Tissue evolution and properties of plasma solid-state surface metallurgical TiCoCrNiWMo high-entropy alloy coatings

Titanium alloy is widely used in aerospace, automobile manufacturing, medical, and health care fields because of its low density, high specific strength, high heat resistance, good corrosion resistance, and excellent biocompatibility. However, the disadvantages of the titanium alloy, such as low hardness, large coefficient of friction, poor wear resistance, and difficult to lubricate, greatly limit its application.^1–3 It is well known that most of the failures of materials such as oxidation, fatigue, wear, and corrosion start from the surface. This is mainly determined by the surface properties of the material rather than the overall properties, so it is especially important to improve the surface properties of the material.^4–6 Surface modification techniques to improve the surface properties of the substrate material to overcome material failure have been proven to be a cost effective and efficient method.^7–9 

At present, the surface treatment technology of the titanium alloy mainly includes thermal spraying, laser cladding, and magnetron sputtering. Although the thermal spraying technology is simple, flexible, and low cost, it has the disadvantages of easy coating flaking and poor uniformity.^10,11 The coating prepared by laser cladding has good bonding properties with the substrate, but the cost of the technology is high and is prone to defects such as porosity and cracks, and there are problems such as compositional segregation and organizational inhomogeneity during the process.^12,13 Although the deposition rate of magnetron sputtering is fast and the resulting coating is dense and uniform, the preparation process and equipment are complicated, the coating is thin, and the bonding strength of the film base is low.^14,15

Plasma solid-state surface metallurgy is a new surface modification technology that has been developed rapidly and widely used in recent years.¹⁶ It is a plasma surface metallurgy method developed on the basis of ion nitriding technology, which uses low-temperature plasma generated by double-layer glow discharge under vacuum conditions to form an alloy coating with special physicochemical properties on the surface of the substrate material by deposition and diffusion.^17,18 The technique is simple, the surface modification is efficient, the thickness of the prepared coating is controllable, and the bonding is good. The process is carried out completely under vacuum conditions, so there is no problem of hydrogen embrittlement and oxidation of the surface alloy coating. It is possible not only to provide high melting point metal elements on the metal surface without pollution and achieve metal surface alloying, but also to prepare high performance surface modified coatings on the surface of various complex shaped workpieces with controllable composition and good organization and structural characteristics. At the same time, it can meet the required performance requirements without adversely affecting the mechanical properties of the substrate.^19–21 It can be seen that plasma solid-state surface metallurgy technology has more advantages compared with the above commonly used surface treatment technology for titanium alloys.

In recent years, plasma solid-state surface metallurgy has been successfully applied to the surface treatment of titanium alloys, and there have been some research results. Wei et al.²² prepared the Cr–Ni alloy coatings with 20, 40, 60, and 80 at. % Ni contents on the TC4 alloy at 730–750 °C using plasma solid-state surface metallurgy technique. It was shown that the alloy coatings with different Ni contents all contained complete deposited layers, interdiffusion layers, and sputtering affected zones. The adhesion of the deposited layers increased with increasing Ni content. When the Ni content was 80%, the bonding force was 87 N and the best bonding performance was achieved, at which time the tensile strength of the TC4 substrate increased by about 2.1% compared with that under the solid solution annealed condition. Zhang et al.²³ used plasma solid-state surface metallurgy to infiltrate alloying element Cr into the Ti-6.5Al-0.3Mo-1.5Zr-0.25Si matrix to form Ti–Cr burn-resistant alloy coatings on its surface. It is shown that the depth of the prepared burn-resistant alloy coating can reach more than 200 μm, and the concentration of alloying elements can reach 90%. When the Cr content in the alloy coating is higher than 14%, the alloy coating plays a good flame retardant effect and can effectively reduce the burning tendency of the titanium alloy substrate. Wei et al.²⁴ prepared Ti–Cr coatings with gradient composition on the TC4 alloy using plasma solid-state surface metallurgy and investigated their isothermal oxidation behavior at 650, 750, and 850 °C. The results show that during the oxidation process, Cr, Ti, and Al diffuse outward to form a multilayer oxide film, which prevents the intrusion of O. The prepared Ti–Cr coating has better oxidation resistance compared with the NiCrAlY thermal barrier coating. Qiu et al.²⁵ deposited a dense CrCoNiAlTiY coating on the surface of the γ-TiAl alloy using plasma solid-state surface metallurgy and investigated its microstructure and wear resistance. It was shown that the coating showed metallurgical bonding with the substrate, with no defects such as pores or cracks on the surface, and the total thickness of the coating was about 5 μm. The overall hardness and wear resistance of the alloy were improved because the coating mainly consisted of high strength and high hardness phases such as AlCr₂, σ-NiCoCr, and γ׳-Ni₃Al. In the wear experiments, the specific wear rate of the coating was reduced by 80% compared with the matrix, and the wear resistance was greatly improved. Wu et al.²⁶ prepared Cr, Cr–W, and Ni–Cr coatings on Ti₂AlNb alloys using plasma solid-state surface metallurgy and investigated their high-temperature oxidation resistance. The results showed that protective Al₂O₃ oxide films or continuous dense NiCr₂O₄ oxide films were formed on the Ni–Cr alloy coatings at 1093 K, and thus, their high-temperature oxidation resistance was better compared with other alloy coatings. It can be seen that plasma solid-state surface metallurgy can be used to form a surface-modified alloy coating on a titanium alloy substrate, thereby effectively improving the properties of the substrate titanium alloy.

However, the surface-modified alloy coatings prepared by plasma solid-state surface metallurgy technology are mainly conventional alloy coatings, and very little research has been conducted on the preparation of new high-entropy alloy (HEA) coatings. In recent years, traditional metallic materials with one or two elements as the main constituents have been difficult to meet production needs, and the research hotspots in the field of materials have gradually shifted to HEAs that break through the traditional alloy design concept.^27–29 As a new type of engineering structural material, the HEA has good development prospects in the fields of aviation, shipping, and construction.^30–32 HEA coatings are generated and developed based on the theory of HEAs and combined with material surface modification technology. It retains the excellent performance of the HEA while reducing the cost of the alloy, which is more conducive to practical applications.^33,34 Many researchers at home and abroad have shown that HEA coatings can substantially improve the surface properties of materials, and their performance is significantly better than that of other coating materials. It has greatly improved the strength, wear resistance, corrosion resistance, and high temperature oxidation resistance of the substrate and is expected to achieve large-scale applications.^35–39 Therefore, a new type of HEA coating can be prepared on the surface of the titanium alloy material by using plasma solid-state surface metallurgy technology to improve other properties of the base material without affecting the original excellent properties of the base material.

Based on the previous research foundation in plasma solid-state surface metallurgy and the HEA, this paper selects five alloy elements, Co, Cr, Ni, W, and Mo, to break the idea of plasma solid-state surface metallurgy technology to prepare traditional alloy-modified coatings on the TC4 substrate and innovatively performs Co–Cr–Ni–W–Mo five-element co-infiltration on the surface of the TC4 substrate. In order to obtain a certain thickness of wear and corrosion-resistant TiCoCrNiWMo HEA coating, a systematic study of its tissue morphology evolution, elemental composition, phase structure, bonding force, hardness, wear resistance, and corrosion resistance was carried out. The aim is to grasp the formation mechanism of HEA coatings, broaden the application scope of plasma solid-state surface metallurgy technology, and provide a research basis for the preparation of HEA coatings by plasma solid-state surface metallurgy.

In this paper, a plasma solid-state surface metallurgy technique (DGLT-15F multifunctional ion chemical heat treatment furnace, Fig. 1) was used to perform Co–Cr–Ni–W–Mo five-member coinfiltration on the surface of TC4 alloy specimens to prepare a new TiCoCrNiWMo HEA coating with metallurgical bonding and a gradient structure. Pure Cr, Ni, Co, W, and Mo bar targets (purity ≥ 99.9%, the size of ∅ 4 × 40 mm²) were selected as source electrode materials to provide deposited elements. The source electrode targets were inserted into circular holes (the size of ∅ 5 mm) uniformly arranged on the surface of the carbon steel hollow cylindrical auxiliary source electrode barrel (the size of ∅ 120 × 160 mm²). Due to the different characteristics of different targets (mainly including atomic radius, sputtering rate, and sputtering yield), and combining the research experience of plasma solid-state surface metallurgy technology and high-entropy alloys as well as the previous experimental explorations, the ratio of targets was selected to be Cr:Ni:Co:W:Mo = 31:31:32:5:5. The size of the TC4 substrate used in this paper was selected to be 130 × 20 × 2.5 mm³ based on prior experimental experience. The substrate is suspended to the center of the auxiliary source barrel by a hook, and the workpiece (substrate) is connected to the auxiliary source barrel by the hook, so that the source and substrate can be controlled simultaneously by a single power source, and the substrate and source can form an equipotential effect. The conventional plasma solid-state surface metallurgy technique for preparing surface alloy coatings (solid solution diffusion layers) on Ti-based alloy surfaces is generally selected at a holding temperature of 800–1000 °C. The preparation of TiCoCrNiWMo gradient HEA coatings on Ti-based alloy surfaces requires a larger sputtering deposition of alloying elements to ensure that the deposited layer can be formed compared with the preparation of conventional surface alloy coatings (solid solution diffusion layer). Therefore, it is proposed in this paper to use a higher holding temperature to ensure a sufficient supply of alloying elements for sputtering deposition. The rest of the process parameters, such as holding time and working air pressure, were determined by previous experience with plasma solid-state surface metallurgy technology. The specific process parameters are shown in Table I. After the test, the furnace temperature was cooled to room temperature before removing the samples. In this paper, different holding temperatures (1000, 1050, and 1100 °C) were set to investigate the organization and properties of TiCoCrNiWMo HEA coatings at different holding temperatures.

A Quanta FEG 450 scanning electron microscope (SEM) equipped with an x-ray energy spectrum analyzer (EDS) was used for morphology, organization, and composition of the HEA coating. A Bruker D8 Advance x-ray diffractometer (XRD) was used to analyze the phases of the HEA coating with a scan range of 20°–100° and a scan step of 0.01°. A WS-2005 coating adhesion automatic scratch tester was used to test the bonding force of HEA coating. The maximum load was 100 N, the loading speed was 100 N/min, the sliding speed was 5 mm/min, and the scratch length was about 5 mm. The HV-1000 micro Vickers hardness tester was used to perform hardness tests on the HEA coatings. To ensure the accuracy of the results, 10 different locations were selected for the same HEA coating for testing, and finally, the average of the hardness values of each group was taken as the final hardness value. A high-speed reciprocating friction and wear tester of type HSR-2M was used to conduct friction and wear tests on HEA coatings, and the abrasive material was small Si₃N₄ balls of 5 mm in diameter. The loading load was 10 N, the running time was 60 min, the reciprocating distance was 5 mm, the speed was 300 rpm, and the sampling frequency was 30 Hz. Five reciprocating friction and wear tests were conducted under the same conditions to ensure the repeatability of the test. Within the error range, the friction coefficient and wear amount of the HEA coating were averaged over the five repeated tests. The samples were cut into 10 × 10 × 2 mm³ using wire erosion, and the HEA coating was electrochemically tested using Shanghai Chenhua CHI660D electrochemical workstation. A standard three-electrode system was used, with the specimen as the working electrode (WE), saturated glycury electrode (SCE) as the reference electrode, and platinum sheet electrode as the auxiliary electrode (CE), and the electrolyte solutions were 1 mol/l H₂SO₄ solution, 3.5 wt. % NaCl solution, and 1 mol/l NaOH solution, respectively. The experiments were started after the open circuit potential stabilized. The scanning interval of the polarization curve test was −1 to 1 V, and the scanning speed was 0.005 V/s. The AC current frequency in the AC impedance test was 10⁵–10⁻² Hz, and all tests were performed three times to improve the accuracy of the test results.

where R denotes the gas constant, n denotes the number of elements, and c_i denotes the mole fraction of component i.

Table II demonstrates the EDS compositions (at. %) of the labeled points in Fig. 2 and the mixing entropy (J K⁻¹ mol⁻¹) of the HEA coatings calculated by Eq. (1). It can be seen that TiCoCrNiWMo HEA coatings meeting the widely accepted definitions of two high-entropy alloys can be successfully prepared on TC4 substrates using plasma solid-state surface metallurgy. However, there are obvious differences in the surface composition of HEA coatings at different holding temperatures. When the holding temperature is 1000 °C, the composition of the HEA coating is uniform, there is no compositional segregation, and the quality of the HEA coating is good. In addition, when the temperature exceeds the (α + β)/β transition temperature of TC4 (Ref. 40) (980–1000 °C) at the holding temperature of 1050 and 1100 °C, the composition of the coating is severely polarized, and the quality of the high-entropy alloy coating is unstable, which leads to obvious differences in the composition of the HEA coating at different morphologies. In general, there are large differences in the composition of HEA coatings with different holding temperatures with respect to the target composition (ratio), which can be explained by the plasma solid-state surface metallurgy sputtering mechanism.

Table III shows the atomic radii and crystal structures of different alloying elements and their sputtering rates under 100 eV Ar⁺ bombardment. Compared with the Ti, Al, and V elements of the matrix, the five target elements of Co, Cr, Ni, W, and Mo have larger differences in the atomic radius size of Cr and Ni elements, and their solid solution in the matrix is relatively low, and their sputtering rates are relatively high, which are easily backsputtered out due to the strong backsputtering effect. Therefore, the target composition of Cr and Ni elements is high, while their composition content in HEA coatings is relatively low. Compared with the Ti, Al, and V elements of the substrate, the Co element has a large difference in the atomic radius size, but its sputtering rate is moderate, while the percentage of the Co element composition in the target material is relatively high, so its composition content in the HEA coating is relatively high. The atomic radius size difference between W and Mo elements and Ti, Al, and V elements of the matrix is small. Meanwhile, in the working holding temperature, W and Mo elements can be infinitely miscible with Ti elements,⁴¹ and their crystal structures are the same as Ti and V elements of the matrix. This leads to good solid solution and penetration properties of W and Mo elements in the matrix TC4 alloy. Therefore, the composition of W and Mo elements in the target material is relatively low, but their composition content in the HEA coating is relatively high. The W element has a lower sputtering rate compared to the Mo element, and the sputtering supply of the element is insufficient, so its composition content in the HEA coating is much lower than that of the Mo element.

Figures 3 and 4 show the cross-sectional SEM images and EDS line scans of the TiCoCrNiWMo HEA coatings at different holding temperatures. From Figs. 3 and 4, it can be seen that the gradient TiCoCrNiWMo HEA coating with a composite reinforced layer structure of deposited layer (1.5 μm) + diffusion layer (3.5 μm) was formed on the surface of the substrate at the holding temperature of 1000 °C. The deposited layer has a columnar crystal structure, but there are gaps, which may be the interstices between the long stripes of the surface organization. The content of Co elements in the diffusion layer decreases with the increase in the depth, showing a clear gradient distribution. It shows that the HEA coating can present metallurgical bonding with the substrate material. When the insulation temperature is 1050 °C, only a deposit layer is formed on the surface of the substrate. Interestingly, the deposited layer is subdivided into an outer layer (2 μm) and a transition layer (4.5 μm), and the alloying element with a distinct gradient process in the transition layer is the Mo element rather than the Co element. This is because the atomic radius of the Co element is small, so it shows a gradient distribution in the diffusion layer when the holding temperature is 1000 °C. The atomic radius size difference between the Mo element and matrix Ti, Al, and V element is small, and it can be infinitely miscible with the Ti element, so its content is higher in the deposited layer and shows a gradient distribution in the transition layer. The (α + β)/β transition temperature of TC4 is 980–1000 °C,⁴⁰ beyond which coarse grain organization is extremely easy to form. Therefore, when the holding temperature is 1050 °C, after the nucleation of the neutral particles first sputtered onto the surface of substrate TC4, the grains rapidly start to grow and coarsen, rapidly forming a diffusion barrier on the surface of the substrate, which prevents the target elements from entering the substrate to form a diffusion layer. On the other hand, because the equipotential mode is used in this experiment, the thickness of the deposited layer does not continue to increase with the disappearance of the vacancy concentration gradient layer and the formation of a dynamic equilibrium between the sputtered and deposited amounts on the substrate surface. However, the whole substrate is still in a high-energy active state, and the outer layer of the deposited layer is constantly bombarded and backsputtered, so the deposited layer gradually forms a delamination phenomenon. At a holding temperature of 1100 °C, a deposited layer (6 μm) was formed on the surface of the substrate only, which was similar to the transition layer formed at a temperature of 1050 °C. At this time, the Mo element content decreases with increasing depth, showing a gradient change. The main reason for the absence of the outer layer of this deposited layer is because the outer grain organization of the deposited layer is further coarsened as the temperature increases, which, together with the significantly enhanced backsputtering effect, causes the loosely bound coarsened organization of the outer layer to be backsputtered out, resulting in the disappearance of the outer layer of the deposited layer.

Figure 5 shows the cross-sectional EDS surface scans of TiCoCrNiWMo HEA coatings at different holding temperatures. It can be seen that at the holding temperature of 1000 °C, Mo elements are enriched in the deposited layer, Co elements are enriched in the diffusion layer, and Ni elements are enriched in the diffusion layer to a shallower extent. When the temperature is 1050 °C, Mo is enriched in the whole sediment layer, W is enriched in the whole sediment layer to a shallow degree, and Co and Ni are enriched in the outer layer of the sediment layer to a shallow degree. At a holding temperature of 1100 °C, Mo is clearly enriched in the whole sediment layer, and W is enriched to a shallow extent in the whole sediment layer. This is consistent with the results of the EDS line scan in Fig. 4.

Figure 6 shows the XRD patterns of TiCoCrNiWMo HEA coatings with different holding temperatures. It can be seen that the effect of holding temperature on the physical phases of the TiCoCrNiWMo HEA coating is small, and only the intensity of diffraction peaks is changed without changing its phase structure. The HEA coatings at different temperatures are the same, and their main alloy phases are Al₁₁V, Al₉Cr₄, (Ni,Ti), AlMoTi₂, and Co_1.3Ni_4.3Mo_4.6 phases.

The bond strength between HEA coating and substrate TC4 is an important indicator of the quality of the alloy coating and also determines the service life of the alloy coating to a certain extent.^42, Figure 7 shows the acoustic emission curves of TiCoCrNiWMo HEA coatings with different holding temperatures during the scratching process and their scratching morphology. During the scratching process, the HEA coating is subjected to a vertical pressure F_V and a tangential stress F_H pointing in the direction of the scratch and generates strain energy, which is released in the form of microcracks within the HEA coating and separation of the HEA coating from the substrate. When the scratch force is small, the interior of the scratch is smooth, and as the load increases, a few cracks begin to develop within the HEA coating, corresponding to a load that is the critical load Fc* for the cohesive failure of the HEA coating. As the load increases, the cracks in the HEA coating gradually become denser until peeling occurs, and even the HEA coating at the boundary of the scratch also peels off, at which time the load is the critical load Fc for the failure of the interface between the HEA coating and the substrate. Fc* and Fc are reflected in the acoustic emission curve (I–F curve) as a sudden enhancement of acoustic emission intensity.^43,44

As can be seen from Fig. 7, the acoustic emission intensity of the HEA coating at different holding temperatures starts to produce some intermittent abrupt changes from about 59–64 N, indicating that the brittleness of the HEA coating itself makes different degrees of fine microcracks within the coating. However, it is worth noting that the critical load Fc for failure does not appear on the figure, indicating that no spalling of the HEA coating occurred. The scratch morphology in Fig. 7 is in good agreement with the I–F curve, and it can be seen that different degrees of transverse cracks can be observed in the direction of the HEA coating perpendicular to the scratch for different holding temperatures, but no seepage spalling occurs inside the scratch or at the boundary of the scratch until the end of the scratch. When the holding temperature is 1000 °C, the HEA coating composition is uniform, there is no compositional segregation, and it can be metallurgically bonded with the TC4 substrate, and the coating quality is good. While the holding temperatures of 1050 and 1100 °C exceed the (α + β)/β transition temperature of TC4, the grain organization of the HEA coating is coarse, and there are interstitial segregation and surface compositional segregation, and the quality of the coating deteriorates. Therefore, the HEA coating with a holding temperature of 1000 °C has the best bonding performance, with a bonding force of about 63.81 N. The bonding force of the film base at holding temperatures of 1050 and 1100 °C is relatively poor, with bonding forces of 61.85 and 59.15 N.

Figure 8 shows the micro-Vickers hardness of the substrate TC4 and TiCoCrNiWMo HEA coatings. It can be seen that the microhardness of HEA coatings with different holding temperatures have all increased to different degrees compared with the substrate. The hardness of the HEA coating with the holding temperature of 1000 °C is the highest, about 541.26 HV_0.3, which is 1.5 times higher than that of the substrate (363.83 HV_0.3). Due to the different atomic radii of each primary element in the preparation of HEA coatings, the lattice dot matrix occupied by them has a random nature. Therefore, the HEA coating has an obvious solid solution strengthening effect, which leads to dislocation movement and difficulty in crystalline surface slip, thus increasing the microhardness of the substrate. The holding temperatures of 1050 and 1100 °C exceed the (α + β)/β transition temperature of the TC4 substrate, the grain organization of the HEA coating is coarse, and there are gaps and surface composition segregation, and the quality of the coating deteriorates. Therefore, the hardness of HEA coatings at holding temperatures of 1050 °C and 1100 °C is relatively low.

Figure 9 shows the frictional wear properties of the substrate TC4 and TiCoCrNiWMo HEA coatings: (a) friction coefficient, (b) friction coefficient and hardness, (c) wear volume, and (d) cross-sectional wear scar profile. The coefficient of friction of a material is an important parameter used to characterize the wear resistance of a material. The higher the average coefficient of friction, the lower the hardness, and the worse the wear performance of the material. In Fig. 9(a), the variation in the friction coefficients of the substrate TC4 and HEA coatings has the same pattern. In the initial stage, the friction coefficient is low because the sample surface is smoother and the roughness and friction resistance are small. After that, with the increase in the friction time, the abrasion marks on the sample surface deepen, causing the roughness to become larger and the hindering effect of the frictional substrate to strengthen. Sharp grinding of the friction sub- and sample surface occurs, resulting in a rapid increase in the friction resistance and friction coefficient. After a short grinding phase, the wear process enters a stable period, the friction coefficient tends to stabilize and enters a stable wear phase.^45,46 From Fig. 9(b), it can be seen that the average coefficients of friction of the substrate TC4 and HEA coatings show a negative correlation with their hardness. The coefficient of friction of a material is an important parameter used to characterize the wear resistance of a material, the higher the average coefficient of friction, the lower the hardness, and the worse the wear resistance of the material.^47,48 The friction coefficients of HEA coatings with different insulation temperatures are smaller than those of the substrate, which shows that HEA coatings have a certain friction reduction effect. As shown in Fig. 9(c), it can be seen that the wear resistance of the HEA coating at different holding temperatures is improved to different degrees compared with the substrate. The HEA coating with an insulation temperature of 1000 °C had the best wear resistance, which improved the wear resistance by 8.9 times compared to the substrate. The wear resistance of the HEA coating at the holding temperatures of 1050 and 1100 °C was average and only slightly improved the wear resistance of the substrate. This is because the holding temperature exceeds the (α + β)/β transition temperature of TC4,⁴⁰ and the grain organization of the HEA coating is coarse, which reduces the plasticity and increases the brittleness of the HEA coating and reduces its wear resistance. As shown in Fig. 9(d), it can be seen that the width and depth of wear marks of HEA coatings with different holding temperatures are smaller than those of the substrate. Among them, the width and depth of abrasion marks of the HEA coating with the holding temperature of 1000 °C are obviously lower, which is consistent with the results shown in Fig. 9(c).

Figure 10 shows the surface wear morphology of substrate TC4 and TiCoCrNiWMo HEA coatings. It can be seen that the surface of the substrate wear marks is distributed with a large number of parallel plow grooves and sporadic abrasive chips, and the surface is torn, forming deep grooves and spalling pits. The surface of the HEA coating with holding temperature of 1000 °C has a shallow separated spalling layer and sporadic abrasive chips. A clear gap between the long strips of tissue can also be observed, indicating that the long strips of tissue are not completely worn away. The surfaces of the HEA coatings with holding temperatures of 1050 and 1100 °C were characterized by the presence of large amounts of abrasive debris and furrows of varying depths parallel to each other. Table IV shows the EDS composition of the marked points in Fig. 10. It can be seen that the wear surfaces of the substrate TC4 and TiCoCrNiWMo HEA coatings with different holding temperatures have a high content of element O on them. This is due to the production of heat during the friction process, which causes high temperature oxidation of the surface, resulting in oxidative wear. The higher elemental O content at points 2 and 3 in Fig. 10 compared to the other points is due to the greater degree of oxidation at these two points. Interestingly, the content of Ti elements on the HEA coating with a holding temperature of 1000 °C is significantly lower, but the content of the five elements Co, Cr, Ni, W, and Mo is relatively high. Combined with surface morphology, it can be seen that the HEA coating at a holding temperature of 1000 °C has good wear resistance. In summary, the wear mechanism of the matrix TC4 is mainly dominated by abrasive wear, accompanied by adhesive wear and oxidative wear. The wear mechanism of the HEA coating at a holding temperature of 1000 °C is mainly adhesive wear, accompanied by oxidative wear and slight abrasive wear. The wear mechanism of the HEA coatings with holding temperatures of 1050 and 1100 °C is dominated by abrasive wear, accompanied by oxidative wear.

Figure 11 shows the electrochemical polarization curves and the AC impedance curves of the substrate TC4 and TiCoCrNiWMo HEA coatings in 1 mol/l H₂SO₄ acidic solution, 3.5 wt. % NaCl neutral solution, and 1 mol/l NaOH alkaline solution. The corresponding electrochemical corrosion parameters are shown in Table V. From Figs. 11(a), 11(c), and 11(e), it can be observed that the electrochemical polarization curves of all samples in different corrosion solutions show a significant passivation, which is due to the generation of passivation films on the surface of the substrate TC4 and HEA coatings, reducing the current density on their surfaces. The passivation film itself is a thin film composed of oxides or hydroxides of metals. The formation process is mainly the process of stable oxides or hydroxides formed by the chemical reaction between Ti, Cr, Ni, Co, Mo, and other metal cations in the sample and the corrosion solution. Typically, the formation of the passivation zone is shown on the polarization curve: as the potential increases, the current density is lower and remains constant. It is noteworthy that for the samples in 1 mol/l NaOH alkaline solution, two obvious current peaks of the activation-passivation transition exist near −0.3 and 0 V in the corresponding anodic polarization curves. A wider secondary passivation zone was formed after 0 V, which indicated that a protective passivation film that could be stable existed at this potential. In contrast, the samples underwent a more pronounced slight self-passivation at a lower polarization potential of −0.2 to 0 V in 1 mol/l H₂SO₄ acidic solution, and the current increased sharply for all samples when the polarization potential increased to about 0.1 V. This was due to the fact that the passivation film generated by the initial passivation at this potential was not sufficient to inhibit the anodic dissolution process caused by the increase in the polarization potential. Relevant studies show that the self-corrosion current density represents the corrosion tendency of the alloy.^49,50 The smaller the self-corrosion current density, the greater the polarization resistance, the smaller the corrosion rate, the more difficult to carry out corrosion, the better the corrosion resistance of the alloy is, and conversely, the worse the corrosion resistance. Combined with Table V, it can be seen that in the active dissolution region under different corrosion conditions, the corrosion current density of the HEA coating with a holding temperature of 1000 °C is the lowest followed by the HEA coating with holding temperatures of 1050 and 1100 °C and the highest corrosion current density of the matrix TC4. It shows that the corrosion resistance of HEA coatings at different holding temperatures is improved to different degrees than the base TC4. In particular, the corrosion current density of the HEA coating with the holding temperature of 1000 °C was one order of magnitude lower than that of the base TC4 in the neutral solution of 3.5 wt. % NaCl and the alkaline solution of 1 mol/l NaOH, showing excellent corrosion resistance.

From Figs. 11(b), 11(d), and 11(f), it can be seen that in different corrosive environments, all samples have similar capacitive arcs of resistance and show similar EIS characteristics. This is due to the charge transfer at the dielectric solution interface and the electrode surface; i.e., the charge transfer occurs at the inhomogeneous interface. Related studies have shown that the radius of the capacitive arc can reflect the polarization resistance of the material in the electrochemical process.^51,52 Different radii of the capacitive resistance rings have different abilities to impede charge transfer. The larger the radius of the capacitive resistance arc, the larger the impedance value, the better the corrosion resistance of the material. As can be seen from Fig. 11, in different corrosive solutions, compared with the radius of capacitive arc resistance of the substrate TC4, the radius of capacitive arc resistance of HEA coatings with different holding temperatures has different degrees of increase. The HEA coating with a holding temperature of 1000 °C has the largest capacitive arc resistance radius, followed by HEA coatings with holding temperatures of 1050 and 1100 °C. Therefore, the corrosion resistance in 1 mol/l H₂SO₄ acidic solution, 3.5 wt. % NaCl neutral solution, and 1 mol/l NaOH alkaline solution in descending order is as follows: HEA coating with a holding temperature of 1000 °C > HEA coating with a holding temperature of 1050 °C > HEA coating with a holding temperature of 1100 °C > substrate TC4 alloy. This is consistent with the results obtained from electrochemical polarization curve tests in Figs. 11(a), 11(c), and 11(e).

Figure 12 shows the corrosion morphology of the substrate TC4 and TiCoCrNiWMo HEA coatings in different corrosive media after electrochemical tests. As can be seen from Fig. 12, the surface of the substrate TC4 uniformly distributed with small corrosion pits, the surface of the corrosion products has been part of the flaking, and residual white products irregularly distributed on its surface. Different shades of black corrosion spots appeared on the surface of the HEA coating for different holding temperatures. The “long stripes” on the surface of the HEA coating with a holding temperature of 1000 °C and the irregular masses on the surface of the HEA coating with a holding temperature of 1050 °C became less pronounced than they were before corrosion. The corrosion mechanisms of all samples in different corrosion solutions were dominated by uniform corrosion with uniform pitting. TiCoCrNiWMo HEA coatings with different holding temperatures have more excellent corrosion resistance than the base TC4 in acidic, alkaline, and salt solutions. This is related to its surface that is more uniform and dense and with less impurities, holes with cracks, and other defects. Interestingly, the corrosion resistance of the HEA coating differs again for different holding temperatures. This may be because the holding temperatures of 1050 and 1100 °C exceeded the (α + β)/β transition temperature of TC4,⁴⁰ HEA coating grain organization is coarse, and there are gaps and surface composition segregation. When the holding temperature is 1000 °C, the composition of the HEA coating is uniform and there is no composition segregation, and it can be metallurgically bonded with the TC4 substrate. In addition, at this temperature, the grains are smaller, and the surface corrosion layer is denser and more uniform than the previous two, which can more effectively impede the penetration of the corrosion solution, thus slowing down the corrosion of the alloy by the corrosion solution. The dense corrosion layer generated in the corrosion increases the bonding force between the corrosion active material and the HEA coating, reduces the shedding of the active material, and also helps to enhance the protective properties of the HEA coating passivation film. Therefore, the HEA coating with a holding temperature of 1000 °C has more excellent corrosion resistance in acidic, alkaline, and salt solutions compared with other samples, which is consistent with the results in the electrochemical polarization curve and AC impedance curve.

In this paper, we used plasma solid-state surface metallurgy to perform Co–Cr–Ni–W–Mo five-member coinfiltration on the surface of the TC4 alloy. The gradient TiCoCrNiWMo HEA coating was successfully prepared. The effects of different holding temperatures on the thickness, organization, composition, morphology, and structure of TiCoCrNiWMo gradient HEA coatings were systematically investigated. The HEA coating was also compared with the base TC4 alloy for bonding properties, mechanical properties, and corrosion resistance. The main conclusions obtained are as follows:

1. At a holding temperature of 1000 °C, the TiCoCrNiWMo gradient HEA coating is a composite reinforced layer structure of deposited layer + diffusion layer with a uniform and dense surface and no composition segregation. When the holding temperature is higher than 1000 °C, the HEA coating only has a deposited layer, and the surface grain organization is coarsened, resulting in interstitial and surface composition segregation.

2. The holding temperature did not change the physical phase structure of the HEA coating, and the main alloy phases were Al₁₁V, Al₉Cr₄, (Ni, Ti), AlMoTi₂, and Co_1.3Ni_4.3Mo_4.6 phases. The bonding performance of the HEA coating with the substrate at a holding temperature of 1000 °C was the best, with a bonding force of about 63.81 N.

3. The mechanical properties of TiCoCrNiWMo HEA coatings with different holding temperatures were all improved to different degrees compared with the base TC4 alloy. Among them, the best performance was obtained for the HEA coating at 1000 °C, with micro-Vickers hardness and wear resistance 1.5 and 8.9 times higher than those of the matrix, respectively.

4. The corrosion resistance of TiCoCrNiWMo HEA coatings with different holding temperatures in 1 mol/l H₂SO₄ acidic solution, 3.5 wt. % NaCl neutral solution, and 1 mol/l NaOH alkaline solution all showed different degrees of improvement relative to the substrate. In particular, the HEA coating with a holding temperature of 1000 °C showed excellent corrosion resistance in salt and alkaline solutions.

The authors acknowledge the National Natural Science Foundation of China (NNFSC) (Grant No. 52161011), the Natural Science Foundation of Guangxi Province (Grant Nos. 2023GXNSFDA026046 and 2020GXNSFAA297060), the Scientific Research and Technology Development Program of Guilin (Grant Nos. 20220110-3, 2020010903, and 20210217-6), the Guangdong Basic and Applied Basic Research Foundation (Grant No. 2020B1515420004), the Guangxi Key Laboratory of Superhard Material (Grant No. 2022-K-001), the Guangxi Key Laboratory of Information Materials (Grant Nos. 221024-Z and 221012-K), the Engineering Research Center of Electronic Information Materials and Devices, Ministry of Education (Grant No. EIMD-AB202009), the Major Research Plan of the National Natural Science Foundation of China (Grant No. 92166112), the Projects of MOE Key Lab of Disaster Forecast and Control in Engineering in Jinan University (Grant No. 20200904006), and the Open Project Program of Wuhan National Laboratory for Optoelectronics (Grant No. 2021WNLOKF010), and Innovation Project of GUET Graduate Education (Grant Nos. 2020YCXS118 and 2022YCXS200) for financial support given to this work.

The authors have no conflicts to disclose.

Xin Li: Writing – original draft (equal); Writing – review & editing (equal). Zixiang Zhou: Writing – review & editing (equal). Chenglei Wang: Writing – review & editing (equal). Haiqing Qin: Writing – review & editing (equal). Jijie Yang: Writing – review & editing (equal). Weijie Liu: Writing – review & editing (equal). Mulin Liang: Writing – review & editing (equal). Chong Liu: Writing – review & editing (equal). Hong Tan: Writing – review & editing (equal). Zhenjun Zhang: Writing – review & editing (equal).

The data that support the findings of this study are available from the corresponding author upon reasonable request.