Investigation into corrosion and wear behaviors of laser-clad coatings on Ti6Al4V

TiAlCoCrxFeNihigh-entropy alloys (HEAs) coatings were fabricated on the surface of Ti6Al4V by laser cladding. Their microstructural evolution with the increase in x value (x = 0, x = 1.0, x = 2.0) was investigated in detail. Besides that, the investigation into the effects of the Cr content on their corrosion behaviors and mechanical properties (in terms of hardness and wear resistance) was also carried out comprehensively. The results indicated that two kinds of phases (a solid solution with thehexagonal close-packed (HCP) structureand Ti2Ni) were synthesized in the coatings, and the HCP content was gradually increased with the increase in x accompanied with the decrease in Ti2Ni content. A HEA coating only composed of single HCP was successfully prepared when x reached 2.0. The electrochemical and immersion tests all confirmed that the coating with x = 2.0demonstrated the most excellent corrosion resistance in a0.1 mol·L−1 HCl solution from different aspects including corrosion tendency and corrosion rate without the applied potential, the formation difficult/stability of the passive film and the dissolution rate in the passive state, and corrosion surface morphology. The averagemicrohardness values of the coatings weregraduallyincreasedfrom 656HV0.2 to 800 HV0.2with increasing xfrom 0 to 2.0, which wereabout double that of the substrate (350 HV0.2). Wear resistance of the coatings also exhibited the upward tendency with increasing the x values (0.562 mm3 at x = 2.0, 0.640 mm3 at x = 1.0, 0.641 mm3 at x = 0 and 1.419 mm3 for the substrate). More Cr addition into the cladding material will contribute to the formation of a HEA coating composed of single HCPwith excellent corrosion and wear resistance.


Introduction
Titanium alloysexhibit excellent corrosion resistance in the static corrosion environment since a thin and compact oxide layer spontaneously formed on their surfaces can protect them from corrosion.However, the thin protective layer easily suffers from the serious damage and peels off from the surfaces when titanium alloys contact with the other components and do a relative motion under the friction condition. The exposedtitanium alloys will be subjected to the severe chemical/electrochemical dissolution, accompanied with the sharp microcutting loss due to their very low hardness of about 350 HV 0.2 . The interaction between corrosion and wear will further accelerate the failure of titanium alloys. Wear and corrosion all stem from the surfaces of titanium alloys, therefore improving the surface properties becomes an urgent problem in the application process of titanium alloys.
Up to now, many surface modification methods have been explored and applied to improve surface properties of titanium alloys [1], including electroplating [2], physical vapor deposition [3], chemical vapor deposition [4], micro-arc oxidation [5], and so on [6][7][8][9]. However, the protective coatings prepared by above methods are usually thin (micrometer scale in thickness) and loose in microstructure, and have a weak bonding with the substrate, whose applications are greatly limited in the harsh corrosion wear environment. Comparatively speaking, laser cladding as a new and promising technique can resolve the above shortcomings and endow titanium alloys with excellent corrosion/wear resistance. The laser-clad coating exhibits an excellent metallurgical bonding with the substrate, and its thickness can be flexibly adjusted from micrometers to several millimeters by controlling the thickness of the pre-placed layer or the feeding rate of the powder, and regulating the processing parameters. Most importantly, the cladding material will undergo a process of rapid melting and solidification underlaserirradiation, resulting in a dense and fine microstructure obtained in averyshortperiodoftime. Many investigations into the laser-clad coating prepared on titanium alloy have been carried out.
A large number of metal-based composite coatings had been synthesized on different metals and alloys. Lei et al [10] prepared carbon fiber reinforced nickel-based composite coatings on the surface of 1Cr13 stainless steel by laser cladding. Compared withthe coating without carbon fibers,the hardness of the coatings with 6 vol.% carbon fibers was increased by 30%. In addition, the corrosion current density and the wear rate were reduced by 93% and 45%, respectively.Shi et al [11]also applied laser cladding to prepare carbon fiber reinforced nickelbased composite coatings on the surface of Q235 steel. The results showed that the distribution of carbon fibers and metal carbides was more uniform, and the combination between carbon fibers and matrix also became better when the content of carbon fiberswas gradually increased from 0 vol.% to 9 vol.%. The average microhardness of the coating with carbon fibers was 678 HV 0.2 , which was 1.7 times higher than that of the coating without carbon fibers. The ultimate tensile strength was evaluated from 164 MPa to 608 MPa when the content of carbon fibers was increased from 0 vol.% to 9 vol.%. The coatings were also fabricated on the other alloys such as copper alloys [12], aluminum alloys [13], magnesiumalloys [14] and titanium alloys [15][16][17][18][19][20] etc With respect to titanium alloys, considerable metal-based composite coatings reinforced by the ceramic particles had been synthesized on their surfaces (TaC-reinforced TiNi/Ti 2 Ni matrix [15], ZrO 2 reinforced TiNi matrix [16], TiC reinforced α-Ti matrix [17], TiB/TiC/TiNreinforced α-Ti matrix [18] and so on [19,20]. Wear resistance of these coatings is all greatly improved when compared with that of titanium alloys due to their very high hardness of about 750-1400 HV resulting from a combination of the refinement/solution/dispersion strengthening effects. However,the intermetallic compounds involved in the above-mentioned composite coatings will bring about some new issues, such as poor machinability and high cracking susceptibility [21,22]. In addition, the further improvement in their comprehensive properties is also greatly restricted by the main components. A novel concept of "multi-component high-entropy alloys (HEAs)" was defined by Yeh et al [23], in which a single solid solution with body-centered cubic (BCC)structure [24], face-centered cubic (FCC)structure [25] and hexagonal close-packed (HCP) structure [26] is synthesized. Owing to the high entropy effect in a multicomponent system, the formation of intermetallic compounds can be greatly inhibited during the solidification. It can be concluded that the HEA coating can effectively resolve the above-mentioned shortcomings involved in the traditional composite coating. However, the present investigations in terms of HEAs mainly focus on the component optimization in different bulk HEAs prepared by the melting method. Some multicomponent systems have been explored (CrMnFeCoNi [27,28], CoCrCuFeNi [29,30], AlCoCrFeNi [31,32], AlCoCrFeNiTi [33,34] and so on [35,36]), among which AlCoCrFeNiTi attracts more attentions due to its higher plasticity/toughness, hardness and excellent corrosion resistance.However, there are few reports about the AlCoCrFeNiTi coatings prepared on titanium alloys (especially by laser cladding). Moreover, Cr as a main corrosion-resistant component in HEAs plays an essential role in improvement of corrosion resistance. Unfortunately, its effect on corrosion resistance of HEAs was not revealed in detail in present studies. The relationship between content of Cr and wear resistance of HEAs was also not established.
In this study, three multicomponent systems with different contents of Cr (TiAlCoCr x FeNi, x=0, 1.0, 2.0) were used to fabricate the HEAs coatings on Ti6Al4V by laser cladding. Microstructural evolution of the coatings with the change in x was investigated. The mechanical properties in terms of wear resistance and microhardness, and corrosion behaviors of the coatings were also analyzed in detail.

Experimental procedures
Ti6Al4V as the substrate was machined into the cylinders with 50 mm in diameter and 10 mm in thickness.Al, Co, Cr, Fe, and Ni powders with high purity (99.5 wt.%) were selected as the cladding materials, which weredried for 1 h at 80°C and weighted in three different mole ratios (x=0, 1.0, 2.0 in AlCoCr x NiFe) using a SartoriousBSA124S analytical balance (0.1 mg resolution), then uniformly mixed by a QM-3SP2 planetary ball milling.
The substrate surface was ground with 150# silicon carbide abrasive paper and ultrasonically cleaned in acetone.Apre-placed layer with a thickness of approximately 1.0 mm was prepared by the modification binding method [37]. Laser cladding was operated using an YLS-5000 fiber laser. The parameters were optimized as follows: applied power of 3 kW, scanning rate of 5 mm·s −1 , and spot diameter of 6 mm. ThreeTiAlCoCr x FeNi HEAscoatings were named asCoating I(x=0), CoatingII (x=1.0)and Coating III (x=2.0).
The macromorphologies and microstructures of the coatingswere characterized using a VHX-600 K optical microscope (OM) and a HITACHI S-3400 scanning electron microscope (SEM). Their chemical compositions were detected by a GENESIS EDAX energy dispersive spectrometer (EDS). Phase constituents were analyzed using a PANalytical X' Pert Pro x-ray diffractometer (XRD) with Cu target radiation (λ=0.154060 nm).
Microhardness distributionacrossthe coatings'crosssectionswas measured by a HXD-1000TMSC/LCD microhardness testerwith a load of 200 gf and a duration time of 15 s. Wear resistance of the coatings and the substrate was evaluated by using the ball-on-disk reciprocating mode on a CFT-1 friction machineat room temperature (25°C). Prior to the tests, the surfaces of the coatings were ground to acquire an approximately same surface roughness. YG6 ceramic balls (5 mm in diameter) were selected as the counterparts. The sliding time was 120 min, the applied load was30 N and the sliding speed was maintained at 0.1m·s −1 . Morphologies and chemical compositions of wear surfaces were characterized by SEMcoupled with EDS. Wear volumes of the samples were measured for four times by a surface mapping profilometer, and the average value was obtained.
The electrochemical tests (including potentiodynamic anodic polarization tests and electrochemicalimpedance tests) were performedonan AUTOLAB PGSTAT302 electrochemical workstation in a0.1 mol·L −1 HCl solution at room temperature (25°C). The traditional three-electrode system was applied, in which the coatings and the substrate were selected as the working electrode, a Pt sheet and a saturated calomel electrode were chosen as the counter electrode and the reference electrode.For the potentiodynamic anodic polarization tests, the potential was swept from −0.8 to 1.0 V with a scanning rate of 1 mV·s −1 . The electrochemical impedance spectrograms were recorded with the frequency from 100 kHz to 0.01 Hz and the amplitude of 10 mV around theopen circuit potential (OCP). In order toinvestigate the corrosion mechanisms of the coatings, an Escalab 250Xi X-ray photoelectron spectroscope (XPS) with a monochromatic Al Kα excitation was applied to detect chemical valence states of metal elements involved in the passive film formed on CoatingIII in the EIS tests. The binding energy scale was calibrated by the C 1s peak (284.67 eV). To further study corrosion resistance of thecoatings, 30 days immersion tests were conducted in a 0.1 mol·L −1 HCl solution at room temperature (25°C). The surface morphologies of the samples after immersion tests were observed by OM.

Results and discussion
3.1. Structural characterization 3.1.1. Macromorphologies and dilution rates Figure 1 shows the cross-sectional macromorphologies of the samples. The samples can be divided into three regions: corresponding to coating, heat-affected zone and substrate. The coatings' surfaces are comparatively smooth and basically parallel to the base surface.The fusion line can be clearly observed between coating and heat-affected zone, which indicates thata good metallurgical bond isformed between them. The coatings are very dense and free from cracks and holes. There are no significant differences in macromorphology of three coatings.
The dilution rate is an essential factor determining the performance of laser-clad coatings. In order to avoid serious degradation of the coating performance, the dilution rate is usually maintained at about 2%-10%. However, the dilution rate should be increased if the elements from the substrate are required to participate in the formation of the phase in the coating. As an important component of high-entropy alloy coatings, Ti comes from the substrate in this study. Therefore, it is necessary to controlthe dilution rateby regulating the processing parameters, so that anappropriateamountTi can participate in the reactions.
The dilution rate can be calculatedas follows [20]: in which S 1 refers to the areaof the coatingbelowthe substrate surface, S 2 signifiesthe area of the coatingabove the substrate surface.
The cross-sectional profile of the coating can be considered as a combination of an ideal arc and a rectangle (figure 2). Therefore, S 1 and S 2 can be calculated according to the geometric relationship [38]: where H 1 and H 2 represent the depths of the coating below and above the surface of the substrate, respectively, W denotesthe width of the coating. Based onthe measured values of H 1 , H 2 and W,the dilution rates of three coatingswere calculated according to equations (1)- (3). As shown in table 1, the dilution ratehas no significant changes with the increase in x (70.9% for x=0, 72.4% for x=1.0 and 70.1% for x=2.0).
The dilution rate mainly depends on the absorption energy of the given substrate, which is closely related to the processing parameters and the pre-placed layer (surface roughness, thickness, compositions and phase constituents and so on). The chemical compositions can be regarded as a most important factor responsible for   . It is clear that the cladding material containing more Cr can absorb more laser energy, which contributesto the melting of the substrate. However, the absorbed energy for melting the cladding material containing more Crwill be enhanced correspondingly due to its higher melting point than the other elements (2178 K for Cr,1768 K for Co, 1726 K for Ni, 933 K for Al, 1809 K for Fe), which weakens the melting of the substrate. The two factors resulting from the addition of Cr will play the role against each other in the dilution rate. No significant change should result from the interplay between the two. Figure 3 shows the X-ray diffraction patterns of thecoatings with different x values (x=0, 1.0, 2.0). The diffraction patterns of Coatings I-III are very similar, in whicheight diffraction peaks are involved. Two strong diffraction peaks are located at 40.2°and 41.4°, besides which the other six weak peaks can also be detected. An interesting phenomenon is observed when x is increased from 0 to 1.0, namely, the intensity ratio between two peaks (40.2°and 41.4°) presents the increasing tendency. When x is further increased to 2.0, the majority of peaks completely disappear and only two peaks are reserved (a strong peak at 40.2°and a weak peaks at 76.2°). Those peaks were fittedby the HighScore Plus software and compared with d values in Joint Committee on Powder Diffraction Standards (JCPDS) cards. The results show that the coating is composed of a solid solution with the HCP structure (JCPDS card: 01-089-2762) and an intermetallic compound of Ti 2 Ni (JCPDS card:01-072-0442). The volume fraction of the two phases can be calculated from the XRD patterns by the Jade 5.0 software. The volume fraction of Ti 2 Ni and HCP is 65% and 35% in Coating I. Ti 2 Ni is gradually reduced to 44% and 7% when x is increased to 1 and 2, accompanied with the increase to 56% and 98% in volume fraction of HCP. That is to say, a HEA coating composed of the single solid solution was successfully synthesized when x reaches 2.0. For the HEAs, only the solution solid is expected to be synthesized in a multicomponent system. However, some unfavorable intermetallic compounds are also easy to be simultaneouslyformed in such a system. Therefore, it is very essential to establish the criterions for effectively predictingphase constituents of the HEAs, which will greatly simplify the design of the multicomponent system and reduce the processing costs. Some criterions in terms of the mixing enthalpy (ΔH mix ), themixing entropy (ΔS mix ) and the atomic size difference (δ) had been established to predict phase constituents. Zhang et al [39] proposed the criterions based on a large number of data obtained from numerous prepared HEAs, namely the simple solid solutions can be synthesized when they satisfy 12ΔS mix 17.5 J·K −1 ·mol −1 , −19ΔH mix 5 kJ·mol −1 , 0δ6.5. The similar criterions were also established by Guo et al [40], which are followed: The above parameters were defined as following [38]:

XRD analyses
where n means the number of components in multi-component alloy system,R denotes the gas constant (R=8.314 J·mol −1 ·K −1 ), C i and C j are the atomic percentage of the i and the j components in thesystem, W =4H i j , AB mix refers to the interaction parameter between the i and j components in the system (H AB mix signifies the mixing enthalpy of a binary liquid alloy system and can be tracked inLiterature [41]), representsthe atomic radius, and R i is the atomic radius of the i component.
It is clear that the calculations of the three parameters are closely related to the contents of different elements in HEAs. The components involved in the laser-clad HEAs coatings are derived from two parts, corresponding to the cladding materials (AlCoCr x NiFe) and the substrate (Ti6Al4V).A model had been established to calculate the contents of those components [38]. Based on that, the contents of all components in the three coatings were calculated (shown in table 2). The values are also measured by EDS (shown in table 3), which are approximately same to the calculated values. Then, ΔS mix , ΔH mix and δ were figured out by Eqs. (4)-(6) (shown in table 4). For the three coatings, the values of ΔS mix and δ can satisfy the criterions proposed by Guo et al However, the values of ΔH mix in Coatings I and II are deviated from the criterions except for that in Coating III. It can be conclude that the simple solid solutioncan only be synthesized in Coating III, which agrees well with the XRD results.
Besides the above-mentioned parameters, a new parameter (Ω) defined as the net driving factor was also put forward to further strengthening the prediction of phase constituents in HEAs. Ω can be calculated by the following equation [42]: ΔS mix and ΔH mix can beregarded as the driving force and the resistance force, respectively. When Ω exceeds 1, the solid solution is inclined to be synthesized owing to the driving force is larger than the resistance force, whereas the unfavorable intermetallic compound tends to be precipitated. As shown in table 4, the values of Ω exhibit the upward tendency with increasing the addition content of Cr in the cladding materials, resulting in the value is approximately equal to 1 in Coating III. The result further validates the conclusion that the solid solution can be formed more easily in Coating III. Figure 4 shows thebackscattered electron (BSE) images of the TiAlCoCr x FeNiHEAs coatings. As shown in figure 4(a 1 ), the island-like structure is uniformly distributed within the amorphous structure in the coating with x=0. When x is increased to 1.0 and 2.0, the amorphous structure is greatly reduced in volume fraction, accompanied with the increase in volume fraction of island-like structure ( figure 4(b 1 ) and (c 1 )). BSE images with a high magnification clearly present the microstructural evolution of the coatings with the change in x ( figure 4(a 2 , b 2 , c 2 )). For the amorphous structure, plenty of fine stripped/equiaxed particles are uniformly scattered within the amorphous matrix. EDS was applied to identify the chemical compositions of three phases  Combined with the XRD results, the island-like phase and the stripped/equiaxed phase in the amorphous structure can be confirmed as α(Ti) with a HCP structure, and the matrix in the amorphous structure is identified as the intermetallic compound of Ti 2 Ni. Obviously, CoatingI is composed of coarse primary α(Ti) grains and the eutectic structure(α(Ti)+Ti 2 Ni). Along the increase in x, the island-like grains become coarser and gradually swallow the amorphous structure, resulting in the amorphous structure mainly distributed along the boundaries of the former. When x reaches 2.0, the coating is almost filled with island-like phase, and the amorphous phase is hardly observed. α(Ti) with a HCP structure presents the increasing tendency in volume fraction together with the opposite change in Ti 2 Ni, and the single solid solution ofα(Ti) is obtained in the coating with x=2.0. Figure 5 shows the potentiodynamic anodic polarization curves recorded on the coatings and the substrate ina0.1 mol·L −1 HCl solution. Some electrochemical parameters were obtained from figure 5 (shown in table 6). The corrosion potential (E corr ) of the coatings is increased from −0.238 V (Coating I) to −0.192 V (Coating II), finally to 0.25 V (Coating III) with increasing the Cr addition content, which is higher than that of the substrate (−0.273 V). The corrosion potential characterizes the electrochemical dissolution tendency of a given material. When the other material is involved in the corrosion environment, the material with a high corrosion potential is more likely to be protected due to its strong ability to acquire electrons. Therefore, the corrosion tendency of the four samples is arranged in the following order: Coating III <Coating II <Coating I<Ti6Al4V. The corrosion current density (I corr ) represents the corrosion rate without the applied current. Clearly, there are no significant differences in corrosion rate for the four samples due to their similar I corr (about 1.5E-7 A·cm −2 ). The samples will step into the active state when the applied potential exceeds E corr , causing the rapid increase in current. When the potential reaches the passive potential (E p ), the samples are transformed from the active state into the passive state due to a thin and compact oxidation film formed on their surfaces. The transformation difficulty is closely associated withthe difference between E p and E corr . Obviously,it iseasier for the samples with  smaller difference value to enterthe passive zone. The difference value of Coating III is the smallest (about 0.059 V), followed by Coating II (0.078 V) andCoating I (0.083 V), which are lower than that of the substrate (0.186 V). The passive current density (I p ) is used to evaluate the stability of the passive film. Coating III exhibits a minimum passive current density of 3.69E-7A·cm −2 , followed by Coating II (5.99E-7 A·cm −2 ), Coating I (8.86E-7 A·cm −2 ). Compared with the coatings, I p of the substrate is increasedalmost one order of magnitude (2.48E-6 A·cm −2 ). It is clear that corrosion resistance of the coatings is significantly improved when compared with that of the substrate. As far as the coatings are concerned, their corrosion resistance is also gradually enhanced with the increase in addition content of Cr. The electrochemical impedance behaviors of the four samples were also investigated in a0.1 mol·L −1 HCl solution by the electrochemical impedance spectroscopy (EIS). Figure 6 indicates the EIS Nyquist curves. Each curve has similar profiles, but the radiiof the capacitance arc are different, which indicates that the corrosion mechanisms of the samples are similar and controlled by the charge transfer resistance. The equivalent circuit model was established by the Zview software and the values of different components in this model are presented in table 7. R s represents the solution resistance between counter electrode and workings electrode, which has no significant changes (about 23.25-25.45 Ω·cm 2 ) in this tests. Q d denotes theinterfacial capacitance, which is used to characterize the degree of the charge concentration around the electrode surface [43]. Among all samples, the substrate demonstrates the highest Q d value, indicating thatthe passive film formed on it is the most unstable and easy to be corroded.n signifiesthe deviation degree ofadouble layer capacitor from an ideal capacitor, which mainly depends on the surface state (roughness, inhibitor adsorption, etc) of the samples [44]. The surfaces of the passive film adhering to all samples' surfaces are similar due to their similar n values.R ct refers tothe chargetransfer resistance, which is regarded asanessential parameters evaluating the resistance of reactions occurring on the samples surface [45,46]. It is clear that the R ct value of Coating III (1.20E6 Ω·cm 2 ) is evidently higher than those obtained at the other samples (9.80E5 Ω·cm 2 for Coating II, 8.56E5 Ω·cm 2 for Coating I, and 9.52E4 Ω·cm 2 for the substrate). Corrosion resistance of the samples is followed:Coating III >Coating II >Coating I>the substrate, which is well in consistent with that acquired in the potentiodynamic anodic polarization tests.

Corrosion behaviors 3.2.1. Electrochemical tests
Corrosion behaviors are mainly related to the inner microstructure with respect to phase constituents, chemical compositions and grain size inside the coatings. Dai et al [47] compared corrosion resistance of asprepared Ti6Al4V by selective laser melting (SLM) with that of commercial Grade 5 sample. The results indicated that the microstructural evolution caused the difference in corrosion resistance of two alloys. The SLM   Ni). Moreover, the total content of Cr involved in the coatings is gradually increased with the increase in x. Cr had been proved to play the positive role in improvement in corrosion resistance of the alloys [38]. Therefore, the evolution in phase constituent and corresponding chemical composition with x can be regarded as an important factor causing the increase in corrosion resistance of the coatings. Besides them, it can be observed that the grain size of HCP is gradually increased with the increase in x. The number in grain boundary is correspondingly reduced. The grain boundary possesses the higher energy than the grain, which indicates that the former is easier to be dissolved than the latter. Therefore, the change in grain size also makes a partial contribution to the improvement in corrosion resistance of the coatings. On the other hand, corrosion behaviors are associated with the passive film formed on the sample. The compositions and chemical valence states of the passive film formed on Coating III were analyzed by XPS. Figure 7 shows the XPS survey and narrow spectra of metal elements (Ti 2p , Ni 2p , Al 2p , Cr 2p , Co 2p and Fe 2p ) involved in the passive film formed on Coating III after the EIS test in a0.1 mol·L −1 HCl solution. The passive film is mainly composed of Ti, Ni, Al, Cr, Co, Fe, and O elements ( figure 7 (a)), implying that some of them have been transformed into oxides during corrosion. For these metal elements, the chemical valence states were further confirmed by high resolution narrow spectra(figures 7(b)-(g)). Figure 7 figure 7(d)). As shown inFig.7(e), the standard values of the metallic state Cr (Cr 2p3/2 at 573.8 eV and Cr 2p1/2 at 583.5 eV) and Cr 2 O 3 (Cr 2p3/2 at 575.9 eV and Cr 2p1/2 at 586 eV) can be detected. Similarly, figures 7(f) and (g) also demonstrate that Co and its oxide (CoO), Fe and its oxide (Fe 2 O 3 ) are all involved in the passive film. It is worth noting that Ti and Al are almost completely transformed into theoxides, and a portion of Co, Fe and Cr are still remained besides their oxides. What is beyond ourexpectation is that Ni doesn't suffer from the oxidization.
These oxides related to the above-mentioned metallic elements can beformed as follows: 2Cr O 12 The formation possibility and difficulty of above oxides can be estimated by the thermodynamics calculation. Based on the data provided in [49], the changes in standard Gibbs free energy (ΔG θ ) of the abovementioned reactions with the temperature are clearly demonstrated in figure 8. The ΔG θ values of all reactions are all negative in the temperature range from room temperature to 1800 K, indicating that those metal oxides can be formed spontaneously. The ΔG θ value related to Reaction (13) is more negative than those of the other reactions, indicating that A1 2 O 3 will be formed preferentially in thermodynamics. On the contrary, Ni is very difficult to be oxidized to NiO due to the more positive ΔG θ value in Reaction (8). The formation difficulty for Cr 2 O 3 and Fe 2 O 3 is between NiO and Al 2 O 3 . The calculated results are well in consistent with those obtained in XPS. However, all Ti elements are oxidized to TiO 2 , contrary to the thermodynamics prediction. The abnormal phenomenon is mainly associated with the dynamics. The thermodynamics mainly clarifies the possibility of a reaction, in which no reaction rate relatedto the dynamics is involved. Although Reaction (10) is difficult to occur in thermodynamics, it may proceed at a higher rate. As a result, no titanium is maintained in the passive film. Reaction (12) is very easy to proceed since its ΔG θ value is only less than that ofReaction (13).This implying that Cr 2 O 3 can be formed preferentially in the passive film. The addition of Cr will contribute to the improvement in corrosion resistance of the coatings due to a dense layer of Cr 2 O 3 with a higher resistance of approximately 10 Ω·cm rapidly formed on the coating. Therefore, Coating III with a highest Cr content can be quickly transformed from the active state into the passive state.
The stability of the passive film is responsible for the passive current density in the passive state. The dissolution process of those metallicoxides can be expressed as follows:  Similarly, thermodynamics calculations were also applied to predict the possibility and difficulty of above reactions. As shown in figure 9, besides TiO 2 and Al 2 O 3 , Cr 2 O 3 is also very difficult to be dissolved in the HCl solution due to the ΔG θ values of Reaction (17), (18) and (19) exceeding zero. It can be concluded that the lowest passive current density of Coating III is closely related to its high Cr content.

Immersion tests
To further verify the effect of Cr on corrosion resistance of the laser-clad coatings, immersion tests were carried out in a0.1 mol·L −1 HCl solution for 30 daysat room temperature (25°C). The surface morphologies of the coatings subjected to the immersion tests were observed by OM ( figure 10). There areno large corrosion pits for Coatings I-III after immersion for 30 days. A close inspection reveals that Coating I suffers from the worst  intergranular corrosion when compared with the other coatings due to a large number of dendrites emerging from the initial polished surface. For Coating II, the dendrites can be clearly observed at local zones.Its surface is very smooth and the dendrites are hardly visible when x is increased to 2.0, implying that the coating demonstrates excellent corrosion resistance. This result is also consistent with the electrochemical tests. Figure 11 shows the microhardness distribution across the coating's cross sections. When x=0, the average microhardness of Coating I is about 656 HV 0.2 . With the increase in content of Cr, the average microhardness of Coating II and Coating III is enhanced to 763 and 800 HV 0.2 . The microhardness of Coating III is improved  about twice as much as that of the Ti6Al4V substrate (350 HV 0.2 ). Obviously, the addition of Cr is beneficial to the improvement in hardness of the coatings.

Microhardness
The improvement in hardness of those coatings is mainly attributed to the solid solution strengthening from HCP and dispersion strengthening from Ti 2 Ni. Ti 2 Ni had been proved to possess a comparatively low hardness of about 616.9 HV [50], which is less than the averagevalueof those coatings. Therefore, the solid solution strengthening from HCP plays a predominant rolein the improvement in hardness of the coatings, which mainly depends on the difference in atomic radius between solvent and solutes, and volume fraction of the solutes. The atomic radii of all elements involved in the coatings are 1.460 Å(Ti), 1.432 Å(Al), 1.310 Å(V), 1.241 Å(Fe), 1.248 Å(Co), 1.246 Å(Ni), 1.249 Å(Cr). The atomic content of V is approximately same in the three coatings, implying that the solid solution strengthening effect of V is roughly the same. The atomic content of Al presents the decline tendency with the increase in x. However, the change in volume fraction of Al does not produce the significant effect on the solid solution strengthening owing to its slight difference from Ti (about 2%in atomic radius). Fe, Co, Ni and Cr possess the approximately same atomic radii of about 1.246 Å (about 15% difference from that of Ti), whose total atomic content is gradually increased with increasing x (25.74 at.%, 28.16 at.% and 28.55 at.% forCoatings I, II and III), accompanied with the decrease in atomic content of Al. It can be concluded that the solid solution strengthening is gradually improved by adding more Cr into the cladding material. This agrees well with the above testing result. Figure 12 showsthe wear profiles of the substrate andthe coatings with different contents of Cr after the sliding tests. The substrate demonstrates a largest wearvolume of about 1.419 mm 3 . For Coatings I and II, their wear volumes are similar (0.640 mm 3 ), decreased obviously by 54.8%compared with that of the substrate. When x is increased to 2.0, the wear volume of the coating is further reduced to 0.562 mm 3 , and about 80% of that obtained in Coatings I and II. Therefore, the increase in contentof Cr in the cladding material can weaken the wear volume.

Wear resistance
In order to reveal the mechanism of the samples, SEM was applied to characterize the surface morphologies of the substrate and CoatingIII. As shown in figure 13(a 1 ) and (b 1 ), the wear track presents an oval shape. However, whether the width or the length in the wear track of CoatingIII are all significantly smaller than those obtained in the substrate, which further confirms that CoatingIII exhibits more excellent wear resistance than the substrate. High-magnification SEM images clearly presents significant differences between their morphologies. A large number of furrows coupled with severe plastic deformation are distributed throughout the whole wear surface of the substrate, indicating that the substrate is subject to severe microcutting. For CoatingIII, its surface is comparatively smooth and no obvious furrows are observed. A portion of zones are covered with a thin layer of products. Other than those, many sheets loosely adhere to the wear surface. EDS results indicate that Area 1 is rich in O (68.08 at.%) and Area 2 is mainly composed of Ti and Al (shown in table 8). It can be concluded that the surface suffers from oxidation during sliding. The sheets can be identified as the debris from the coating, which peelsoff from the surface and is oxidized, and finally rolled into the sheets. The oxidation layer and flaky debris will protect the coating from microcutting to a certain extent. Wear loss of the substrate and CoatingIII mainly results from the microcutting effect. Wear loss of the coatings presents the slight difference, which should be mainly attributed to the change in phase constituent of the coatings with x. Besides HCP, the comparative content of Ti 2 Ni (65 vol.%)is involved in the coating when x is zero. Along with the increase in x, the content of Ti 2 Ni was gradually reduced to 44vol.% (x=1.0) and 7 vol.% (x=2.0). As far as the two phases are concerned, Ti 2 Ni possesses the weaker resistance to plastic deformation than HCP due to its lower hardness than that of HCP. Moreover, Ti 2 Ni with a complex cubic structure also has more slip systems than HCP with a close-packed hexagonal structure, which implies the former possesses better ductility than the latter. When the friction pairs contact at an applied load and do a relative motion, the sharp protrusions on YG6 ceramic balls are easier to penetrate into the soft Ti 2 Ni with weak resistance to plastic deformation, and producemore serious microcutting. Therefore, wear loss of the coatings presents the downward tendency with the increase in x due to the reduction in content of Ti 2 Ni. Accompanied with the cyclic reciprocating, stripped debris gradually peels off from the surface, leaving many furrows. The oxidation layer and the flaky debris can also be formed on the substrate surface, but cannot exist steadily since the soft substrate is unable to provide the enough structural support for them. On the contrary, they can adhere to the surface of the coatings with a high hardness, resulting in the significant reduction in wear loss.

Conclusions
(1) TiAlCoCr x FeNi HEAs coatings were prepared on surface of Ti6Al4V by laser cladding.The volume fraction of Ti 2 Ni was decreased accompanied with the increase in volume fraction of HCP along the increase in x (0, 1.0, 2.0). When x=2.0, the coating was only composed of the single solid solution of HCP. The evolution in phase constituent was closely associated with the change in Ω.Ωwas close to 1when x reached 2.0, implying that the unfavorable intermetallic compound was avoided due to the high resistance force in the system.
(2) The coating with x=2.0 exhibited the optimum corrosion resistance in electrochemical and immersion tests carrying out ina0.1 mol·L −1 HClsolution, which demonstrated a stronger corrosion tendency, a higher  passivation tendency and a lower active dissolution rate when compared with the substrate and the other coatings (x=0, 1.0).It should be attributed to the rapid formation of oxides containing Cr, which shielded the coating from further dissolution. Moreover, the oxides containing wit a high stability also made a partial contribution to the improvement in corrosion resistance of the coating.
(3) The hardness of the substrate was significantlyimproved by preparing the TiAlCoCr x FeNiHE as coatings. A high Cr content contributed to the improvement in hardness of the coatings (enhanced about 20% with x increased from 0 to 2.0). Correspondingly, wear resistance of the coatings was improved than that of the substrate, and their wear resistance presented an increasing tendency with increasing x values.The change mainly resulted from the evolution in phase constituent of the coatings. The content of HCP with a high hardness and a low ductility was gradually increased, accompanied with the decrease in content of Ti 2 Ni with a low hardness and a high ductility. This improved the resistance to microcutting, resulting in the decrease in wear loss.