Comparative Investigation on Wear Properties of Composite Coatings with Varying CeO₂ Contents

Abstract

Several innovative mixed powders of Ti6Al4V and NiCr-Cr₃C₂ with different CeO₂ contents (0, 1, 2, 3, and 4 wt.%) were designed, and Ti₂C-reinforced CrTi₄-based composite coatings were prepared on the Ti6Al4V surface via laser cladding technology. The effects of CeO₂ amount on the forming quality, microstructure, hardness, and wear resistance of the composite coatings were studied. The results showed that the CeO₂ amount had a significant influence on the forming quality of the composite coatings. The cracks were eliminated completely when the CeO₂ content was 2 wt.%; furthermore, the lowest porosity was obtained with the addition of 3 wt.% CeO₂. The primary phase constituents of the coatings were non-stoichiometric Ti₂C and a β-type solid solution (CrTi₄) as the reinforcement and matrix, respectively. CeO₂ and a low quantity of Ce₂O₃ were re-precipitated at the Ti₂C/CrTi₄ interface and CrTi₄ grain boundary in the coatings with CeO₂ addition. In addition, the average hardness of the composite coatings was 1.28–1.34 times higher than that of the Ti6Al4V substrate. The wear resistance of the composite coatings was significantly higher than that of the substrate. However, both the composite coatings and the Ti6Al4V substrate exhibited a mixed-wear mode, i.e., abrasive and adhesive wear.

1. Introduction

The dual-phase (α + β) Ti alloy Ti6Al4V is extensively used in aerospace components, oceanographic engineering, and biomedical engineering [1,2]. However, the applications of this alloy in tribology-related engineering components are limited owing to its low hardness and wear resistance [3]. Laser surface modification can increase the surface hardness and wear resistance of Ti6Al4V while retaining its desirable bulk characteristics, and has thus attracted considerable attention in recent years [4]. Furthermore, the use of a ceramic material with high elasticity and hardness as the reinforcement phase increases the wear resistance of the matrix phase. Metallic carbides, borides, and nitrides, such as TiC, WC, TiB/TiB₂, and TiN, are commonly used in the synthesis of ceramic-reinforced composite coatings [5,6,7,8]. Titanium carbide is among the most extensively used reinforcement materials for Ti6Al4V owing to its high elastic modulus, hardness, and wettability; furthermore, its thermal expansion coefficient is similar to that of Ti6Al4V [9]. Wang et al. [10] found that in situ synthesis could alleviate the cracking and spalling of the coating induced by the direct addition of a ceramic reinforcing phase. Qu et al. [11] demonstrated that the forming quality and microstructure of the cladding coating were strongly dependent on the thickness of the preplaced powder; however, this influence was difficult to control. The in situ synthesis of a titanium carbide-reinforced phase via synchronous powder feeding is therefore considered the optimal technique for the development of a titanium carbide-reinforced composite coating.

The metal ceramic NiCr-Cr₃C₂ is commonly used in powdered form during thermal spraying; however, the coating fabricated via spraying comprises multiple microdefects, including pores and cracks, while the bond between the coating and substrate is weak [12,13]. Research concerning the laser cladding of NiCr-Cr₃C₂ powder on Ti-alloy surfaces is limited. Fan et al. [14] prepared TiC- and Cr₇C₃-reinforced laser-clad coatings on the surface of Ti6Al4V via preplacement of NiCr-Cr₃C₂ powder. Aghili et al. [15] observed numerous pores within NiCr-Cr₃C₂ coatings on the surface of a titanium aluminide substrate prepared via laser cladding using a coaxial power supply.

The addition of an optimal quantity of rare-earth oxides, such as Y₂O₃, CeO₂, and La₂O₃, prevents the formation of defects, including cracks, pores, and inclusions, during laser cladding [16]. Zhang et al. [17] reported that the addition of Y₂O₃ powder resulted in the complete elimination of pores in laser-clad Ti-6Al-4V coatings. The addition of rare-earth oxides creates an effective balance between the hardness and toughness and prevents crack formation within the coatings, thereby optimizing the tribomechanical properties [18]. Zhang et al. [19] demonstrated that CeO₂ showed greater efficacy in preventing microcracking and increasing hardness than Y₂O₃ and La₂O₃. CeO₂ was therefore selected to increase the forming quality and wear resistance of the coatings in the present study. The influence of the CeO₂ content on the wear resistance of a Ti6Al4V/NiCr-Cr₃C₂ coating currently remains unexplored, and will be examined in the present study.

The previous study [20] demonstrated the effect of the CeO₂ contents on the forming qualities and microstructure of titanium carbide-reinforced Ti-based composite coatings on the surface of Ti6Al4V using multitrack overlapping laser cladding in detail. The forming qualities of the composite coatings, including crack, porosity, and dilution rate, were firstly evaluated by penetrant inspection and optical microscopy (OM). Then, the phase composition, microstructure, and element distribution were characterized by X-ray diffraction (XRD), scanning electron microscopy (SEM), energy-dispersive spectroscopy (EDS), backscattered electron (BSE) images, electron probe microanalysis (EPMA), transmission electron microscopy (TEM), and electron backscattered diffraction (EBSD). Finally, the present study employed Vickers hardness tests and wear testing to characterize the hardness and wear resistance of the Ti-based composite coatings for Ti6Al4V. The wear mechanism was revealed by analyzing the wear morphologies and composition of the wear debris using SEM, EDS, and BSE.

2. Materials and Methods

2.1. Materials and Laser Cladding

An annealed α/β dual-phase Ti alloy, Ti6Al4V (executive standard: GB/T3621; Chinalco Shenyang Nonferrous Metal Processing Co. Ltd., Shenyang, China), was used as the substrate, as shown in Figure 1. The substrate plate was cut into dimensions of 12 mm × 10 mm × 10 mm via electro-discharge machining for laser cladding.

Five groups of Ti6Al4V/NiCr-Cr₃C₂/CeO₂ composite powders were designed to investigate the influence of the CeO₂ content of the composite powders on the tribological performance of the coatings. The CeO₂ content of each of the five groups was 0, 1, 2, 3, and 4 wt.%, while all groups contained 25 wt.% NiCr-Cr₃C₂; furthermore, the Ti6Al4V content was balanced. Mechanical mixing of the powders was performed via planetary ball milling (MSK-SFM-1; MTI Corporation, Richmond, CA, USA) for 1 h prior to the cladding process. The mixed powders were dried in an oven at 70 °C for at least 2 h, until cladding, to increase their fluidity.

The composite coatings were prepared using a modular laser processing system (Trulaser cell 7040, Trumpf China Co. Ltd., Nanjing, China) equipped with a fiber optic laser (TruDisk 4002, TRUMPF). The mixed powders were supplied using a scraper-spray-type powder feeder at flow rates dependent on the rotational speed of the powder supply channels. A stream of inert He gas was used to continuously feed the powder particles into a coaxial deposition nozzle that generated a powder particle stream flowing into the spot of the focused laser beam. The nozzle exhibited a concentric alignment with the laser beam (diameter: 3 mm). The cladding process was performed under an atmosphere of Ar to prevent any reaction between the liquid cladding materials and atmospheric contaminants such as O₂ or H₂O. The optimal process parameters were as follows: power, 1100 W; scanning speed, 0.4 m/min; powder feeding, 4–5 g/min; overlapping rate, 50%; feeding gas flow rate, 7 L/min; protective gas flow rate, 8.0–15 L/min.

2.2. Hardness and Wear Resistance Testing

The porosity and dilution rate were investigated via OM (Olympus Corporation, Tokyo, Japan) in conjunction with image analysis software, Image-Pro Plus 7.0 (IPP; Media Cybernetics, Inc., Rockville, MD, USA). The microstructures and chemical compositions were characterized via SEM (S-3000, Hitachi High-Technologies Corporation, Tokyo, Japan) in conjunction with EDS (INCAPentaFET-x3, Oxford Instruments plc, Abingdon, UK). The small precipitates in the coatings were identified using TEM (Tecnai G2 F30, FEI Company (Thermo Fisher Scientific), Hillsboro, OG, USA) in conjunction with EDS. The microhardness was measured along the depth direction of the coatings using a Vickers microhardness tester (KB 750-MHSR, KB Prüftechnik GmbH, Rhineland-Palatinate, Germany) with a load of 0.5 kgf and a dwell time of 10 s. Each measurement was repeated at least 5 times, and the average value was calculated. The dry-sliding friction test was performed using a block-on-ring friction and wear testing machine (M-2000Z, Jinan Fangyuan Testing Instrument Co., Ltd., Jinan, China) at ambient temperature (25 °C) in air for 50 min, with an applied load and sliding speed of 100 N and 200 rpm, respectively. A quenched bearing steel, GCr15, with a hardness of approximately 735.1 HV_0.5, was used as the counterpart and replaced after each sliding friction test. The weight of the samples was measured before and after each test using a high-precision electronic balance. Each test was repeated at least 3 times, and the average value was calculated. The surface morphologies were observed after wear testing using SEM and BSE, and the chemical composition of the wear debris was analyzed using EDS.

3. Results and Discussion

3.1. Microstructure

Ti6Al4V/NiCr-Cr₃C₂ composite coatings with different CeO₂ contents were prepared on the Ti6Al4V substrate by laser cladding with coaxial powder feeding. The surface, as well as the cross-sectional SEM macrographs of the various coatings, is presented in Figure 2. With the increase in CeO₂, the cracks on the coating surface displayed the tendency to decline at the beginning and rise in later stages. There were no cracks on the surface when the addition of CeO₂ was 2 wt.%.

The addition of CeO₂ exerted a significant impact on the cracking behavior, porosity, and dilution rate of the composite coatings. The dilution rates and the porosity of the five coatings revealed by the metallographic statistics are listed in Figure 3. It can be seen that the porosity decreased first and then increased with an increase in the CeO₂ content (0 → 4 wt.%), and the minimum porosity was observed in the coating with 3 wt.% CeO₂. In addition, the pores were mainly distributed near the transitional zone, between the former and latter cladding tracks, and exhibited an obvious aggregation phenomenon, as indicated by the red line in Figure 2. The CeO₂ addition can lower the coating porosity, which is due to the following factors. (a) The active Ce atom tends to combine with the O atom to generate rare oxides, which contributes to reducing gases in the molten pool. (b) The CeO₂ increases the absorption rate of laser energy, and further increases the peak temperature of the molten pool. The molten pool can maintain liquid for longer periods, and thus the bubbles have more time for escaping. CeO₂ can improve the laser energy absorption mainly because of its unique electronic structure. Generally, rare earth elements (such as Ce, Y, and La) have abundant electronic energy levels and diverse spectra, which can increase the absorption rate of laser light through electronic transitions. (c) CeO₂ promotes more sufficient convection in the molten pool, which can restrain the formation of pores effectively. On the contrary, the dilution rate was positively correlated with the amount of CeO₂. However, the fluidity of the molten pool was attenuated by the presence of superfluous CeO₂ (4 wt.%). This resulted in a decrease in the dilution rate.

Figure 4 presents the XRD patterns of the laser-clad coatings. All the coatings present a similar XRD peak position. Therefore, the CeO₂ addition had an insignificant influence on the phase type of the composite coating. The primary phases in the coatings were face-centered cubic (FCC) titanium carbide and BCC CrTi₄. The BSE images for the cladding zone of the composite coatings are presented in Figure 5. The EDS results of the matrix in the coatings are shown in

Table 1. The EDS results of matrix in the coatings with different CeO₂ contents.

CeO2 Content	Ti/(at.%)	C/(at.%)	Al/(at.%)	V/(at.%)	Cr/(at.%)	Ni/(at.%)
0 wt.%	74.33	3.82	8.86	2.05	8.53	2.41
	74.22	4.29	9.01	2.19	8.39	1.90
1 wt.%	73.20	3.89	8.80	2.13	9.55	2.43
	73.13	4.22	9.32	2.14	9.22	1.97
2 wt.%	73.7	4.04	8.82	2.10	9.04	2.30
	73.3	4.08	8.95	1.97	9.55	2.15
3 wt.%	73.16	5.66	8.98	2.09	8.22	1.89
	72.90	6.25	8.83	2.02	7.90	2.10
4 wt.%	73.07	2.71	9.37	2.20	10.28	2.37
	72.90	3.72	9.10	2.09	10.01	2.18

. The microstructure revealed the development of numerous dendrites without a specific orientation. Combined with the results of XRD, EDS, and TEM, the matrix phase was determined to be CrTi₄, while the dendrites were determined to be non-stoichiometric titanium carbide (TiC_x). The dendrites of TiC_x in the coating were subjected to statistical analysis through the metallographic method. The results revealed that the contents of TiC_x in the coatings with 0, 1, 2, 3, and 4 wt.% CeO₂ were 55, 45, 40, 35, and 41%, respectively. An increase in the addition amount of CeO₂ induced an initial increase and a subsequent decrease in the content of titanium carbide. This variation trend was opposite to that observed for the change in the dilution rate.

There was a positive correlation between the proportion of the white particles and the CeO₂ content, as shown in Figure 5. Namely, the more CeO₂ was added to the coating, the more particles were precipitated. It was inferred that these highlighted particles were composed of Ce oxides. The EDS analysis revealed that these particles were mainly enriched in Ce and O elements. The element distribution of the composite coatings with 2 wt.% CeO₂ also showed that Ce was mainly present in the form of oxides in the coating, which supported the previous EDS results, as shown in Figure 6. Moreover, the result of TEM further verified that these Ce oxides were CeO₂ and Ce₂O₃ (Figure 7). In addition, it also can be observed from Figure 6 that the dendritic phases were mainly enriched in Ti and C elements, and the alloying elements Al, V, Ni, and Cr as the solute were mainly retained in the matrix (CrTi₄). However, the Al and V were uniformly distributed in the matrix phase, while Ni and Cr showed obvious segregation, and Ni presented a more serious level.

Furthermore, the addition of CeO₂ exerted no discernible effect on the phase composition of the composite coatings, except for the precipitation of new Ce oxides. The added CeO₂ powders initially decomposed and subsequently precipitated as CeO₂ and Ce₂O₃ during laser cladding (Figure 7). Figure 8 presents the phase map of the composite coating with 2 wt.% CeO₂. The matrix phase (red area) was identified as β-Ti, and CrTi₄ was confirmed as a solid solution phase rich in Cr and Ti and whose crystal structure was identical to that of β-Ti. The dendritic and granular non-stoichiometric titanium carbide (TiC_x) was confirmed as Ti₂C (green phase), which was consistent with the C/Ti atomic ratio obtained via EDS analysis. The rare-earth oxides (blue phase) were identified as primarily CeO₂, and they were mainly distributed at the Ti₂C/CrTi₄ interface and CrTi₄ grain boundary. The influence of CeO₂ addition on the average hardness of the composite coatings resulted from two competitive factors. (a) The re-precipitated Ce oxides can hinder the growth of grains, and thus refine grains. Generally, the finer the Ti₂C dendrite and CrTi₄ grain were, the higher hardness the composite coating exhibited. (b) However, the hardness of the Ce oxide was lower than that of titanium carbide (reinforcement phase). In addition, the amount of the re-precipitated Ce oxide was relatively low. Thus, there was little effect on the increase in hardness if only considering the mechanical performance of the Ce oxide. Therefore, based on the above two aspects, the re-precipitated Ce oxide played a positive role in improving the hardness.

3.2. Hardness and Tribological Performance

3.2.1. Microhardness

The average hardness values of the substrate and various composite coatings are presented in Figure 9. The hardness of the substrate was 363.2 ± 12.8 HV_0.5. The coatings with 0, 1, 2, 3, and 4 wt.% added CeO₂ exhibited hardness of 488.2 ± 6.2 HV_0.5, 472.4 ± 23.9 HV_0.5, 464.2 ± 3.0 HV_0.5, 487.2 ± 5.7 HV_0.5, and 472.9 ± 9.6 HV_0.5, respectively. All of the coatings exhibited hardness significantly higher (1.28–1.34 times) than that of the substrate. The increase in the hardness of the coatings relative to that of the substrate was attributed to several factors.

First, the presence of the titanium carbide (TiC_x) reinforcement phase, with extremely high hardness of 29–34 GPa, contributed to the high hardness of the coatings [21]. The hardness of TiC_x increased with x, i.e., the C/Ti stoichiometric atomic ratio [22]. Similarly, the hardness of the coating increased with the quantity of the reinforced phase. The content of dendritic TiC_x decreased with an increase in the CeO₂ content. A maximum dendritic TiC_x content of 55% was observed in the coating without added CeO₂. The dendritic TiC_x content decreased to 35% with the addition of 3 wt.% CeO₂ and increased upon the addition of further CeO₂.

Second, the dissolution of numerous alloying elements, including Ni, Cr, Al, and V, in the matrix of the composite coatings induced a certain degree of lattice distortion, thereby increasing the strength of the coatings by solution strengthening. Ni, Cr, and V elements all have a similar atomic size, crystal structure, electronegativity, and valences to Ti elements; therefore, Ni, Cr, and V can be solvable in Ti. Compared with the Ti6Al4V substrate, which only included Al and V alloy elements in addition to Ti elements (as shown in Figure 1), a large number of Ni, Cr atoms were solidly dissolved in the matrix phase β-Ti of composite coatings when 25 wt.% NiCr-Cr₃C₂ was added, and this led to the lattice distortion of β-Ti and the increase in the strength as well as the hardness of the matrix. Furthermore, the segregation of C in the dendrite, as well as Ni and Cr in the matrix, induced significant distortion, thereby increasing the hardness of the coatings.

Third, the presence of Ce oxides at the matrix/TiC_x phase boundary inhibited grain growth. Generally, a decrease in the grain size results in an increase in the number of grain boundaries per unit volume. Consequently, dislocation movement is inhibited, and this results in higher hardness.

Finally, the dilution rate exerted a significant impact on the hardness of the coatings. Strong adhesion between the coating and substrate was achieved with optimal dilution, whereas excessive dilution weakened the strengthening effect of the reinforced phase on the coating. An increase in CeO₂ induced an initial increase and a subsequent reduction in the dilution rate. CeO₂ content of 2 and 3 wt.% resulted in the highest dilution rate, while the coating with 3 wt.% CeO₂ exhibited the highest hardness, which was mainly due to the more prominent grain refinement within the coating. The hardness of the coatings was reduced by CeO₂ contents over 3 wt.% because excess CeO₂ formed impurities in the coating, thereby reducing the hardness [23].

In addition, there was an unexpected phenomenon in which the composite coating with 3 wt.% CeO₂ addition had the lowest titanium carbide content, but it also had the highest average hardness. This is mainly due to the following factors. (a) Compared with the coating without CeO₂, more Al, V, Ni, and Cr atoms were solid-dissolved in the matrix phase, so the solution strengthening the coating with 3 wt.% CeO₂ had a stronger effect. (b) The re-precipitated Ce oxides were distributed at the Ti₂C/CrTi₄ interface and CrTi₄ grain boundary of the coating with CeO₂. The Ce oxides can hinder the growth of grain, and thus refine grains. Generally, the finer the Ti₂C dendrite and CrTi₄ grain were, the higher hardness the composite coating exhibited. (c) Compared with the coating without CeO₂, there were fewer cladding defects (such as pores) within the coating with CeO₂. In general, the cladding defects can decrease the coating hardness. Therefore, although the TiC_x content of the coating with 3 wt.% CeO₂ was only 35%, which was obviously lower than that of the coating without CeO₂ (55%), its hardness was only slightly lower than the coating without CeO₂.

3.2.2. Friction Coefficient

Figure 10 presents the friction coefficients of the substrate and the cladding coatings with different CeO₂ contents. The friction coefficient of the substrate (0.359) was lower than those of the cladding coatings, which ranged from 0.508 to 0.587, and this was attributed to several factors. First, Ti6Al4V exhibited low hardness and was susceptible to plastic deformation while grinding against GCr15. Second, the formation of an oxide film on the substrate during wear was facilitated by its high chemical activity and strong affinity for O. Sumitomo et al. [24] reported that Ti oxides are easily sheared and deformed, and their presence reduced the friction coefficient of a tribo-film. Third, the composite coating comprised a substantial quantity of protruding titanium carbide, the hardness of which is significantly higher than that of GCr15. These hard protuberances were embedded in the GCr15 surface, thereby generating high friction resistance. Furthermore, the hard titanium carbide destroyed the stable and continuous lubricating film on the wear surface, which also increased the friction coefficient of the coatings.

The substrate exhibited a stable friction coefficient that rapidly increased after approximately 1500 s, indicating a change in the characteristics of the oxide film on the Ti6Al4V surface. An O-rich α layer was formed owing to the strong affinity of Ti for O at elevated temperatures in air [25]. Deshmukh et al. [26] reported that the presence of the hard and brittle α layer induced a possible deterioration in the mechanical properties of the Ti alloy. Therefore, substantial heat of friction was considered to accumulate between the friction pairs with continuous wear, which gradually transformed the TiO_x film into the α layer, resulting in an increase in the friction coefficient of the substrate.

Figure 10a reveals a simultaneous, slight reduction in the friction coefficients of the composite coatings after approximately 1800 s, except for the coating with 1 wt.% CeO₂. It may be the consequence of the coaction of many factors. (a) Compared with the coating with 1 wt.% CeO₂, the more re-precipitated Ce oxide in the other coatings exerted a significant grain refinement effect and coarse titanium carbide dendrites were significantly reduced. When the friction and wear experiment was carried out for a long time, the smaller abrasive particles in the friction interface were discharged, resulting in the weakening of the ploughing effect of abrasive particles on the surface of the coating, so the friction coefficient decreased. (b) The hardness of each of the composite coatings was significantly higher than that of the substrate, resulting in substantial heat of friction in the coatings. A large amount of the matrix phase (CrTi₄) and a small amount of titanium carbide were oxidized to titanium oxides (TiO_x) under the high-temperature and oxygen-rich conditions, thereby reducing the friction coefficient [27]. (c) The generated heat of friction softened the matrix, increasing the adhesion between the coating and GCr15 and thereby increasing the friction coefficient.

For coatings with 2 wt.%, 3 wt.%, and 4 wt.% CeO₂, the effect of factors (a) and (b) was significantly greater than that of factor (c); therefore, the friction coefficients of these composite coatings showed an obvious downward trend in the later stages of the friction and wear experiment. However, for the coating with 1 wt.% CeO₂, the factors contributing to the increase or decrease in the friction coefficient were in a process of “dynamic balance”, so the friction coefficient changed smoothly with time. On the contrary, for the substrate and the coating without CeO₂, factor (c) was dominant in the coating without CeO₂; therefore, the friction coefficient of this coating tended to consistently increase.

3.2.3. Wear Morphology

The wear morphologies on the surfaces of the substrate and composite coatings with various CeO₂ contents are presented in Figure 11. Deep furrows were formed on the wear surface of Ti6Al4V, which exhibited severe plastic and extrusion deformation (Figure 11a). These observations corresponded to the abrasive and adhesive wear failure modes, which were attributed to several factors. First, the surface of Ti6Al4V was subjected to the combined action of normal and tangential force while sliding against GCr15. The micro-protrusions on the GCr15 surface were easily pressed into and cut the substrate owing to the low hardness of Ti6Al4V. Numerous deep, wide furrows were formed parallel to the direction of sliding by continuous wear on the surface of the substrate. Second, the soft Ti6Al4V adhered to the surface of GCr15 during relative sliding. When the bonding strength between the sticky nodes exceeded the shear strength of the substrate, the adhesive part was torn from the substrate. Third, localized deformation occurred in the wear area during periodic sliding.

Conversely, the composite coatings exhibited extremely rough surfaces, and the existing furrows exhibited low depths and widths owing to a combination of factors. First, the hardness of the TiC_x-reinforced composite coatings was higher than that of the substrate, making the coatings highly resistant to ploughing by the micro-protrusions on the GCr15 surface. Second, the uniformly distributed dendritic and acicular TiC_x in the composite coatings facilitated load transfer to the ductile matrix phase and released the stress via coordinating deformation, thus avoiding the brittle fracture of the hard and brittle TiC_x. Third, the solid solution of Ni, Cr, and other alloying elements in the matrix phase induced lattice distortion, thereby increasing the strength and hardness of the coating and increasing its deformation resistance. Finally, the rare-earth oxides (CeO₂/Ce₂O₃) were distributed at both the phase and grain boundaries between TiC_x and the matrix, which inhibited the motion of grain boundaries, resulted in grain refinement, and hindered deformation, thus improving the ploughing resistance of the coatings [28].

The wear morphology of the coating without added CeO₂ exhibited discontinuous furrows with distinct adhesive deformation (Figure 11b). The presence of numerous titanium carbide dendrites induced significant micro-vibrations on the coating surface during sliding, resulting in fretting wear. Furthermore, the titanium carbide dendrites prevented the continuous ploughing of hard abrasive particles, thereby forming discontinuous furrows. The generation of substantial heat of friction on the contact surface between the composite coating and GCr15 softened the surface of the coating and increased the mutual adhesive tendencies. The coatings with added CeO₂ each exhibited a lower average hardness than that of the coating without added CeO₂, resulting in the generation of significant heat of friction during wear. The wear morphology of the coating without added CeO₂ exhibited distinct adhesion characteristics. The addition of CeO₂ reduced both the dendritic titanium carbide content and the hardness of the coating in conjunction with an increase in the content of the matrix phase. These changes effectively alleviated the micro-undulations on the coating surface, thereby forming continuous scratches. However, the adhesive strength of the coatings was lowered, and multiple continuous scratches were observed.

The chemical composition of the wear debris was analyzed using EDS (Figure 12) to elucidate the wear mechanism. The EDS spectrum of the wear debris of the substrate (Figure 12a) revealed the presence of Ti and O, demonstrating the formation of a TiO_x film during wear. Fe, originating in GCr15, was also detected in the wear debris, confirming the material transfer arising from the adhesion wear between the frictional materials. The EDS spectra of the coatings (Figure 12b–f) revealed that the wear debris on the different composite coatings contained more Fe and Cr than that on the substrate (Figure 12a), suggesting that GCr15 sustained significant damage. The hardness of TiC_x was higher than that of GCr15. Therefore, the contact between TiC_x and GCr15 during periodic sliding induced the reverse ploughing of GCr15, resulting in the desquamation of the constituents of GCr15 as wear debris. Minor quantities of Ti, Al, V, and Ni were also detected in the wear debris, indicating ploughing of the matrix. The Ni content in the debris of the coating without CeO₂ was minimal because this coating contained the most dendritic titanium carbide. The major loads were undertaken by these hard, uniformly distributed reinforcement phases, which lowered the degree of ploughing in the matrix phase. Reducing the dendritic titanium carbide content of the coatings by adding CeO₂ resulted in pronounced exposure and ploughing of the matrix phases.

BSE images of the coatings (Figure 13) were used to characterize their morphologies and thus elucidate the influence of the rare-earth oxide, CeO₂, on the wear mechanism of the composite coatings. The wear surface exhibited numerous bright particles, which were found to comprise primarily Ce and O by EDS, and thus these particles were identified as Ce oxides. Ding et al. [29] reported that abrasive CeO₂ particles induce the formation of wide, deep furrows. The Ce oxides were distributed at the front of the furrow (Figure 13). Here, the continuous furrows were interrupted, or the original ploughing routes were changed, indicating that the presence of Ce oxide in the coating inhibited ploughing.

The wear resistance of the composite coatings was evaluated based on the wear rate, as calculated by Equation (1):

$$W_s=\frac{W_l}{F\cdot t}$$

(1)

where W_s is the wear rate, W_l is the weight loss owing to wear, F is the frictional load, and t is the time.

The results revealed that the wear resistance of the composite coatings was significantly higher than that of the substrate, despite their high friction coefficients (Figure 14). The wear rate of the substrate was calculated as 9.36 (×10⁻⁸ g/(N·s)), while those of coatings with 0, 1, 2, 3, and 4 wt.% CeO₂ were 0.42, 1.44, 4.05, 5.17, and 1.99 (×10⁻⁸ g/(N·s)), respectively. The coatings with 0, 1, 2, 3, and 4 wt.% CeO₂ exhibited wear resistances 22.3-, 6.5-, 2.3-, 1.8-, and 4.7-times, respectively, higher than that of Ti6Al4V. The wear resistance of a given material typically exhibits a positive correlation with its hardness [30]. The hardness of Ti6Al4V was lower than that of the composite coatings; therefore, the wear resistance of the composite coatings was superior to that of the substrate.

However, the hypothesis that high hardness results in high wear resistance is not supported by the observed wear rates of the coatings with different CeO₂ contents. Therefore, no simple positive correlation was observed between the microhardness and wear resistance, which was because the wear resistance of the materials was strongly associated with their microstructure and forming quality. Several factors may contribute to this phenomenon. (a) Stress concentration occurs at the pores, cracks, inclusions, and other defects, and functions as a source of failure that promotes crack initiation and propagation until fracture. Therefore, the forming quality of the composite coating exerts a significant influence on the wear resistance. When the addition of CeO₂ is 3 wt.%, it has the highest hardness, but there are many defects, such as cracks, resulting in a higher wear rate under the comprehensive action. (b) Furthermore, a high quantity of hard titanium carbide was tightly embedded in the CrTi₄ matrix phase. Consequently, the matrix phase exhibited high toughness, with high resistance to the ploughing action of the friction pair and hard debris. The friction load was undertaken by the ceramic reinforcing phase, and stress was released from the matrix phase via plastic deformation. Thus, brittle fracture of the ceramic reinforcing phase was avoided and numerous hard, abrasive particles would not be produced by the brittle fracture of the ceramic reinforcement phase, which can weaken the grinding effect of abrasive particles on the coating. Therefore, the wear rate of the composite coating is lower than that of the Ti6Al4V substrate. (c) The appropriate amount of Ce oxides at the phase boundary could impart the composite coating with high wear resistance by preventing dislocation motion during wear, thereby suppressing plastic deformation and surface ploughing. (d) However, excess rare-earth oxides behaved as impurities to induce brittleness in the coating [23]. It was inferred that the hardness exerted the dominant influence on the wear resistance of the composite coatings. The nature of the microstructure and the presence of defects also influenced the wear resistance of the coatings. The in-situ-synthesized hard Ti₂C bound strongly to the matrix, imparting high resistance to plastic deformation and ploughing. Therefore, the Ti₂C-reinforced Ti-based composite coatings exhibited high wear resistance. The hardness of the composite coatings with 3 wt.% CeO₂ is approximately the highest of all coatings, but the cracks on the coating surface seriously reduced its wear rate. Conversely, although the composite coating with 2 wt.% CeO₂ presented the lowest hardness compared with other composite coatings, its forming quality is the best, with no cracks and few pores. As a confluence of microstructure, forming quality, and hardness, its wear rate is not the lowest.

4. Conclusions

Ti₂C-reinforced CrTi₄-based composite coatings were successfully prepared on the Ti6Al4V surface via laser cladding technology. The effects of CeO₂ amount on the forming quality, microstructure, hardness, and wear resistance of the composite coatings were systematically studied. The main conclusions are as follows:

(1) Ti6Al4V/NiCr-Cr₃C₂ composite coatings with different CeO₂ contents were successfully prepared on a Ti6Al4V substrate using laser cladding technology with coaxial powder feeding. The addition of CeO₂ exerted a significant impact on the cracking behavior, porosity, and dilution rate of the composite coatings.

(2) The hardness of the composite coatings was 1.28–1.34-times higher than that of the substrate owing to the precipitation of hard Ti₂C and solid solution strengthening originating from the alloying elements. The coatings with added CeO₂ exhibited lower hardness than those without added CeO₂. Moreover, the CeO₂ content exerted a significant influence on the hardness of the composite coatings.

(3) The wear resistances of the composite coatings were significantly higher than that of the substrate. Both the substrate and composite coatings exhibited a mixed-wear mode with both abrasive and adhesive wear. Furthermore, ploughing of the coating was prevented by the presence of Ce oxides in the composite. CeO₂ content of 2 wt.% in the composite coating was determined to be optimal, as these presented no cracks, low porosity, and high wear resistance.