Effect of Chromium Carbide Addition on the Microstructures and Properties in Dual Carbide Phases Reinforced Ni-Based Composite Coatings by Plasma Cladding

Cr3C2-modified NiCr–TiC composite coatings were prepared using the plasma spraying technique for different Cr3C2 contents on the microstructure and the properties of the Ni-based TiC cladding layer were investigated. The microstructures of the coatings were characterized using scanning electron microscopy, and the friction and wear performance of the coating was evaluated by the wear tests. The results revealed that the surfaces of the Cr3C2-modified NiCr–TiC composite coatings with varying Cr3C2 contents were dense and smooth. TiC was uniformly distributed throughout the entire coating, forming a gradient interface between the binder phase of the Ni-based alloy and the hard phase of TiC. At high temperatures, Cr3C2 decomposes, with some chromium diffusing and forming complex carbides around TiC, some chromium solubilizes with Fe, Ni, and other elements. An increase in chromium carbide content leads to an upward trend in hardness. The measured hardness of the coatings ranged from 600 to 850 HV3 and tended to increase with increasing Cr3C2 content. When the mass fraction of Cr3C2 reached 30%, the hardness increased to 850 HV3, and the cracks and defects were observed in the coating, resulting in a wear resistance decline.


Introduction
Particle-reinforced metal matrix composites significantly enhance the wear resistance of coatings. By leveraging the advantageous physical, chemical, and mechanical properties of the reinforcing materials, these coatings compensate for the deficiencies of the substrate material, resulting in a superior overall performance. These coatings are mainly applied in industries requiring surfaces with excellent wear resistance and components with long lifespans, such as in the engineering of mixed-material mixing drums for new energy equipment, marine engineering, and automotive turbochargers [1,2]. The type and amount of the reinforcing phase added to the composite determines the operating performance of the coating.
Typically, the reinforcing phases used in composite coatings are metal carbides such as WC, TiC, TiN, and CrC. The addition of carbides can be achieved through ex situ and in situ methods. In situ generated carbides tend to be uniformly distributed and well bonded to the matrix; however, their content is remarkably limited, restricting the potential for improvements in hardness and wear resistance. For example, composite coatings with different TiC/Co50 contents were produced using a laser-cladding technique [3]. As the TiC content increased, the hardness of the cladding layer increased, reaching approximately

Methods
The composite powders were mixed using an ultrasonic vibration mixer for 1 h. The cladding experiments were conducted using Duomu A300 plasma cladding equipment. During the cladding process, a coaxial carrier gas powder feeding method was used for single-layer multiple-pass cladding, with argon gas protection. The parameters are presented in Table 3. The test samples were polished by the ZXQ-5HS inlay machine gold sandpaper (60 #to2000#), and then on the MP-3S-2 with 15 micron, 8 micron, 1 micron automatic polishing machine, each polishing time is 3 min. Phase analysis of the cladding layers was performed using a Bruker D8 X-ray diffractometer. The parameters were as follows: operating voltage of 40 kV, operating current of 40 mA, Cu Kα radiation, step scan speed of 2 • /min, and scanning range of 20~90 • . Microstructural observation and analysis of the cladding layers were carried out using a Tescan Mira4 field-emission scanning electron microscope. The microhardness of different positions along the cross-section of the laser cladding layer at the same horizontal level was measured using a Huayin 200HV-5 Vickers hardness tester. A five-point average was taken, with a test load of 3000 g (30 N) applied for 10 s. The friction and wear performance of the cladding layers were evaluated using an MM-U10G pin-on-disk tribometer at room temperature with a load of 50 N, rotational speed of 50 r/min, and wear time of 1 h, using a silicon nitride ball with a diameter of ϕ 9.425 mm as the counterface. After the wear test, the mass loss owing to friction was determined using a FA2004 electronic balance with a precision of 0.1 mg.

Results and Discussion
3.1. Phase Structure of the Cladding Layer Figure 1 shows the X-ray diffraction (XRD) pattern of the composite coating. The results indicate that the microstructure of the cladding layer is composed of solid solution phases such as FeNi/NiCrFe, TiC, (Fe, Cr)C/(Fe, Cr, Ni)C, Cr 7 C 3 , and complex carbides. According to the literature [23], the enthalpy of formation (∆rH) for Cr 3 C 2 is -0.114 eV/atom. The enthalpies of formation for the solid solution phases were calculated using Materials Studio software (Version 8.0); the enthalpies for (NiCr 11 )C 8 , (Ni 2 Cr 10 )C 8 , and (Ni 3 Cr 9 )C 8 were calculated to be −0.124, −0.0758, and +0.1298 eV/atom. This indicates that when Ni ≥ 3, the enthalpy of formation for the substituted solid solution phases becomes positive, and thus solid solution phases cannot be formed. Cr 3 C 2 decarburizes at high temperatures to produce Cr 7 C 3 . The dissolution of chromium carbides into TiC causes changes in the lattice parameters of the δ phase. Wenlong et al. [20] found that when the chromium carbide content was TiC, TiC-6.7% Cr 3 C 2, TiC-19.5% Cr 3 C 2 , and TiC-30% Cr 3 C 2 , respectively, the lattice parameters changed at 1500 • C to 4.33 −10 , 4.31 −10 , 4.27 −10 , 4.25 −10 m, which remain essentially unchanged. Moreover, for the Cr 3 C 2 solid solution in TiC, a strong and stable structure was formed.   Table 2). The coatings are designated 1# for the Ni-15Cr3C2-45TiC coating with 15% Cr3C2 content, 2# for the Ni-19Cr3C2-45TiC coating with 19% Cr3C2 content, 3# for the Ni-23Cr3C2-45TiC coating with 23% Cr3C2 content, 4# for the Ni-26Cr3C2-45TiC coating with 26% Cr3C content, and 5# for the Ni-30Cr3C2-45TiC coating with 30% Cr3C2 content, respectively Figure 2 shows that the cladding layer can be divided into three distinct regions: the heat During the plasma cladding of nickel-based carbide composite coatings on an iron base plate, melting of the matrix provides the Fe source, and Cr 3 C 2 pyrolysis provides the Cr and C sources. Under the high-temperature conditions of plasma cladding, a compound FeNi/NiCrFe and (FeNi)C/(FeCrNi)C solid solution was produced, with a significant peak value [24][25][26][27][28]. With increasing Cr 3 C 2 content, the main phase types of the different composite coatings did not change. They were composed of TiC, Cr 3 C 2 , FeNi/NiCrFe, (FeNi)C/(FeCrNi)C solid solution and other phases. The main diffraction peaks of each phase were enhanced and sharpened to varying degrees; however, the positions of the diffraction peaks did not shift. The XRD spectrum revealed that, owing to the role of the high-energy plasma beam, Cr 3 C 2 added to the nickel-based alloy undergoes a large amount of decomposition, which significantly increases the Cr and C contents in the culture pool. During the rapid solidification process, a large number of solid solutions and carbides will be generated, followed by the formation of (FeNi)C/(FeCrNi)C solid solution.  Table 2). The coatings are designated 1# for the Ni-15Cr 3 C 2 -45TiC coating with 15% Cr 3 C 2 content, 2# for the Ni-19Cr 3 C 2 -45TiC coating with 19% Cr 3 C 2 content, 3# for the Ni-23Cr 3 C 2 -45TiC coating with 23% Cr 3 C 2 content, 4# for the Ni-26Cr 3 C 2 -45TiC coating with 26% Cr 3 C 2 content, and 5# for the Ni-30Cr 3 C 2 -45TiC coating with 30% Cr 3 C 2 content, respectively. Figure 2 shows that the cladding layer can be divided into three distinct regions: the heat-affected zone, the bonding zone, and the cladding zone. The bonding zone appears as a prominent "bright" band, which was formed by the mutual dilution of the substrate metal and the cladding powder during the plasma cladding. The bonding zone acts as a transition region between the substrate material and coating, indicating a metallurgical bond between them.   Table 2). The coatings are designated 1# for the Ni-15Cr3C2-45TiC coating with 15% Cr3C2 content, 2# for the Ni-19Cr3C2-45TiC coating with 19% Cr3C2 content, 3# for the Ni-23Cr3C2-45TiC coating with 23% Cr3C2 content, 4# for the Ni-26Cr3C2-45TiC coating with 26% Cr3C2 content, and 5# for the Ni-30Cr3C2-45TiC coating with 30% Cr3C2 content, respectively. Figure 2 shows that the cladding layer can be divided into three distinct regions: the heataffected zone, the bonding zone, and the cladding zone. The bonding zone appears as a prominent "bright" band, which was formed by the mutual dilution of the substrate metal and the cladding powder during the plasma cladding. The bonding zone acts as a transition region between the substrate material and coating, indicating a metallurgical bond between them.

Microstructure of the Cladding Layer
SEM images of the cross-sectional micromorphology of the plasma cladding layer with various mass fractions of Cr 3 C 2 are presented in Figure 3. The black phase represents the TiC phase, which is uniformly distributed in the coating. In the coating with 15% Cr 3 C 2 , the second phase is present in smaller quantities and has smaller particle sizes, with a relatively uniform distribution. However, when the mass fraction of Cr 3 C 2 is increased to 30%, the particle TiC size remains similar and relatively unchanged, whereas the size of the second phase significantly increases. With increasing distance from the bottom of the molten pool, the microstructure of the cladding layer transitions from cell grains to dendritic grains, and at the top of the cladding layer, it transforms into a fully disordered dendritic microstructure. This is likely influenced by the polycrystalline orientation of the base material and possible nucleation at the solidification front of the molten pool.
with various mass fractions of Cr3C2 are presented in Figure 3. The black phase represents the TiC phase, which is uniformly distributed in the coating. In the coating with 15% Cr3C2, the second phase is present in smaller quantities and has smaller particle sizes, with a relatively uniform distribution. However, when the mass fraction of Cr3C2 is increased to 30%, the particle TiC size remains similar and relatively unchanged, whereas the size of the second phase significantly increases. With increasing distance from the bottom of the molten pool, the microstructure of the cladding layer transitions from cell grains to dendritic grains, and at the top of the cladding layer, it transforms into a fully disordered dendritic microstructure. This is likely influenced by the polycrystalline orientation of the base material and possible nucleation at the solidification front of the molten pool.  The microstructure of the cladding layer consists of a dendritic solid solution, eutectic structure, undissolved chromium carbides, as well as rod-and block-like carbide phases. As the Cr3C2 content increases, the dendritic grains in the cladding layer transform into equiaxed grains, resulting in a finer microstructure and a higher number of block-like carbide phases, as confirmed by XRD analysis. Additionally, a significant amount of undissolved titanium carbides are uniformly distributed in the coating, which can be attributed to the rapid cooling rate during the plasma fusing process.
The high-magnification image of the 1# cladding layer, as shown in Figure 4, reveals the presence of numerous black TiC particles, as well as a significant amount of grey elongated and block-like structures. Energy-dispersive spectroscopy (EDS) analysis (Table The microstructure of the cladding layer consists of a dendritic solid solution, eutectic structure, undissolved chromium carbides, as well as rod-and block-like carbide phases. As the Cr 3 C 2 content increases, the dendritic grains in the cladding layer transform into equiaxed grains, resulting in a finer microstructure and a higher number of block-like carbide phases, as confirmed by XRD analysis. Additionally, a significant amount of undissolved titanium carbides are uniformly distributed in the coating, which can be attributed to the rapid cooling rate during the plasma fusing process. The high-magnification image of the 1# cladding layer, as shown in Figure 4, reveals the presence of numerous black TiC particles, as well as a significant amount of grey elongated and block-like structures. Energy-dispersive spectroscopy (EDS) analysis (Table 3) indicates that Point A contains a high concentration of Fe, Cr, and Ni, while Point B is primarily composed of Fe and Ni, with lower Cr content. Combined with the XRD results ( Figure 2), it can be concluded that the phases at Point A are (Fe, Cr)C and (Fe, Cr, Ni)C, while Point B contains (Fe, Cr) and (Cr, Fe, Ni) solid solutions owing to the incorporation of Cr. The EDS results of the TiC region in the 1# cladding layer (Table 4) reveal that during the cladding process, a significant amount of iron from the substrate melted into the coating [29]. The contents of other elements may vary depending on the specific region selected for analysis. undissolved titanium carbides are uniformly distributed in the coating, which can be attributed to the rapid cooling rate during the plasma fusing process.
The high-magnification image of the 1# cladding layer, as shown in Figure 4, reveals the presence of numerous black TiC particles, as well as a significant amount of grey elongated and block-like structures. Energy-dispersive spectroscopy (EDS) analysis (Table  3) indicates that Point A contains a high concentration of Fe, Cr, and Ni, while Point B is primarily composed of Fe and Ni, with lower Cr content. Combined with the XRD results ( Figure 2), it can be concluded that the phases at Point A are (Fe, Cr)C and (Fe, Cr, Ni)C, while Point B contains (Fe, Cr) and (Cr, Fe, Ni) solid solutions owing to the incorporation of Cr. The EDS results of the TiC region in the 1# cladding layer (Table 4) reveal that during the cladding process, a significant amount of iron from the substrate melted into the coating [29]. The contents of other elements may vary depending on the specific region selected for analysis.  Coating #1 was further using EDS, revealing the formation of multicomponent carbides (NixCryFezTih)C with a thickness of about 0.5-1 µm around the TiC phase. The values of x, y, z, and h gradually transitioned as the thickness increased, until the outermost titanium (Ti) disappeared completely. The results for P1-P5 (Table 5) show that  Coating #1 was further using EDS, revealing the formation of multicomponent carbides (NixCryFezTih)C with a thickness of about 0.5-1 µm around the TiC phase. The values of x, y, z, and h gradually transitioned as the thickness increased, until the outermost titanium (Ti) disappeared completely. The results for P1-P5 (Table 5) show that the decreasing Ti content is accompanied by a further transition to multicomponent carbides, such as (FeNi)C and (FeCrNi)C occurs. The elemental distribution at point P4 indicates that the melted Cr 3 C 2 provides sufficient carbon for the precipitation of a second phase, namely (Fe, Cr, Ni)C. The results from point P5 indicate the formation of an FeNi solid solution enabled by the high Fe and Ni contents in the substrate [30]. During the melting process, the melted TiC reacts with decomposed Cr 3 C 2 and Fe to form ternary carbides, as shown in Tables 4-6, with precipitation of a second phase observed at Point P6. Meanwhile, the unmelted large TiC particles remain intact. At Point P7, a network-like binary carbide consisting of Fe and Cr is formed. With an increase in the Cr 3 C 2 content, both the quantity and size of the formed complex carbides increase. Pokhmurska et al. [31] showed that Cr 3 C 2 dispersed in the solid solution can effectively enhance the hardness of the coating. Figure 5 compares the average surface and crosssectional hardness values of the cladding layer, revealing that the hardness of the composite coating generally increases with the increase in Cr 3 C 2 content, reaching a maximum of 825.4 HV3. The analysis results indicate that the morphology of the TiC particles in the coating remains largely unchanged, while the primary changes occur in the binder-phase nickelbased alloy, Cr3C2, and in the melted Fe. Owing to the low melting point of Cr3C2, it dissolves in the molten pool under the action of a high-temperature plasma arc, providing a significant source of Cr and C. For one thing, Cr and C are supersaturated in the solid solution, enhancing the solid-solution strengthening effect of the alloying elements. For another thing, during the cooling process, they can further react with C, B, and other elements to form more compounds, thus further improving the performance of the coating. Partial melting of Cr3C2 and subsequent cooling results in the precipitation of Cr3C2 particles as non-uniform nuclei that promote nucleation and refining of the microstructure of the coating, thus providing a grain refinement strengthening effect. In addition, the area surrounding the black TiC particles is mainly composed of (NiCrFe)C and (FeCr)C along with dispersed fine carbides such as TiC. Among them, the γ-Ni dendritic solid solution, which contains dissolved C, Cr, Si, Fe, and other elements, possesses a eutectic structure with compound phases. Based on the morphology and Xray analysis of the coating, the region with higher Cr content is considered the primary Cr-rich phase and is believed to contain a certain amount of dissolved Ni, Fe, and other The analysis results indicate that the morphology of the TiC particles in the coating remains largely unchanged, while the primary changes occur in the binder-phase nickelbased alloy, Cr 3 C 2 , and in the melted Fe. Owing to the low melting point of Cr 3 C 2 , it dissolves in the molten pool under the action of a high-temperature plasma arc, providing a significant source of Cr and C. For one thing, Cr and C are supersaturated in the solid solution, enhancing the solid-solution strengthening effect of the alloying elements. For another thing, during the cooling process, they can further react with C, B, and other elements to form more compounds, thus further improving the performance of the coating. Partial melting of Cr 3 C 2 and subsequent cooling results in the precipitation of Cr 3 C 2 particles as non-uniform nuclei that promote nucleation and refining of the microstructure of the coating, thus providing a grain refinement strengthening effect. In addition, the area surrounding the black TiC particles is mainly composed of (NiCrFe)C and (FeCr)C along with dispersed fine carbides such as TiC. Among them, the γ-Ni dendritic solid solution, which contains dissolved C, Cr, Si, Fe, and other elements, possesses a eutectic structure with compound phases. Based on the morphology and X-ray analysis of the coating, the region with higher Cr content is considered the primary Cr-rich phase and is believed to contain a certain amount of dissolved Ni, Fe, and other elements. During the solidification process, the harndess of the cladding layer is enhanced via solid solution, dispersion, and grain refinement effects. The composite coating contains a significant amount of TiC and hard phases of complex carbides, which serve as the supporting framework of the entire cladding layer, further improving the hardness of the coating [32].

Nanoindentation
Nanoindentation testing was performed on sample 4# ( Figure 6). Based on the results of the XRD and SEM-EDS analysis, Point 1 can be concluded to correspond to TiC particles; points 2-4 correspond to complex carbides; and Points 5 and 6 correspond to NiFe/FeNi solid solution. The microhardness results are shown in Table 7.  The large TiC particles exhibit a high hardness of 32 GPa. The (TiCr)C/(TiCrFe)C carbides have hardness values exceeding 20 GPa, indicating high hardness. The NiFe/FeNi solid solution exhibits a hardness above 6 GPa. This further confirms that Cr3C2 dissolves significantly during the high-temperature process of thermal spraying, forming multicomponent carbides around the TiC particles, including (TiCr)C/(TiCrFe)C multicomponent carbides with other elements. With increasing Cr3C2 content, the volume fraction of multicomponent carbides in the coating increases, which enhances the hardness and elastic modulus. The NiFe and FeNi phases exhibit high hardness and toughness. The high-volume fraction of hard phases helps to reduce thermal stresses and improve the coating's defect structure, thereby enhancing the overall performance of the coating. Figure 7 compares the friction coefficient and weight loss due to sliding wear for composite coatings with different Cr3C2 contents. The coating with 26% Cr3C2 addition exhibits a lower friction coefficient and the least amount of wear. Although the coating labeled 1# has a lower friction coefficient, it also has a lower hardness, resulting in a higher weight loss compared to coating 4#.  The large TiC particles exhibit a high hardness of 32 GPa. The (TiCr)C/(TiCrFe)C carbides have hardness values exceeding 20 GPa, indicating high hardness. The NiFe/FeNi solid solution exhibits a hardness above 6 GPa. This further confirms that Cr 3 C 2 dissolves significantly during the high-temperature process of thermal spraying, forming multicomponent carbides around the TiC particles, including (TiCr)C/(TiCrFe)C multicomponent carbides with other elements. With increasing Cr 3 C 2 content, the volume fraction of multicomponent carbides in the coating increases, which enhances the hardness and elastic modulus. The NiFe and FeNi phases exhibit high hardness and toughness. The highvolume fraction of hard phases helps to reduce thermal stresses and improve the coating's defect structure, thereby enhancing the overall performance of the coating. Figure 7 compares the friction coefficient and weight loss due to sliding wear for composite coatings with different Cr 3 C 2 contents. The coating with 26% Cr 3 C 2 addition exhibits a lower friction coefficient and the least amount of wear. Although the coating labeled 1# has a lower friction coefficient, it also has a lower hardness, resulting in a higher weight loss compared to coating 4#.  Figure 8 shows the cross-section of the composite coating after friction testing. The wear coefficients of the samples are shown in Table 8. The 1# coating exhibits a lower friction coefficient, but larger wear dimensions in terms of the width and depth of the wear marks. This is mainly due to its lower hardness, resulting in higher wear compared to the 4# coating. The friction coefficient of sample 4# was lower than that of sample 5#. Although the hardness of sample 4# was lower than that of sample 5#, the width and depth of the wear marks were smaller in sample 4#, indicating excellent wear resistance performance. For sample 5#, the width and depth of the wear marks were measured to be 924.43 µm and 35 µm, respectively. For sample 4#, the width and depth were measured to be 845.23 µm and 8 µm, respectively. During friction and wear testing, the friction coefficient curve was directly obtained using software (origin2021), and the wear rate (WR; mm 3 /N·m) was calculated using Equation (1).

Wear Morphology
Here, V is the wear volume (mm 3 ), F denotes the normal force applied during the wear process (in N), L represents the total friction stroke (m), and WR represents the wear rate (mm 3 /N·m), Therefore, WR indicates the wear volume per unit wear distance under a unit load.  Figure 8 shows the cross-section of the composite coating after friction testing. The wear coefficients of the samples are shown in Table 8. The 1# coating exhibits a lower friction coefficient, but larger wear dimensions in terms of the width and depth of the wear marks. This is mainly due to its lower hardness, resulting in higher wear compared to the 4# coating. The friction coefficient of sample 4# was lower than that of sample 5#. Although the hardness of sample 4# was lower than that of sample 5#, the width and depth of the wear marks were smaller in sample 4#, indicating excellent wear resistance performance. For sample 5#, the width and depth of the wear marks were measured to be 924.43 µm and 35 µm, respectively. For sample 4#, the width and depth were measured to be 845.23 µm and 8 µm, respectively.

Wear Morphology
During friction and wear testing, the friction coefficient curve was directly obtained using software (origin2021), and the wear rate (WR; mm 3 /N·m) was calculated using Equation (1).
Here, V is the wear volume (mm 3 ), F denotes the normal force applied during the wear process (in N), L represents the total friction stroke (m), and WR represents the wear rate (mm 3 /N·m), Therefore, WR indicates the wear volume per unit wear distance under a unit load.  Figure 8 shows the cross-section of the composite coating after friction testing. The wear coefficients of the samples are shown in Table 8. The 1# coating exhibits a lower friction coefficient, but larger wear dimensions in terms of the width and depth of the wear marks. This is mainly due to its lower hardness, resulting in higher wear compared to the 4# coating. The friction coefficient of sample 4# was lower than that of sample 5#. Although the hardness of sample 4# was lower than that of sample 5#, the width and depth of the wear marks were smaller in sample 4#, indicating excellent wear resistance performance. For sample 5#, the width and depth of the wear marks were measured to be 924.43 µm and 35 µm, respectively. For sample 4#, the width and depth were measured to be 845.23 µm and 8 µm, respectively. During friction and wear testing, the friction coefficient curve was directly obtained using software (origin2021), and the wear rate (WR; mm 3 /N·m) was calculated using Equation (1).   Table 8 indicate that the wear resistance of the composite coating is exceptional. This improvement in wear performance aligns with Archard's wear theory [13], which suggests that the wear volume (Vw) is directly proportional to the wear load, but inversely proportional to the material's hardness ( Figure 5). The composite coating exhibits higher hardness and thus superior wear resistance. Additionally, the lower friction coefficients contribute to further the wear performance of these coatings. Figure 9 shows that the widths of the wear marks on coatings 2# and 4# are remarkably small. In terms of friction coefficient and wear volume, sample 4# exhibits the best wear resistance. Microscopically, numerous branching cracks can be observed on the surface of the hard phase. This may be due to crack deflection around the two-phase particles and the bending of crack fronts between particles, resulting in the reduced driving force for crack propagation and increased fracture toughness. The plastic deformation zone ahead of the crack absorbs some energy, hindering crack propagation [33][34][35].   Table 8 indicate that the wear resistance of the composite coating is exceptional. This improvement in wear performance aligns with Archard's wear theory [13], which suggests that the wear volume (Vw) is directly proportional to the wear load, but inversely proportional to the material's hardness ( Figure 5). The composite coating exhibits higher hardness and thus superior wear resistance. Additionally, the lower friction coefficients contribute to further the wear performance of these coatings. Figure 9 shows that the widths of the wear marks on coatings 2# and 4# are remarkably small. In terms of friction coefficient and wear volume, sample 4# exhibits the best wear resistance. Microscopically, numerous branching cracks can be observed on the surface of the hard phase. This may be due to crack deflection around the two-phase particles and the bending of crack fronts between particles, resulting in the reduced driving force for crack propagation and increased fracture toughness. The plastic deformation zone ahead of the crack absorbs some energy, hindering crack propagation [33][34][35]. The wear surface contains hard phases of TiC, (FeCr)C, (FeCrNi)C, FeNi, FeNiCr, oxide, and bonded Si3N4, among others. Scanning electron microscopy (Table 9) reveals The wear surface contains hard phases of TiC, (FeCr)C, (FeCrNi)C, FeNi, FeNiCr, oxide, and bonded Si 3 N 4 , among others. Scanning electron microscopy (Table 9) reveals that the wear mechanism of the samples is abrasive wear, with some accumulation of wear debris during the wear process. This accumulation of wear debris increases the roughness of the wear surface, which is one of the main reasons for the relatively high friction coefficient of the material under room temperature conditions. The micromorphology and energy spectrum chemical composition analysis of the wear scars after the wear test of the composite coating are shown in Figures 9 and 10, respectively. The composite coatings exhibit adhesive wear characteristics, with the formation of thick oxide films on the friction surfaces [36,37]. The morphology of the wear debris at room temperature is shown in Figure 9. The energy spectrum reveals that the white particles are mainly composed of titanium carbide, chromium carbide ceramic particles, FeCrNi, and some SiN dual-phase debris. This is primarily because, during the wear process, ceramic phases fracture and accumulate on the surface, forming nanosized ceramic particles after repeated rolling and fragmentation by the abrasive balls. The gray portion consists of a mixture of Cr 3 O 2 , FeO, and NiO oxides. In addition, the energy spectrum indicates that the upper oxide layer is mainly composed of Ti, Cr, and Ni oxides, while the lower oxide layer is primarily composed of Ti oxides. Similar results are observed in the wear scars of the composite coating. These oxides are formed owing to the high temperature generated during the friction process, resulting in rapid surface oxidation and effectively improving the wear resistance of the coating. The Gibbs free energy diagrams of the elemental oxidation reactions reveal that the oxidation of TiC into TiO 2 and CO 2 has the lowest energy compared with the other oxidation reactions [38]. This suggests that TiC preferentially forms oxides during the friction process and accounts for the significant amount of Ti oxide in the oxide films generated in the composite coating during friction. The formation of TiO 2 and other metal oxides during the friction process is advantageous because they are relatively soft and possess lubricity [39]. This allows for the development of a continuous friction-induced oxide film, which exhibits lubricating properties. The presence of such a lubricating and continuous oxide film further enhances the wear resistance performance of the coating. that the wear mechanism of the samples is abrasive wear, with some accumulation of wear debris during the wear process. This accumulation of wear debris increases the roughness of the wear surface, which is one of the main reasons for the relatively high friction coefficient of the material under room temperature conditions. The micromorphology and energy spectrum chemical composition analysis of the wear scars after the wear test of the composite coating are shown in Figures 9 and 10, respectively. The composite coatings exhibit adhesive wear characteristics, with the formation of thick oxide films on the friction surfaces [36,37]. The morphology of the wear debris at room temperature is shown in Figure 9. The energy spectrum reveals that the white particles are mainly composed of titanium carbide, chromium carbide ceramic particles, FeCrNi, and some SiN dual-phase debris. This is primarily because, during the wear process, ceramic phases fracture and accumulate on the surface, forming nanosized ceramic particles after repeated rolling and fragmentation by the abrasive balls. The gray portion consists of a mixture of Cr3O2, FeO, and NiO oxides. In addition, the energy spectrum indicates that the upper oxide layer is mainly composed of Ti, Cr, and Ni oxides, while the lower oxide layer is primarily composed of Ti oxides. Similar results are observed in the wear scars of the composite coating. These oxides are formed owing to the high temperature generated during the friction process, resulting in rapid surface oxidation and effectively improving the wear resistance of the coating. The Gibbs free energy diagrams of the elemental oxidation reactions reveal that the oxidation of TiC into TiO2 and CO2 has the lowest energy compared with the other oxidation reactions [38]. This suggests that TiC preferentially forms oxides during the friction process and accounts for the significant amount of Ti oxide in the oxide films generated in the composite coating during friction. The formation of TiO2 and other metal oxides during the friction process is advantageous because they are relatively soft and possess lubricity [39]. This allows for the development of a continuous friction-induced oxide film, which exhibits lubricating properties. The presence of such a lubricating and continuous oxide film further enhances the wear resistance performance of the coating. The abrasion debris exhibits larger delamination in a layered form. In Figure 9, it can be observed that under room temperature conditions, the composite coating predominantly undergoes abrasive wear with abrasive particles when in contact with a silicon nitride counterpart [40], and a minor amount of oxidative wear. Fracture of titanium carbide and multicomponent carbide particles occurs in the sub-surface region, with cracks primarily appearing at the interface between the multicomponent carbide The abrasion debris exhibits larger delamination in a layered form. In Figure 9, it can be observed that under room temperature conditions, the composite coating predominantly undergoes abrasive wear with abrasive particles when in contact with a silicon nitride counterpart [40], and a minor amount of oxidative wear. Fracture of titanium carbide and multicomponent carbide particles occurs in the sub-surface region, with cracks primarily appearing at the interface between the multicomponent carbide particles and the metal binder phase.

Wear Morphology
SEM and EDS analysis were performed on the wear debris of coating 4#, as shown in Figure 11, respectively. The wear resistance of the composite coating is mainly attributed to the presence of titanium carbide (TiC), as revealed by EDS analysis of the wear debris. As the Cr 3 C 2 content varies in the range of 15% to 26%, the increase in FeCr carbide (FeCr)C content gradually leads to higher friction coefficients, increased hardness, and improved wear resistance. However, when the Cr 3 C 2 content exceeds 30%, the hardness increases, but cracks penetrating the coating appear, resulting in more coating defects and a decrease in wear resistance. The abrasion debris exhibits larger delamination in a layered form. In Figure 9, it can be observed that under room temperature conditions, the composite coating predominantly undergoes abrasive wear with abrasive particles when in contact with a silicon nitride counterpart [40], and a minor amount of oxidative wear. Fracture of titanium carbide and multicomponent carbide particles occurs in the sub-surface region, with cracks primarily appearing at the interface between the multicomponent carbide particles and the metal binder phase.
SEM and EDS analysis were performed on the wear debris of coating 4#, as shown in Figure 11, respectively. The wear resistance of the composite coating is mainly attributed to the presence of titanium carbide (TiC), as revealed by EDS analysis of the wear debris. As the Cr3C2 content varies in the range of 15% to 26%, the increase in FeCr carbide (FeCr)C content gradually leads to higher friction coefficients, increased hardness, and improved wear resistance. However, when the Cr3C2 content exceeds 30%, the hardness increases, but cracks penetrating the coating appear, resulting in more coating defects and a decrease in wear resistance.

In Situ Formation of the Hard Phase
To explain the formation mechanism of FeNi/NiCrFe,(Ti,Cr)C and (FeNi)C/(FeCrNi)C, and their influence on the coating properties, the solidification process of the molten tank was considered. During the plasma cladding process, when the melting pool temperature exceeds 2000 °C, Cr3C2 is decomposed into CrC molecules to form a composite carbide transition layer around TiC and form a stable binding body with

In Situ Formation of the Hard Phase
To explain the formation mechanism of FeNi/NiCrFe,(Ti,Cr)C and (FeNi)C/(FeCrNi)C, and their influence on the coating properties, the solidification process of the molten tank was considered. During the plasma cladding process, when the melting pool temperature exceeds 2000 • C, Cr 3 C 2 is decomposed into CrC molecules to form a composite carbide transition layer around TiC and form a stable binding body with the matrix. The melting of Fe, Ni, Cr, and C is generated in situ (FeNi)C/(FeCrNi)C. The melting point of the nickel-based alloy was approximately 1300 • C, of TiC was approximately 3067 • C, and of Cr 3 C 2 was approximately 1890 • C. Cr 3 C 2 and TiC will preferentially nucleate during the pool solidification. The study showed that (Ti, Cr) C formed, even though the melting pool temperature was below the melting point of the nickel-iron alloy. According to the B-C-Cr ternary phase diagram, the melting point of Cr 3 C 2 is higher than that of Cr 23 C 6 . The addition of Ti and Cr consumed some C atoms and increased the Cr/C ratio, favoring Cr 3 C 2 formation in the B-C-Cr system, while Cr 23 C 6 was inhibited. This was the primary reason for the presence of numerous stable carbides in the coating and improved wear resistance [41].
Compound carbide particles formed in situ have high hardness and are uniformly distributed in the coating, which improves the coating wear resistance (see Figure 6). Additional hot melting of Cr 3 C 2 caused a sharp increase in the contents of carbide in (TiCr)C and (FeNi)C/(FeCrNi)C in the coating. The high hardness of the carbides in the coating can hinder the dislocation movement and deformation of the substrate and play a significant role in pinning the substrate. Thus, the strength of the substrate is enhanced, and the wear resistance of the coating is improved. In addition, Fe in the Fe matrix, Ni form Ni25, and Cr atoms form Cr 3 C 2 form a solid reinforcement phase that increases the hardness of the matrix and improves the wear resistance of the coating.

1.
Plasma-sprayed Ni-based TiC composite coatings with various Cr 3 C 2 contents exhibit a smooth surface morphology, dense structure, high bonding strength, and high microhardness. The composite coatings are composed of TiC, (FeCr)C, and (Fe, Cr, Ni)C multicomponent carbides, as well as (FeCr) and (Fe, Cr, Ni) solid solutions.

2.
As the mass fraction of Cr 3 C 2 increases, the microhardness of the plasma-sprayed coating also tends to increase. When the mass fraction of Cr 3 C 2 reaches 30%, the overall hardness of the coating reaches its maximum at 825.4 HV3. Decarburization of Cr 3 C 2 provides a source of carbon for the formation of multicomponent carbides of Fe, Ni, and other elements, resulting in the formation of ring-shaped phases (Ti, Cr, Fe, Ni)C around TiC particles that exhibit excellent wetting properties with the nickel-based binder phase.

3.
The main wear mechanisms of the composite coating are abrasive wear and oxidative wear. With an increasing mass fraction of Cr 3 C 2 , both the friction and wear volume initially decreases and then increases, while the wear volume follows a similar trend. When the mass fraction of Cr 3 C 2 reaches 26%, the friction coefficient is relatively low, the wear marks are narrower, and the wear volume is the lowest. However, at a mass fraction of 30% Cr 3 C 2 , the thermal stress increases, and the number of defects and cracks penetrating the coating surface increases, decreasing the wear resistance. Institutional Review Board Statement: Not applicable.
Informed Consent Statement: Not applicable.

Data Availability Statement:
The data presented in this study are available on reasonable request from the corresponding author.