Study on the Organization and Wear Resistance of In718/(2Nb + 1SiC) Tool Composite Coating Based on Laser Cladding Technology

Abstract

As an advanced surface modification technology, laser cladding technology can significantly improve the hardness, wear resistance, and corrosion resistance of the base material by forming a high-performance composite coating on the surface of the substrate. In this paper, In718 nickel-based high-temperature alloy was used as the substrate, and laser cladding technology was used to prepare In718/(2Nb + 1SiC) composite coatings, and the effects of different laser powers (800 W, 1200 W, 1600 W, and 2000 W) on the microstructure and wear-resistant properties of the coatings were systematically investigated. Through SEM morphology observation, EDS elemental distribution analysis, wear behavior characterization, wear mechanism discussion, the regulatory mechanism of laser power on the microstructure evolution, the uniformity of reinforced phase distribution, and the coating properties were revealed. The results show that different melting powers significantly affect the phase structure and the generation of reinforcing phases in the In718/(2Nb + 1SiC) composite coatings. A 1200 W laser input power can make the coatings reach the optimal phase structure equilibrium. Both properties of the SiC-reinforced phase are retained, and a large number of NbC- and MoC-reinforced phases are generated. This results in excellent properties for the coating.

1. Introduction

IN718 alloy is a well-known nickel-based high-temperature alloy, which is widely used in aerospace, navigation, and chemical industries due to its excellent corrosion resistance and high-temperature mechanical properties [1,2,3]. However, its wear resistance is relatively limited, often leading to severe surface wear and failure. Laser cladding technology, as an efficient and precise surface strengthening technology, has been widely used in aerospace, mold making, and tool making in recent years [4,5,6]. By adding suitable reinforcing materials to the surface of the substrate and melting them rapidly at high laser temperatures, a surface coating with dense organization and excellent performance is formed [7,8,9]. This not only can significantly improve the material’s wear resistance, corrosion resistance, and high-temperature oxidation resistance, but can also extend the service life of the base material.

At present, the focus of IN718 laser cladding is mainly on repair and surface performance improvement [10,11,12,13,14]. Lin et al. [15] investigated the influence of different Y₂O₃ contents on the microstructure evolution and high-temperature corrosion resistance of Y₂O₃/IN718 composite coatings at 700 °C. The results showed that 1.0 wt.% Y₂O₃ could effectively reduce crack initiation, refine grain size, and improve high-temperature corrosion resistance. Liu et al. [16] used CrCoNi TiC/SiC powder for laser cladding on IN718 alloy to enhance wear resistance. The addition of TiC and SiC powders not only generates dispersed fine carbides in the coating but also changes the grain and substructure morphology by altering the solidification thermodynamics of the cladding zone. Mojtaba et al. [17] successfully prepared IN718/SiC composite coatings with better hardness and wear resistance through laser cladding. Compared with pure In718, adding up to 5 wt.% SiC nanoparticles to In718 increased wear resistance and microhardness by 51% and 50%, respectively, and reduced the friction coefficient by 129%. Thus, it can be seen that in the surface strengthening of the nickel-based, high-temperature alloy In718, the reasonable addition of reinforcing phases, such as carbide (SiC), can further enhance the mechanical properties and wear resistance of the coating. In addition, Nb plays an important strengthening role in nickel-based alloys. J et al. [18] prepared Inconel 625/Nb composite coatings with different Nb contents using laser cladding technology. With the increase of Nb content, the surface structure of the coating changed from dendritic to cellular crystals, and the formation of NbC was promoted. Hong et al. [19] investigated the effect of adding the strong carbide-forming element Nb on the microstructure and properties of Ni60/WC composite coatings. The increase of Nb promoted the increase of NbC content, and the morphology of NbC changed from coarse strips to fine needle-like aggregates and then to block-like structures. When the Nb content reached 3 wt.%, the needle-like aggregated NbC was evenly distributed, and the coating had the best wear resistance.

In summary, both SiC and Nb have a promoting effect on the performance improvement of nickel-based alloys. However, there is currently no research on the composite strengthening mechanism of SiC and Nb on IN718. Based on our previous references, this article intends to explore the laser process of In718/(2Nb + 1SiC) composite coatings, laying the foundation for further research in the future. In this study, In718/(2Nb + 1SiC) composite coatings were prepared with different microstructural properties by modulating the laser power using laser cladding technology. The effects of laser power on the microstructure, enhanced phase distribution, and wear resistance of the coatings were systematically investigated. Through microstructure observation, wear behavior analysis, and element distribution study, the intrinsic correlation between coating properties and laser process parameters is revealed. It provides theoretical guidance for optimizing the design of laser cladding coatings and, at the same time, provides a useful reference for the practical application of tool surface strengthening.

2. Materials and Methods

2.1. Experimental Materials

The substrate material was In718 high-temperature alloy provided by Shenzhen Baishun Metal Materials Co. The size of the specimen was 70 mm × 50 mm × 3 mm, and the composite powder consisted of Nb, SiC, and In718 with an average particle size of 53–150 μm, provided by New Life Alloys R&D Laboratory (see Figure 1), and the specific material compositions are shown in

Table 1. The chemical composition of In718 powders (wt.%).

C	Cr	Si	Nb	Mo	Co	Fe	Ti	Al	Mn	P	Ni
0.01	20.18	0.01	5.15	2.94	0.91	17.28	0.68	0.54	0.25	0.003	Bal

below. Nb and SiC powders with a particle size of 1 μm were supplied by the CMT Jingyao Alloy Powder Research Institute. In718, Nb and SiC were mixed homogeneously. The total mass of each group was 100 g, and the proportions of composite powders are In718:Nb:SiC = 97:2:1. The first group was prepared by 800 W process parameters, and the remaining three groups were prepared by 1200 W, 1600 W, and 2000 W process parameters, respectively [20,21].

2.2. Experiments and Characterization

The laser cladding experiment was conducted at Nanjing Huirui Optoelectronics Technology Co., Ltd. (Nanjing/China), using a BFSC-6000 laser with a length of 20 m and a 600 µm core diameter power transmission fiber. After the laser cladding was completed, 15 × 15 × 3 mm multi-pass cladding standard specimens and 5 × 5 × 3 mm single-pass cladding standard specimens were prepared using a wire-cutting machine [22,23,24,25]. The multi-pass cladding specimens were surface ground on 80–1000 grit SiC sandpaper for friction wear experiments. For single-pass fused cladding samples, the cross-sections were ground on 100–2500 grit sandpaper. A velvet polishing cloth was used to polish to a mirror finish and then etched with aqua regia (HCl:HNO₃ = 3:1) for 15 s. Finally, it was cleaned with alcohol and dried with a hair dryer. BRUKER X-ray Diffractometer (XRD) with CuK α radiation was chosen for the identification of the phase composition. Scanning electron microscopy was used to characterize the microstructure using Zeiss equipment from Germany. The element distribution and composition were analyzed by energy dispersive spectroscopy (EDS, XMX1003, Oberkochen, Bavaria, Germany). The specific parameters of the cladding process were set to a scanning speed of 12 mm/s, a powder feed rate of 15 g/min, and a spot diameter of 3 mm, as shown in

Table 2. Experimental parameters of laser cladding.

Number	Laser Power	Composite Powders	Scanning Speed	Feeding Rate	Spot Diameter
S1	800 W	Nb 2 wt.% + SiC 1 wt.% + IN718 97 wt.%	12 mm/s	15 g/min	3 mm
S2	1200 W
S3	1600 W
S4	2000 W

below. Each sample was tested three times during the testing process, and the average value was taken [26,27,28,29,30].

3. Results and Discussion

3.1. Microstructure

Figure 2 shows the X-ray diffraction (XRD, Bruker D8 Advance, Bruker Corporation, Billerica, MA, USA) patterns of In718/(2Nb + 1SiC) composite coatings at different powers. By looking at the In718 matrix, it can be seen that all samples (S1–S4) show typical In718 diffraction peaks for Ni-based alloys. Examples are the (111), (200), and (220) crystal plane diffraction peaks of Ni. The intensity and number of characteristic peaks of the enhanced phase varied in different samples. This suggests that the laser power significantly affects the precipitation behavior of the enhanced phase and the grain evolution mechanism by modulating the cooling rate of the molten pool (in synergy with the scanning speed). The reason for this is that the power increase (800 W → 1600 W) reduces the cooling rate of the melt pool (from 3.8 × 10³ K/s in S1 to 1.2 × 10³ K/s in S4). This thermodynamic difference directly contributes to the formation of microstructures. In some samples, S2 and S3, characteristic peaks of NbC and MoC appeared. This may be the carbides generated by chemical reaction during the cladding process. For sample S1, the intensity of diffraction peaks of matrix Ni is high. No obvious characteristic peaks of the new phase (e.g., NbC or MoC) were observed. It can be explained that the lower power resulted in insufficient melting temperature. The diffraction peaks of the S2 sample matrix Ni had moderate intensities, and those of the SiC reinforcing phase were significantly enhanced. At the same time, the diffraction peaks of NbC and MoC appeared, indicating that the reinforcing phase reacted effectively with the substrate to generate refined carbides. At 1200 W (S2), the high melt temperature and the melt duration of 4.2 s work together to provide sufficient thermodynamic conditions and a time window for the diffusion of activated carbon. In this process, niobium (Nb) and molybdenum (Mo) elements combine with carbon to form nanoscale carbides, and the strengthening mechanism is mainly due to the synergistic effect of fine crystallization, dispersion strengthening, and solid solution strengthening. Under the condition of low power (S1), the carbon activity is low due to the insufficient melt temperature, and the short diffusion time limits the full formation of carbides. Although the melt energy can be increased at high power, the turbulent effect of the molten pool will induce the decomposition of carbides (such as NbC), resulting in insufficient carbide generation or structural instability. Neither of the two extreme power conditions can achieve stable carbide strengthening phase regulation.

Figure 3 demonstrates the building direction and the microstructure of the cross-section of the In718/(2Nb + 1SiC) composite at different laser powers. Figure 3a shows the histograms of the top, middle, and bonded regions of the fused cladding layer at a laser power of 800 W. The top and middle regions of the fused cladding layer are shown in Figure 3b. From the figure, it can be seen that both columnar grains and a few equiaxed grains are present in the fused cladding layer. In the picture, columnar grains are predominant, and the growth direction of columnar grains is clearer, growing along the direction of the thermal gradient. Equiaxed crystals are distributed between the columnar crystals, showing local grain refinement. Due to the lower power, the melt pool size is smaller, and the cooling rate is faster, at which time the columnar crystals dominate. In the middle region, columnar crystals are mixed with a few equiaxed crystals. Compared to the top, the orientation of the columnar crystals is slightly weakened, a transition region appears, and some of the grains show a tendency to refine. The cooling rate in the middle region is slightly lower than at the top, but since the power is still low, the crystal growth is mainly dominated by the thermal gradient. At this point, columnar crystals are still predominant, and a limited number of equiaxed crystals are generated. This mixed distribution of microstructures suggests that in the intermediate region, grain growth is influenced by changes in melt pool size and cooling rate. The bottom of the molten cladding is dominated by equiaxed crystals with a more uniform and relatively fine grain distribution. Under lower power conditions, the cooling rate in the bottom region is the highest, leading to rapid solidification of the grains. At the same time, the columnar crystals are broken and transformed more into equiaxed crystals. This organization improves the densification of the bottom region to some extent, but due to the overall low power. The degree of grain refinement is still limited, and the organization uniformity is not satisfactory. It can be seen that the cooling rate of the fusion cladding layer is high at 800 W power, and the transition effect between columnar and equiaxed crystals is weak. The degree of grain refinement is average, resulting in the inability to achieve a high level of wear resistance. The smaller melt pool volume also limits the metallurgical bonding and organizational homogeneity of the material.

Figure 3b shows the organization of the top, middle, and bonded regions of the fused cladding layer at a laser power of 1200 W. As can be seen from the figure, the microstructure of the top region is dominated by columnar crystals at 1200 W power. However, compared with group a, the grains are more orientated and uniform in size. The experimental evidence demonstrates that the moderate power makes the melt pool size in the top region expand, the thermal gradient is stable, and the grains can grow in an orderly manner. In addition, the appearance of equiaxed crystals is slightly increased compared to group a. This may be due to a slight decrease in the cooling rate, which caused the columnar crystals to partially break up and begin to transform into equiaxed crystals. The crystals in the top region are better refined with well-defined grain boundaries. This indicates that this power regime helps to obtain a dense equiaxed structure in the top region. The microstructure in the middle region reaches an optimal balance between columnar and equiaxed crystals. The number of columnar crystals has decreased, while the proportion of equiaxed crystals has increased significantly. The overall grain refinement and distribution are more homogeneous, showing that the cooling rate and thermal gradient in the intermediate region are optimally matched at this power. The grain refinement effect in the intermediate region is enhanced, combined with the mechanical properties perspective. This refined organization contributes to the improvement of wear resistance and overall metallurgical bond quality. The bottom region is dominated by equiaxed crystals, which are very fine and uniformly distributed. The 1200 W power resulted in a moderate cooling rate in the bottom region, and the columnar crystals were completely broken, forming a large number of equiaxed crystals that were uniformly distributed. The better densification and homogeneity of the tissue in the bottom region indicate that the 1200 W power has a very significant effect on the bottom of the cladding layer.

Figure 3c shows that the top region of group c is still dominated by columnar crystals, but the grain size increases significantly, and the orientation is weakened. At 1600 W, the melt pool temperature becomes too high, and the cooling rate is reduced. The growth rate of columnar crystals is slowed down, resulting in coarsening of the grains. In addition, the number of equiaxed crystals in the top region decreased compared to group b. This demonstrates that the high power weakened the top region. This indicates that the high power weakened the refining effect of the grains in the top region. The grains in the middle region are coarser, and the proportion of equiaxed crystals decreases. Compared to group b, the thermal gradient in the middle region is reduced. However, the increase of the melt pool temperature makes the columnar crystals dominate, and the refinement effect is not as good as that of group b. The phenomenon of grain coarsening indicates that the high power limits the grain fragmentation and the generation of equiaxed crystals, leading to the deterioration of the uniformity of the microstructure. The grain coarsening phenomenon is more significant in the bottom region, and the number of columnar crystals increases. At the same time, the blurring of grain boundaries shows the uncontrolled direction of grain growth. The high power reduces the cooling rate in the bottom region, resulting in the failure of columnar crystals to break up sufficiently and a decrease in the number of equiaxed crystals.

Figure 3d shows blurred boundaries between grains with poor orientation. This suggests that the high power of 2000 W causes the temperature of the molten pool to be excessively high. As a result, the cooling rate is further diminished, and the controlled growth of columnar crystals weakens. The microstructure in the top region is severely coarsened, with almost no equiaxed crystals being formed. In the middle region, the columnar crystals continue to grow, and obvious grain defects emerge. For example, fracture or porosity between grain boundaries. This indicates that the 2000 W power caused excessive heat input to the melt pool, leading to internal stress concentration and structural instability during crystal growth. Equiaxed crystals almost disappeared, and the organizational inhomogeneity was further exacerbated. The bottom region was accompanied by obvious metallurgical defects, such as holes and cracks. The high power resulted in an extremely low cooling rate at the bottom, the columnar crystals were not sufficiently broken; the grain growth was out of control; and the presence of defects further reduced the densification and wear resistance of the fusion cladding.

SEM analysis of the In718/(2Nb + 1SiC) coatings at different powers in Figure 4 shows that the 1200 W (S2) parameter significantly optimizes the microstructure and elemental distribution of the coatings. Point 1 and point 2 indicate the formation of Nb/MoC, with point 1 indicating the enrichment of carbides at grain boundaries, which plays a role in grain boundary strengthening. Compared with the conventional nickel-based composite coatings (e.g., In718/TiC system), the surface flatness of the S2 specimens was higher (roughness Ra = 0.8 μm). At the same time, the crack density (0.2 bars/mm²) is only one-fifth of that of the laser-fused In718/WC coating. This difference stems from the synergistic strengthening mechanism between the NbC/MoC reinforcing phase and the substrate in this study: on the one hand, the uniform dispersion of nanocarbides (82 nm) suppresses the grain boundary slippage, which reduces the wear rate by 58% compared with that of the pure Ni-based coatings; on the other hand, the in situ decomposition of SiC generates the reactive carbon combined with Nb/Mo efficiently, resulting in a better elemental distribution uniformity. Notably, the porosity of the coatings at the 1200 W parameter (0.6%) is close to the level of single-crystal high-temperature alloys (e.g., 0.4% for CMSX-4), suggesting that the process has the potential to rival expensive preparation techniques (e.g., e-beam melting) in achieving high densification. It provides a new idea for the industrialization of low-cost, high-performance tool coatings.

The distribution of basic elements and reinforcing elements on the surface of S3 samples is generally uniform, with local accumulation of particles. The reason for this is that the high laser power input leads to the high temperature of the molten pool, which enhances the fluidity. However, it also triggers local instability, and the particles tend to accumulate in the local area. At the same time, the surface of the morphology has porosity and slight roughness. This is due to the reduced cooling rate of the cladding layer due to the high laser power, which makes it difficult for the gases to escape completely during solidification. Point 3 and point 4 indicate the formation of Nb/MoC, which is enriched within the crystal and plays a role in dispersion strengthening. The distribution of matrix and reinforcement elements in the S4 samples is extremely heterogeneous and agglomerated. This is due to the increased convection in the melt pool under high power conditions. However, the mobility is out of control, and the reinforcing-phase particles are aggregated in localized areas during the cooling process. Point 4 indicates the enrichment of Nb/MoC within the crystal, followed by more severe segregation of Nb and Mo. In addition, the high power input overheats the melt pool, leading to the concentration of internal stresses during the cooling of the cladding layer. At the same time, the gas trapping phenomenon was intensified, and cracks and holes were formed.

3.2. Microhardness

Figure 5 illustrates the hardness display of the coating cross-section, where Figure 5a shows the microhardness distribution plot division of the coating cross-section into three regions. Substrate, heat-affected zone (HAZ), fusion zone (FZ). and coating region (coating), respectively. Where the substrate section is located near the lower region, with a gradual transition to the hardness of the coating. The heat-affected zone (HAZ) is located between the substrate and the fusion cladding and has a definite increase in hardness. The hardness of the coating area reaches the highest value and shows different distribution characteristics. The hardness distribution pattern shows that all the samples (S1, S2, S3, and S4) located in the substrate region have a low and nearly uniform hardness of about 236.5 HV, which indicates that the substrate hardness is not significantly affected during the coating melting process. The hardness begins to gradually increase as the coating is approached. In the HAZ region, a significant increase in hardness occurs, indicating that the thermal input has resulted in partial grain refinement and phase transformation of the matrix organization. The hardness reaches its maximum value rapidly after entering the coating region. This indicates that the microstructure and reinforcing phase particles of the composite coating play a significant strengthening role. The S2 coating has the highest hardness and a smooth hardness distribution, indicating the best coating quality and organization homogeneity at this power. The hardness of the S3 and S4 coatings is slightly lower than that of the S2, and the distribution fluctuates. This may be related to the insufficient mobility of the melt pool or the local organizational inhomogeneity caused by too high power. S1 has the lowest hardness, which may be related to the insufficient dispersion of the reinforcing phase particles or the coarsening of the grains due to the insufficient melting power.

By comparing the average microhardness in Figure 5b, it can be seen that the average hardness of the substrate for all the samples is 236.5 HV, which is much lower than the hardness of the coated area. This indicates that the coating significantly increased the hardness of the material surface. Among the multiple sets of coating hardness data, the S2 specimen had the highest hardness. The S3 specimen was slightly lower than the S2 specimen, which is a preliminary indication that the higher power may have resulted in a slight decrease in tissue uniformity. The S4 specimen showed a significant decrease in hardness, which may be due to coarsening of the tissue or an increase in the number of defects caused by the high power. The S1 specimen was close to the S3 specimen but was slightly lower than the S2 specimen, which suggests that the low power may have resulted in the formation of a denser coating, but that there may not have been sufficient dispersion of the reinforcing phase, resulting in a limited increase in the hardness. resulting in limited hardness enhancement. Therefore, in terms of the margin of error, the S2 sample has the smallest fluctuation in hardness, indicating that the coating organization is the most homogeneous at this power, and the margin of error for the S4 sample is slightly larger, suggesting that the high power conditions may lead to localized areas of uneven organization or increased defects within the coating. The main reason for the formation of the above-mentioned hardness enhancement of the coatings can be attributed to the Nb- and SiC-reinforced phase particles in the composite coatings acting as diffuse reinforcement in the coating region. The overall hardness of the coating was significantly increased. The fluctuation of the hardness in the coating area may be related to the distribution of the reinforcing phase particles and the flow stability of the molten pool during the melting process. The fluctuation of S2 is the smallest, which indicates that the organizational homogeneity is the best under the melting conditions. While S4 fluctuated more, which may be due to the defects and organizational inhomogeneity triggered by the melt pool instability under high power conditions.

3.3. Tribological Properties

Figure 6 demonstrates the specimen wear data for different specimens located in the sliding and static wear states, respectively. From Figure 6a, it can be seen that the substrate has the highest overall coefficient of friction with an average value of 0.609. The curve fluctuates widely. The COF rises rapidly in the initial stage and then fluctuates continuously in the range of 0.5 to 0.6. It shows a large instability. This indicates that the surface hardness and wear resistance are relatively low. During the sliding process, more serious adhesive wear occurs on the surface, resulting in a higher and unstable coefficient of friction. The COF value of the S1 specimen is significantly lower than that of the substrate, reaching 0.503. However, there are still some fluctuations. The curve rises rapidly in the initial stage and then tends to stabilize. The friction coefficient of the S2 specimen rises rapidly in the initial stage to nearly 0.7 and then stabilizes at about 0.407 with less fluctuation. The curve shows high smoothness, indicating stable friction performance. It indicates that when the input power is 1200 W, the heat input and cooling rate of the molten pool reach a better balance, and the coating densification and hardness are effectively improved. Meanwhile, the hardness of the coating surface is increased, which reduces the inhomogeneity during sliding contact and makes the friction coefficient more stable. The friction coefficient of the S3 specimen rises rapidly to about 0.4 in the initial stage. However, with the increase of sliding time, the curve fluctuation increased, and the average friction coefficient was slightly higher than that of S2. At this time, the coating densification and hardness were still high. However, there may be local agglomeration of reinforcing phase particles or fusion coating defects, which may have a certain adverse effect on the overall friction performance. The friction coefficient of the S4 specimen rises rapidly to close to 0.5 in the initial stage, and then the curve fluctuates greatly. The overall performance showed low stability, and the average COF rose back to 0.52. At this time, the decrease in surface hardness and tissue defects triggered adhesive wear during sliding, leading to an increase in the coefficient of friction. From Figure 6b, it can be seen that the microstructure, surface homogeneity, and hardness of the coatings significantly affected the friction performance under different power conditions. The average friction coefficient of the S1 specimen is 0.503, which indicates the poor densification of the coating. The lowest average friction coefficient of the S2 specimen is 0.407. The average friction coefficient of the S3 specimen rises back to 0.444, which shows a certain performance degradation.

The average friction coefficient of the S4 specimen rises back to 0.52, which indicates the worst friction performance. The curves of S1 and S4 have large fluctuations, which reflect the high surface inhomogeneity of the coating. The higher The frictional contact during sliding is unstable. The curves of S2 and S3 are smoother, especially S2, which shows the best stability. This indicates that the homogeneity and densification of the coating organization are optimal under this power condition.

Figure 7 demonstrates the three-dimensional morphological wear state of different specimens. From the figure, it can be seen that the wear surface of the substrate exhibits the deepest wear depth (−67 μm). And the wear surface has a rough contour with more pits. The wear area has the largest extent, and the wear width is larger. The substrate lacks a reinforcing phase and has a low surface hardness. It is prone to adhesive wear and plastic deformation during friction, resulting in a larger wear depth and width. Meanwhile, the surface roughness indicates that the substrate fails to effectively resist the external force during friction, resulting in a large amount of material flaking. The hardness of the coating prepared at 800 W power is improved, but the melt pool temperature is low. The densification and organizational homogeneity of the coatings were poor, and there were still more microscopic defects. The increase in hardness reduces the wear depth, but the wear resistance is not yet optimal. At 1200 W, the wear depth further reduces to −35 μm, and the wear surface is flatter and more uniform. This shows that the wear resistance of the coating has been significantly improved. It can be seen that the densification and organizational homogeneity of the coatings are optimal at 1200 W. The distribution of the reinforcing phases is uniform, which improves the wear resistance of the coatings. The uniform distribution of reinforcing phases improves the hardness and wear resistance of the coating. This indicates that the optimization of the surface organization effectively reduces the adhesive and cutting wear, thus reducing the depth and width of wear. The depth of wear of specimen S3 increases slightly to −40 μm, and the flatness of the wear surface decreases compared to that of S2. Specimen S4 shows a further increase in wear width. The wear depth increases to −50 μm, and the wear surface is uneven, and the roughness increases. Figure 7b wear profile analysis shows the deepest and widest wear profile of the substrate In718. The wear depth and width of the S1 specimen are reduced compared to the substrate, but the profile is still steeper on both sides, indicating an improvement in the wear resistance of the coating. S2 has the shallowest and narrowest wear profile, showing the smoothest ‘V’ type depression and the best wear resistance. Samples S3 and S4 have deepened wear profiles compared to S2, and the wear resistance of the coatings decreased significantly. In Figure 7d, the wear rate comparison shows that the substrate has the highest wear rate (about 1.8 × 10⁻³ mm³/N-m), which indicates the worst wear resistance, while S2 has the lowest wear rate (about 0.8 × 10⁻³ mm³/N-m), which is the best wear resistance, and the wear rates of S1, S3, and S4 specimens are significantly lower but still higher compared with that of the substrate. The base material is still high. The trend of wear rate is consistent with the trend of wear depth and width. This indicates that the microstructure uniformity and hardness of the coating are the key factors determining the wear resistance.

According to Figure 8, there is a significant amount of oxidation wear debris and grooves in the In718 matrix. There are also significant cracks and delamination on the surface. Particle build-ups, surface irregularities, and damage are localized. It is also worth noting that the base material is subjected to high friction and high temperatures. The surface undergoes a serious oxidation reaction, forming oxidized abrasive debris. At the same time, due to the low hardness of the surface, adhesive wear and plastic deformation are prone to occur, resulting in large areas of material spalling from the surface. For the sample S1 surface, there are large wear particles (wear debris), as well as a certain degree of spalling (delamination). The more pronounced distribution of furrows and particles indicates the presence of friction particle cutting during the wear process. Cracks and rougher surfaces are still present in the wear area. The reason for this phenomenon is that the surface is still susceptible to friction particle impact due to the low hardness of the coating at 800 W. This results in the formation of furrows and abrasive chips. The occurrence of spalling indicates that the surface wear resistance has improved, but the coating is not sufficiently dense. The surface of the S2 sample was the smoothest, and the number of peeling pits was significantly reduced. At this time, the furrow becomes shallow, and the particle wear phenomenon is reduced. The surface homogeneity is better, indicating that the densification of the coating tissue is significantly improved. S3 specimens have deeper furrows, and particles and wear debris are more obvious. Therefore, it is suggested that the higher power of 1600 W may have led to agglomeration or coarsening of the reinforcing phase particles in the coating, which in turn reduced the homogeneity of the coating. The deepened furrows indicate a decrease in surface hardness and an increase in the cutting action of the abrasive particles, leading to increased spalling of the material during friction.

The surface of the S4 specimen was the most severely damaged, with large areas of cracks. Furrows are deep and dense, and particles and wear debris are present in large numbers. There is a significant increase in spalling, and the surface roughness is at its maximum. This indicates that the 2000 W power is too high, and defects, such as cracks, may appear inside the coating tissue, leading to a significant decrease in the wear resistance of the coating. At the same time, the expansion of surface damage led to increased crack extension. The formation of large spalling pits indicates a decrease in the coating bond strength. It can be seen that the coating densification and uniformity are best when the laser power is 1200 W. Cracking and spalling are significantly reduced, and the wear resistance is optimized. The wear mechanism was dominated by slight abrasive wear, and the coating bond strength and hardness were excellent. Inputs of 1600 W and 2000 W resulted in uneven local tissue distribution, significant cracking and spalling phenomena, and a significant decrease in wear resistance. For this reason, the introduction of coatings significantly improved the wear resistance. In particular, the coating under 1200 W conditions can effectively inhibit crack extension and spalling phenomena and prolong the service life of the material.

In Figure 9, there are a large number of cracks, spalling areas, and oxidation wear debris on the surface. It shows that there is a significant amount of oxygen (O) on the surface, and the oxidative wear phenomenon is severe. The high concentration distributions of O and C indicate that the main wear mechanisms are oxidative wear and adhesive wear. Meanwhile, due to the insufficient hardness, the surface is prone to forming deep furrows and large material spalling, resulting in the worst overall wear resistance. The hardness of the coatings was improved at 800 W power, and the cracks and spalling were reduced, but the uneven distribution of the reinforcing phases led to the more pronounced granular wear and furrowing phenomena. The distribution of O and C indicates that oxidation wear is still dominant, and the coating wear resistance is improved but still poor. The 1200 W power shows the best wear resistance. The surface is flat, and cracks, spalling, and furrows have largely disappeared. The reinforcing phases (Nb and Mo) are uniformly distributed, oxidative wear is significantly reduced, and the coating hardness and densification reach an optimal state so that the wear mechanism is mainly dominated by slight abrasive wear. The 1600 W power coating begins to show crack expansion and spalling pits, and the surface uniformity decreases. The furrows deepen. Localized agglomeration of the reinforced phase and increased distribution of O and C indicate that the wear resistance of the coating begins to decline. The wear mechanism is a combination of abrasive and oxidative wear. The most serious damage to the surface of the coating is observed at 2000 W, with a large number of cracks and spalling pits. The distribution of reinforcing phases was uneven or even missing, and oxidative and adhesive wear increased significantly. The surface roughness increases, and the furrow depth is the largest, leading to a significant decrease in the overall wear resistance of the coating, close to the performance of the substrate. As a result, the wear resistance of the coatings showed a tendency to increase and then decrease with the power change. Among them, the 1200 W coating (S2) showed the best performance, indicating that the medium power condition can achieve uniform distribution of the reinforcing phase and high densification of the coating. While both low power (S1) and too high power (S3 and S4) resulted in a decrease in wear resistance due to tissue inhomogeneity or increased defects.

4. Conclusions

In this paper, the laser cladding technique was selected to study the organization and wear resistance of In718/(2Nb + 1SiC) composite coatings for tools under different laser process parameters. In this study, a comprehensive analysis of morphology, crystal structure, X-ray diffraction pattern, microhardness, and friction properties of the composite material of In718/(2Nb + 1SiC) composite coating material was obtained by selecting different process parameters. The main conclusions are as follows:

(1)

The micro-morphological and XRD analyses showed that the SiC/NbC/MoC multiphase synergistic strengthening effect was achieved at the 1200 W (S2) parameter. The coating porosity under this power parameter was reduced by 62% compared with the baseline value, and the volume fraction of the reinforced phase reached 35.7%. This finding provides a theoretical basis for the parameter setting of industrial laser cladding equipment. By precisely controlling the input power in the range of 1150–1250 W, the in situ synthesis of the reinforced phase can be achieved while ensuring the densification of the coating. It is suitable for aerospace turbine blade repair and other scenarios that require strict coating densities.

(2)

When the laser input power is 800 W and 1600 W, respectively, the surface inhomogeneity of the composite coating is high, and the frictional contact in the sliding process is unstable. Their mean friction coefficients were 0.503, 0.407, 0.444, and 0.52, respectively. When the input power is 1200 W, the wear curve is smoother when the specimen coating organization uniformity and densification are optimal.

(3)

The trend of wear rate is consistent with the trend of wear depth and width, indicating that the microstructure uniformity and hardness of the coating are the key factors determining the wear resistance.

(4)

By analyzing the cross-section hardness values of the coating, the average microhardness of S1, S2, S3, and S4 samples are 353.1 H, 375.4 H, 355.2 H, and 337.8 H.