Effect of Nb on MC (M:Ti,Nb,Cr) precipitation and tribological behavior of TiC/Ni-based coating fabricated via plasma transferred arc

Figure 2(a) shows the XRD diffraction pattern of the coatings. The diffraction peaks of the XRD pattern corresponding to γ-(Fe, Ni) [PDF:33-0397], TiC [PDF:32-1383], (Nb, Ti)C [PDF:47-1418] and M₂₃C₆. The diffraction peak of phase Nb was not detected in the diffraction pattern. This confirms that Nb melts and reacts with other elements in the molten pool during PTA cladding. With the increase of Nb content from 0 wt% to 10 wt%, the diffraction peak of M₂₃C₆ changes dramatically. At the diffraction angle of about 38° in figure 2(a), there is a diffraction peak of M₂₃C₆ in sample P2, which is not detected in sample S1 and S3. Generally, the forming elements of M₂₃C₆ mainly come from Ni50 alloy powder. Hence, the Nb content affects the behavior of alloying elements precipitated from the molten pool (discussed later).

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Figure 2(b) shows the diffraction pattern of the local magnification mode of XRD from 43° to 45°, and the diffraction peak corresponds to γ-(Fe, Ni). According to $2{\rm{d}}\,\sin \,\theta ={\rm{n}}\lambda ,$ the smaller 2θ values indicate a larger interplanar crystal distance of the corresponding crystal planes [24]. This means that the γ-(Fe, Ni) crystal lattice is distorted during non-equilibrium metallurgical process. Table 2 shows the 2θ location and FWHM of the γ-(Fe, Ni) peak. It can be seen that there is a large deviation in the γ-(Fe, Ni) diffraction peak of sample S3. In other words, with the increase of Nb content in the coating, the matrix was transformed correspondingly. One of the most important roles in the phase evolution of coatings is the precipitation and solid solution of Cr₂₃C₆. From the XRD pattern, Cr is present in the coating in two forms. When the coating contains 5 wt% Nb, Cr is precipitated as carbide (Cr₂₃C₆). When the coating contains 10 wt% Nb, Cr is mainly dissolved in the austenitic matrix.

The SEM images of the cross-sectional of the coatings are shown in figure 3 and EDS results corresponding to feature areas in figure 3 are shown in table 2. The unmelted TiC particles, primary dendritic TiC, and spherical TiC are observed in figures 3(a)–(a1). According to previous literature, the formation of primary dendritic formed by recrystallization of TiC was harmful to the mechanical properties of MMCs [19]. In the current work, one goal is to avoid the formation of coarse primary dendritic formed by recrystallization of TiC by adding Nb element. The sizes of in situ TiC particles are approximately 3 ∼ 5 μm shown in figure 3(a1). As shown in figures 3(b)–(b1), the number of coarse primary dendrites decreased and two kinds of carbide with different contrast appeared when the Nb content is 5 wt%. Furthermore, a reaction layer formed at the interface between TiC and matrix. Combined with EDS results (EDS points C, D, and E), the content of each element in regular geometry particles is basically the same as that in the reaction layer. It is worth noting that, the carbides with low contrast are carbides with high Cr content, which can be determined as M₂₃C₆ according to EDS analysis results. This corresponds to the diffraction peak at about 38 degrees in the XRD pattern. Figures 3(c)–(c1) shows the SEM images of sample S3. It can be seen that the characteristic microstructure of sample S3 underwent further evolution compared with sample S1 and S2. In addition to the carbides generated in situ, a network microstructure distributed along grain boundaries is also found in figure 3(c). The mechanism of Nb content on the microstructure evolution of the original coating (sample S1) is discussed as follows:

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As known, the microstructure of carbide reinforced metal matrix composite materials can be affected by carbon content strongly [25]. In this study, the source of carbon provided by Ni-based alloy powder and partially recrystallized TiC. The increase of Nb content means more carbon consumption in the molten pool. Zhao et al studied the relationship between carbon content and tribological behavior of NbC/Fe-based coatings. They found that the evolution of NbC microstructure was closely related to C content [25].

Based on the analysis of microstructure features, the solidification and precipitation process in the molten pool is as follows: In the initial powder system, the melting point of Nb is lower than that of TiC but higher than that of Ni-based powder. TiC with a high melting point (3200 °C) did not melt completely in the molten pool, but partially dissolved at the liquid-solid interface. This is confirmed by the SEM image in figures 3(a)–(a1). Dissolved TiC provides a Ti source and a C source. TiC and NbC have similar lattice lattice constants, both Ti and Nb are strong carbide forming elements [26]. In the spontaneous nucleation process of MC alloy carbide, niobium and titanium atoms can replace each other in their respective positions in the lattice, thus precipitating out in the form of (Nb, Ti)C carbides. Hence, regular geometry (Nb, Ti)C particles were precipitated out first in the first stage. As the molten pool temperature decreases, the element Cr formed by carbide with the lower melting point was precipitated by M₂₃C₆ phase at the liquid-solid interface. Finally, unprecipitated carbide-forming elements remain in the solid solution, and a small amount of eutectic may be precipitated by the eutectic reaction. As can be seen from figure 3(a1), point B), a small amount of eutectic is distributed around the matrix.

It should be noted that the phase of M₂₃C₆ is significantly present in sample S2 shown in Figure 4, but not observed in sample S1 and sample S3. This corresponds well with the XRD results. It is well known that the presence of V, Cr, Mo elements often reduces the Ms point [27]. Similarly, the addition of Nb elements will also cause changes in the precipitation behavior of various substances in the system. With the increase of Nb content to 10 wt%, Nb and Ti atoms consume a large number of C atoms, resulting in carbon-poor behavior in the molten pool. Therefore, even if the pool system has a lower Ms point, the carbide-forming elements remain in the matrix or are expelled from the grain boundary after austenitizing due to insufficient nucleation of carbide. However, the addition of too much Nb into the Ni-based melt results in an increase in the viscosity of the molten pool [28]. This greatly reduces the overall rheological property of the molten pool. At the same time, in the process of plasma transferred arc scanning, there is a large temperature gradient at the top of the molten pool, which leads to the surface tension caused by Marangoni convection [9]. This, in turn, creates a thermal capillary flow on the surface of the molten pool, further destabilizing the rapid solidification process. Figure 3(c1) shows a high-power SEM image of in situ (Nb, Ti)C particles. It can be seen that carbide particles have a core–shell structure. EDS results show that the core has a high content of Ti elements, and the shell has basically the same content as the reaction layer.

There are three distribution forms of Nb atoms in the molten pool. ①: Nb is a strong carbide forming element, forms nucleation with Ti, C atoms which from the unmelted TiC to form (Nb, Ti)C in the molten pool. ②: The surface of the unmelted TiC particles in the molten pool provides an attachment point for free Nb atoms. Subsequently, a solid solution composed of Ti, Nb, Cr, and C is formed by a dissolution/precipitation mechanism [17]. It should be emphasized that EDS results in table 3 (EDS point: C-D, H-F,) show that the content of elements in the reaction layer is very similar to that in MC particle. There is basically no difference between the composition of the reaction layer and MC particles. It is also shown from the side that the formation of the reaction layer and the particles precipitate out at the same solidification stage. And the distribution of the reaction layer on the surface of the unmelted TiC particles is random. The formation of in situ film has a strong interface with the matrix, which enhances the bonding force between the TiC and the matrix (discussed later). Theoretically, both the microstructure of the remelted TiC and the interfacial adhesion of the unmelted TiC are improved. ③: The coating prepared by plasma transferred arc is formed by rapid non-equilibrium solidification process. Some carbide-forming element atoms fail to fully nucleate, but are solids dissolved in the matrix. This will cause certain component segregation. As the temperature is reduced to phase transition point, in order to reduce the shear resistance during the phase transition, alloying elements is discharged to the matrix [29]. Hence, a small amount of Nb exists in the matrix or grain boundary in the form of a solid solution.

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Figure 5 shows the SEM image of the carbide characteristic morphology. In order to obtain the fine microstructure image of the coating, the etching time of the sample was 120 s before SEM observation. It can be seen from figure 5(a) that the in situ formation of (Nb, Ti)C particle is octahedral structure. Figure 5(b) depicts the independently grown petal-shaped M₂₃C₆. Figure 5(c) shows the core–shell structure and in situ morphology of (Nb, Ti)C particles. The addition of Nb strongly regulates the microstructure of the original coating. Fine dispersion distribution of (Nb, Ti)C particles, petal-like M₂₃C₆ and unmelted TiC particles with reaction layer are used as the reinforcing phase of the coating. The austenitic matrix provides good toughness, which is conducive to the construction of perfect metal matrix composites. The goal of making full use of the performance of each phase is achieved as far as possible.

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Figure 6 shows the average microhardness values of the coatings with different Nb contents. It can be seen that sample S2 has the highest microhardness in all samples. The main enhancement contribution of microhardness is derived from the formation of the hard phase [4]. According to XRD result and SEM image, finely dispersed (Nb, Ti)C and M₂₃C₆ particles are the main source of contribution to the increase in microhardness of sample S2. When the Nb content reaches 10 wt% in the coating (sample S3), the microhardness is lower than that of the sample S1. Based on the above description, the reason for the decrease in microhardness of sample S3 is mainly attributed to the insufficient carbon content, which leads to insufficient nucleation of carbide forming elements.

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Generally speaking, the friction coefficient of the metal materials varies with the normal load. Hashempour et al studied the relationship between different normal loads and friction coefficients of metal materials. They found that the friction coefficient of the material increases as the load decreases [30]. The obvious COF decrease by load increase can be explained by the formation of mechanically mixed layer (MMLs) [31, 32]. When the loading force is applied on the material surface, the plastic deformation of the material results in the work hardening behavior. When the load is not enough to induce the random deformation of the material, the friction coefficient between the pairs is larger. When the load causes the material to undergo plastic deformation, the existence of the plastic deformation zone makes the pair of pairs relatively smooth, so it shows a lower friction coefficient. Cao et al also studied the relationship between different normal loads and friction coefficients of metal materials. They found that as the normal load increased from 20 N to 100 N, the corresponding coefficient of friction decreased from about 0.8 to 0.65 [33]. Figure 7 shows the friction coefficient (COF) of different samples. It can be seen from figure 7 that the three coatings exhibit different friction coefficient variation behaviors under 25 N loading. As the 25 N load cannot make sample S1 and sample S2 have enough plastic deformation behavior, they finally show a higher friction coefficient behavior. However, the friction coefficient of sample S3 decreases sharply. It is possible that work hardening behavior occurred in the friction process with S3, and the coating surface was compacted, thus showing a lower friction coefficient.

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Figure 8 shows the wear volume loss of each samples. With the increase of Nb content to 10 wt%, the wear volume and friction coefficient increase, which is unfavorable to the tribological properties of the coating. The main reason for this result is the carbon poor behavior in the coating, so the carbides formed are metastable structures, which can be expressed as (Nb, Ti)C_(1−x), x < 1.

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In order to discuss the tribological behavior of coatings more specifically, figure 9 shows the SEM morphology of the wear scar after dry sliding friction. The most two important factors that influence the mechanical property of the MMCs are the matrix microstructure and the carbide particles. It is well known that carbide reinforcing phase carries and resist the effects of applied loads during wear test due to their high hardness until failure occurs. Meanwhile, the soft matrix binds the carbide particles and transfers the external load to individual carbide particles [17]. Therefore, the good compatibility between the matrix and the hard carbide particles is the key to the material's full play. The matrix will be quickly damaged when the TiC peels off and fails to exist the effects of applied loads. This result is harmful to improving wear resistance. Figure 9(a) shows that the unmelted TiC particles remain relatively shape at the junction of the worn and unworn areas, but cracks around the TiC particles appear at the junction of TiC and the matrix. This indicates that under the action of the load, the interface with low bonding strength fails before the TiC particles fail. The end result will cause the TiC particles to flake off. Figure 9(b) shows the TiC particles in the worn area are broken but no cracks appear at the interface. Figure 9(c) further shows that when the unmelted TiC is broken, there is still a continuous binder phase at the interface between TiC and the matrix. Figure 9(A) shows the EDS results of marked points. The results show that the continuous binder phase is a solid solution composed of Ti, Nb, Cr, and C. This also proves that the film formed surrounding TiC particles has the effect of enhancing the interface bonding force. In this work, the microcracks first originated from the stress concentration tip of the dendritic structure and the defect of the reinforcement/matrix, and it is more likely to induce microcracks than the dendritic structure at the defect of the reinforcement/matrix. The solid solution film attached around the TiC particles has the function of the in situ method, that is, the interface of the reinforcement/matrix is a stronger interface. It is worth mentioning that, as far as the authors know, this method has a certain reference value for laser cladding technology.

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