Microstructure and Tribological Properties of Lubricating-Reinforcing Laser Cladding Composite Coating with the Ti₂SC-Ti₂Ni Mosaic Structure Phase

Abstract

Lubricating-reinforcing composite coatings were successfully prepared on Ti6Al4V using laser-clad Ti6Al4V/Ni60/Ni-MoS₂ mixed powders with different Ni-MoS₂ contents (25, 35, and 45 wt.%), and their microstructure and tribological properties were studied. The reinforcing phase TiC, Ti₂Ni, and the lubricating phase Ti₂SC were in situ precipitated while Ti₂SC and Ti₂Ni formed a mosaic coherent structure within the above three coatings. In the 25 and 45 wt.% Ni-MoS₂ coatings, the microstructure distribution uniformity of the coatings was not effectively improved by the Ti₂SC-Ti₂Ni mosaic structure phase due to the lower or higher content of Ti₂SC. In the 35 wt.% Ni-MoS₂ coating, the forming quality of the coating was the best due to an appropriate amount of the uniformly distributed Ti₂SC-Ti₂Ni mosaic structure phase. Furthermore, the microhardness of the coatings gradually decreased as the amount of Ni-MoS₂ increased. In the 35 wt.% Ni-MoS₂ coating, due to the uniformly and diffusely distributed Ti₂SC-Ti₂Ni mosaic structure phase, the stable lubricating-reinforcing mosaic structure transfer composite films were formed during the progress of the friction and wear tests, which led to the optimal worn surface evenness and quality, the anti-friction and the wear resistance properties compared with the Ti6Al4V, 25 and 45 wt.% Ni-MoS₂ coating.

1. Introduction

Titanium alloys are extensively applied in aerospace, national defense, petrochemical, and other key industrial fields due to their low density, high strength performance, and excellent corrosion resistance [1,2]. As a commonly used material for aircraft and its engines, Ti6Al4V is widely applied in the 4th and 5th flap tracks of Boeing 737NG series aircraft and in the fan blades, compressor disc, low compressor blades of CFM56 engines, which is helpful to improve the fuel economy of aircraft engines and the transient response capability of aircraft in the air [3,4]. However, the service life of components made of Ti6Al4V alloys is greatly shortened under harsh friction and wear conditions for its low hardness and poor wear resistance, which restricts the application of titanium alloys [5].

In recent years, laser cladding technology has been widely used in the preparation of wear-resistant and self-lubricating coatings on the surface of Ti6Al4V alloys as a new surface modification method [6,7]. Compared with other thermal deposition methods such as atmospheric plasma spraying (APS), flame spraying (FS), high-velocity oxygen-fuel (HVOF), detonation spray coating (DSC), etc., laser cladding provides a superior metallurgical bonding with the substrate, which means less spalling or cracking behavior and denser coatings [8,9,10,11,12,13,14,15]. In the study of Zhou et al. [16], self-lubricating anti-wear composite coatings were fabricated by laser cladding with Ni60/TiC/WS₂ mixed powders. It was found that the tribological properties of the coatings were decreased due to the fact that the intermetallic compound phase was accumulated at the top region of the coatings. Ke et al. [17] successfully fabricated Ti₂SC self-lubricating wear resistant coatings while the results showed that the lubricating phase Ti₂SC was segregated in the coatings, resulting in serious fluctuation on the mechanical properties of the coatings. Liu et al. [18] prepared Ni-based self-lubricating wear resistant composite coatings. The study indicated that the lubricating phase and reinforcing phase in the coating were separated, which led to significant fluctuation in the tribological properties of the coating. From the above studies, it could be found that the lubricating phase and reinforcing phase were agglomerated and segregated in the coatings, which resulted in an uneven distribution of the microstructure and serious fluctuation in the mechanical performance. However, the interface relationship between the lubricating phase and reinforcing phase and its effect on the microstructure and tribological properties of coatings have not been studied.

In conclusion, the lubricating-reinforcing composite coatings with three different proportions of Ti6Al4V/Ni60/Ni-MoS₂ mixed powders on Ti6Al4V alloy were fabricated using the coaxial powder-fed laser cladding technology. The microstructure, phase structure characteristic, and the tribological performance of the coatings were emphatically researched. The mosaic growth mechanism between the reinforcing phase Ti₂Ni and the lubricating phase Ti₂SC in the coatings was revealed according to the Bramfitt two-dimensional lattice misfit theory. This research is expected to lay a theoretical and experimental foundation for the extended application of the self-lubricating wear resistance laser-clad layer on a Ti6Al4V alloy.

2. Materials and Methods

2.1. Materials

Ti6Al4V was selected as the substrate material and cut into cuboid specimens with the dimensions of 60 mm × 40 mm × 10 mm. The optimized design for the coating material system is listed in

Table 1. The design of the coating cladding material system.

No.	Coating	Composition of the Cladding Materials/wt.%
		Ti6Al4V	Ni60	Ni-MoS2
1	25 wt.% Ni-MoS2 coating	35	40	25
2	35 wt.% Ni-MoS2 coating	35	30	35
3	45 wt.% Ni-MoS2 coating	35	20	45

, and the selection criteria of the coating materials were as follows: Ti6Al4V is beneficial to the improvement in the wettability between the precipitation phases and the matrix [19]. Ni, B, C elements offered by Ni60 can react with the Ti element to form intermetallic compounds and ceramic reinforcing phases with high hardness that can effectively increase the microhardness and tribological performance of the coatings [8]. MoS₂ is a typical thermal-decomposed lubricant that can offer the S element to react with the Ti element and in situ synthesize various lubricating phases, which leads to a better anti-friction performance [20].

The coaxial powder-fed laser cladding technology requires the fine flowability of the cladded powders to achieve the synchronous supply of cladding materials. Therefore, spherical Ti6Al4V, Ni60, and quasi-spherical Ni-MoS₂ (the mass ratio of Ni to MoS₂ was 3:1, and the wrapping rate of Ni coated MoS₂ is 100%) were selected in this material system, and all powder sizes ranged from 30 to 120 μm with the SEM morphologies shown in Figure 1. The chemical compositions of Ti6Al4V and Ni60 are listed in

Table 2. The chemical compositions of Ti6Al4V and Ni60.

Materials	Elements (wt.%)
	Ti	Al	V	Fe	C	N	O	H	Ni	Cr	B	Si
Ti6Al4V	Bal.	5.5–6.75	3.5–4.5	≤0.3	≤0.08	≤0.05	≤0.02	≤0.015	-	-	-	-
Ni60	-	-	-	≤17.0	1.0–2.0	-	-	-	Bal.	14–18	2.5–4.5	3.5–4.5

.

2.2. Preparation and Laser Cladding Process

Before the laser cladding experiment, the Ti6Al4V specimen was subjected to sandblasting treatment to remove the contaminants and oxide layer on the surface, roughen the specimen surface, and enlarge the absorption rate of the Ti6Al4V to the laser energy. The sandblasted specimen was ultrasonically cleaned in absolute ethanol for 5 min. The proportional cladding powders were mixed by mechanical stirring for 12 h, then placed in a constant-temperature vacuum dryer and heated to 80 °C for 10 h.

The laser cladding process was conducted on a German TruDisk4002 coaxial powder feeding laser system. Helium was adopted as the feeding gas and argon was used as the shielding gas to create an inert atmosphere and avoid oxidation. The process parameters of the laser cladding are listed in

Table 3. The process parameters of the laser cladding.

Process Parameter	Value	Process Parameter	Value
Power	1100 W	Feeding gas flow rate	7.0 L/min
Scanning speed	400 mm/min	Shielding gas flow rate	11.0 L/min
Spot diameter	3.0 mm	Laser focus	16.0 mm
Powder flow rate	1.4 r/min	Overlap ratio	50%

, and the schematic diagram of the laser cladding process is shown in Figure 2.

2.3. Microstructure Characterization and Property Tests

The cladded specimen was cut into 13 mm × 10 mm × 4 mm samples along the cross-section perpendicular to the cladding direction. The samples were prepared as metallographic samples with a diameter of 22 mm, then eroded for 15 s with Kroll’s reagent (with the volume ratio of HF:HNO₃:H₂O = 1:1:40). The test equipment and its parameters are listed in

Table 4. The test equipment and test parameters.

No.	Test Items	Test Equipment	Test Parameters
1	Phase composition	X’Pert-Pro MPD X-ray diffraction (XRD)	Scanning velocity: 6°/min, diffraction range: 20–80°
2	Microstructure observation	Hitachi S-3000N scanning electron microscopy (SEM)	-
3	Phases analysis	Oxford INCAPentaFET-X3 energy dispersive spectrometry (EDS)	-
4	Element distribution	JXA-8530F field-emission electron probe X-ray microanalyzer (EPMA)	Electronic optical system resolution: 3 nm
5	Microhardness	KB30SR-FA digital microhardness tester	Test load: 500 g, dwell time: 12 s
6	Tribological properties	RTEC MFT-5000 tribometer	Normal load: 50 N, WC counterpart diameter: 6 mm, sliding time: 90 s
7	Coating wear 3D profiles	Solarius AOP non-contact white-light interferometer	Scanning speed: 50 μm/s

.

3. Results and Discussion

3.1. Cross-Sectional Morphology

Figure 3 shows the penetrant test results of the coatings with different Ni-MoS₂ additions (25, 35, and 45 wt.%). As shown in Figure 3, the 35 wt.% Ni-MoS₂ coating showed the best forming quality. The cracks presented an interlaced netlike structure in the 25 and 45 wt.% Ni-MoS₂ coatings and the number of them was significantly more than that of the 35 wt.% Ni-MoS₂ coating. The possible reason for the above phenomenon is that under the same process conditions, the microstructure uniformity of the 35 wt.% Ni-MoS₂ coating was the best. Moreover, the cracks gradually decreased along the cladding direction, which could be explained by the fact that the thermal stress caused by the temperature difference was higher in the initial melting zone of each track than that in the final melting zone during solidification of the molten pool [21].

The cross-sectional SEM morphologies of the coatings are illustrated in Figure 4. It is clear in Figure 4 that the coatings with different Ni-MoS₂ additions presented a wavy transition. The above experimental phenomena can be explained as follows: the laser energy density at the Gaussian heat source center is higher than that of the surrounding area, which leads to the fact that the coating at the Gaussian heat source center was thickened and the surrounding area became thinner. As can be seen from the figures, there were few pores in each of the three coatings and an obvious longitudinal through crack appeared in the 25 wt.% Ni-MoS₂ and 45 wt.% Ni-MoS₂ coatings. The 35 wt.% Ni-MoS₂ showed the best forming quality, which matched the penetrant test results of the coatings (Figure 3).

3.2. Phase Composition

Figure 5 shows the XRD results of the coatings. As shown in Figure 5, the composite coatings mainly consisted of TiC, Ti₂Ni, Ti₂SC, and the matrix α-Ti. The ceramic phase TiC and the intermetallic compound phase Ti₂Ni are generally deemed as reinforcing phases that can significantly increase the microhardness and the wear resistance of the laser-clad layers [22,23]. Ti₂SC is a ternary layered ceramic M_n+1AX_n phase with high elastic modulus, shear modulus, and melting point as well as superb yield strength and thermal stability [24]. Furthermore, Ti₂SC is often used as a self-lubricating phase in the coating because of its typical layered structure [25]. According to the XRD results, the diffraction peak strength of Ti₂SC gradually increased with the rise in the Ni-MoS₂ addition in the coatings, indicating that the content of Ti₂SC was increased, which was beneficial to improving the anti-friction properties of the coating. Furthermore, the XRD diffraction pattern did not show the corresponding MoS₂ diffraction peak, which implied that most of MoS₂ had been thermally decomposed during the laser cladding process.

3.3. Microstructural Characterization

Figure 6 shows the microstructure of the coatings with different Ni-MoS₂ additions. Figure 6(a1–a5,b1–b5,c1–c5) shows the microstructure morphologies of the 25, 35, and 45 wt.% Ni-MoS₂ coatings with different magnification, respectively. It can be seen from Figure 6(a5,b5,c5) that the coatings mainly included the petal-like and granular phase A1, matrix phase A2, strip-like phase A3, and irregular bulk-like phase A4. It should be noted that with the increase in Ni-MoS₂ in the cladding material system, the number of the strip-like phases gradually increased (as shown in Figure 6(a3,b3,c3)). In the 45 wt.% Ni-MoS₂ coating, the strip-like phases were agglomerated and segregated, which aggravated the cracking tendency and matched the experimental phenomenon in Figure 3c.

The EDS results of the aforementioned phases are listed in

Table 5. The EDS results of the phases in the coatings.

Phase	Proportion	Ti	S	C	Ni	Al	Cr	V	Mo	Si
A1 (TiC)	Wt.%	58.68	-	12.30	6.52	3.63	-	5.02	13.85	-
	At. %	45.52	-	39.96	4.13	1.37	-	3.66	5.36	-
A2 (α-Ti)	Wt.%	70.15	-	-	6.84	5.66	3.04	4.86	9.45	-
	At. %	72.12	-	-	5.89	10.61	2.96	4.06	4.36	-
A3 (Ti2SC)	Wt. %	69.45	19.05	5.74	5.02	-	-	-	-	0.74
	At. %	54.85	22.65	18.22	3.26	-	-	-	-	1.02
A4 (Ti2Ni)	Wt. %	58.14	1.15	2.66	35.27	2.78	-	-	-	-
	At. %	55.81	1.65	10.18	27.62	4.74	-	-	-	-

. From Table 5, it can be deduced that the petal-like and granular phase A1 mainly consisted of Ti and C elements with an atomic ratio of approximately 1:1. The continuous matrix phase A2 mainly consisted of Ti with a mass ratio over 70%. The strip-like phase A3 mainly consisted of Ti, S, and C elements with an atomic ratio of approximately 2:1:1. The irregular bulk-like phase A4 primarily consisted of Ti and Ni elements with an atomic ratio of approximately 2:1. Thus, combined with the XRD analysis, it could be verified that the petal-like and granular phase A1 was the ceramic reinforcing phase TiC, the continuous matrix phase A2 was α-Ti, the strip-like phase A3 was the lubricating phase Ti₂SC, and the irregular bulk-like phase A4 was the intermetallic compound phase Ti₂Ni.

It can be seen from the growth characteristics and distribution law of the synthetic phase in the coating (Figure 6(a3,c3,c3)) that a low amount of precipitated Ti₂SC was distributed sporadically in the 25 wt.% Ni-MoS₂ coating. With the continuous rise in the Ni-MoS₂ content, the amount of Ti₂SC in the 35 wt.% Ni-MoS₂ coating was increased significantly and the distribution of Ti₂SC was more uniform than that in the 25 wt.% Ni-MoS₂ coating. For the 45 wt.% Ni-MoS₂ coating, the precipitation amount of Ti₂SC reached the peak, which led to large-area agglomeration of the precipitated Ti₂SC.

Due to a large number of the Ni element provided by Ni60 and Ni-MoS₂ in the molten pool, the brittle intermetallic compound Ti₂Ni was synthesized in large quantities and mainly distributed in the form of a continuous large-area irregular block phase in the coatings, which substantially increased the cracking risk of the coatings. It is worth noting that in the three coatings prepared with different additions of Ni-MoS₂, the continuous growth behavior of Ti₂Ni was blocked by in situ synthesized Ti₂SC (Figure 6(a4,b4,c4)), which greatly reduced the cracking sensitivity of the coatings. In the 25 wt.% Ni-MoS₂ coating, the continuous large-area growth of Ti₂Ni was not effectively destroyed due to the low precipitation amount of Ti₂SC. In the 35 wt.% Ni-MoS₂ coating, the continuous distribution of Ti₂Ni was effectively obstructed due to the uniform distribution of an appropriate amount of Ti₂SC. In the 45 wt.% Ni-MoS₂ coating, the segregation growth of Ti₂Ni was not effectively blocked due to the agglomeration behavior of excessive Ti₂SC.

Figure 6(a4,a5,b4,b5,c4,c5) demonstrates that the Ti₂SC-Ti₂Ni mosaic structure phases were generated by the behavior of Ti₂SC blocking the continuous growth of Ti₂Ni, which were formed in the 25, 35, and 45 wt.% coatings. Combining the structural characteristics of the mosaic structure phase, it can be inferred that if Ti₂SC and Ti₂Ni were the relation of coherent attachment growth, the synergistic stirring effect of Ti₂SC and Ti₂Ni could be formed during the cyclic convection process of the molten pool to prevent the segregation and agglomeration behavior of Ti₂SC and Ti₂Ni due to the density difference, which effectively increased the distribution uniformity of the coating microstructure and reduced the performance fluctuations of the coatings at the same time. Therefore, this paper would lucubrate the formation mechanism and the phase boundary relationship of the Ti₂SC-Ti₂Ni mosaic structure phase as well as the coating tribological properties.

Based on the aforementioned analysis, it was speculated that the following reactions (1)–(5) might occur in the molten pool [26,27].

Ti + C = TiC

(1)

Ti + S = TiS

(2)

TiC + TiS = Ti₂SC

(3)

Ti + Ni = TiNi

(4)

Ti + TiNi = Ti₂Ni

(5)

During the reaction process aforementioned, the ceramic reinforcing phase TiC with the highest melting point would first precipitate from the molten pool L. According to the Ti–C binary alloy phase diagram, TiC precipitated through a uniform grain of L → residual molten pool (L1) + TiC. Attributed to the high liquidus slope of TiC, the molten pool was easily supercooled, leading to the fact that the primary TiC grew up in the way of dendrite and formed the residual molten pool (L1) + dendritic TiC while the residual molten pool (L1) offered TiC with abundant space for growth, thus the TiC finally grew to become coarse dendrites. However, in the present experiment, TiC mainly existed in the form of petal-like or granular phases, demonstrating a noticeable refinement in the ceramic phase TiC in the coating, which was conducive to the reduction in the coating cracking susceptibility and had a fine grain strengthening effect in the coating. The refinement of TiC can be explained as follows: TiC would precipitate preferentially from the molten pool for its high melting point. As the temperature decreased, the lubricating phase Ti₂SC would synthesize in the molten pool when the saturation of Ti and S elements on the TiC surface met the conditions of Ti₂SC precipitation (reaction Equation (3)). It could be seen from the reaction synthesis of the lubricating phase Ti₂SC that the ceramic reinforcing phase TiC was swallowed and consumed as a reactant. Thus, the growth of TiC was inhibited, which caused TiC to exist in the form of petal-like and granular phases in the coating.

In terms of the precipitation process of Ti₂SC, it is believed that the lubricating phase Ti₂SC was synthesized by the reaction between the ceramic reinforcing phase TiC in the solid state and Ti and S elements in the liquid state (i.e., TiC(solid) + TiS(liquid) = Ti₂SC, rather than the solid-solid reaction of TiC(solid) + TiS(solid) = Ti₂SC). The main reasons were as follows. First, in the progress of the reaction between the solid-state TiC and solid-state TiS, there was a relatively high energy barrier for the synthesis of Ti₂SC [28]. Second, Ti and S elements in the liquid state were diffused and converged easily on the surface of the solid-state TiC, creating a favorable reaction environment for the precipitation of Ti₂SC. Furthermore, in the Ti-rich molten pool, once solid-state TiS was precipitated, it would then react with the Ti element enriched in the surrounding and form interfering phases of Ti₂S. Nevertheless, no corresponding diffraction peaks for TiS and Ti₂S appeared in the XRD, which demonstrated that TiS and Ti₂S did not effectively synthesize in the coating. A similar phenomenon was also found in the work of Zhai et al. [25] during the preparation process of Ti₂SC using laser-cladding technology, where they discovered that Ti₂SC was synthesized by the reaction between solid-state TiC and the liquid-state Ti and S elements in the Ti–S–C material system.

As the intermetallic compound, Ti₂Ni is a typical coating reinforcing phase with high microhardness and a complex face-centered cubic (FCC) structure [29]. According to the Ti–Ni binary alloy phase diagram, TiNi can produce Ti₂Ni in a Ti-rich environment [30]. The precipitation process of Ti₂Ni is as follows. TiNi begins to precipitate in the molten pool when the temperature drops to 1310 °C, then undergoes a peritectic reaction with Ti to produce Ti₂Ni when the temperature drops to 984 °C. The aforementioned precipitation process of Ti₂Ni is shown in reactions (4) and (5).

As shown in Figure 7, the micro-area of the 35 wt.% Ni-MoS₂ coating was analyzed by EPMA due to the better forming quality.

As seen in Figure 7a,b,d,j, the distribution of Ti, S, and C elements was consistent in the strip-like phase of the coating. Combined with the XRD results, it was determined that the strip-like phase was Ti₂SC. It can be seen from Figure 7c that a small amount of the Si element was enriched in Ti₂SC, which was consistent with the EDS results of Ti₂SC. From Figure 7a,e,j and the XRD analysis results, it can be confirmed that the irregular bulk-like phase in the coating was mainly Ti₂Ni, which was in accordance with the EDS results of the Ti₂Ni. Moreover, the strip-like lubricating phase Ti₂SC was embedded in Ti₂Ni, forming a Ti₂SC-Ti₂Ni mosaic phase that matched to the aforementioned SEM (Figure 6(a4,b4,c4)). It can be seen from Figure 7a,d,j that TiC mainly existed in the form of the granular phase and petal-like phase with a size of approximately 1–3 μm, which was in accordance with the structural characteristics of TiC (Figure 6). It can be seen from Figure 7a,f–i that Al, Mo, V, and Cr dissolved in the matrix as solid solvable elements and played the role of solid solution strengthening in the coating.

3.4. Formation Mechanism of the Ti₂SC–Ti₂Ni Mosaic Structure Phase

To further explore the interface relationship of the Ti₂SC-Ti₂Ni mosaic structure phase, the two-dimensional lattice misfit theory was adopted. It was defined as Equation (6) and the corresponding variable definitions are listed in

Table 6. The variable definition of the Bramfitt lattice misfit.

No.	Variable	Definition
1	(hkl)s	a low-index crystal face of the nucleation substrate
2	[uvw]s	a low-index crystal direction of (hkl)s
3	(hkl)n	a low-index crystal face of nucleating phase
4	[uvw]n	a low-index crystal direction of (hkl)n
5	d[uvw]n	an interatomic spacing of [uvw]n
6	d[uvw]s	an interatomic spacing of [uvw]s
7	θ	an angle between the [uvw]s and [uvw]n

. The two-dimensional lattice misfit theory states that the composite structure phase can be formed when the lattice misfit is under 12%. Moreover, when the lattice misfit is lower, the consumed two-phase mismatch energy is lower and more atoms are matched on the two-phase mismatch interface, which leads to the stronger binding force of the two phases in the composite structure phase [31].

$${\mathsf{\delta }}_{{\left(hkl\right)}_n}^{{\left(hkl\right)}_s}=\frac{1}{3}{\displaystyle \sum }_{i=1}^3\left[\frac{\left|d_{{\left[uvw\right]}_s}{}^i\cos \mathsf{\theta }-d_{{[uvw]}_n}{}^i\right|}{d_{{\left[uvw\right]}_n}{}^i}\right]\times 100$$

(6)

Due to the fact that the melting point of Ti₂SC is higher than that of Ti₂Ni [29], Ti₂SC would precipitate before Ti₂Ni. Therefore, for the Ti₂SC-Ti₂Ni mosaic structure composite phase, Ti₂SC was the substrate phase and Ti₂Ni was the nucleation phase. The unit cell models of Ti₂SC and Ti₂Ni are illustrated in Figure 8 with their lattice parameters [27,32] shown in

Table 7. The lattice parameters of Ti₂SC and Ti₂Ni.

Crystal	Crystal System	Space Group	Lattice Parameters/Å
			a	b	c
Ti2SC	Hexagonal	P63-MMC	3.2100	3.2100	11.2000
Ti2Ni	Cubic	FD-3M	11.3193	11.3193	11.3193

.

To obtain a better interfacial matching relationship of Ti₂Ni and Ti₂SC, the low-index crystal faces were selected for the calculation of the two-dimensional lattice misfit. The schematic of the crystallographic relationship and the results of the lattice misfit between Ti₂SC and Ti₂Ni are shown in Figure 9 and

Table 8. The calculated lattice mismatch between Ti₂SC and Ti₂Ni.

Matching Face	${\mathbf{Ti}}_2\mathbf{SC}\left(01\overline{1}0\right)/{\mathbf{Ti}}_2\mathbf{Ni}\left(001\right)$
[uvw] Ti2SC	[$\overline{1}$000]	[$\overline{1}$001]	[0001]
[uvw] Ti2Ni	[110]	[020]	[$\overline{1}$10]
θ (°)	0	0.653	0
$d{\mathrm{Ti}}_2\mathrm{SC}$(Å)	3.210	11.651	11.200
$d{\mathrm{Ti}}_2\mathrm{Ni}$(Å)	3.025	11.437	11.029
δ (%)	3.18

, respectively. As shown in Table 8, the lattice misfit between the crystal-face (001) of Ti₂Ni and the crystal-face ($01\overline{1}0$) of Ti₂SC was 3.18%, indicating that Ti₂SC was a significantly effective heterogeneous-nucleus substrate phase for the Ti₂Ni, which resulted in the formation of a coherence interface between Ti₂SC and Ti₂Ni. The above precipitation behavior and formation mechanism of the Ti₂SC-Ti₂Ni mosaic structure composite phase is shown in Figure 10.

It can be seen from the above analysis that an appropriate amount of in situ precipitated Ti₂SC in the 35 wt.% Ni-MoS₂ coating resulted in the fact that the continuous growth behavior of Ti₂Ni was effectively blocked by the Ti₂SC-Ti₂Ni coherent mosaic structure phase, greatly increasing the distribution uniformity and reducing the cracking sensitivity of the coating, which was exactly the main reason for the optimal forming quality of the 35 wt.% Ni-MoS₂ coating.

3.5. Microhardness

Figure 11 shows the microhardness distribution of the coatings with different Ni-MoS₂ additions. The microhardness of the 25%, 35%, and 45% Ni-MoS₂ coatings was increased by 108.41%, 96.23%, and 80.04%, respectively, compared with Ti6Al4V (374.26 HV_0.5). The fluctuation range of the 35 wt.% Ni-MoS₂ coating microhardness was significantly smaller than that of the 25 and 45 wt.% Ni-MoS₂ coatings, which can be explained by the microstructure distribution uniformity of the coating being the best in the 35 wt.% Ni-MoS₂ coating.

Studies have shown that the main factors influencing the microhardness of the coatings are the content and particle size of the reinforcing phase TiC, Ti₂Ni, and the addition amount of Ni60 [33,34]. Furthermore, the microhardness of TiC was significantly higher than that of Ti₂Ni [35]. It can be seen from the above analysis and the material system in this paper that TiC was progressively refined, and its precipitation amount as well as the addition amount of Ni60 gradually decreased with the increase in the Ni-MoS₂ content, which contributed to the phenomenon that the microhardness of the coatings illustrated a trend in the gradual decrease.

3.6. Tribological Properties

Figure 12a shows the average friction coefficients of Ti6Al4V and the 25, 35, and 45 wt.% coatings at room temperature. Under the same wear test conditions, the average friction coefficient of the 35 wt.% Ni-MoS₂ coating was 8.7%, 0.9%, and 4.8% lower than that of Ti6Al4V and the 25 and 45 wt.% Ni-MoS₂ coatings, leading to the conclusion that the 35 wt.% Ni-MoS₂ coating exhibited the best anti-friction performance. The above phenomenon can be explained as follows. In the 25 wt.% Ni-MoS₂ coating, a small amount of the Ti₂SC-Ti₂Ni composite phases were generated, which cannot form a long-term effective lubricating film in the process of pair grinding, so the reduction in the friction coefficient was limited. In the 35 wt.% Ni-MoS₂ coating, an appropriate amount of Ti₂SC-Ti₂Ni composite phases were evenly distributed, forming a stable lubricating film between the worn surface and the grinding ball, which led to a further decrease in the friction coefficient. In the 45 wt.% Ni-MoS₂ coating, a viscoelastic resistance was formed during the wear process due to the segregation of an excessive amount of Ti₂SC, which resulted in an increase in the friction coefficient [17]. Studies by [16,25] revealed that the self-lubricating phase Ti₂SC could effectively improve the anti-friction performance of the coatings, while combined with the above analysis, it can be found that the anti-friction effect of Ti₂SC was weaker than that of graphite and MoS₂ and so on [6].

Figure 13 shows the worn surface roughness measured perpendicular to the rubbing direction. It can be seen that the worn surface roughness of the 25 wt.% Ni-MoS₂ coating was obviously decreased and there were less convex peaks and pits in the worn surface compared with Ti6Al4V. In the 35 wt.% Ni-MoS₂ coating, the roughness was further decreased and its fluctuation was the least, which led to the optimal worn surface evenness. In the 45 wt.% Ni-MoS₂ coating, the amplitude fluctuation of the roughness was increased and the worn surface evenness declined.

The wear rate is a frequently-used parameter to evaluate the material wear resistance [36], which is defined as:

W = V/(L × S)

(7)

where W = wear rate (mm³·N⁻¹·m⁻¹); V = wear volume (mm³), L = normal load (N); S = sliding distance (m). In this paper, the friction and wear tests were carried out in the configuration of reciprocation with a total sliding distance of 3 m and a normal load of 50 N. Measured by the white-light interferometer, the wear volume of Ti6Al4V and the 25 wt.%, 35 wt.%, and 45 wt.% Ni-MoS₂ coatings were 8.66 × 10⁻² mm³, 6.15 × 10⁻² mm³, 6.14 × 10⁻² mm³, and 8.07 × 10⁻² mm³, respectively. According to Equation (7), the wear rate of the 35 wt.% Ni-MoS₂ coating was 4.09 × 10⁻⁴ mm³·N⁻¹·m⁻¹, approximately 29.09%, 0.20%, and 23.40% lower than that of Ti6Al4V (5.77 × 10⁻⁴ mm³·N⁻¹·m⁻¹), the 25 wt.% Ni-MoS₂ coating (4.10 × 10⁻⁴ mm³·N⁻¹·m⁻¹), and the 45 wt.% Ni-MoS₂ coating (5.38 × 10⁻⁴ mm³·N⁻¹·m⁻¹) (Figure 12b). In conclusion, the wear resistance of the 35 wt.% Ni-MoS₂ coating was the optimal.

Morphologies of the worn surfaces at different magnifications of the coatings are shown in Figure 14. It can be seen from Figure 14(a1,a2) that due to the poor wear resistance of Ti6Al4V, large flaking and plastic deformation appeared on the worn surface, showing typical adhesive wear and fatigue wear. There were some visible bulk-like spalls and grooves as well as several debris with sizes of approximately 5 μm distributed on the worn surface of the 25 wt.% Ni-MoS₂ coating, as shown in Figure 14(b1,b2). Combined with the morphology and the roughness value of the worn surface, the 25 wt.% Ni-MoS₂ coating presented a better worn surface evenness than that of Ti6Al4V due to the self-lubricating effect of Ti₂SC.

For the 35 wt.% Ni-MoS₂ coating shown in Figure 14(c1,c2), it was noted that the worn surface of the coating was smoother, there were a few granular debris and slender furrow scratches but no bulk-like spalls on it. This was mainly because the uniformly distributed Ti₂SC-Ti₂Ni mosaic structure composite phases had a dragging and binding effect during the wear process. On one hand, as a high-hardness “skeleton”, the intermetallic compound Ti₂Ni could effectively prevent the plastic deformation and block spalling of the coating. Meanwhile, in the progress of the friction and wear tests, the coherent interface Ti₂SC-Ti₂Ni was extruded and crushed by WC concurrently, which formed the Ti₂SC-Ti₂Ni lubricating-reinforcing mosaic structure transfer composite films. The Ti₂Ni debris and Ti₂SC debris with a typical layered structure would extend and fill the whole coating surface simultaneously, which reduced the coating friction coefficient and prevented further wear of the coating.

The worn surface morphology of the 45 wt.% Ni-MoS₂ coating is shown in Figure 14(d1,d2). There was a lot of debris agglomerated on the coating, which caused an obvious decrease in the worn surface quality compared with the 35 wt.% Ni-MoS₂ coating. The possible reason was that the local stress concentration was caused by a large amount of Ti₂SC agglomeration, which accelerated the wear failure of the coating. In addition, it can be seen from Figure 14 that the wear mechanism of the 25 wt.% Ni-MoS₂ coating was abrasive wear and slight adhesive wear, and the 35 and 45 wt.% Ni-MoS₂ coating presented typical characteristics of abrasive wear.

In summary, the anti-friction, the wear resistance, the worn surface evenness and its fluctuation, and the worn surface quality of the 35 wt.% Ni-MoS₂ coating were the best compared with the Ti6Al4V and the 25 and 45 wt.% Ni-MoS₂ coatings.

4. Conclusions

(1) Ti-based lubricating-reinforcing composite coatings were fabricated on Ti6Al4V using Ti6Al4V/Ni60/Ni-MoS₂ mixed powders with different additions of Ni-MoS₂. The phase composition of the coatings with different Ni-MoS₂ contents remained unchanged, which mainly consisted of the ceramic phase TiC, the intermetallic compound phase Ti₂Ni, the lubricating phase Ti₂SC, and the matrix phase α-Ti, meanwhile Ti₂SC and Ti₂Ni presented a mosaic structure.

(2) The Bramfitt two-dimensional lattice misfit between the crystal-face (001) of Ti₂Ni and the crystal-face ($01\overline{1}0$) of Ti₂SC was 3.18.%, indicating that the coherence interface between Ti₂SC and Ti₂Ni was effectively formed in the coatings.

(3) The 35 wt.% Ni-MoS₂ coating showed the optimal forming quality because an appropriate amount of the diffusely distributed Ti₂SC-Ti₂Ni mosaic structure phase significantly increased the microstructure distribution uniformity and reduced the cracking sensitivity of the coating.

(4) The microhardness of the coatings gradually decreased with the increasing Ni-MoS₂ content from 25 to 45 wt.%, and the fluctuation range of the 35 wt.% Ni-MoS₂ coating microhardness was the smallest compared with the other coatings. Because the Ti₂SC-Ti₂Ni lubricating-reinforcing mosaic structure transfer composite films were formed in the 35 wt.% Ni-MoS₂ coating, the coating showed the optimal worn surface evenness, the worn surface quality, the anti-friction, and the wear resistance properties. In addition, the experimental results showed that the anti-friction effect of Ti₂SC was weaker than that of graphite and MoS₂.