Effect of V Content on Microstructure and Properties of TiNbZrV_x Medium-Entropy Alloy Coatings on TC4 Substrate by Laser Cladding

Abstract

In order to improve the wear resistance of titanium alloy and apply it to the high-speed train brake disc, TiNbZrV_x (x = 0, 0.2, 0.4, 0.6, 0.8) refractory medium-entropy alloy coatings were prepared on Ti-6Al-4V (TC4) substrate. The effect of V content on the microstructure, mechanical properties, and friction and wear properties of the coatings was studied. TiNbZrV_x coatings achieved good metallurgical bonding with the substrate, forming BCC and B2 phases and AlZr₃ intermetallic compound (IMC). From TiNbZr coating to TiNbZrV_0.8 coating, V promotes element segregation and new phase formation, which decreased the average grain size from 85.055 μm to 56.515 μm, increased the average hardness from 265.5 HV to 343.4 HV, and reduced the room temperature (RT) wear rate by 97.8%. However, the ductility of the coatings decreased from 15.7% to 5.8% because the grain boundary precipitates changed the dislocation arrangement, and the tensile fracture mode changed from ductile fracture to brittle fracture. Abrasive wear was the main wear mode at RT, and adhesive wear and oxidation wear were the main wear modes at elevated temperature. The COF at elevated temperature was lower than that at RT, because a large number of friction pair components were transferred to the coating surface at high temperature and were repeatedly rolled to form a dense film, which played a certain lubricating role.

1. Introduction

The brake disc is a key component of a high-speed train, and the lightweighting of brake discs plays a crucial role in enhancing train performance [1]. Traditional cast steel or forged steel brake discs are relatively heavy, making it difficult to meet the lightweighting requirements for high-speed train brake discs. Titanium alloys have a density of only 60% that of cast steel, offering low density, high specific strength, and good corrosion resistance [2], making it an ideal material for lightweight brake discs. However, its low thermal conductivity, poor wear resistance, and inadequate high-temperature oxidation resistance limit its application on brake discs [3]. Surface modification of titanium alloys can enhance their surface wear resistance and high-temperature performance [4], extending their service life [5] and realizing the application of titanium-based brake discs in high-speed trains.

Laser cladding, characterized by high precision, rapid cooling rates, and strong bond strength between coating and substrate [6], has found extensive application in titanium alloy surface modification. However, Ti forms brittle IMCs with most metals, such as Al, Cr, and Ni [7,8,9], at elevated temperature, resulting in poor bonding with the substrate and susceptibility to cracking. This represents one of the challenges faced in laser cladding on titanium alloys.

Primary methods for producing wear-resistant, high-temperature coatings on titanium alloy surfaces via laser cladding include the following: adding solute elements such as Cu and Mo [10,11]; incorporating rare earth elements and oxides like CeO₂ and Y₂O₃ [12,13]; introducing hard particles like TiC, WC, and TiB [14,15,16]; preparing medium-entropy alloy (MEA) or high-entropy alloy (HEA) coatings [17,18]; and combining these approaches [19,20]. Among these methods, MEAs and HEAs composed of three or more metals can form single FCC, BCC, or HCP structures [21], achieving metallurgical bonding with the substrate and reducing cracking. They also exhibit excellent mechanical properties, corrosion resistance, wear resistance, and oxidation resistance due to the combined effects of the high-entropy effect, hysteresis diffusion effect, and lattice distortion effect [22]. Refractory medium/high-entropy alloys (RMEA/RHEA) formed by high-melting-point elements such as Ti, Nb, Zr, V, Mo, Ta, and W show outstanding high-temperature resistance [23]. Among the above elements, V and Nb belong to β-phase stabilizing elements [24] of titanium alloy, which can reduce the α↔β phase transformation temperature and promote the formation of β-phase. β-Ti has a low elastic modulus, high tensile strength, better strength-toughness balance, and can maintain high strength at elevated temperatures [25].

In recent years, numerous researchers have employed laser cladding technology to fabricate RHEA coatings on titanium alloy surfaces, achieving notable progress in enhancing wear resistance and high-temperature oxidation resistance. Zhang [26] prepared AlNbTiVB coating on TC4 substrate by laser cladding, increasing the coating hardness to 1.34 times that of the substrate while reducing the wear rate by 29.76%. Deng [27] laser-clad CoCrFeNiMo_0.2 coatings on TC4 substrate, achieving a microhardness of 900 Hv. A dense Cr₂O₃ oxide layer formed on the coating surface, enhancing corrosion resistance and high-temperature oxidation resistance. Lu [28] prepared NbMoTaVTi coating on TC4 substrate. The phase composition of the coating was found to be BCC+HCP. The hardness, elevated temperature oxidation resistance, and corrosion resistance of the coating were higher than those of the substrate. In summary, most current studies on MEA/HEA coatings for titanium alloys focus on enhancing the mechanical properties, such as hardness, tensile properties, and wear resistance, with limited investigation into the formation mechanism of solid solution and precipitated phases, and the transition between them.

In this study, TiNbZrV_x (x = 0, 0.2, 0.4, 0.6, 0.8) coatings were prepared on TC4 substrate by laser cladding. The effects of V content on the microstructure, especially the transformation of phases and IMCs, and their effect on mechanical properties were studied. This study provides essential data support for the surface modification of the titanium alloy brake disc.

2. Experiment

2.1. Materials and Methods

In this study, TC4 plates with dimensions of 100 × 100 × 40 mm were used as the laser cladding substrate. The atomic ratio was converted into a weight fraction as shown in

Table 1. Chemical composition of substrate and powder (wt.%).

Samples	Ti	Nb	Zr	V	Al
Ti-6Al-4V	Bal	-	-	4	6.02
TiNbZr	20.63	40.05	39.32	0	-
TiNbZrV0.2	19.76	38.36	37.67	4.21	-
TiNbZrV0.4	18.96	36.80	36.14	8.1	-
TiNbZrV0.6	18.23	35.39	34.74	11.64	-
TiNbZrV0.8	17.55	34.07	33.44	14.94	-

, and the spherical Ti, Nb, Zr, and V powders (99%, 53–150 μm) were weighed using a high-precision balance (0.1 mg) and uniformly mixed by a dual motion mixer for 4 h with a rotation speed of 120 rpm to obtain TiNbZrVx (x = 0, 0.2, 0.4, 0.6, 0.8) powders.

Before experiments, powders were placed in a vacuum drying oven and dried at 80 °C for 2 h to remove moisture. The surface of the substrate was polished with 400# and 600# sandpapers and washed with absolute ethanol. Laser cladding experiments were conducted in the Ar gas glove box, with water and oxygen content controlled below 50 ppm to prevent high-temperature oxidation of powders. Process parameters were as follows: laser power was 3 kW, spot diameter was 4 mm, scanning speed was 210 mm/min, powder feeding speed was 12 g/min, lap ratio was 50% (4.5 mm), and shielding gas flow was 30 L/min. The thickness of the single cladding layer was about 0.85~0.9 mm. After each cladding layer was finished, they were cooled for 10 min, and then the next layer was clad. Four layers of TiNbZrV_x coatings were deposited onto the substrate using orthogonal deposition directions to reduce deformation of the substrate. All the coatings had no cracks on the surface and inside, and had good bonding with the substrate, with a total thickness of approximately 3.5 mm.

2.2. Microstructure Characterization

The coatings were processed into 10 mm×10 mm samples by wire cutting, the observation surface was polished with sandpapers and SiO₂ (50 nm) polishing solution until the surface was free of scratches, and etched with Keller’s reagent (HNO₃:HCl:HF:H₂0 = 2.5:1.5:1:95) for 30~40 s.

Microstructure and elemental composition were characterized by SEM and EDS. Phase analysis was performed by XRD at a scanning angle of 20°~80° and a scan step of 2°/min. Phase structure and crystal orientation were determined by EBSD and TEM testing.

2.3. Mechanical Property Characterization

Microhardness of the coatings was tested using a Vickers microhardness tester with a load of 9.8 N and a holding time of 15 s. Tests were conducted from the top to the bottom of the coatings at 0.1 mm intervals in the vertical direction.

The top surfaces of the coatings were machined into 48 mm standard tensile specimens with a thickness of 1.6 mm. The gauge length and width were 18 mm and 3 mm, respectively. Tensile properties were evaluated using a tensile testing machine at a tensile speed of 0.48 mm/min. Sampling locations and specimen dimensions are shown in Figure 1a,b. Each group of experiments was repeated three times.

2.4. Friction and Wear Performance Testing

Process the coating into standard friction block specimens with a diameter of 32 mm, as shown in Figure 1c. The friction pairs are copper-based powder metallurgy pins with a diameter of φ4.75 × 14 mm. The chemical composition is shown in

Table 2. Chemical composition of friction pair (wt.%).

Composition	Cu	Fe	CrFe	Graphite	MoS2	SiC	Others
Content	40~45	15~20	5~8	15~20	2~5	2~5	2~5

. Before the experiment, the contact surface of the pins was polished smoothly with 800# sandpaper. According to the provisional technical conditions of the EMU brake disc “TJ-CL3102019”, RT and 700 °C were chosen as test temperatures. The COF versus time curves of the coatings at RT and 700 °C were tested using a friction and wear testing machine.

To simulate the real braking process, pressure between the brake disc and brake pad, three friction pair pins were installed on the rotation shaft as shown in Figure 1d, and the foundation was raised so that the pins can contact the friction block with a load of 61 N, a contact pressure of 1.14 MPa.

Start the device and rotate the rotation shaft with a rotational speed of 300 r/min, and a test time of 30 min. After the test, the total sliding distance of the three pins is 2034.72 m. The experiment was carried out in an atmospheric environment. When the high temperature test was carried out, the specimens were placed in the heating furnace, heated by electric heating wire, and the system automatically measured the temperature. After reaching the specified temperature, the test began. After the test, the time–COF curve was obtained. 0~5 min belongs to the running-in stage, the COF wave was large, and the friction pair was consumed too much in 25~30 min, which may not reach the set contact pressure. Therefore, the data in these intervals were not considered, and the average COF in 5~25 min was counted. The profile of wear scars was measured by a white light interferometer, and the wear rate was calculated from the following equation:

$$\mathsf{\omega } = {\displaystyle \frac{\mathrm{V}}{\mathrm{N}·\mathrm{L}}}$$

(1)

where ω is the wear rate, V is the wear volume, N is the contact pressure between the friction pad and the friction pair, and L is the total distance the friction pair slides.

3. Results

3.1. Phase Analysis

Figure 2 shows the XRD patterns (XOY plane in Figure 1a) and binary phase diagram of TiNbZrV_x coatings. The coatings were mainly composed of the BCC phase. From the intensity change in diffraction peaks, it can be seen that the amount of BCC increased first and then decreased with the increase in V content, and the BCC content of TiNbZrV_0.4 was the highest. A small amount of B2 and AlZr₃ were detected in TiNbZrV_0.4~TiNbZrV_0.8. HCP and Laves phases calculated theoretically in Figure 2c did not appear. Figure 2b is the enlarged view of the main diffraction peak. The interplanar spacing of the crystal can be calculated using the Bragg equation:

$$\mathrm{d} = {\displaystyle \frac{\mathsf{\lambda }}{2\sin \mathsf{\theta }}}$$

(2)

where d is the interplanar spacing, λ is the wavelength of the X-ray, and θ is the diffraction angle.

As V content increased, diffraction peaks shifted to the right, increasing the diffraction angle θ. With the X-ray wavelength λ remaining constant, the interplanar spacing d decreased, indicating lattice contraction. This occurred because the atomic radius of V was smaller than that of Ti, Nb, and Zr, leading to a reduction in the lattice constant upon V incorporation.

3.2. Microstructure

Figure 3a shows the overall microstructure of the TiNbZrV_0.2 coating cross-section (YOZ plane in Figure 1a). Good metallurgical bonding is observed between the coating layers and between the coating and substrate with distinct layer boundaries. A small amount of unmelted Nb powder is present at the top.

According to rapid solidification theory [29], the microstructure of laser cladding coating at different positions is related to temperature gradient (G) and crystal growth rate (R), and the G/R ratio represents the stability factor at solid–liquid interface. Figure 3b presents the microstructure of the top of the coating, where the melt pool cooled rapidly, and the crystals grew fast, forming equiaxed crystals with a small size. Figure 3c,d shows the microstructure in the middle section of the coating. This region is composed of columnar dendrites, which grew in the vertical direction. Closer to the base, the temperature gradient increased, allowing dendrites to undergo repeated heating and cooling cycles for a longer time [30]. This extended thermal cycling facilitated crystal growth, resulting in increased crystal length and width. At the junction of coating and substrate, the temperature gradient is the greatest, and the crystal growth rate is the lowest. Crystal in this region grew in a planar crystal manner [31]. The microstructure of other coatings is similar to TiNbZrV_0.2.

Figure 3e shows the elemental composition variation in the TiNbZrV_0.2 coating from bottom to top along the deposition direction. Ti content decreased layer by layer, while Nb and Zr contents increased, with V content remaining essentially constant. Melt pool flow is the main reason for this variation. During laser cladding, part of the laser energy acted on the substrate, melting its surface to form a melt pool. Under the combined effects of the Marangoni effect and laser forces, convection occurred within the melt pool. Ti atoms, abundant in the substrate, migrated into the coating, while Nb and Zr were exchanged into the substrate. Al existed in the substrate rather than in the cladding powder, so it flowed into the substrate along with the melting pool during the laser cladding process, and the content decreased layer by layer. As the number of cladding layers increased, the concentrations of all three elements stabilized. Since the mass fraction of V in both the substrate and TiNbZrV_0.2 powder was approximately 4%, and the extent of elemental exchange was similar, the V content remained essentially unchanged.

Figure 4 shows SEM microstructures and element distributions of TiNbZrVx coatings. As shown in Figure 4(a1–e1), with the increase in V content, the morphology of dendrites changed, the primary dendrites became shorter and wider, while the number of secondary dendrite arms increased and the dendrite arms spacing decreased.

Figure 4(a2–e2) is the EDS mapping images of TiNbZrV_x coatings. Ti and V were uniformly distributed throughout the coatings, while Nb and Zr exhibited dendritic segregation. Zr content was higher in the inter-dendrite (ID) than within the dendrite regions (DR), while Nb primarily aggregated within the DR and was less distributed between them. This phenomenon arose because rapid cooling inhibited atomic diffusion, leading to non-equilibrium solidification [32]. During cooling, the high-melting-point Nb (2468 °C) first reached its solidification temperature and crystallized at the dendrite trunks. The low-melting-point Zr (1852 °C) crystallized more slowly and, thus, enriched the dendrite interstices [33].

Element contents of energy spectrums (Sp1–Sp6) shown in Figure 4(a2–e2) were counted, and the result is shown in

Table 3. Element content of ID and DR (wt.%).

Position	Samples	Ti	Nb	Zr	V	Al
ID	TiNbZr	27.19 ± 0.45	28.67 ± 2.26	42.53 ± 1.78	0.58 ± 0.11	1.03 ± 0.09
	TiNbZrV0.2	26.56 ± 0.60	24.91 ± 2.55	43.17 ± 2.48	4.19 ± 2.35	1.17 ± 0.15
	TiNbZrV0.4	25.37 ± 0.37	23.78 ± 0.73	41.53 ± 1.06	8.50 ± 0.30	0.82 ± 0.07
	TiNbZrV0.6	23.05 ± 0.32	23.61 ± 1.44	40.63 ± 1.66	12.15 ± 0.29	0.56 ± 0.08
	TiNbZrV0.8	20.31 ± 3.08	20.16 ± 1.78	42.46 ± 1.61	16.35 ± 3.49	0.72 ± 0.17
DR	TiNbZr	26.54 ± 1.06	37.71 ± 2.73	34.64 ± 1.91	0.50 ± 0.17	0.61 ± 0.09
	TiNbZrV0.2	26.22 ± 0.42	35.49 ± 0.39	34.24 ± 0.71	3.25 ± 0.19	0.80 ± 0.05
	TiNbZrV0.4	26.36 ± 0.49	36.52 ± 1.02	29.89 ± 0.51	6.73 ± 0.05	0.50 ± 0.04
	TiNbZrV0.6	23.56 ± 0.27	37.59 ± 2.25	28.63 ± 1.43	9.86 ± 0.69	0.36 ± 0.11
	TiNbZrV0.8	22.53 ± 0.24	37.46 ± 1.10	26.75 ± 0.86	12.94 ± 0.41	0.32 ± 0.05

. According to Table 1, the Nb and Zr content should be reduced with the increase in V content. However, in fact, the inter-dendritic Zr and the dendritic Nb contents were maintained at about 42% and 37%, respectively, and the inter-dendritic V and Al contents were also higher than those within dendrites. It can be seen that Zr, V, and Al segregated to ID, Nb segregated to DR during solidification, and V promoted the degree of segregation.

Figure 5(a1–e1) shows the inverse pole figures (IPF) of the coatings cross-section (YOZ plane in Figure 1a). It can be observed that the grain size decreased gradually from the bottom to the top of the coatings. As shown in Figure 5(a1), vertical columnar crystals with small orientation differences between adjacent grains were dominant in the central region (Layers 2–3) of TiNbZr, and a few equiaxed crystals appeared at the interlayer boundaries. These columnar crystals extended through both layers along the deposition direction. The top (Layer 4) of the coatings predominantly consisted of columnar crystals grown in a horizontal direction, which is the same direction as the laser cladding at the top. As shown in Figure 5(b1,c1), with increasing V content, grain size in Layers 2–3 decreased, and the grains were still mainly columnar crystals. The top layer exhibited a mixed structure of equiaxed and columnar grains, with larger orientation differences between adjacent grains. When V content exceeded 0.6, as shown in Figure 5(d1,e1), grain size further decreased significantly. Layers 2–4 of the coatings were predominantly composed of equiaxed grains, with grain sizes less than 250 μm.

Figure 5(a2–e2) and 5f shows the grain diameter distribution and average grain diameter variation in the coatings as a whole and in each layer. With increasing V content, the average grain diameter of the coatings decreased from 85.055 μm to 56.515 μm, representing a 33.55% reduction. The difference in grain size of each layer was also gradually reduced.

Figure 6 shows TEM test results of TiNbZr, TiNbZrV_0.4, and TiNbZrV_0.8 coatings. Element contents of grain boundaries (sp1) and crystals (sp2) were measured, and the results are shown in

Table 4. Element content of grain boundaries and crystals (at.%).

Samples	Position	Ti	Nb	Zr	V	Al
TiNbZr	sp1	28.56	29.79	39.49	0.33	1.82
	sp2	30.34	27.74	39.46	0.29	2.17
TiNbZrV0.4	sp1	31.74	9.80	49.07	7.75	6.17
	sp2	29.56	27.21	35.78	6.17	1.28
TiNbZrV0.8	sp1	9.82	14.94	49.71	23.48	2.05
	sp2	22.01	35.11	31.64	10.13	1.11

. As shown in Figure 6(a1), in the bright field image of TiNbZr, the contrast distribution was uniform, and no precipitate was observed. The SAED images in directions [110] and [100] showed no diffraction spots of precipitated phases. The contrast of the crystal was also relatively uniform, and no obvious element segregation occurred, indicating that the coating consisted of a single BCC phase. Aggregated precipitates appeared on grain boundaries of TiNbZrV_0.4, and SAED images of the precipitates are shown in Figure 6(b3). Two sets of diffraction spots appeared, and the weaker diffraction spot was the B2 phase. When V content increased to 0.8, as shown in Figure 6(c1,c4), grain boundaries widened, V, Zr, and Al segregated at grain boundaries, and SAED images at grain boundaries also appeared two sets of diffraction spots, indicating the existence of AlZr₃ IMC.

Element segregation at grain boundaries was similar to dendrite in Figure 4. With the increase in V content, Zr, V, and Al are segregated at grain boundaries. Combined with the binary phase diagram in Figure 2c, it can be determined that Zr segregation destroyed the order of the B2 phase and transformed it into AlZr₃ IMC [34,35].

3.3. Microhardness

Figure 7 shows the microhardness variation curves of TiNbZrV_x coatings along the vertical direction (YOZ plane in Figure 1a). With the increase in V content, the overall hardness of the coatings gradually increased. The average hardness of TiNbZr~TiNbZrV0.8 coatings was 265.5 ± 14.15 HV, 293.4 ± 18.50 HV, 299.1 ± 23.76 HV, 312.3 ± 27.43 HV, and 343.4 ± 35.81 HV, respectively.

The increase in hardness of the coatings was related to lattice shrinkage. The smaller the interplanar spacing is, the larger the energy required for dislocation movement is. With the increase in V content, the slip resistance of the coatings increased with the decrease in interplanar spacing, and finer grains hindered dislocation movement, resulting in higher microhardness [36].

For the same coating, the hardness gradually increased from the bottom to the top, and the average hardness of each layer from TiNbZr~TiNbZrV_0.8 coatings was shown in Figure 7b. The bottom of the coatings was subjected to cyclic heating and cooling during the upper cladding process, resulting in a heat treatment effect similar to normalizing, leading a certain softening of the microstructure, and the hardness was reduced to below 300 HV.

3.4. Tensile Properties

Figure 8 shows the stress–strain curves and corresponding mechanical properties of TiNbZr~TiNbZrV_0.8 tensile specimens. TiNbZrV_0.2 exhibited the highest ultimate tensile strength (UTS) at 985.2 MPa, while the tensile strengths of other specimens ranged between 890 and 960 MPa. The elongation at break of TiNbZr and TiNbZrV_0.2 specimens was 15.7% and 14.5%. For other specimens, elongation decreased with increasing V content. When V content reached 0.8, elongation dropped to 5.8%. Overall, TiNbZrV_0.2 shows the best mechanical properties with the highest UTS and a slightly lower elongation than TiNbZr.

Figure 9 displays the fracture morphology of the specimens. As shown in Figure 9(a3,b3), TiNbZr and TiNbZrV_0.2 exhibited ductile fracture, characterized primarily by dimples and microvoids. Tear ridges appeared on the fracture surfaces of TiNbZrV_0.2 and TiNbZrV_0.4 specimens, indicating localized plastic deformation [37]. The longer tear ridges on the TiNbZrV_0.4 specimen suggested enhanced shear tearing during fracture. As shown in Figure 9(c2–e2,c3–e3), fractures at TiNbZrV_0.4-TiNbZrV_0.8 exhibited both dimples and river patterns. With increasing V content, dimples decreased while river patterns increase. TiNbZrV_0.6 and TiNbZrV_0.8 displayed rock candy patterns and cleavage steps, showing intergranular fracture characteristics. The fracture mode was a mixed fracture of intergranular fracture and cleavage fracture.

Figure 10 shows EBSD images of the XOY plane in Figure 1a near the tensile fracture. The IPF images shown in Figure 10(a1,b1) revealed that the grains near the fracture of TiNbZr and TiNbZrV_0.2 were elongated along the tensile direction, and stress concentration occurred in the grains. The stress concentration resulted from plastic deformation of crystals through slip, resulting in residual stress with uneven distribution. As V content increased, as shown in Figure 10(c1–e1), both the extent to which the grains were elongated and the extent to which stress was concentrated decreased, indicating a decrease in the ability of the coating to plastically deform.

The KAM diagram displays orientation differences at grain boundaries and within crystals, which are related to dislocations and facilitate plastic deformation [38]. As shown in Figure 10(a2,b2), high-density dislocations appeared near the tensile fracture surfaces of TiNbZr and TiNbZrV_0.2, with average KAM values of 1.563° and 1.311°. The KAM values at grain boundaries exceeded those within grains, indicating that dislocations primarily accumulated at grain boundaries [39], which exhibited higher local plastic deformation capacity. As shown in Figure 10(c2,d2), the KAM values of TiNbZrV_0.4, TiNbZrV_0.6, and TiNbZrV_0.8 were lower than those of TiNbZr and TiNbZrV_0.2. High-density dislocations were concentrated near the fracture surface and decreased with increasing V content, indicating reduced dislocation density and diminished plastic deformation capacity.

3.5. Friction and Wear Behaviour

The instantaneous coefficient of friction (COF) can be calculated using the following equation:

$$\mathsf{\mu } = {\displaystyle \frac{\mathrm{F}}{\mathrm{N}}}$$

(3)

where μ is the COF, F is the sliding friction force between the friction pair and friction surface, and N is the contact pressure between the friction pair and friction surface.

Figure 11a shows the time-dependent instantaneous COF curves of TiNbZrVx coatings at RT. It can be observed that the COF of TiNbZrV_0.2, TiNbZrV_0.4, and TiNbZrV_0.8 exhibited significant fluctuations during the running-in stage (0–5 min). This phase represented the break-in process between the coating and the friction pair. After the break-in completion, the COF gradually stabilized. The curves of TiNbZr and TiNbZrV_0.6 showed a continuous increase throughout the process, with the COF rising by over 0.1. This increase was attributed to excessive wear of the friction pair, where the contact pressure N failed to reach the set value, leading to an elevated COF.

Figure 11b shows the time-dependent variation in the instantaneous COF of TiNbZrVx coatings at 700 °C. Similar to the friction experiment at RT, the friction curve exhibited significant fluctuations in the initial stage. Between 5 and 25 min, the overall fluctuation of the COF curve decreased as the V content increased. After exceeding 25 min, the curves of TiNbZrV_0.2 and TiNbZrV_0.6 exhibited an upward trend, with the increase in COF exceeding 0.1. The curves of TiNbZr, TiNbZrV_0.4, and TiNbZrV_0.8 also remained relatively stable in the later friction stage, indicating that these coatings caused minimal wear to the friction pair during the friction process.

The average COFs of each sample are calculated and shown in Figure 11c. The provisional technical conditions of the EMU brake disc “TJ-CL3102019” stipulate that the COF of the brake disc shall be within the range of 0.28~0.44 during braking. The average COFs of TiNbZr~TiNbZrV_0.6 at RT ranged from 0.28 to 0.44, meeting braking requirements. The COF of TiNbZrV_0.8 was 0.255 ± 0.013, which was below the normal operating requirement for brake discs. The COF of the coatings decreased gradually with the increase in V content at RT, which was caused by the increase in coating hardness and the decrease in plastic removal [40,41].

At 700 °C, the average COFs of the coatings were lower than at RT, initially increasing and then decreasing with increasing V content. The average COFs of TiNbZrV_0.2, TiNbZrV_0.4, and TiNbZrV_0.6 were 0.282 ± 0.038, 0.315 ± 0.014, and 0.288 ± 0.024, respectively, meeting braking requirements. However, the COFs of TiNbZr and TiNbZrV_0.8 were both below 0.28, failing to meet braking requirements. In summary, TiNbZrV_0.2~TiNbZrV_0.6 coatings met braking requirements for both ambient and elevated temperatures, with TiNbZrV_0.4 exhibiting the smallest overall coefficient fluctuation.

Figure 12 shows the surface morphology and the EDS spectra of the coatings and the friction pair’s main elements after the friction wear test at RT. As seen in Figure 12(a1–e1) and Figure 12(a2–e2), the wear surface morphology at RT was dominated by grooves and debris, which was the typical characteristic of abrasive wear [42]. Grooves on the coating surface were formed by the cutting action of hard particles such as Cr-Fe, SiC, and SiC from the friction pair. Debris appeared as irregular particles, and EDS images showed higher Cu content in debris areas, indicating that debris primarily originated from the friction pair. Figure 12(c2–e2) revealed delamination structures near the plow grooves in TiNbZrV_0.6 and TiNbZrV_0.8 coatings. These structures formed due to repeated compaction of debris.

Table 5. Element content of wear surface at RT (wt.%).

Samples	O	Cu	Ti	Nb	Zr	V
TiNbZr	9.47	11.76	26.56	25.63	26.58	-
TiNbZrV0.2	11.08	10.98	23.10	24.62	26.73	3.49
TiNbZrV0.4	10.35	9.6	19.73	25.86	28.12	6.34
TiNbZrV0.6	12.06	6.53	18.59	26.54	27.15	9.13
TiNbZrV0.8	10.42	5.49	15.28	27.68	28.61	12.52

shows the major element contents on the wear surfaces at RT. The mass fraction of O fluctuated around 10%, indicating minor oxidative wear at RT. With increasing V content, the mass fraction of Cu decreased, suggesting that the degree of adhesive wear decreases. This trend was related to the hardness change in the coatings. The higher the hardness, the stronger the adhesion resistance of the coatings [43].

Figure 13 shows the surface morphology of the coatings after the friction wear test at 700 °C, along with EDS spectra of the coatings and primary elements in the friction pair. Compared to RT, the worn surface appeared relatively smooth with reduced plowing and debris. Only shallow plowing and minor debris were observed on the TiNbZr surface. Surfaces from TiNbZrV_0.2~TiNbZrV_0.8 predominantly exhibited flake-like tribofilms. EDS images revealed that the tribofilm contained the friction pair element Cu but lacked the coating element Ti, indicating the tribofilm originated from the friction pair. At elevated temperature, the friction pair softened, allowing its components to adhere more readily to the coating surface, which accorded with the characteristics of adhesive wear. The transferred friction pair components rapidly oxidized at elevated temperature and were repeatedly compacted into dense films. These films partially functioned as lubricants, reducing the COF and wear rate [44]. This explains the data in Figure 11c, in which the average COF at elevated temperature was lower than that at RT. As V content increased, the tribofilm area expanded and the structure transitioned from dispersed to continuous.

Table 6. Element content of wear surface at 700 °C (wt.%).

Samples	O	Cu	Ti	Nb	Zr	V
TiNbZr	31.67	21.94	9.84	19.51	17.04	-
TiNbZrV0.2	33.72	25.47	9.13	15.32	14.84	1.52
TiNbZrV0.4	36.53	29.91	7.64	12.30	11.25	2.37
TiNbZrV0.6	35.52	27.91	7.82	12.25	12.83	3.67
TiNbZrV0.8	39.64	22.19	6.83	13.18	13.62	4.54

shows the content of major elements on the wear surface at 700 °C. Compared to RT, the O and Cu content on the wear surface was higher, indicating higher levels of oxidative wear and adhesive wear at elevated temperature. As V content increased, the mass fraction of Cu first increased and then decreased, indicating that the degree of adhesive wear first increased and then decreased. The mass fraction of Cu in the worn surface of the TiNbZrV_0.4 coating reached 29.91%, exhibiting the highest degree of adhesive wear.

Figure 14 shows the wear scar profile and wear rate of TC4 and TiNbZrV_x coatings at RT and 700 °C. As shown in Figure 14a,b, the wear scars of TC4 and TiNbZr coating at RT were composed of multi-channel grooves, showing a deep middle and shallow two sides, and the wear scar depth of TiNbZr reaches 363.6 μm. After V addition, the wear scar depth decreased obviously, the boundary between grooves became clearer, and the wear scar depth of TiNbZrV_0.8 decreased to 14.9 μm.

Figure 14c,d are profiles of wear scars at 700 °C. TC4 was more easily softened than the coatings at elevated temperature [45], leading to the friction pair cutting more into it and forming grooves. However, the tribofilm played a role in protection and lubrication, and the grooves were shallower than those at room temperature. So the scar surface of TC4 was still dominated by furrows, but the wear depth decreased to 266.4 μm, which was lower than the wear depth at RT but far greater than the coating’s depth. The wear scars of TiNbZrV_x coatings were flattened, and the friction films adhering to the coating surface increased with increasing V content.

According to the Archard formula, the wear rate is inversely proportional to the hardness of the material [46]. Measure the microhardness of the contact surface between the friction block and the friction pair, respectively, at room temperature, measure 5 times for each sample, and calculate the average value. Calculate the wear rates of TC4 and TiNbZrV_x coatings at RT and elevated temperature. The relationship between wear rate and microhardness of the contact surface is shown in Figure 15. TC4 is composed of HCP and BCC phases [47] and has higher hardness than single BCC. Therefore, TC4 showed higher microhardness and lower wear rate than TiNbZr. The wear rate of TC4 and TiNbZr at RT were 148.92 × 10⁻⁸ mm³/(N∙m) and 231.44 × 10⁻⁸ mm³/(N∙m). The addition of V can cause lattice distortion and hinder dislocation movement, and improve the plastic deformation resistance of the coating. Therefore, although the hardness of TiNbZrV_0.2 and TiNbZrV_0.4 was lower than that of TC4, the wear rates were greatly reduced. As V content increased, wear rate decreased continuously with increasing hardness and precipitates, because higher hardness values increased the bearing capacity of the coating to some extent [48]. The wear rate of TiNbZrV_0.8 decreased to 5.03 × 10⁻⁸ mm³/(N∙m), which was 97.8% lower than that of TiNbZr.

The wear rate of TC4 at 700 °C decreases to 123.13 × 10⁻⁸ mm³/(N∙m), which is 17.32% lower than that at RT. For TiNbZrV_x coatings, the decrease in wear rates was more obvious. The wear rates of the coatings fluctuated from 0.95 to 2.52, which were 99.64%, 91.93%, 92.97%, 80.86%, and 81.11% lower than those at RT, respectively.

4. Discussion

4.1. Phase Structure Control

Thermodynamic parameters of multi-principal element alloys (MPEAs) used to predict the phase composition include atomic radius difference (δ), valence electron concentration (VEC), mixing entropy (ΔSmix), mixing enthalpy (ΔHmix), melting temperature (Tm), etc. In order to understand the effect of interdendritic element segregation on the phase stability of coatings, the thermodynamic parameters of n-principal element alloys are calculated as follows [49]:

$$\mathsf{\delta } =\sqrt{{\sum }_{\mathrm{i}=1}^{\mathrm{n}}{\mathrm{x}}_{\mathrm{i}}{(1-{\displaystyle \frac{{\mathrm{r}}_{\mathrm{i}}}{\overline{\mathrm{r}}}})}^2};\overline{\mathrm{r}} ={\sum }_{\mathrm{i}=1}^{\mathrm{n}}{\mathrm{x}}_{\mathrm{i}}{\mathrm{r}}_{\mathrm{i}}$$

(4)

$$\mathrm{VEC}={\sum }_{\mathrm{i}=1}^{\mathrm{n}}{\mathrm{x}}_{\mathrm{i}}{\left(\mathrm{VEC}\right)}_{\mathrm{i}}$$

(5)

$${∆\mathrm{S}}_{\mathrm{mix}}=-\mathrm{R}{\sum }_{\mathrm{i}=1}^{\mathrm{n}}{\mathrm{x}}_{\mathrm{i}}\ln {\mathrm{x}}_{\mathrm{i}}$$

(6)

$${∆\mathrm{H}}_{\mathrm{mix}}=4{\sum }_{\mathrm{i}=1;\mathrm{i}\ne \mathrm{j}}^{\mathrm{n}}{∆\mathrm{H}}_{\mathrm{ij}}^{\mathrm{mix}}{\mathrm{x}}_{\mathrm{i}}{\mathrm{x}}_{\mathrm{j}}$$

(7)

$${\mathrm{T}}_{\mathrm{m}}={\sum }_{\mathrm{i}=1}^{\mathrm{n}}{\mathrm{x}}_{\mathrm{i}}{\mathrm{T}}_{\mathrm{i}}^{\mathrm{m}}$$

(8)

$$\Omega ={\displaystyle \frac{{\mathrm{T}}_{\mathrm{m}}{∆\mathrm{S}}_{\mathrm{mix}}}{\left|{∆\mathrm{H}}_{\mathrm{mix}}\right|}}$$

(9)

where x_i, r_i, (VEC)_i, ${\mathrm{T}}_{\mathrm{i}}^{\mathrm{m}}$ are the atom percent, atomic radius, VEC, and melting temperature of component i, respectively, and ${∆\mathrm{H}}_{\mathrm{ij}}^{\mathrm{mix}}$ is the mixing enthalpy between component i and component j. R is the gas constant valued at 8.314 J/K·mol. Ω is the enthalpy–entropy ratio, reflecting the formation ability of the solid solution phase.

According to the Hume–Rothery criterion, solid solution tends to form when δ < 6.6% and −15 kJ/mol < ΔHmix < −5 kJ/mol, while IMCs are easy to form when δ ≥ 6.6%. According to the VEC criterion, MPEAs tend to form the BCC phase when VEC < 6.8. Moreover, MPEAs tend to produce a single-phase solid solution structure when Ω >1 [49].

The thermodynamic parameters of inter-dendrite regions of TiNbZrVx coatings were calculated, and the results are shown in

Table 7. Phase stability parameter of inter-dendrite of TiNbZrV_x coatings.

Parameters	Criterion	TiNbZr	TiNbZrV0.2	TiNbZrV0.4	TiNbZrV0.6	TiNbZrV0.8
δ (%)	<6.6	4.64	5.61	6.16	6.64	7.11
VEC	<6.8	4.20	4.22	4.27	4.33	4.35
ΔHmix (kJ/mol)	−15~−5	−1.58	−2.68	−2.13	−1.73	−2.80
ΔSmix (J/K·mol)	-	10.01	10.97	11.39	11.59	11.76
Tm (K)	-	1810.89	1792.80	1807.37	1824.59	1811.50
Ω	≥1.1	11.49	7.33	9.69	12.21	7.62

. VEC and Ω values accorded with the condition to form a single BCC solid solution, but the values of ΔHmix were not in the range, and δ values of TiNbZrV_0.6 and TiNbZrV_0.8 exceeded 6.6. With the increase in V content, δ and VEC values increased, while Ω decreased, indicating an increasing tendency to form new phases.

4.2. Mechanical Property

The B2 phase has an ordered lattice structure, which is not easily sheared by dislocation. The amount of the B2 phase affects the arrangement of dislocations at grain boundaries. Figure 15 shows the dislocation distribution at grain boundaries of TiNbZr, TiNbZrV_0.4, and TiNbZrV_0.8 coatings. As shown in Figure 16a, the TiNbZr coating was a single BCC solid solution structure, and dislocations near grain boundaries were dominated by dislocation bundles [50] arranging in parallel, which reduced dislocation resistance and was beneficial to improve the plasticity of the material. When V content increases, as shown in Figure 16b, the amount of B2 phase increases and gathers at grain boundaries; the dislocation structure changes from parallel dislocation bundles to dislocation loops, resulting in local lattice deformation and stress concentration, which reduces the plasticity and tensile strength of the material. Furthermore, when V content increased to 0.8, the B2 phase transformed into AlZr₃ IMC, which improved the tensile strength by second-phase particle strengthening. However, the resistance to dislocation movement was increased in this process, resulting in a continuous reduction in plasticity.

4.3. Friction and Wear Mechanism

Figure 17 is the friction and wear mechanism diagram of TiNbZr and TiNbZrV_0.8 coatings at RT and elevated temperature. At RT, the wear mode is mainly abrasive wear. The hardness of the friction pair is 81.2 Hv, significantly lower than that of the coatings. Under contact pressure and high rotational speeds, it readily adheres to the coating surface, leading to material migration and debris formation. This suggests that, in addition to abrasive wear, minor adhesive wear also occurred. V element, on the one hand, reduced the grain size and, on the other hand, promoted Zr segregation at grain boundaries and formed AlZr₃ IMC. Smaller grain size and IMC increased the hardness of the coatings. As the coating hardness rose, the grinding effect of hard particles in the friction pair reduced, resulting in flatter plow grooves and fewer debris particles.

At elevated temperatures, friction films improved wear resistance by reducing the frictional force. On the other hand, according to Figure 2c, a part of the BCC phase transforms into the HCP phase at 700 °C [51]. Compared with BCC, the HCP structure has fewer slip systems, which is conducive to resisting plastic deformation during the wear process and improving the wear resistance of the coatings.

5. Conclusions

This study prepared TiNbZrV_x (x = 0, 0.2, 0.4, 0.6, 0.8) refractory MEA coatings via laser cladding. The effects of V on coating microstructure, mechanical properties, and friction/wear behaviour were investigated, yielding the following conclusions:

(1) The microstructures at the top and middle regions of the coatings were equiaxed crystals and dendritic crystals. V promoted element segregation, leading to B2 phase and AlZr₃ IMC formation. These phases fixed grain boundaries and inhibited grain growth, reducing grain size from 85.055 μm to 56.515 μm.

(2) The grain boundary precipitates changed the dislocation arrangement and increased the strength of the coatings by hindering the dislocation movement but also led to a decrease in the plasticity. The fracture mode transitioned from ductile fracture to brittle fracture. The TiNbZrV_0.2 coating exhibited the highest UTS, while the tensile properties of other coatings decreased with increasing V content.

(3) At RT, the coatings exhibited abrasive wear, while at elevated temperature, they demonstrated adhesive and oxidative wear. The TiNbZrV_0.8 coating demonstrated the lowest wear rate, whose wear rate was reduced by 97.8% at RT and formed a friction film that exceeded the wear volume at elevated temperature compared to the TiNbZr coating.