Using titanium powder, aluminum powder and boron carbide powder as raw materials, a Ti-C-B-Al composite coating containing in-situ self-generated TiC, TiB and TiAl intermetallic compounds reinforced titanium matrix was prepared on the surface of Ti6Al4V titanium alloy by synchronous coaxial laser cladding system. The microstructure, phase composition, microhardness and micro-nano mechanical properties of the composite coating were studied. The results show that the three powders fully reacted under the action of high-energy laser beam and generated TiC, TiB and TiAl intermetallic compound reinforcement phases. The composite coating has a smooth surface, no pores and cracks inside the coating, and a good metallurgical bond is formed between the coating and the substrate. The reinforcement phase is evenly distributed in the coating. Compared with the substrate, the cladding coating has better mechanical properties. The hardness of the coating is between 460HＶ0.3 and 510HＶ0. 3, and the nanomechanical properties of the coating have a good correspondence with the microhardness. When m(Ti)∶ m(Al)∶ m(B4C)＝84∶12∶4, the obtained coating has better mechanical properties.

Titanium alloy is widely used in aerospace, chemical industry, medical equipment and other fields due to its light weight, good biocompatibility, fatigue resistance and corrosion resistance. However, due to the low hardness, poor high-temperature oxidation resistance and wear resistance of titanium alloy, and easy adhesion, its large-scale use as high-temperature friction moving parts is limited. Surface enhancement technology can improve the mechanical properties of the titanium alloy surface, improve the friction and wear resistance and high-temperature oxidation resistance of titanium alloy. Among them, the titanium-based composite coating obtained by in-situ synthesis technology is closely combined with the substrate, the reinforcement is evenly distributed inside the coating, the size is controllable, and the thermodynamics is stable. The obtained coating has better performance, which effectively solves the defects of composite coatings prepared by traditional processes. In recent years, the use of magnetron sputtering, cold spraying, laser cladding and other technologies to obtain in-situ self-generated reinforcement phases to improve the surface properties of titanium alloys has received widespread attention. Among them, the high-energy laser beam of laser cladding can melt the powder and the substrate surface material at the same time, and quickly solidify to form a metallurgical bonding coating with a low dilution rate and dense interface structure. Therefore, laser cladding technology is an effective and feasible surface modification technology to improve the surface properties of titanium alloys.

In the laser cladding process, the coating performance obtained by selecting different cladding materials will vary greatly. Studies have shown that in the preparation of titanium-based composite coatings by in-situ synthesis, TiC, TiB, and TiB2 reinforcement phases obtained on the surface of titanium alloys using hard ceramic reinforcement phases such as Ti-C system and Ti-B system can improve the wear resistance of the titanium alloy surface; TiN and Ti2Ni obtained on the surface of titanium alloys using Ti-N system and Ti-Ni system can improve the corrosion resistance of the coating; TiAl, Al3Ti, and Ti5Si3 obtained on the surface of titanium alloys using Ti-Al system and Ti-Si system can improve the oxidation resistance of the titanium alloy surface; and the addition of refractory metals such as Nb and Zr can improve the biocompatibility of the coating. In-situ synthesis technology can achieve metallurgical bonding between the coating and the substrate, greatly improving the performance of titanium-based coatings.

Titanium alloy reinforcement phases enhance the single performance of the coating, but the service conditions of titanium alloy parts are complex, and it is urgent to integrate multiple excellent properties to improve their utilization rate. Due to the differences in the thermal expansion coefficient, melting point and other physical properties of the coating reinforcement phase, the research on in-situ synthesis of high-performance titanium-based composite coatings by laser cladding is still in its infancy. Therefore, this work uses different proportions of titanium powder, aluminum powder and boron carbide powder to prepare three different composite powders, and then uses synchronous cladding to prepare in-situ Ti-C-B-Al composite coatings. The phase composition, microstructure, microhardness and nanoindentation mechanical properties of the coating are analyzed, and the feasibility of laser cladding to prepare in-situ enhanced titanium-based multifunctional composite coatings is explored.

1. Experiment
The experimental substrate is Ti6Al4V hot-rolled titanium alloy plate, and the substrate size is L50 mm×H30 mm×T10 mm. Before the experiment, sandblasting is used to remove the oxide layer on the substrate surface, and the substrate surface is cleaned with petroleum ether. Powder system mainly verifies the effect of aluminum addition on the performance of the cladding layer. The components of three different agglomerated powders are prepared according to the mass ratio, S1 (m (Ti) ∶ m (Al) ∶ m (B4C) = 92 ∶ 4 ∶ 4), S2 (m (Ti) ∶ m (Al) ∶ m (B4C) = 88 ∶ 8 ∶ 4), S3 (m (Ti) ∶ m (Al) ∶ m (B4C) = 84 ∶ 12 ∶ 4). Figure 2 is a typical SEM image of the agglomerated powder obtained. As shown in Figure 2, the prepared agglomerated powder is nearly spherical, which meets the requirements of the synchronous coaxial powder feeding laser cladding test.

The single-layer single-pass coating was prepared by synchronous coaxial powder feeding RH-A3000D flexible fiber laser remanufacturing forming system. To ensure good metallurgical bonding and low dilution rate between the cladding layer and the substrate, the optimized process parameters were used: laser power of 2.4 kW, scanning speed of 24 mm/s, powder feeding speed of 1.2 r/min, defocus of -2 mm, and argon flow rate of 8 L/min.

When preparing metallographic specimens, the specimens obtained after laser cladding were firstly subjected to wire cutting sampling and hot mounting treatment, and then the cross-section of the specimens was ground and polished with SiC sandpaper of different meshes, and the specimens were corroded for 15 s using Keller reagent (Kroll reagent, a mixed aqueous solution of 2% HF + 5% HNO3 (volume fraction)). The coating cross section was analyzed by using a Bruker D8 Advance X-ray diffractometer; the cross section morphology of the laser cladding coating was analyzed by a FEI NovaSEM 450 field emission scanning electron microscope.

The microhardness of the coating cross section was measured by a Huayin 200HVS-5 digital display small load Vickers hardness tester. The test method was as follows: the loading load was 300 g and the loading time was 10 s. Each sample was measured along the depth direction of the cross section. Three parallel measurements were performed at each depth, and the average value was taken as the hardness value at that depth. The Agilent Nano Indenter G200 nanoindenter was used to perform a single nanoindentation mechanical property test on the cladding layer. The indenter was a standard triangular pyramid diamond indenter. Test method: Use the maximum depth test mode with a loading depth of 2 000 nm. Select four different positions on the cladding coating for measurement to obtain the nanohardness (HIT) and elastic modulus (EIT) of the cladding coating.

2 Results and Analysis

2.1 Analysis of Coating Phase and Microstructure
Figure 3 shows the SEM photos of the surface layer, middle layer and transition layer of the composite coating obtained by cladding the three powders. As shown in Figure 3, there are no defects such as cracks and pores inside the coatings obtained by cladding the powders of different compositions. The overall thickness of the S1-S3 cladding layers is 700 μm, 750 μm and 790 μm, respectively. As can be seen from Figures 3a1-a3 and b1-b3, a large number of granular precipitates, needle-shaped precipitates and lath-shaped precipitates are dispersed in the coatings, and the types of precipitates obtained by powders of different compositions are quite different: with the increase of Al content in the powder, the content of granular precipitates in the surface layer of the coating increases, and the content of needle-shaped and lath-shaped precipitates decreases. The size of the granular precipitates in the composite coatings obtained by cladding the three powders (S1-S3) is 0.6-2 μm, 0.4-1. 4 μm, 0.2-1.2 μm, the size of needle-shaped precipitates is 2-10 μm, 3-10 μm, 2-7 μm, and the size of lath-shaped precipitates is 3-8 μm, 3-7 μm, 2.6-6 μm. It can be seen from Figure 3c1-c3 that the precipitates in the transition zone formed by the cladding coating and the substrate are short fibers. The lengths of the short fibers in S1-S3 are about 16 μm, 10 μm and 8 μm, and the precipitates diffuse into the matrix phase, indicating that a good metallurgical bond is formed between the cladding layer and the substrate.

Comparing the coating structures obtained under different compositions, it can be seen that there are great differences in the types and sizes of precipitates inside the coating. With the increase of Al content in the powder, the granular precipitates in the coating increase continuously, the needle-shaped and lath-shaped precipitates decrease continuously, and the size of all reinforcing phases decreases continuously. The above results show that the increase of Al content will reduce the size of the precipitated phase in the coating and affect the occurrence of in-situ reaction.

Figure 4 is the electron probe microanalysis (EPMA) result of the typical area of ​​the coating of specimen S1. It can be seen that the element with the highest content in the coating is Ti, and other characteristic elements of the coating are evenly distributed in the coating, indicating that the reinforcing phase generated by the in-situ synthesis reaction is evenly distributed in the coating.

Figure 5 is the XRD diagram of the cladding coating. It can be seen that the XRD diagram of the composite coating obtained after the three powders are clad has the characteristic peaks of TiC, TiB and a small amount of TiAl, and the characteristic peaks of B4C, Al and TiB2 are not seen, indicating that during the laser cladding process, the titanium powder reacts completely with the boron carbide powder and the aluminum powder to generate three reinforcing phases, and due to the high Ti content, the TiB2 generated during the reaction is completely converted into TiB. At the same time, by comparing the XRD patterns of the three coatings, it can be seen that with the increase of Al content, the diffraction peak intensity of TiB and TiC in the coating increases continuously, while the diffraction peak intensity of Ti decreases continuously. This indicates that during the laser cladding process, the in-situ synthesis reaction inside the coating will be affected by the Al content, that is, with the increase of Al content, the in-situ reaction between powders is more sufficient, which increases the content of the newly generated reinforcement phase.

Combined with the above X-ray diffraction phase qualitative analysis results, it can be seen that the laser in-situ chemical reaction that occurs when Ti and Al powders, Ti and B4C interact with laser is:
3Ti+B4C→2TiB2+TiC (1)
5Ti+B4C→4TiB+TiC (2)
TiB2+Ti→2TiB (3)
Ti+Al→TiAl (4)

From the above reactions, it can be seen that during the laser cladding process, the agglomerated powder melts rapidly under the laser energy, and multiple in-situ chemical reactions occur simultaneously in the melt to generate new reinforcement phases, which are then cooled to form a cladding layer. Figure 6 shows the principle of laser in-situ reaction chemistry. First, Ti powder, Al powder and B4C powder are separated into active [Ti], active [Al], active [B] and active [C]. Then active [Ti] and active [B] are combined to form TiB ceramic phase, active [Ti] and active [C] are combined to form TiC ceramic phase, and excess active [Ti] and active [Al] are combined to form TiAl intermetallic compound.

2.2 Analysis of coating microhardness Figure 7 shows the curve of microhardness of the composite coating section with depth. As shown in Figure 7, the hardness change trend of the composite coating obtained after the three powders are clad is roughly the same. Among them, the hardness of the cladding coating (CZ) is the highest, followed by the heat affected zone (HAZ). The hardness of the molten pool where the coating transitions to the heat affected zone shows a downward trend. This is because the grain size at the bottom of the molten pool gradually increases and the reinforcement phase decreases, resulting in a decrease in the hardness of the coating. The reason why the hardness of the heat affected zone (HAZ) is higher than that of the substrate (Subb) is that the material of the substrate surface layer undergoes phase transformation hardening due to the heat of the molten pool during the laser cladding process, resulting in a difference in hardness from the substrate. The hardness of the substrate (Subb) is about 310HV0.3, and the average hardness of the coatings S1, S2, and S3 are 482.5HV0.3, 489.2HV0.3, and 503.1HV0.3, respectively. Compared with the substrate, the hardness of the cladding composite coating increased by 55.6%, 57.8%, and 62.3%, respectively, which is due to the large number of in-situ generated small-sized reinforcement phases dispersed inside the coating.

Comparing the hardness of S1-S3 specimens, it can be seen that with the increase of Al content, the microhardness of the composite coating increases. When the Al content is 12%, the hardness of the composite coating is the highest. According to the results of coating structure analysis, as the Al content in the agglomerated powder increases, the size of the reinforcement phase generated in the coating becomes smaller, that is, the volume fraction of the reinforcement phase increases, and the microhardness of the composite coating gradually increases under the strengthening effect of multiple reinforcement phases. However, the hardness of the coating gradually decreases in the molten pool. This is because during the cladding process, the powder and the substrate surface are melted and mixed at the same time, so that the reinforcement phase in the molten pool is diluted by the coating matrix phase, and the reinforcement phase in the molten pool is reduced, resulting in a decrease in the comprehensive mechanical properties of the molten pool.

2.3 Mechanical properties of coatings The performance of the composite coating is affected by the mechanical properties of the TiC, TiB, TiAl intermetallic compounds synthesized in situ inside the coating and the matrix phase. Therefore, the single nanoindentation mechanical properties of the composite coating were studied using a nanoindenter. Figure 8 shows the relationship between the load and the indentation depth obtained by the nanoindentation test of the cladding coating of specimens S1-S3. From the curve, it can be seen that a continuous and smooth load-displacement curve is formed when the triangular pyramid diamond indenter is pressed into the coating. There is no “step” phenomenon in the curve, indicating that the coating does not crack during the nanoindentation test. When the indenter reaches the preset depth, the greater the required load, the better the mechanical properties of the coating. After comparison, it is found that the applied loads of each coating are different when it reaches the preset depth, which indicates that the mechanical properties of each coating are different.

Figure 9 shows the mechanical properties of S1-S3 cladding coatings obtained by a single nanoindentation test. As shown in Figure 9, the average indentation hardness and elastic modulus of S1-S3 coatings have the same change trend, and are consistent with the change trend of microhardness. The average indentation hardness and elastic modulus of S3 coating are the largest, which are 5.6 GPa and 137.2 GPa, respectively.

According to the conversion relationship between indentation hardness and microhardness (HV = 94.5HIT), the Vickers hardness of each coating can be calculated. After calculation, the Vickers hardness of S3 coating is 529.2HV. It can be seen that there is an error between the Vickers hardness value converted from nanoindentation hardness and the measured hardness value, and the error value is about 26HV, which is converted to indentation hardness of 0.275 GPa. This error is caused by the fact that the area covered by the indenter in the nanoindentation test experiment is smaller than the area covered by the indenter in the microhardness test, that is, when performing nanoindentation testing, the indenter may be more concentrated in the reinforcement phase area, so the obtained indentation hardness value is larger. On the whole, the mechanical properties of nanoindentation can form a good correspondence with the microhardness results.

3 Conclusions
(1) Using titanium powder, aluminum powder and boron carbide powder as raw materials, a Ti-Al-B-C composite coating with smooth surface, no internal defects such as pores and cracks, and good metallurgical bonding was successfully prepared on the surface of TC4 alloy by laser cladding. The coating is mainly composed of TiC, TiB and TiAl intermetallic compounds.
(2) The microhardness of the Ti-Al-B-C composite coating is 480HV0. 3 ~ 510HV0. 3, among which the S3 cladding coating has the highest microhardness of 503. 1HV0. 3, which is 62. 1% higher than that of the substrate. The Al content in the agglomerated powder affects the in-situ reaction of the coating. The grain size is refined with the increase of Al content, which in turn affects the mechanical properties of the coating.
(3) The change trend of the nanoindentation hardness and elastic modulus of the coating is basically consistent with the change trend of the microhardness. When m(Ti): m(Al): m(B4C) = 84:12:4, the prepared coating has the best mechanical properties, with the average indentation hardness and elastic modulus being 5.6 GPa and 137.2 GPa, respectively.