Microstructure and properties of in-situ Ti5Si3-TiC composite coatings by reactive plasma spraying

Applied Surface Science 508 (2020) 145264

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Microstructure and properties of in-situ Ti5Si3-TiC composite coatings by reactive plasma spraying

T

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Xuan Sun, Wei Li, Jihua Huang , Jian Yang, Shuhai Chen, Xingke Zhao School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing 100083, PR China

A R T I C LE I N FO

A B S T R A C T

Keywords: Ti5Si3 TiC Wear Oxidation Reactive plasma spraying Coating

Ti powders, Si powders, and sucrose as a carbonaceous precursor were used to prepare Ti-Si-C composite powders by a spray-drying/precursor-pyrolysis technology. Ti-Si-C composite coatings were in-situ synthesized by reactive plasma spraying of Ti-Si-C composite powders. The microstructure and properties of the composite coatings were investigated. The results showed that the composite coatings consisted of in-situ formed Ti5Si3, TiC and Ti3O. The TiC particles with a diameter of hundreds of microns formed into clusters. The Ti5Si3 and Ti3O with sizes of about 20–40 nm coexisted in the lamellae. The microhardness of Ti5Si3-TiC composite coatings is up to 1906 HV0.1, about 4 times higher than that of TC4 alloys because of the nano Ti5Si3 and the submicron TiC grains. The composite coatings showed excellent wear resistance at room temperature and high temperature (600 °C), about 170 times and 40 times that of TC4 alloys, respectively. The oxidation tests suggested that the Ti5Si3-TiC composite coatings improved the oxidation onset temperature and have an oxidation weight change half of that of the TC4 alloys.

1. Introduction Titanium alloys are widely used in fields of aerospace, military, medical devices, sports equipment due to their low density, high strength-to-weight ratio, high yield strength and toughness, excellent biocompatibility and corrosion resistance [1–3]. However, titanium alloys have poor tribological properties and high-temperature oxidation resistance, which limit their application in wear environment [4–6] and high-temperature parts over 823 K [7–9]. Wear and oxidation are essentially a surface failure process that concerns only the composition and microstructure of the near-surface region. Therefore, coating technology has been considered as a promising approach to simultaneously improve the wear and oxidation resistance of titanium alloys by preparing appropriate coatings. The ternary Ti-Si-C system has attracted significant interest with high wear resistance, good oxidation resistance, and excellent electrical and thermal conductivity [10–12]. It contains a series of possibly formed high-hard or/and high-temperature ceramic phases such as TiC, Ti5Si3, SiC, Ti3SiC2 and thus is a promising protective coating material for titanium alloys. Currently, Ti-Si-C composite coatings have been fabricated and studied via various methods, such as plasma enhanced chemical vapor deposition (PECVD) [11], magnetron sputtering [13,14], gas tungsten arc welding (GTAW) deposition [15,16], double

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cathode glow discharge technique [17], self-propagating high-temperature sintering (SHS) [18] and reactive plasma spraying [19,20]. In the applied methods, reactive plasma spraying, combining plasma spraying with in-situ synthesis technology, has unique advantages for manufacturing ceramic composite coatings [21–23]. The in-situ synthesis allows the formation of submicron-sized or nanosized ceramic particles and clean interfaces between the particles, which makes the coatings high mechanical properties. Furthermore, the exothermic reactions as a second heat source, along with the heat of the plasma jet, prompt the melting of the powders and improve the bond strength between the coatings and the substrates. Agarwal et al. reported Ti-Si-C composite coatings prepared by reactive plasma spraying using tribochemically mixed Ti–SiC–C powders and discussed the reaction mechanisms under high heating and rapid cooling rate conditions [20]. Li et al used the Ti-SiC-graphite agglomerate powders by spray drying for reactive plasma spraying of Ti-Si-C composite coatings and high hardness was obtained [19]. However, in both Ti-Si-C composite coatings obtained, the reaction products were far from the initially designed composition, which meant that the reactions were inadequate in the plasma spraying process. As was reported by Agarwal et al., the unevenly mixed raw materials in the tribochemically blended powders were partly the reason [20]. Li et al. prepared composite powders with uniformly distributed raw particles

Corresponding author at: School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing 100083, PR China. E-mail address: [email protected] (J. Huang).

https://doi.org/10.1016/j.apsusc.2020.145264 Received 28 September 2019; Received in revised form 11 December 2019; Accepted 2 January 2020 Available online 07 January 2020 0169-4332/ © 2020 Elsevier B.V. All rights reserved.

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by spray drying while the phenomenon of inadequate reactions was still existing [19]. That can be explained by the separation of reactive constituent particles in the plasma spraying process [24–26]. In the traditional process of composite powders, the raw materials were bonded by mechanical mix or by a small amount of binder, which led to weak bonds in the composite powders. The reaction components were liable to be separated due to the dramatic plasma jet, which led to the incomplete reactions of the raw materials and the phase composition diverged from the initial design. Therefore, composite powders with high cohesive strength are the key factor for high-quality coatings. In this paper, a kind of powders with high cohesive strength was prepared using sucrose as a carbon precursor and the pyrolyzed carbon as a binder by a spray-drying/precursor-pyrolysis technology [27]. The feature of the process was that sucrose was used as the origin of carbon, which retained the high adhesion even after the pyrolysis. The mixture of Ti, Si and sucrose were first spray-dried into spherical agglomerate powders and then were pyrolyzed to carbonize the sucrose into carbon. By reactions among Ti, Si and C, Ti-Si-C composite coatings were in-situ synthesized by reactive plasma spraying, and the microstructure and properties of the coatings were investigated.

Table 2 Experimental parameters of the plasma spraying. Power 45 kW

Commercial pure Ti powders (80.18 wt%, 99% in purity, 5–20 μm, General Research Institute for Nonferrous Metals), Si powders (5.41 wt %, 99.99% in purity, 1–10 μm, Zhongnuo New Materials Co., Ltd) and sucrose (14.41 wt%, 99.9% in purity, Sinopharm Chemical Reagent Co., Ltd) were used as raw materials. The calculated theoretical composition of the composite powders is 89 wt% Ti, 6 wt% Si and 5 wt% C, which takes the pyrolysis loss of C into account. The Ti powders, Si powders and sucrose were blended by wet milling in absolute ethanol for 24 h at 120 rev min−1. Then the slurry was dried for 12 h. The obtained mixture was spray-dried into spherical agglomerate powders. The spray drying was conducted using a centrifugal spray dryer (LGZ 8, Wuxi Dongsheng) with the optimized process parameters listed in Table 1. The spray-dried powders were pyrolyzed in flowing argon gas at 250 °C for 2 h and 350 °C for 1 h to carbonize the sucrose into carbon. After pyrolysis, the powders were mechanically sieved to gain powders in the size range of 40–90 μm. Commercial TC4 alloys, of which the chemical composition was Ti6Al-4V (wt.%), were used as substrate materials with a size of 50 mm × 50 mm × 8 mm. Before plasma spraying, the substrate surfaces were activated by grit blasting with aluminum oxide particles to improve the adhesive strength between the coatings and the substrates. The feedstock was deposited onto the substrates by atmospheric plasma spraying using a UniCoatPro Plasma system (F4MB90-XL spraying gun, Oerlikon Metco) with Ar-H2 as working gases. Table 2 shows the experimental parameters of the plasma spraying.

Centrifugal atomization Co-current airflow 200 °C 90 °C 11000 rpm 50 wt% A15 (Hantai Chemical) 10% PVA (Polyvinyl alcohol) solution

100 mm

Argon flow −1

60 L min

Hydrogen flow 12 L min−1

2.3. Performance tests The microhardness tests were conducted on the cross-sections of the Ti-Si-C composite coatings by a Vickers hardness tester (VMHT30M, Leica) with applied loads of 100 g, 200 g and 300 g and a dwell time of 15 s. The Vickers microhardness value was the average of 10 measurements. The linearly reciprocating sliding wear tests were carried out on samples (Φ24 mm × 8 mm) of the TC4 alloys and the Ti-Si-C composite coatings using a ball-on-disk tribometer (Optimal SRV-4, Germany). The tests were conducted under a nonlubricated condition at room temperature (RT) and high temperature (600 °C). The friction counterpart was Si3N4 balls with a diameter of 10 mm. Table 3 shows the detailed parameters of the wear tests. The wear volume loss was the average value of three samples. The coefficient of friction (COF) as a function of the wear time was obtained after tests. The high-temperature wear tests started after the temperature of the samples and the worktable rose to 600 °C. The thermogravimetric (TG) analysis was used for oxidation tests by measuring the oxidation weight change of the samples in the heating process. The samples (Φ5 mm × 200 μm) of TC4 alloys and Ti-Si-C composite coatings were heated from 20 °C to 1000 °C in alumina crucibles at the rate of 10 °C min−1 in flowing air. The TC4 alloys and the coatings (10 mm × 10 mm × 2 mm) were held in a box muffle furnace at 800 °C under a flowing air atmosphere for 100 h to study their isothermal oxidation behavior. The weight of the samples was measured at 1 h, 4 h, 7 h, 10 h, 30 h, 50 h, 70 h, and 100 h using analytical balances to calculate the oxidation weight increment per unit area. After each time point, the XRD results of the coatings were tested to examine the phase change.

Table 1 Optimized process parameters of the spray-drying process.

Atomization method Spray-air contact mode Inlet temperature Outlet temperature Disk rotating velocity Solid content Dispersive medium Binder

75 V

Spray distance

The composition of the Ti-Si-C composite powders and the Ti-Si-C composite coatings were examined by X-ray diffraction (XRD, SmartLab, Rigaku) with Cu Kα target (λ ~ 1.5406 Å). The microstructure of the powders and the coatings were observed by the field emission scanning electron microscopy (FE-SEM, Zeiss Merlin VP compact) equipped with an energy-dispersive X-ray spectrometer (EDS) and the transmission electron microscopy (TEM, TECNAIG2 20) equipped with a high-resolution mode. To prepare the TEM samples, the Ti5Si3-TiC coatings with a diameter of 6 mm were cut from the substrates by wire-electrode cutting and grounded to 50 μm by hand using SiC sandpapers. Then, the samples were cut with a diameter of 3 mm by a punching press and glued onto a copper ring with a diameter of 3 mm. The center thickness of the samples was reduced to 20 μm by dimpling grinder (GATAN 656). Finally, the samples were thinned by ion-beam thinning with ion beam incident at 10° and ion energy of 6 kV. The worn surfaces of the TC4 alloys and the Ti-Si-C coatings were observed by FE-SEM after the tests.

2.1. Experimental materials and coating process

Values

600 A

Voltage

2.2. Characterization

2. Experimental details

Process parameters

Current flow

Table 3 Process parameters of the wear tests.

2

Normal force

Stroke length

Oscillating frequency

Test duration

Temperature

10 N

2000 μm

10 Hz

20 min

RT /600 °C

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Fig. 1. FE-SEM images and XRD pattern of Ti-Si-C composite powders: (a) top view, (b) cross-section and (c) XRD results of the powders during different processes.

3. Results and discussion 3.1. Microstructure of the powders and the coatings Fig. 1 shows the morphology and internal microstructure of the TiSi-C composite powders. The composite powders are spherical in a diameter range of 40–90 μm. In the backscattered electron micrograph of the cross-section in Fig. 1b, the powder exhibits a dense structure and homogeneous distribution of Ti, Si, and C. In the preparation process of the composite powders, the spray drying process produced powders with high sphericity and the uniform reactant distribution. That good features ensure the uniformly heating and the thorough reactions of the molten droplets. The Ti and Si were wrapped by the carbon, which provides high bond strength. Fig. 1c shows the XRD patterns of the composite powders in the stages of milling, spray drying and pyrolysis. The result shows that the mixture consists of Ti, Si, and sucrose after the milling. The sucrose disappeared when the mixture was spray-dried into spherical powders, indicating that the sucrose became amorphous owing to the rapid drying process. The amorphous sucrose was then carbonized to amorphous carbon in the pyrolysis process, which was also not present in the XRD result. The peaks of Ti and Si have no change in the preparation process, suggesting that there is no reaction or oxidation occurring during the preparation process of the powders. Fig. 2 shows the phase composition of the Ti-Si-C composite coatings. The composite coatings consist of Ti5Si3, TiC and Ti3O. The peaks of Ti5Si3 appear and the peaks of Si disappear in the coatings compared with the powders, suggesting that the Si and Ti react to form Ti5Si3 and the Si is exhausted. The TiC peaks exist in the composite coatings, demonstrating that TiC forms as a result of the reaction between Ti and amorphous C. The existence of the Ti3O implies that the Ti is residual after reacting with Si and C. The residual Ti reacted with O of the atmosphere to form Ti3O in the coatings owing to the high solubility of O element in Ti liquid phase, especially at high temperature. The oxidation phenomenon of Ti in atmospheric plasma spraying is difficult to avoid and has been reported in other reactive plasma spraying researches [28,29]. According to designed element composition in Ti-Si-C ternary phase diagram, the phase composition by equilibrium reaction

Fig. 2. XRD pattern of the Ti-Si-C composite coatings.

process is Ti5Si3, TiC and Ti, which is consistent with the actual results except for the oxidation of the residual Ti. That means the reaction components fully reacted due to high cohesive strength. In addition, an obvious phenomenon in the coatings is wider diffraction peaks compared to the powders, which means finer grains are formed in the reactive plasma spraying process. The top-right inset of Fig. 2 shows a more detailed XRD pattern of the three phases, which is obtained by scanning from 30° to 50° at a lower speed. The TiC peaks shift to larger angles compared to the standard peaks, which may be attributed to the high compress stress and the change of the phase structure [30]. Fig. 3a shows the cross-sectional structure of the Ti5Si3-TiC composite coatings. The composite coating has a thickness of about 200 μm and shows a dense structure with little pores and cracks. The dense coatings indicate that the powders were fully heated and melted into droplets in the plasma jet, which can be confirmed by the thin splat spreading on the surface of the coatings in Fig. 3b. Besides the heat input of the plasma jet, another heat source for the complete melting of the powders is the exothermic reactions. The two reactions of Ti5Si3 and TiC released a lot of heat and promoted the melting of the powders. The formation of pores in the coatings may be attributed to three reasons. 3

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Fig. 3. FE-SEM images of (a) the cross-section of the composite coatings and (b) the splat on the coating surface.

about tens of nanometers are observed. It is hard to distinguish the Ti5Si3 and the Ti3O due to the limited resolution of EDS as 1 μm. It will be realized by TEM in a later section. Fig. 4f, magnified E area in Fig. 4d, shows a high O content layer, which is supposed to form on the surfaces of the in-flight droplets or on the surfaces of the coatings in the air before the torch returns. To confirm the Ti5Si3 and the Ti3O in the lamellae of the composite coatings, the TEM images of the lamellae are shown in Fig. 5. Fig. 5a shows the microstructure of the Ti5Si3-Ti3O lamellae, the whole area consisting of particles with a size range of 20–40 nm. The nanograins are identified as Ti5Si3 and Ti3O by high-resolution images and shown in Fig. 5b and d, respectively. The insets in the two pictures show the Fast Fourier Transform (FFT) results of the grains, which are consistent with the Ti5Si3 and Ti3O crystal structure. In most of the area, the Ti5Si3 and Ti3O are uniformly mixed in the lamellae according to the microstructure observation and phase identification. The nano sizes of the two phases also suggest that the Ti5Si3 and Ti3O are formed in a fast cooling process, which should be the impacting process of the droplets. Therefore, the Ti5Si3-Ti3O lamellae are two-phase coupled growth microstructure of Ti5Si3 and Ti3O, which are simultaneously formed in the impacting droplets. Fig. 5c shows an area mainly composed of Ti5Si3, which is confirmed by the electron diffraction rings. The clear diffraction rings also suggest that the Ti5Si3 particles formed in the coatings are with small grain size in different crystallographic orientations. In addition, a small area of TiC is found in the lamella as shown in Fig. 5e even though most of the TiC particles exist in TiC accumulation areas in the coatings.

First, a small number of unmelted or partially melted powders had a poor binding with deposited coatings and resulted in pores. Second, the molten droplet impacted on the substrate and formed rugged splat. When the next droplet impacted the splat, the high area of the splat blocked the droplet from closely covering the splat in the low area. Then the pore formed between the two splats [22]. Third, the contraction of the splats and the gas with no time to expel during the cooling process also formed pores. The majority of the cracks are along the direction perpendicular to the substrate, suggesting that the cracks originate from the release of in-plane stresses in the cooling process. For further investigating the microstructure, the cross-sections of the Ti5Si3-TiC composite coatings were etched using a solution of HF: HNO3: H2O in the ratio 1: 1: 2 by volume. As shown in Fig. 4a, the composite coating was severely etched to expose the internal structure. Fig. 4b, magnified A area from Fig. 4a, shows the typical layered structure of thermal spraying coatings. The thickness of the lamellae is less than 5 μm, suggesting that the droplets impacting onto the substrates have a well-melted state. Two kinds of main microstructure can be found in Fig. 4b: B area of TiC accumulation and C area of Ti5Si3Ti3O layered structure. Fig. 4c, magnified B area in Fig. 4b, shows the TiC accumulation area, which is composed of hundreds of TiC particles in sizes of hundreds of microns. The accumulation of the TiC particles has been observed in the reactive thermal spraying of Fe-Ti-C powders in the previous studies [24,26,27]. The TiC formed in the flight process of the droplets, aggregated to TiC clusters, and finally remained in the coatings. Fig. 4d, magnified C area in Fig. 4b, shows the lamellae in the composite coatings. The lamellae are composed of Ti, Si and O according to the EDS analysis. The surface of the lamella is shown in Fig. 4e, magnified D area in Fig. 4d, in which particles with sizes of

Fig. 4. FE-SEM images of the etched cross-section of the composite coatings: (a) the etched cross-section, (b) the layered structure after etching, (c) the TiC accumulation area, (d) the magnified images of the lamellae, (e) the nanostructured Ti5Si3-Ti3O, and (f) the Ti3O lamella. 4

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Fig. 5. TEM images of the Ti5Si3-TiC composite coatings: (a) the microstructure of Ti5Si3-Ti3O area, (b) the high-resolution image of the Ti5Si3, (c) the polycrystal diffraction of the Ti5Si3, (d) the high-resolution image of Ti3O, and (e) the high-resolution image of TiC.

3.2. Thermodynamic calculation of the reactions As Ti5Si3 and TiC have been synthesized in the Ti-Si-C composite coatings, it is still necessary to clarify the actual reactions involved in the reactive plasma spraying process for the further optimization of the coatings. The thermodynamic predictions combined with the microstructure analysis of the coatings are used to explain the real reactions in the reactive plasma spraying process. The reaction of Ti-Si-C system is a complicated process and the possible reaction products include Ti5Si3, TiC, Ti3SiC2, TiSi, TiSi2, SiC and so on. The reactions with O2 and N2 also need to be taken into account as the plasma spraying is conducted in the atmosphere. Given the literature and the actual results, the following reactions are thought to be possible in the plasma spraying [20,31]:

5Ti + 3Si = Ti5Si3

(1)

Ti + C= TiC

(2)

Si + C= SiC

(3)

Ti + Si = TiSi

(4)

Ti + 2Si = TiSi2

(5)

Ti + (1/2)O2 = TiO

(6)

Ti + (1/2)N2 = TiN

(7)

3Ti + Si + 2 C= Ti3SiC2

(8)

Fig. 6. Gibbs free energy of formation of different phases in the plasma spraying process.

inorganic substances [32]. The Gibbs free energy for Ti3SiC2 formation was based on the work done by Y. Du [33]. From Fig. 6 the Gibbs free energy of all the reactions is negative from 298 K to 3000 K according to the thermodynamic calculation, suggesting that all the reactions are possible. The formation of Ti5Si3 has the lowest Gibbs free energy among the eight reactions, and then Ti3SiC2 and TiO. It can be inferred that the Ti element has different affinities with different elements, the bond trend for synthesizing a new phase, in the equilibrium reaction condition. The Si element is first, O element the second, C element the third. Therefore, Ti5Si3 is the most easily formed and stable phase in the

The Gibbs free energy of the reactions except for Ti3SiC2 in Fig. 6 was calculated using the handbook of thermodynamic data of practical 5

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equilibrium Ti-Si-C system. And then, the Ti3SiC2 should be the second easily formed phase. However, the synthesis of Ti3SiC2 requires the involvement of three elements, which makes the reaction difficult. The formation of Ti3SiC2 is limited due to the rapid non-equilibrium process, especially in the reactive plasma spraying. The difficulty to form Ti3SiC2 has been reported in the efforts to in situ synthesize Ti3SiC2 coatings by A. Agarwal [20]. The next possible phase is TiO, the Gibbs free energy of which is only higher than Ti5Si3 and Ti3SiC2. The XRD and TEM results show there is no TiO but Ti3O in the coatings, which can be explained by the protective effect of the plasma working gases, Ar and H2. During the plasma spraying process, the O element will be dissolved in the liquid droplets in the flight process because of the high solubility of O element in Ti liquid phase. The Ar and H2 protected the droplets from the atmosphere so that the O element dissolved in the droplets was not enough for the formation of TiO. Therefore, the Ti liquid phase with dissolved O element eventually formed Ti3O in the impact process. Above the Gibbs free energy of TiO are TiN and TiC, whose curves cross at about 1900 K. TiC has a stronger formation trend above 1900 K and thus is more possible to form at the high temperature of the plasma jet. Therefore, the Ti5Si3, Ti3SiC2, TiO, and TiC are the main possible phases as a result of the thermodynamic predictions, which is partly in accord with the real results of Ti5Si3, TiC, Ti3O.

Fig. 8. Wear volume loss of the TC4 alloys and the Ti5Si3-TiC coatings in the room-temperature wear tests.

3.3.2. Room-temperature wear tests Fig. 8 shows the wear volume loss of the TC4 alloys and the Ti5Si3TiC coatings. The wear volume loss of the composite coatings is 1/170 of that of the TC4 alloys, demonstrating that the Ti5Si3 and TiC in the coatings dramatically improve the wear resistance. Fig. 9 shows the coefficient of friction of the TC4 alloys and the Ti5Si3-TiC coatings. The wear process of the TC4 alloys, as shown in the COF curve, consists of two stages: the run-in stage and the stable friction stage. The run-in stage is from 0 s to about 600 s, in which the COF has an increasing trend along with the increase of wear time. In this stage, the contact surface expanded from a point to a circular arc between the Si3N4 balls and the TC4 substrates. The COF increases as a result of the increasing area of the contact surface. When it comes to about 600 s, the contact area has been constant up to its maximum and the wear process enters a stable state. After 600 s is the stable friction stage. In this stage, the COF is stable, of which the average value is about 0.9 during 600 s to 1200 s. It can be seen that the stable value of the COF is also the maximum because the contact area reaches its peak in the stable friction stage. The COF of Ti5Si3-TiC coatings exhibits an opposite trend. Once starting, the COF of the composite coatings increases to a stable value in just about 60 s. The short run-in period demonstrates that the composite coatings rapidly enter the stable friction stage due to the high microhardness. The average COF of the composite coatings is about 0.8 from 60 s to 1200 s, less than that of the TC4 alloys, suggesting that the working friction force is smaller in the wear surface of the composite coatings. It can also be seen that the COF of the composite coatings has a slightly decreasing trend in the stable friction stage, which is attributed to the formation of oxide films on the coating surface. To explain the different anti-wear performance between the TC4

3.3. Mechanical properties 3.3.1. Microhardness tests Fig. 7 shows the microhardness of the TC4 alloys and Ti5Si3-TiC coatings under loads of 100 g, 200 g, and 300 g. The composite coatings have higher microhardness under each load compared to the TC4 alloys. The highest microhardness reaches 1906 HV0.1 at the load of 100 g, about 4 times higher than that of the TC4 alloys. The high microhardness is attributed to the nano Ti5Si3 and submicron TiC distributed in the composite coatings. The microhardness of the composite coatings follows the Rule of Mixtures, which means the high percentage and high hardness of the Ti5Si3 and the TiC are the main sources of the high microhardness [34]. On the other hand, the finer size of the grains also plays an important role in promoting the hardness. The microhardness of the Ti5Si3-TiC coatings has a decreasing trend along with the increase of the loads. As the loads increase from 100 g to 300 g, the indentation size increases. The larger the indentation size, the stronger influence of the pores and the cracks on the hardness measurement, which leads to the decrease of the microhardness. In addition, the standard deviation of the microhardness values also has a decreasing trend with the increasing loads. When the low load is applied, the microhardness has big differences between the areas with and without defects. With the high load, the indentation size is large enough to reflect the influence of the defects on the microhardness, which is thus more stable and more exact.

Fig. 7. Microhardness of the TC4 alloys and the Ti5Si3-TiC coatings under various loads.

Fig. 9. Coefficient of friction of the TC4 alloys and the Ti5Si3-TiC coatings in the room-temperature wear tests. 6

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Fig. 10. FE-SEM images of the worn surfaces of the TC4 alloys and the Ti5Si3-TiC coatings after the room-temperature wear tests: (a) the worn surface of the TC4 alloys, (b) the worn surface of the coatings, (c) magnified area in a, and (d) magnified area in b.

reveals that the wear of the composite coatings follows the oxidation wear mode. As discussed above, the in-situ synthesized Ti5Si3 and TiC endow the composite coatings comparable hardness with the Si3N4 balls. This hard-to-hard wear mode makes it hard for the micro-protrusions of the Si3N4 balls to cut off the coating surface. With the wear time increasing, the temperature of the contact surfaces rose and the coating surface was oxidized due to the continuous friction. The oxide layers have a smooth surface, reducing the force of friction, responsible for the decreasing COF of the Ti5Si3-TiC composite coatings. Fig. 10d presents the wear scratches on the oxide layer, which have a width of about 1.5 μm and a very shallow depth. The tiny scratches show that the oxide layers also have high resistance to the wear, which improves the anti-wear performance of the composite coatings. The uncovered coating in Fig. 10b has a flat and smooth surface without scattered oxide layers on it, suggesting that the oxide layers are not worn by crushing or cutting. A large area of bare coatings means that the oxide layers peel off in large-flake form from the interfaces between the oxide layers and the composite coatings. The peeling of the oxide layers may be attributed to the different coefficient of thermal expansion between the oxide layers and the composite coatings. After the oxide layers peeled off, the exposed coating surface contacted with the wear balls and oxidized to form new oxide layers in the wear situation again. Therefore, the wear process of the Ti5Si3-TiC coatings in the roomtemperature condition is an oxidation wear mode, relating to a cycle of the oxide layers: forming-peeling off-forming.

alloys and the composite coatings, the worn surfaces were observed and shown in Fig. 10. Fig. 10a shows the worn surface of the TC4 alloys, which presents a furrow structure in the linear-reciprocating wear direction. The furrows are composed of grooves and plowed ridges with wear debris on them, suggesting that the wear mechanism of the TC4 alloys is abrasive wear. As the Si3N4 balls have a higher hardness (1600–2000 HV) than the TC4 alloys, the TC4 alloys are the main wear parts. The wear of the surfaces is carried out by micro-cutting. The micro-protrusions on the surface of the Si3N4 balls, acting as a hardabrasive grain, cut the TC4 surface to form the grooves in the wear direction. Simultaneously, the normal force induces plastic deformation to form the plowed ridges on both sides of the grooves. The rough surface formed by plastic deformation is shown in Fig. 10c. When the grooves and the ridges were repetitively formed, the surface materials were crushed to form debris. It can be seen in the inset of Fig. 10c that the debris has apparent curly shape, demonstrating that plastic deformation occurs in the micro-cutting process. Therefore, the wear process of the TC4 alloys is an abrasive wear process, substantially a micro-cutting process along with plastic deformation. A typical worn surface of the Ti5Si3-TiC coatings is shown in Fig. 10b. The worn surface consists of two areas, which are the oxide layer on the left side and the coating surface on the right side according to the element distribution analysis in Fig. 11. The oxide layer is also composed of two areas: the dark grey area and the light grey area as shown in Fig. 10b. By comparing the element distribution of the two areas in Fig. 11, the dark grey area has a high content of Si while the light grey area has a high content of Ti. From the EDS analysis of the point A and the point B in Fig. 10b shown in Table 4, it can be inferred that the main oxide phase in the dark grey area is SiO2 and the main oxide phase in the light grey area is TiO2. The Si content in the oxide layer is higher than the nominal compositions of Si in the raw powders, which means part of the Si element is transferred onto the surface of the composite coatings from the Si3N4 balls. The coating surface has a low concentration of O element according to Fig. 11 and the EDS result of point C in Fig. 10b shown in Table 4, indicating that this area is “fresh” and is just worn by the Si3N4 ball for a short time. The existence of the oxide layer and the “fresh” coating surface

3.3.3. High-temperature wear tests Fig. 12 shows the wear volume loss of the TC4 alloys and the Ti5Si3TiC composite coatings at room temperature and 600 °C under the same wear parameters. The wear volume loss of the TC4 alloys at 600 °C is less than that at room temperature, exhibiting a slightly higher antiwear performance at high temperature, which has been reported [35]. This promotion is related to the surface oxide layers, which form as a result of the high temperature. The wear volume loss of the composite coatings slightly increases at 600 °C. The Ti5Si3-TiC composite coatings still show an excellent wear performance with 45 times volume loss less than the TC4 alloys. 7

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Fig. 11. Element distribution of the worn surface of the Ti5Si3-TiC coatings by EDS mapping. Table 4 EDS analysis of the different areas in the worn surface of the Ti5Si3-TiC coatings in Fig. 10b. Area

A B C

Element (at%) C

O

Si

Ti

2.76 3.01 4.38

56.44 62.79 15.46

25.52 4.29 6.01

15.28 29.91 74.15

Fig. 13 shows the COFs of the TC4 alloys and the composite coatings at 600 °C. The TC4 alloys had a shorter run-in period at 600 °C than that at room temperature, indicating that the high temperature accelerated the run-in process. The wear entered the stable friction stage more rapidly. The average COF of the TC4 alloys in the stable friction stage is about 0.9, close to that at room temperature. The similar COFs imply that the TC4 alloys have the same wear mode at the two temperatures. The Ti5Si3-TiC composite coatings have a run-in state of about 200 s, longer than that at room temperature, demonstrating a tougher run-in process at 600 °C. The average COF of the composite coatings in stable friction process has an obvious rise to about 1.5, higher than that of the TC4 alloys and that of the composite coatings at room temperature. The higher COF suggests that the wear mechanism has changed as the temperature rises from room temperature to 600 °C.

Fig. 12. Wear volume loss of the TC4 alloys and the Ti5Si3-TiC coatings in the room-temperature and high-temperature wear tests.

Fig. 14 shows the worn surfaces of the TC4 alloys and the Ti5Si3-TiC composite coatings after the wear tests at 600 °C. The worn surface of TC4 alloys at high temperature is similar to that at room temperature, exhibiting the grooves and the plowed ridges with a mass of debris on them. The similar surface morphology indicates that the high-temperature wear of the TC4 alloys is also abrasive wear mode with the micro-cutting mechanism. However, the high temperature still brings 8

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surface morphology compared with the worn surface at room temperature. The flat area shows the typical oxidation wear mode with oxide layers on the surface of the composite coatings. This area endured the same wear process with the worn surface at room temperature, a cycle of the oxide layers: forming-peeling off-forming. As shown in Fig. 14e, most surface regions of the flat area have not been oxidized even at 600 °C, which means that the existence of Ti5Si3 slows down the oxidation at high temperature. However, the pits beside the flat areas are the main reason for the wear volume loss at 600 °C. From Fig. 14f, smooth fracture planes are found in the pits, suggesting that the pits are formed via brittle fracture. The brittle fracture has a close relationship with the high-temperature induced oxidation. At room temperature, the in-situ synthesized Ti5Si3 and TiC make the coatings high hardness and thus high wear resistance. Ti5Si3 and TiC also lead to high brittleness. The high brittleness results in cracks, especially under the high load at the high reciprocating frequency in the wear process. The cracks are observed both on the surface of the worn surface at room temperature in Fig. 10d and high temperature in Fig. 14e. However, the cracks became the fast channel of the oxidation at high temperature, which can be confirmed by the darker oxidized cracks in Fig. 14e. The oxidation, in turn, decreased the cohesive strength of the coatings, resulting in the brittle fracture along the cracks under the cyclic loads. Therefore, the high-temperature wear of the Ti5Si3-TiC coatings mainly follows a brittle fracture mode, which is attributed to the oxidation along the cracks.

Fig. 13. Coefficient of friction of the TC4 alloys and the Ti5Si3-TiC coatings in the high-temperature (600 °C) wear tests.

some differences. As shown in Fig. 14b, the grooves at 600 °C have a relatively smoother surface compared to that at room temperature. Moreover, the groove surfaces at 600 °C show obvious plastic deformation, which exists in morphology of plastic flow. The softened metal of the surface at 600 °C was pushed ahead by the wear balls along the wear direction and shaped to stripes as a result of plastic flow. This change in surface slightly facilitates the wear process because it is easier for the micro-protrusions to cut the surface metal of the TC4 alloys. Another change brought by high temperature can be concluded to oxidation behavior. Oxide films formed rapidly as the temperature rose up to 600 °C and reduced the coefficient of friction, resulting in a less wear volume loss than that at room temperature. An obvious evidence of the oxide films is the debris shown in the inset of Fig. 14c, which are formed by brittle fracture and have a high O content of 54.25 at% by EDS analysis. Therefore, the wear process of the TC4 alloys in the hightemperature condition is the abrasive wear mode by the micro-cutting mechanism, in which the plastic flow and the oxide films exist. Fig. 14d shows the worn surface of the Ti5Si3-TiC coatings after the wear tests at 600 °C. The worn surface is composed of flat areas and lots of pits, which is clearly different from the worn surface at room temperature. Fig. 14e shows the flat area in Fig. 14d, which has a similar

3.3.4. high-temperature oxidation tests To compare the anti-oxidation performance of the Ti5Si3-TiC coatings to the TC4 alloys, the samples of the coatings and the TC4 alloys were heated in flowing air to measure the oxidation weight change. As shown in Fig. 15, the Ti5Si3-TiC coatings and the TC4 alloys both have three stages in the heating process, namely I unoxidized stage, II initial oxidation stage, and III rapid oxidation stage. The initial oxidation stage of the TC4 alloys started at about 600 °C while the initial temperature of the coatings was about 700 °C. The higher initial temperature indicates that the Ti5Si3-TiC coatings can withstand higher operating temperature. In the initial oxidation stage, both the Ti5Si3-TiC coatings and the TC4 alloys started to oxidize and the oxidation rates were low. The initial oxidation stage lasted about 200 °C of temperature-rising, up to

Fig. 14. Worn surfaces of the TC4 alloys and the Ti5Si3-TiC coatings after the high-temperature (600 °C) wear tests: (a) the worn surface of TC4 alloys, (b), (c) the magnified images of the A, B areas in a, (d) the worn surface of the coatings, (e) the magnified back-scattered electron image of the C area in d, and (f) the magnified image of the D area in d. 9

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between the oxide layers and the substrates. With the increase and propagation of the cracks, the oxide layers peeled [36]. However, there was no oxide layer spalling in the oxidation process on the surface of the Ti5Si3-TiC coatings comparing Fig. 16d and e, which can be attributed to the similar coefficient of thermal expansion between the oxide layers and the coatings and the small thickness of the oxide layers. Fig. 17 shows that the TiO2 is the main oxide of the TC4 alloy and the coatings after 100 h. From the XRD patterns of different time points of the coatings in Fig. 17b, the TiC, Ti5Si3, and Ti3O gradually decrease and TiO2 increases with the increase of the oxidation time. The TiC is still existing after 100 h in the XRD patterns even the height of the peaks decreases gradually. In the Ti-Si-C system, TiO2 and SiO2 are the main oxides, and TiO2 is easier to form than SiO2 [36]. When the coatings were put in air at 800 °C, the TiO2 rapidly formed on the surface of the coatings to an oxide layer. Then, the oxidation process began to process in the interface between the oxide layers and the coating surface and the oxide layers grew [36,37]. In the interface, the formation of new oxides needed the O to diffuse from the outside into the interface and the Ti and Si to diffuse outwards. The SiO2 formed as an adherent oxide layer to prevent the diffusion of O and retard the oxidation progress, which is the main reason for the high oxidation resistance of the Ti5Si3-TiC coatings. Moreover, the Ti5Si3 in the coatings can retard the oxidation by inhibiting the diffusion of Ti needed for oxide layer growth [36]. Therefore, the oxidation resistance of the coatings was improved with the addition of the Si element.

Fig. 15. Oxidation weight change of the TC4 alloys and the Ti5Si3-TiC composite coatings as a function of oxidation temperatures.

800 °C for the TC4 alloys and to 900 °C for the Ti5Si3-TiC coatings, after which the oxidation rate obviously increased. The rapid oxidation stage started from 900 °C suggests that the Ti5Si3-TiC coatings still possessed good oxidation resistance below 900 °C while at the same temperature the TC4 alloys had worse oxidation. In the rapid oxidation stage, the diffusion of the O element was fast so that both the coatings and the TC4 alloys bore severe oxidation. After the oxidation until 1000 °C, the oxidation weight gain of the Ti5Si3-TiC coatings is less than half of that of the TC4 alloys, suggesting that the Ti5Si3-TiC coatings have better anti-oxidation performance than TC4 alloys, and are promising for the high-temperature protection of titanium alloys. To further study the isothermal oxidation behavior, the mass change of the TC4 alloys and the Ti5Si3-TiC coatings at 800 °C for 100 h were measured and shown in Fig. 16. The oxidation temperature of 800 °C was chosen because this temperature was in the initial oxidation stage of both the TC4 alloys and the Ti5Si3-TiC coatings. The result shows that the mass change of the Ti5Si3-TiC coatings is less than that of the TC4 alloys and is about half of that of the TC4 alloys after 100 h. Therefore, the Ti5Si3 improves the isothermal oxidation resistance of the Ti5Si3-TiC coatings compared with the TC4 alloys, which is in accordance with the non-isothermal oxidation results above. Fig. 16a, b, and c show the surface of the TC4 alloy before and after the oxidation, and the oxide layer peeling from the TC4 alloy. In the oxidation process, the oxide layers of the TC4 alloys rapidly formed and grew with the high temperature. With the increasing thickness, cracks generated due to the mismatch in the coefficient of thermal expansion

4. Conclusions Ti-Si-C composite powders were prepared by spray-drying/precursor-pyrolysis technology using Ti powders, Si powders and sucrose as raw materials. Ti5Si3-TiC composite coatings were in-situ synthesized by reactive plasma spraying of Ti-Si-C composite powders. The microstructure, mechanical properties and anti-oxidation performance were investigated. 1. The Ti5Si3-TiC composite coatings have a dense structure and consist of Ti5Si3, TiC, and Ti3O phases. The TiC particles have sizes of hundreds of microns and form TiC clusters while Ti5Si3 and Ti3O particles with sizes of 20–40 nm coexist in the coatings. 2. The Ti5Si3-TiC coatings exhibit high microhardness, which is about 4 times higher than that of the TC4 alloys. 3. The Ti5Si3-TiC composite coatings process excellent room-temperature and high-temperature wear performance, which are about 170 times and 40 times higher than that of TC4 alloys, respectively. The nanosized Ti5Si3 and the submicron-sized TiC are the main reason for the high hardness and wear resistance. 4. The Ti5Si3-TiC composite coatings have better anti-oxidation performance than TC4 alloys at high temperature.

CRediT authorship contribution statement Xuan Sun: Conceptualization, Methodology, Validation, Formal analysis, Investigation, Data curation, Writing - original draft. Wei Li: Investigation. Jihua Huang: Conceptualization, Methodology, Supervision, Project administration, Funding acquisition. Jian Yang: Writing - review & editing. Shuhai Chen: Writing - review & editing. Xingke Zhao: Writing - review & editing.

Declaration of Competing Interest Fig. 16. The oxidation weight gain and surface images of the TC4 alloys and the Ti5Si3-TiC coatings for 100 h at 800 °C: (a) the TC4 alloy, (b) the TC4 surface after the peeling of oxide layer, (c) the peeled oxide layer of TC4 alloy, (d) the Ti5Si3-TiC coating, (e) the oxidation surface of the coating.

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. 10

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Fig. 17. the XRD patterns of (a) the TC4 alloy after 100 h, and (b) the Ti5Si3-TiC coatings after the isothermal oxidation for 0 h, 1 h, 4 h, 7 h, 10 h, 30 h, 50 h, 100 h at 800 °C.

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