Ti2Ni matrix composite coatings with TaC addition prepared on Ti6Al4V by laser cladding

Applied Surface Science 402 (2017) 478–494

Contents lists available at ScienceDirect

Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc

Full Length Article

High-temperature wear and oxidation behaviors of TiNi/Ti2 Ni matrix composite coatings with TaC addition prepared on Ti6Al4V by laser cladding Y.H. Lv, J. Li ∗ , Y.F. Tao, L.F. Hu School of Materials Engineering, Shanghai University of Engineering Science, Shanghai 201620, China

a r t i c l e

i n f o

Article history: Received 21 July 2016 Received in revised form 4 January 2017 Accepted 12 January 2017 Available online 17 January 2017 Keywords: Coating Laser cladding High temperature Wear resistance Oxidation resistance

a b s t r a c t TiNi/Ti2 Ni matrix composite coatings were produced on Ti6Al4V surfaces by laser cladding the mixed powders of Ni-based alloy and different contents of TaC (0, 5, 10, 15, 20, 30 and 40 wt.%). Microstructures of the coatings were investigated. High-temperature wear tests of the substrate and the coatings were carried out at 600 ◦ C in air for 30 min. High-temperature oxidation tests of the substrate and the coatings were performed at 1000 ◦ C in air for 50 h. Wear and oxidation mechanisms were revealed in detail. The results showed that TiNi/Ti2 Ni as the matrix and TiC/TiB2 /TiB as the reinforcements are the main phases of the coatings. The friction coefficients of the substrate and the coatings with different contents of TaC were 0.431 (the substrate), 0.554 (0 wt.%), 0.486 (5 wt.%), 0.457 (10 wt.%), 0.458 (15 wt.%), 0.507 (20 wt.%), 0.462 (30 wt.%) and 0.488 (40 wt.%). The wear rates of the coatings were decreased by almost 83%-98% than that of the substrate and presented a decreasing tendency with increasing TaC content. The wear mechanism of the substrate was a combination of serious oxidation, micro-cutting and brittle debonding. For the coatings, oxidation and slight scratching were predominant during wear, accompanied by slight brittle debonding in partial zones. With the increase in content of TaC, the oxidation film better shielded the coatings from destruction due to the effective friction-reducing role of Ta2 O5 . The oxidation rates of the substrate and the coatings with different contents of TaC at 1000 ◦ C were 12.170 (the substrate), 5.886 (0 wt.%), 4.937 (5 wt.%), 4.517 (10 wt.%), 4.394 (15 wt.%), 3.951 (20 wt.%), 4.239 (30 wt.%) and 3.530 (40 wt.%) mg2 cm−4 h−1 , respectively. The oxidation film formed outside the coating without adding TaC was composed of TiO2 , NiO, Cr2 O3 , Al2 O3 and SiO2 . When TaC was added, Ta2 O5 and TaC were also detected, which effectively improved the oxidation resistance of the coatings. The addition of TaC contributed to the improvement in high-temperature wear and oxidation resistance. © 2017 Elsevier B.V. All rights reserved.

1. Introduction Titanium alloy is an important material in the fabrication of components used in aerospace, automotive, medical, marine, and other industries due to its low density, high strength-to-weight ratio, excellent corrosion resistance and other superior properties [1–3]. With the rapid development of the aircraft industry, a higher thrust-weight ratio is required to improve aircraft performance of aircraft engine. Titanium alloys have attracted much attention due to their high specific strength at high temperature. However, titanium is very active at high temperature, which results in titanium alloys being easily oxidized under high-temperature conditions

∗ Corresponding author. E-mail address: jacob [email protected] (J. Li). http://dx.doi.org/10.1016/j.apsusc.2017.01.118 0169-4332/© 2017 Elsevier B.V. All rights reserved.

[4,5]. Therefore, titanium alloys are mostly used at lower temperature (below 600 ◦ C). For example, the permitted maximal used temperature of Ti6Al4V cannot exceed 350 ◦ C under the condition of long service, which makes it only suitable for blades and the first and second blades of aircraft engines [6]. At present, mature titanium alloys used at higher temperature (about 600 ◦ C) belong to the Ti-Al-Sn-Zr-Mo-Si system (such as IMI384 and Ti1100 alloys) [7]. For instance, the IMI384 alloy developed by IMI and Rolls Royce has been used in the Trent700, EJ200 and PW350 engines. Many turboprop engine centrifugal impellers are also manufactured using IMI384 alloy. The Ti1100 alloy developed by Timet has been used to produce the engine valves of automobiles and motorcycles, and has also been applied in the T55-712 engine redesigned by Lycoming [8]. However, at present, the permitted maximal used temperature of titanium alloys still does not exceed 600 ◦ C.

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

479

Fig. 1. The schematic drawing of the preparation process of the preplaced layer.

The shortcoming of being easily oxidized restricts the further application of titanium alloys in high-temperature parts. Currently, the nickel-based superalloys, with excellent resistance to oxidation, are still the main material used to prepare high-temperature parts in modern aero engines. However, since the thrust-weight ratio of aviation engines cannot be improved effectively, it is extremely essential to improve the oxidation resistance of titanium alloys to widen their application in the aircraft industry. Some effective surface-modification methods [such as thermal spraying (TS) [9], physical or chemical vapor deposition (PVD or CVD) [10,11] and plasma surface alloying [12,13]] have been studied to fabricate thermal barrier coatings (TBCs) to improve the oxidation resistance of titanium alloys. However, the coatings produced by these methods tend to be loose and porous and are weakly bonded with the substrate [14,15]. As a promising surface-modification technology, laser cladding is widely used in many fields due to its high efficiency and stability. TBCs fabricated by laser cladding have fewer pores, good wear and corrosion resistance, and strong metallurgical bonding with the substrate [16–19]. Many researchers have prepared coatings to improve the oxidation resistance of titanium alloys. Liu et al. [20] produced intermetallic composite coatings on TA2 titanium alloy by laser cladding Ni-based alloy and evaluated the oxidation resistance. It was shown that these composite coatings mainly consisted of TiNi, Ti2 Ni and Ni3 Ti. After oxidation tests at 800 ◦ C, the oxidized surfaces of these coatings were much denser than that of the substrate due to the formation of Al2 O3 . Huang et al. [21] fabricated a coating consisting of (Ti, V)5 Si3 and a bcc solid solution on Ti6Al4V alloy by laser cladding a mixture of Ti, V, Cr, Al and Si elemental powders. The oxidation test at 800 ◦ C showed that the weight gain of the coating was decreased by almost 77% than that of the substrate after oxidation for 50 h. This change was attributed to the formation of a dense oxide layer consisting of TiO2 , Al2 O3 , SiO2 , Cr2 O3 , and a small amount of V2 O5 outside the coating. Liu et al. [22] reported a study on the preparation of TiN/Ti3 Al intermetallic composite coatings on Ti6Al4 V surface by laser cladding a mixture of Ti and AlN. An isothermal oxidation test at 600 ◦ C indicated that the oxidation resistance of the coatings was improved by almost six times that of the substrate due to the formation of TiN, Al2 O3 and TiO2 . From the above-mentioned examples, it can be seen that some metallic elements are frequently added into laser cladding materials to form a dense oxide layer under high-temperature conditions, which can effectively inhibit further inward diffusion of oxygen. However, these present studies mainly focus on oxidation behaviors of the laser-clad coating at high temperature (less than 800 ◦ C). There are few studies about oxidation behaviors of the coating at higher temperature (for example 1000 ◦ C). It is well known that the

increase in service temperature will significantly accelerate the oxidation of the coating. Moreover, more complicated reactions may occur, which may further promote the oxidation of the coating. Therefore, it is very essential to investigate oxidation behaviors of the coating at higher temperature, which contributes to the development of a new coating used at higher temperature and widens its application fields. On the other hand, besides oxidation behaviors of the coating, no high-temperature wear behaviors were involved in those studies. The coating may contact with the other components and do a relative motion when it is in service, which will result in the occurrence of wear. Wear will remarkably reduce the service performance of the coating. For example, the counterpart may produce the micro-cutting on the coating, the coating may be separated from the substrate. The coating may also peel off from the substrate due to the initiation and propagation of cracks during wear. Wear causes the destruction of the coating and accelerates the oxidation of the substrate. Therefore, high-temperature wear resistance of the coating is as important as its oxidation resistance. Many scholars also prepared laser-clad coatings to improve the wear property of titanium alloys at high temperature. Liu et al. [23] and Lu et al. [24] fabricated coatings on Ti6Al4V by laser cladding WS2 , CaF2 , h-BN, and other materials with a self-lubricating function to reduce the friction coefficient of the coatings. These laser-clad coatings presented better wear resistance than titanium alloys at 600 ◦ C. Among those studies, the friction-reducing role of self-lubrication phases (such as WS2 , CaF2 , h-BN) was investigated in detail. Besides those phases in the coating, a thin oxidation film formed on the coating surface during high-temperature wear will also significantly affect wear behaviors of the coating. However, the effect of the oxidation film on high-temperature wear behaviors was not involved in those studies. Moreover, oxidation behaviors of the coating were not investigated. Based on the above analyses, it can be found that present researches do not consider both oxidation behaviors and wear behaviors of the laser-clad coating at high temperature. The investigations into oxidation mechanism of the coating at higher temperature (for example 1000 ◦ C) are also lack. Moreover, the effect of the oxidation film on high-temperature wear mechanism needs to be revealed. This study aimed to prepare a new type of coating by laser cladding on the titanium alloy (Ti6Al4V) surface, which possesses not only outstanding oxidation resistance but also excellent wear resistance at high temperature. Ni-based alloy (NiCrBSi) was selected as the cladding material. The laser cladding parameters were optimized to obtain a smooth coating surface, good interface bonding and a uniform microstructure without pores and cracks. Moreover, suitable content of the substrate can be melted to react

480

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

with the cladding material, resulting in the in situ synthesis of some expected phases. Ti2 Ni and TiNi were designed as the matrix, and TiB, TiB2 , TiC were designed as the reinforcements. TiB, TiB2 and TiC as the ceramic particles have very high microhardness, which can prevent the coating from being cut by the counterpart. TiNi with body-center cubic structure and Ti2 Ni with face-center cubic structure possess more slip systems (12) than ␣(Ti) (3 slip systems) with close-packed hexagonal structure [␣(Ti) is the main phase in Ti6Al4 V]. This illustrates that TiNi and Ti2 Ni have better plasticity and toughness than ␣(Ti). They will endow the coating with good resistance to brittle debonding. The coating with good resistance to cutting and brittle debonding will present excellent wear resistance. Mostly importantly, TaC with high melting point (3880 ◦ C) is added into the cladding material (0, 5, 10, 15, 20, 30 and 40 wt.%) to further improve oxidation and wear resistance of the coating. Tantalum carbide (TaC) is widely used to fabricate components used at high temperature due to its high melting point (3880 ◦ C) and good high-temperature properties. Barbatti et al. [25] fabricated WC-Co cemented carbides using state-of-the-art powder sintering techniques and investigated the effect of the partial replacement of WC by cubic refractory carbides (TaC and NbC) on the oxidation resistance of the carbides at 600 ◦ C. The results showed that the oxidation kinetics obeyed a linear function and that increasing (Ta, Nb)C content caused a reduction of the oxidation rate. Yao et al. [26] prepared the Ta2 O5 -TaC inner layer and SiC outer layer on the surface of a C/C composite by pack cementation and chemical vapor deposition, respectively. Oxidation resistance of the composite coated by Ta2 O5 -TaC/SiC was compared to that only coated by SiC in a static air atmosphere at 1500 ◦ C over 100 h. The results showed that the mass loss of the former was only a quarter of the latter, illustrating that TaC can effectively improve the oxidation resistance of carbon/carbon composites. Yang et al. [27] prepared Ti3 SiC2 and 20 wt.% TaC/Ti3 SiC2 compounds using the hot pressing method. Their oxidation resistance was investigated at 1100–1500 ◦ C for 20 h. The results showed that the compound with TaC addition had a comparatively low mass gain rate at all temperatures (only about 7% at 1500 ◦ C). However, the mass gain rate of the compound without TaC addition was dramatically increased with increasing the temperature (about 50% at 1500 ◦ C), illustrating that the addition of TaC can effectively protect the material from being seriously oxidized. The above examples prove that TaC indeed has good oxidation resistance at high temperature. However, there are few reports about the application of TaC in laser cladding. In this study, the microstructure evolution of the coatings with different TaC contents was characterized. The effect of the addition content of TaC on oxidation behaviors of the coatings at very high temperature (1000 ◦ C) was investigated in detail. The investigation into high-temperature (600 ◦ C) wear behaviors of the coatings with different TaC contents was also carried out. Besides the microstructure of the coatings, the effect of the oxidation film on wear mechanism was highlighted.

2. Experimental procedures 2.1. Preplaced layer preparation A columned Ti6Al4 V alloy with a chemical composition of Ti6.5Al-4.26V-0.1C (in wt.%) was used as the substrate. The substrate was cut into a circular disk with a diameter of 50 mm and a thickness of 10 mm, and was then ultrasonically cleaned in alcohol for 15 min to remove surface stains. The cladding materials were mixtures of NiCrBSi alloy powder (Ni-16Cr-3.5B-4.5Si-1C in wt.%) with different contents of TaC powder (0, 5, 10, 15, 20, 30 and 40 wt.%). The mixtures were dried at 80 ◦ C for 15 min prior to being ground in a mechanical ball mill for 8 h.

Fig. 2. The treating process of the coatings for wear tests.

Prior to laser cladding, the preplaced layer needs to be prepared on the substrate. The bonding method is commonly used to prepare the preplaced layer. The mixed powder is homogeneously mixed with an organic binder such as acetyl cellulose alcohol, sodium silicate, etc. [28,29] and the slurry is formed. Then the slurry is pasted onto the substrate surface to form a preplaced layer. The method is very simple and flexible. However, a large amount of binder used in the method may be decomposed into gas during laser cladding. Those gas may has no enough time to overflow from the molten pool due to the rapid heating and cooling characteristics of laser cladding. The remaining gas will deteriorate the microstructure and reduce mechanical properties of the coating. Moreover, the thickness of the preplaced layer cannot be controlled precisely, which results in poor microstructural reproduction of the coating. In order to solve the shortcomings mentioned above, a new method was invented to prepare a high-quality preplaced layer with precise thickness and less binder in this study. The preparation process is illustrated in Fig. 1. The substrate surface was brushed with a thin layer of binder (4% polyvinyl alcohol) and placed into a hollow cylinder with 10.8 mm in height and 50.2 mm in inner diameter. The mixture was allowed to fill the empty space with a height of 0.8 mm above the substrate. Finally, the layer was pressed by a tablet machine at 30 MPa with a holding time of 3 min to make it dense, resulting in a preplaced layer with a thickness of approximately 0.8 mm. 2.2. Laser cladding Laser cladding was carried out in an YLS-5000 fiber laser system. The laser cladding parameters were optimized to obtain a smooth coating surface, good interface bonding, and a uniform microstructure without pores and cracks. Moreover, suitable content of the substrate can be melted to react with the cladding

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

481

Fig. 3. The topographies of the coatings with different contents of TaC addition before wear tests: (a) 0 wt.%; (b) 20 wt.%; (c) 30 wt.%; (d) 40 wt.%.

material, resulting in the in situ synthesis of some expected phases, such as Ti2 Ni, TiNi, TiB2 , TiB and TiC. The laser processing parameters were selected as follows: output power, 3 kW; spot diameter, 6 mm; scanning speed, 5 mm/s. After laser cladding, the singletrack and overlapped coatings with an overlap rate of 50% were cooled in air. The single-track coatings were used for X-ray diffraction tests, microstructural analysis, and oxidation tests, designated groups 1, 2 and 3, respectively. The overlapped coatings, designated group 4, were used for wear tests. The substrate was also used as the reference sample in the oxidation and wear tests. 2.3. Microstructure characterization The surfaces of the group 1 samples and the substrate were cut into flat shapes using an electrical discharge wire-cutting machine and polished using 150-grit SiC abrasive papers. A PANalytical X’Pert Pro X-ray diffractometer (XRD) with Cu K␣ radiation (␭=0.1540560 nm) was used to identify the phase constituents of the coatings, and their microstructures and chemical compositions were examined using a Hitachi S-3400 scanning electron microscope (SEM) equipped with a Genesis EDAX energy-dispersive spectrometer (EDS). Prior to the test, the cross-sections of the group 2 samples were successively ground by 150-, 600- and 1200- grit SiC abrasive papers, polished, and finally eroded in a mixed solution consisting of 4 ml water, 6 ml HNO3 , and 8 drops of HF for 30 s. The microhardness along the depth direction of the coatings was measured by a HXD-1000TMSC/LCD Vickers microhardness tester with a load of 500 gf and a dwell time of 15 s; the microhardness was measured repeatedly and an average was adopted. 2.4. Wear test The dry sliding friction test was carried out using a GHT-1000E high-temperature tribometer at 600 ◦ C in air for 30 min. The applied load was 10 N and the sliding speed was 0.188 m/s. Hard Si3 N4 ceramic balls with a diameter of 4 mm were used as counterparts and replaced after every sliding friction test. The friction coeffi-

Fig. 4. X-ray diffraction patterns of the coatings: (a) without TaC addition; (b) with 30 wt.% TaC addition.

cients were recorded in real time by the test system during testing. The wear volumes of the substrate and the coatings were measured using a CFT-1 ultrafunctional wear-test machine. The topographies and chemical compositions of wear surfaces were examined using the SEM and EDS. In order to reveal the wear mechanisms of the substrate and coatings with different TaC contents, the oxide films formed during the wear tests were evaluated by X-ray photoelectron spectroscopy (XPS) equipped with a Kratos Axis Ultra spectrometer using a monochromatic Al K␣ source. The energy steps for spectra and narrow-scan spectra were 1 and 0.1 eV, respectively. Considering the adventitious carbon on the analyzed sample surface, the binding-energy scale was corrected in accordance with the C1s peak located at 284.8 eV. Relative humidity will strongly affect the wear behaviors of material surfaces due to the presence of adsorbed films of water

482

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

vapor or hydrocarbons [30,31]. Takadoum et al. [32] investigated the effect of relative humidity in air on friction and wear of Al2 O3 , Si3 N4 , and partially stabilized zirconia (PSZ) sliding on various metals (copper, nickel, aluminum, and titanium). The results indicated that the increase in humidity resulted in a decrease in friction coefficient (␮) for copper and nickel. In contrast, the friction coefficient for aluminum and titanium was almost unaffected by the relative humidity level. Surface roughness is the other important factor affecting the properties of the materials (especially wear properties) [33–35], as it will significantly change the friction force during dry friction, further affecting the wear loss of the material. One main aim of this study is to investigate the effect of the addition of TaC on the high-temperature wear behaviors of the coatings. Therefore, in order to avoid interference from humidity and surface roughness, wear tests were carried out under approximately the same humidity and surface roughness. Prior to wear testing, the coatings underwent the procedure shown in Fig. 2. The surfaces of the coatings present a convex shape after laser cladding, which cannot be directly applied for wear tests. Therefore, the substrate surface was taken as the reference surface, and the part of the coating above the reference surface with a thickness of approximately 0.7 mm was removed using an electro-spark cutting machine. In order to further reduce the surface roughness, the surfaces of the coatings were ground successively with 150-, 600-, and 1200-grit SiC abrasive papers in an M-2 pregrinder. Finally, the coatings were ultrasonically cleaned in acetone to remove contaminants adhering

Fig. 6. Typical microstructure morphologies of the coating without TaC addition.

to the surfaces, after which the microroughness of the coatings was measured using a Bruker Dimension Icon scanning probe microscope (SPM). As shown in Fig. 3, the surfaces of the coatings with different TaC contents are extremely level and their roughness is approximately the same (48.9, 49.1, 41.6, and 46.9 nm). The applied temperature for the wear tests was controlled at 600 ◦ C, which can effectively remove water vapor adhering to the surfaces.

Fig. 5. BSE images of the cross section from the upper part to the lower part of the coating without TaC addition.

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

483

2.5. Oxidation test The substrate and coatings for oxidation testing were cut into cubes with dimensions of 5 mm × 4 mm × 1.5 mm. The surfaces of the cubes were ground by abrasive papers to ensure that the surface was extremely smooth. The isothermal oxidation test was carried out at 1000 ◦ C for 50 h in air using a muffle furnace with a temperature accuracy of ±0.5 ◦ C. The masses of these samples were measured every hour during the entire duration of the test using an analytical balance with an accuracy of 0.1 mg. In order to reveal the oxidation mechanisms of the substrate and coatings with different TaC contents, the oxide films formed during the oxidation tests were also evaluated by XPS.

3. Results and discussion 3.1. XRD results The XRD patterns of the coatings without and with 30 wt.% TaC are shown in Fig. 4. It can be seen from the figure that the diffraction pattern of the coating with 30 wt.% TaC is similar to that of the coating without TaC. All of the diffraction peaks are indexed in terms of the Joint Committee on Powder Diffraction Standards (JCPDS) cards. According to the indexed results, the main peaks of the coating without TaC are in accordance with the following JCPDS cards: No. 03-065-4572 for TiNi, No. 00-005-0687 for Ti2 Ni, No. 3-065-

Fig. 8. Typical microstructure morphologies of the coating with 30 wt.% TaC addition.

8805 for TiC, No. 01-075-1045 for TiB2 , and No. 01-089-3922 for TiB. This illustrates that the coating mainly consists of TiNi, Ti2 Ni, TiC, TiB, and TiB2 , all of which are Ti-rich compounds. These compounds are synthesized in situ between Ti resulting from the partial melting substrate and the other elements from the cladding material. When 30 wt.% TaC is added, some new peaks located at 35.26◦ , 39.9◦ , and 70.54◦ are found. They are conformed as the peaks of TaC in terms of the relevant JCPDS card (No. 01-089-2870 for TaC). The

Fig. 7. BSE images of the cross section from the upper part to the lower part of the coating with 30 wt.% TaC addition.

484

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

Table 1 Chemical compositions of the phases with different morphologies in the coatings with 0 wt.% and 30 wt.% TaC. Content of TaC

Zones

Ti

Al

V

Ni

Cr

B

Si

C

Ta

Phase

0 wt.%

a b c d e f

15.61 16.62 14.82 65.15 43.73 41.88

0.35 – – – 8.08 4.09

– – – – 2.75 0.68

1.17 – – – 24.74 41.56

1.53 1.56 1.62 – 9.03 4.37

80.81 59.96 61.55 – – –

0.42 – – – 7.24 2.91

0.11 21.86 22.01 34.85 4.43 4.51

– – – – – –

TiB2 TiB TiB TiC Ti2 Ni TiNi

30 wt.%

g h i j k

37.28 43.82 51.35 6.43 9.86

6.84 11.52 – – –

0.83 4.11 – 0.46 –

42.66 19.45 – – –

3.93 8.54 – 0.36 –

– – – 53.97 81.27

1.09 5.93 2.41 1.54 1.56

3.92 3.37 42.68 34.02 3.25

3.45 3.26 3.56 3.22 4.06

TiNi Ti2 Ni TiC TiB TiB2

Fig. 9. Morphology of the initial TaC particles.

TaC detected may be the original TaC added into the cladding materials or the TaC in situ synthesized during laser cladding. Moreover, the peaks corresponding to the coating with TaC exhibit a slight shift to the left, which indicates that the lattice distance is slightly larger. This may be due to the dissolution of Ta atoms into these compounds. 3.2. Microstructural characterization Fig. 5 shows the back-scattering electron (BSE) images from the upper to lower part of the cross-section of the coating without TaC. The evolution of the microstructure can be inferred from these images. A great number of reinforcements, including black blocky phase, grey blocky/needlelike phase, and equiaxed dendrites, are uniformly and dispersively distributed in the matrix. Generally speaking, the microstructure is uniform and similar throughout the entire coating, which should be related to Marangoni flow effects during the heat-transfer process. However, a slight change can still be observed, namely that with the increase in distance from the coating surface, the volume fraction of the black blocky phase presents a decreasing tendency, accompanied by an increase in the amount of grey blocky/needlelike phase. The black blocky phase almost disappears at the zones close to the interface, which is totally substituted by the grey blocky/needlelike phase. A transition area with a thickness of approximately 60 ␮m can be clearly observed at the interface between the coating and the substrate, indicating that the coating has a strong metallurgical bonding with the substrate. A BSE image (Fig. 6) at a higher magnification reveals that the matrix is composed of two phases, corresponding to a poriferous blocky phase and a grey honeycomb phase. EDS was used to identify the chemical compositions of these morphologically different phases (Table 1). The black blocky phase and grey blocky/needlelike phase are mainly composed of B and Ti, which can be considered

titanium boride. Based on the theory of back-scattering electron imaging, the image of one phase with a higher average atom number is brighter than that with a lower average atom number. Thus, the grey blocky/needlelike phase and the black blocky phase can be identified as TiB and TiB2 , respectively. The equiaxed dendrite is rich in Ti and C, which can be regarded as TiC. The poriferous blocky phase and grey honeycomb phase mainly consist of Ti and Ni, accompanied by a small amount of Al, V, Cr, Si, and C, and thus should be Ti-Ni compounds. Since a Si atom has a similar radius and electronegativity as a Ni atom, the lattice sites of Ni might be replaced by Si atoms in the Ti-Ni compounds. Similarly, the lattice sites of Ti may be substituted by Cr, Al, and V. For the poriferous blocky phase, the atomic ratio of Si and Ni (31.98 at.% in total) to Ti, Cr, Al, and V (63.59 at.% in total) is approximately 1:2. For the grey honeycomb phase, the ratio is about 1:1. Therefore, the poriferous blocky phase and the grey honeycomb phase can be confirmed as Ti2 Ni and TiNi, respectively. The cross-section of the coating with 30 wt.% TaC is shown in Fig. 7. The microstructure is also very uniform in the entire coating. When compared to the coating without TaC, a significant difference is that the black blocky phase and the grey blocky/needlelike phase are replaced by a white blocky/needlelike phase. Five phases marked in the BSE image (Fig. 8) with a higher magnification were analyzed by EDS. Combined with XRD analysis, phases g and h should be TiNi and Ti2 Ni, respectively, and phases i and j are TiC and TiB, respectively. Phase k with polygonal shape should be TiB or TiB2 . TaC identified in the XRD result is not clearly observed due to its low content and small size. Compared to the coating without TaC, the phase constituents have no significant changes. However, these phases all consist of different Ta contents, and these Ta atoms replace the lattice sites belonging to Ti atoms in those compounds during laser cladding. Owing to Ta’s atomic number being higher than that of Ti, these phases (especially TiB2 and TiB) look brighter in the BSE images. Based on the above analyses, it can be found that the addition of TaC mainly causes two changes in microstructure, corresponding to microstructural refinement and the solid solution of Ta into the intermetallic compounds (Ti2 Ni, TiNi, TiB2 , and TiC). Other than that, the XRD analyses also confirm that a small quantity of TaC is in the coatings. However, TaC particles cannot be observed clearly in the SEM images (shown in Fig. 5 and 7), and the reasons for this will be discussed later. The TaC existing in the coatings may be the incompletely dissolved TaC or the TaC in situ synthesized during laser cladding. Fig. 9 shows the morphology of the TaC initially added into the cladding material, in which the aggregation of a large number of fine equiaxed particles with an average size of approximately 0.5 ␮m can be seen. TaC particles with similar morphology are not observed in the coatings, and therefore it can be concluded that TaC in the coatings should be that in situ synthesized during laser cladding. Owing to the very high processing temperature and fine size of the TaC particles, TaC particles initially added into the

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

cladding material are completely melted during laser cladding, and TaC will be first precipitated in the molten pool due to its very high melting point (3880 ◦ C). The three reinforcements with high melting points (3067 ◦ C of TiC, 2980 ◦ C of TiB2 , and 2200 ◦ C of TiB) are then preferentially nucleated by the heterogeneous nucleation mode and grow on the surfaces of the TaC particles, which significantly refines the microstructure due to the increase in nucleation rate. In situ synthesized TaC particles will quickly be enclosed by TiB, TiB2 , and TiC, which further inhibits the growth of TaC. As a

485

result, fine TaC cannot be identified in the microstructure, and the other Ta atoms will enter into the crystal lattices of the compounds of TiB2 , TiB, TiC, Ti2 Ni and TiNi in solute form. More addition of TaC means that more heterogeneous nucleation centers are formed and more Ta dissolves into the solvents in solute form. As a result, the microstructure is further refined and the solid solution of Ta in intermetallic compounds is also further strengthened when more TaC is added into the coating.

Fig. 10. Friction coefficient of the substrate and coatings with different contents of TaC addition.

486

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

Fig. 11. Average friction coefficients of the substrate and coatings with different contents of TaC addition.

3.3. High-temperature wear resistance Fig. 10 shows the changes in friction coefficient of the substrate and coatings against a Si3 N4 ball at 600 ◦ C. Their average friction coefficients at the stable wear stage were also calculated (seeing Fig. 11). The changes in friction coefficient of the coating without TaC and the substrate with sliding time are very similar. The process of the sliding friction can be divided into two stages, the initial wear stage and the stable wear stage. At the initial wear stage (approximately 5 min), the friction coefficient presents an increasing trend with increasing sliding time due to the wear occurring between the protuberant parts of the samples’ surface and the counter ball. With the increase in contact area between the counterparts, the sliding friction enters a stable stage. Comparatively speaking, the friction coefficient of the coating without TaC fluctuates more dramatically with sliding time than that of the substrate. Moreover, the average friction coefficient (about 0.554) of the coating is higher than that of the substrate (about 0.431). When 5 wt.% TaC is added, it is clear that the friction turns into the stable stage in a very short time. Meanwhile, the average friction coefficient is slightly decreased from 0.554 (the coating without TaC) to 0.486; however, the friction coefficient still obviously fluctuates. When the addition content of TaC is further increased to 10 wt.%, the average friction coefficient is reduced to 0.457, and the fluctuation in friction coefficient with the sliding time is also weakened. When the addition content of TaC is increased to 15 wt.%, the relation curve between the friction coefficient and the sliding time becomes more smooth when compared to that (the coating with 10 wt.% TaC). The average friction coefficient (0.458) is similar to that (0.457 for the coating with 10 wt.% TaC). However, with the further increase to 20, 30 and 40 wt.% in addition content of TaC, the stability in friction coefficient with increasing sliding time is not further improved, and the average friction coefficient does not present the decreasing trend (0.507, 0.462 and 0.488 for the coatings with 20, 30 and 40 wt.% TaC). It can be concluded that the suitable addition of TaC (less than 15 wt.%) contributes to decreasing friction force and improving stability of the contact surfaces between the Si3 N4 ball and the coatings. The positive effect can be maintained with the further increase in addition content of TaC, but cannot be improved significantly. Fig. 12 illustrates the wear profiles of the substrate and the coatings after wear tests at 600 ◦ C for 30 min. The substrate has a much larger profile with a wear volume per millimeter of 1.15 × 10−5 mm3 . For the coating without TaC, the wear rate is greatly decreased, by approximately 83.0% (about 1.95 × 10−6 mm3 ). When 5 wt.% TaC is added, the wear rate is fur-

ther decreased by nearly 94.8% (about 5.99 × 10−7 mm3 ). When more TaC is added into the coating, the wear rate is further decreased by 95.7% (4.97 × 10−7 mm3 ), 95.9% (4.69 × 10−7 mm3 ), 96.6% (3.92 × 10−7 mm3 ), 97.2% (3.26 × 10−7 mm3 ) and 97.8% (2.49 × 10−7 mm3 ). It is observed that the wear rates of these coatings are greatly decreased when compared to that of the substrate. The addition of TaC will further improve the wear resistance of the coatings; in this case, when different contents of TaC are added, the wear rate is decreased in the range 69–87%. This indicates that the increase in the TaC content contributes to the reduction in wear rate during a high-temperature sliding friction test for 30 min. There are no similar reports about the effect of TaC addition on the wear behaviors of laser-clad coatings at high temperature. However, two reports had confirmed that TaC contributed to improving the wear resistance of laser-clad coatings at room temperature. Chao et al. [36] prepared a Ni-based composite coating reinforced by in situ synthesized TaC on mild steel by laser cladding the mixture of Ni60 powder with 20 wt.% (Ta2 O5 + C). As a reference, a pure Ni-based coating was also fabricated by laser cladding the Ni60 powder. The wear resistance of the two coatings was investigated by dry sliding wear tests for a duration of 20 min at ambient temperature. The results showed that the wear mass loss (2.2 mg) of the coating with TaC was only one-fifth that (11.1 mg) of the coating without TaC. Yu et al. [37] also prepared two coatings using Ni60 with 7 wt.% Ta powder and pure Ni60 powder on medium carbon steel plate surface by multilayer laser cladding. The tribological properties of the coatings were studied by pin-ondisc tests with a load of 25 N for a duration of 5 min at ambient temperature. For the coating prepared by laser cladding pure Ni60 powder, the weight loss was 0.061 g. The weight loss was significantly reduced to 0.015 g for the coating prepared by laser cladding Ni60 with 7 wt.% Ta powder, which can be attributed to the formation of TaC during cladding. However, the above studies only proved that the existence of TaC contributed to the improvement in wear resistance. The relationship between the content of TaC and wear resistance was not established. That is to say, with the increasing in content of TaC, whether wear resistance of the coating can be further improved was not verified. This study proves that the addition of TaC is also very efficient in improving the wear resistance of the laser-clad coating at high temperature. Moreover, it is also confirmed that the coatings with higher contents of TaC exhibit more excellent wear resistance. The change in wear resistance of the coatings is closely related to the microstructural evolution resulting from the addition of TaC. As discussed in Sec. 3.2, the addition of TaC can not only lead to microstructural fineness, but may also result in the solid solution of some Ta atoms entering into the intermetallic compounds (TiNi, Ti2 Ni, TiC, TiB and TiB2 ). With the increase in content of TaC, the above effects become more prominent, which further results in the improvement in strength/hardness of the coatings. Fig. 13 shows the microhardness distribution across the cross-section of the coatings with different contents of TaC. The microhardness does not change significantly along the direction from the surface to the interface, which indicates that the microstructure of the coatings is very uniform (as shown in Fig. 5 and 7). For the coating without TaC added, its average microhardness (855 HV0.5 ) is almost 2.5 times of that (345 HV0.5 ) of the substrate, which should be attributed to the in situ synthesis of high-hardness reinforcements, among which TiB2 (34 GPa) and TiC (30 GPa) uniformly distributed in the matrix play the predominant role. With the increase in the content of TaC, the average microhardness of the coatings presents an increasing trend as analyzed above (912 HV0.5 for 5 wt.%, 925 HV0.5 for 10 wt.%, 934 HV0.5 for 15 wt.%, 951 HV0.5 for 20 wt.%, 963 HV0.5 for 30 wt.%, and 983 HV0.5 for 40 wt.%). The enhancement in hardness can effectively improve the cutting resistance of the coat-

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

487

Fig. 12. The wear profiles of the substrate and coatings with different contents of TaC addition after wear tests at 600 ◦ C for 30 min.

Fig. 13. Microhardness distribution across the cross-section of the coatings.

ings, further endowing them with more excellent wear resistance, which was further verified by the roughness-testing results. Fig. 14 illustrates the surface topographies of the coatings after wear tests.

The surface roughness of the coatings before wear tests is about 45 nm (seeing Fig. 3). The surface roughness of the coatings with 0, 20, 30 and 40 wt.% TaC addition after wear tests were 109 nm, 66.4 nm, 63.4 nm and 59.5 nm, respectively. Owing to the improvement in resistance to micro-cutting by increasing TaC content, the worn surfaces of the coatings become more leveled. In addition to microstructure, oxide film is the other crucial factor affecting the wear resistance of the coatings. Fig. 15(a) and (b) illustrate the morphologies of the wear surface of the Ti6Al4 V alloy substrate after a wear test at 600 ◦ C for 30 min. The wear surface is very rough, and some long and narrow plough grooves parallel to the sliding direction can be clearly observed in some zones (A), accompanied by serious plastic deformation around them. The plough grooves are not very obvious in the other zones (B). However, a large number of debonding pits with different sizes and the white debonding edges around them are visible. Moreover, some micro-cracks and scratches can also be found. Fig. 15(b) is the back scattering electron (BSE) image of the wear surface. According to the BSE imaging mechanism, it can be inferred that zones A have higher atomic number than zones B. EDS was used to identify the chemical compositions in different zones, and the results are listed in Table 2. It can be seen that zones A are mainly composed of Ti

Fig. 14. The topographies of the coatings with different contents of TaC addition after wear tests: (a) 0 wt.%; (b) 20 wt.%; (c) 30 wt.%; (d) 40 wt.%.

488

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

Fig. 15. BE and BSE images of wear morphologies of the substrate and the coatings with different contents of TaC: (a) BE image of the substrate; (b) BSE image of the substrate; (c) BSE image of the coatings without TaC addition; (d) a partial zone of the former figure with a higher magnification; (e) and (f) BSE images of the coatings with TaC addition of 30 wt.% and 40 wt.%. Table 2 Chemical compositions of difference zones marked in Fig. 15. Zone

O

Ti

Al

V

Ni

Si

Cr

C

Ta

A B C D E F

24.11 58.93 13.87 29.54 9.51 24.46

65.29 35.16 52.69 42.66 48.02 49.11

7.92 4.45 9.83 6.45 5.79 5.17

2.68 1.46 3.90 1.93 3.46 –

– – 11.30 13.17 13.32 8.47

– – 2.03 2.46 6.35 5.98

– – 5.67 3.21 3.20 2.30

– – 0.71 0.58 1.57 1.35

– – – – 8.78 3.15

(65.29 at.%), O (24.11 at.%), Al (7.92 at.%) and V (2.68 at.%). Zones B mainly consist of Ti (35.16 at.%), O (58.93 at.%), Al (4.45 at.%) and V (1.46 at.%). The content of oxygen in zones B is about 2.4 times as much as that in zones A. However, the contents of metal elements are decreased (in particular, the content of Ti is reduced by almost half of that in zones A). Thus, zones A can be identified as exposed titanium alloy with mild oxidation, and zones B should be the oxide film formed on the substrate surface at high temperature. Fig. 15(c)–(f) illustrate the morphologies of the wear surfaces of the coatings with different TaC contents. It can be seen from the figures that the wear surfaces are smoother than that of the substrate. The long and deep plough grooves are not visible, and the number of debonding pits is greatly reduced. However, slight scratch traces can be found in different zones. According to these BSE images [Fig. 15(c)–(f)], it is observed that the bright zones with

high atomic number and the dark zones with comparatively low atomic number are discontinuously distributed within the wear surface. The majority of the dark zones are very smooth and dense, while the other zones are rough and loose due to the agglomeration of a large number of fine particles. In addition, the number of brighter zones is decreased with the increase in content of TaC. The EDS analysis results of the zones labeled C-F in Fig. 15(c) and (e) are listed in Table 2. The brighter zone C in the coating without TaC is mainly composed of Ti (52.69 at.%), Ni (11.30 at.%), O (13.87 at.%), and Al (9.83 at.%). The dark zone D is also rich in Ti (42.66 at.%), Ni (13.17 at.%), O (29.54 at.%), and Al (6.45 at.%). Small amounts of V, Si, Cr, and C were also detected. Ta is found in Zones E and F in the coating with 30 wt.% TaC. By comparison, it can be seen that the dark zones contain a higher content of O than the bright zones. Similarly, it can be inferred that the dark and bright zones should be the oxide film and the coating with slight oxidation, respectively. The oxide film has a significant effect on the wear resistance of these samples due to the change of the essence of sliding surfaces. In order to reveal the wear mechanism of the coatings, XPS was used to identify the constituents and these elements’ chemical valence state. Fig. 16(a) shows the survey spectrum of the oxide film formed on the coating with the addition of 40 wt.% TaC, which further verifies the existence of oxides containing Ti, Ni, Cr, Si, Al, and Ta. Fig. 16(b) and (c) show the high-resolution narrow XPS spectra of the Ti2p and Al2p states. The Ti2p peak can be fitted into two peaks

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

489

Fig. 16. XPS spectra of the oxide film formed on the surface of the coating with 40 wt.% TaC addition after a wear test at 600 ◦ C in a 30 min duration.

located at 458.1 and 463.9 eV, which are in good agreement with those of Ti4+ in TiO2 . The Al2p peak located at 74.0 eV indicates the existence of Al2 O3 . Fig. 16(d)–(g) show the spectra of Si, Ni, Cr, and Ta. TheSi2p peak located at 103.3 eV agrees with that of Si in SiO2 . The Ni2p peaks located at 855.6 and 872.87 eV are in good agreement with those of Ni2+ in NiO. The Cr2p peaks can be fitted at the positions with binding energies of 576.8 and 586.0 eV, which coincide with the standard values of Cr3+ in Cr2 O3 . The Ta4f peaks can be fitted into four peaks located at 22.2 eV and 24.1 eV (corresponding to Ta4f7/2 and Ta4f5/2 in TaC), 25.5 eV and 27.4 eV (corresponding to Ta4f7/2 and Ta4f5/2 in Ta2 O5 ).

Based on the above analysis, the formation mechanism of wear debris of titanium alloy can be described as follows (seeing Fig. 17). The oxide film formed on the surface of the titanium alloy substrate prior to the sliding friction test. In the initial wear stage, wear mainly occurs between the hard Si3 N4 ceramic ball and the oxidation film. Some micro-protrusions on the oxide film may shear off and become debris from the violent collisions with the hard ball. The sharp protrusions on the Si3 N4 ball’s surface may also be pressed into the surface of the oxide film under normal load and produce the mild scratches on the oxidation film upon subsequent sliding. Some sharp protrusions may even penetrate the

490

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

Fig. 17. The schematic drawing of different formation mechanisms of wear debris.

film and are impressed into the soft substrate. A strip of substrate coupled with the oxide film is microcut. Then the substrate, with its many grooves, is exposed and slightly oxidized. The final important formation mechanism of the debris is closely related to the mismatch in the elasticity modulus (E) between the substrate and the oxide film. Compared to the sharp protrusions on the Si3 N4 ball, some blunt protrusions may not be impressed into the surface of the oxide film, which only results in elastic deformation. Owing to the significant difference in the elasticity modulus between the soft substrate with low E and the oxide film with comparatively high E, the substrate cannot provide sufficient support for the oxide film. Thus, the oxide film acts as a cantilever beam, in which the zone subject to the force is like the deformation end and the other zones, without the applied force, are like the fixed end. The stress concentration will be generated at the interface between them, resulting in separation of the zone (deformation end) from the oxide film. The debris will further accelerate wear loss, and oxide particles, in particular, will produce the slight scratching on the oxidation film and exposed substrate. The other fixed particles may cause the stress concentration in the local zones of the oxidation film and result in the initiation and extension of the cracks. In addition to wear debris from the coatings, wear debris from the Si3 N4 balls may adhere to the surfaces of the coatings and produce the micro-cutting on the coatings. Enomoto et al. [38] investigated the characterization of wear behavior of steel-Si3 N4 and Si3 N4 -Si3 N4 pairings in the Versailles Project on Advanced Materials and Standards (VAMAS) round-robin tests, the results of which indicated that some particles from Si3 N4 were detached from the sliding surface and further oxidized into SiO2 debris. These wear particles may act as “free abrasives” in the Si3 N4 -Si3 N4 pairing or can be embedded on the soft steel surface as “fixed abrasives” in the Si3 N4 -steel pairing, which will abrade the surfaces of the coatings and cause their wear loss. Because the coatings also consist of Si element, which is also subjected to oxidization, it is very difficult to distinguish SiO2 wear particles from the Si3 N4 ball on worn surfaces. Therefore, oxidation, micro-cutting, and brittle debonding are predominant over the course of the entire wear process for titanium alloy. The wear mechanism of the coatings should be a combination of oxidation and micro-cutting, which is significantly different from that of the substrate. Wear debris mainly comes from the microcutting produced on the oxidation film by the counterpart of the Si3 N4 ceramic ball. Brittle debonding resulting from the mismatch between the oxidation film and the coating, and the serious cutting produced on the coatings, make almost no contribution to wear

Fig. 18. The mass gain of the substrate and coatings in an oxidation test at 1000 ◦ C in air for 50 h.

debris formation. This can be attributed to the strong support provided by the hard substrate (the coating without oxidization) to the oxide film. The high hardness endows the coating with excellent resistance to micro-cutting from the counterpart. Moreover, the similar elasticity modulus between the coating and the oxide film inhibits crack initiation and propagation within the oxidation film; that is, the oxide film can effectively shield the coating from destruction due to the support provided by the coating. The result reveals that the addition of TaC can increase the coverage scale of oxide film and protect the coating more effectively. It should be closely related to the formation of Ta2 O5 . A previous study reported that Ta2 O5 has a lower friction coefficient than TiO2 against a Si3 N4 ceramic ball [39]. Thus, the oxide film consisting of Ta2 O5 can obviously weaken the shear effect of the counterpart, and further reduce the wear loss. 3.4. High-temperature oxidation resistance Fig. 18 illustrates the mass gain data of the substrate and coatings with different TaC contents after an oxidation test at 1000 ◦ C in air for 50 h. The mass of these samples is increased with oxidation time as a whole. In the initial stage of the oxidation test, the mass gain of the samples is increased rapidly with oxidation time. In the later period, these samples’ mass tends to be stable with oxidation time increasing due to the formation of the stable oxidation film. For the same oxidation-time duration, the mass gain of tita-

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

nium alloy is obviously much higher than that of these coatings, illustrating that they can effectively protect the substrate from suffering significant oxidization. The relationship between mass gain and oxidization is in accordance with the parabolic rule, and the curves are fitted according to the growing tendency of mass gain. As shown in Fig. 19, the oxidation rates of the substrate and these coatings with TaC contents of 0, 5, 10, 15, 20, 30 and 40 wt.% are 12.170, 5.886, 4.937, 4.517, 4.394, 3.951, 4.239 and 3.530 mg2 cm−4 h−1 , respectively. With an increase in the content of TaC, the mass gain of the coatings exhibits a decreasing trend. After the entire oxidation test, the mass gain of titanium alloy and these coatings are 83.27, 45.41, 37.99, 35.19, 33.53, 30.81, 31.17, and 26.10 mg cm−2 . The mass gain of these coatings is decreased by nearly 45.5-68.7% when compared to that of titanium alloy. When TaC is added, the mass gain can be further decreased effectively, indicating that the addition of TaC can effectively protect the coatings from being significantly oxidized due to the formation of Ta2 O5 . However, with the further addition of TaC, the effect on mass decrease is not very obvious. After the oxidation test at 1000 ◦ C in air for 50 h, the surface of titanium alloy becomes loose and coarse. Moreover, a large area of oxidation film is detached from the surface. However, the surfaces of these coatings are very compact and smooth. In order to reveal the oxidation mechanism, the surfaces of the coatings without TaC and with 40 wt.% TaC are analyzed via XPS.

491

Fig. 19. Oxidation rates of the substrate and coatings with different contents of TaC addition.

Fig. 20(a) shows the XPS survey spectrum of the coating without TaC addition. The peaks of O, Ti, and Al can be observed quite obviously with small amounts of Si, Ni, and Cr. Corresponding high-resolution spectroscopy is employed to detect these elements’ chemical states [shown in Fig. 20(b)–(f)]. The binding energy of

Fig. 20. XPS spectra of the oxide film formed on the surface of the coating without TaC addition after an oxidation test at 1000 ◦ C in air for 50 h.

492

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

Fig. 21. XPS spectra of the oxide film formed on the surface of the coating with 40 wt.% TaC addition after an oxidation test at 1000 ◦ C in air for 50 h.

Ti2p3/2 is 458.0 eV, which indicates the presence of TiO2 in the oxide film. One Al2p peak of is located at 74.3 eV, which corresponds well to that of Al2 O3 . The Ni2p peaks located at 855.6 eV corresponding to Ni2p3/2 and 873.5 eV corresponding to Ni2p1/2 accord well with those in NiO. The peaks located at 577.5 eV corresponding to Cr2p3/2 and 103.2 eV corresponding to Si2p reveal the presence of Cr2 O3 and SiO2 , respectively. Fig. 21(a) is the survey spectrum of the oxide film formed on the surface of the coating with 40 wt.% TaC. Similarly, the oxide film

formed on the surface of the coating with 40 wt.% TaC added has similar oxides according to Fig. 21(b)–(f). Owing to the addition of TaC, two new compounds corresponding to TaC and Ta2 O5 are clearly detected according to the Ta4f peaks located at 22.6 and 26.0 eV, respectively [shown in Fig. 21(g)]. TiNi, Ti2 Ni, TiC, TiB2 , TiB, and a portion of TaC in the coatings reacted with oxygen and were transformed into TiO2 , NiO, B2 O3 , Ta2 O5 , CO, and CO2 by the reactions as follows: 2TiNi + 3O2 → 2TiO2 + 2NiO

(1)

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

2Ti2 Ni + 5O2 → 4TiO2 + 2NiO

(2)

2TiC + 3O2 → 2TiO2 + 2CO

(3)

TiC + 2O2 → TiO2 + CO2

(4)

2TiB2 + 5O2 → 2TiO2 + 2B2 O3

(5)

4TiB + 7O2 → 4TiO2 + 2B2 O3

(6)

4TaC + 7O2 → 2Ta2 O5 + 4CO

(7)

4TaC + 9O2 → 2Ta2 O5 + 4CO2

(8)

The possibility of these reactions occurring can be predicted by thermodynamic calculations. The changes in the standard Gibbs free energy (G) of these reactions at 1000 ◦ C are all calculated based on data from Ref. [40], and the results are listed in Table 3. It is obvious that all of the G values are negative at approximately 1000 ◦ C, which illustrates that these reactions can occur spontaneously in the oxidation test. Among all products, B2 O3 can be evaporated at an elevated temperature above 723 K due to its low melting point [41]. Moreover, C escapes from the coatings in the forms of CO and CO2 . Therefore, B and C can nearly be undetected. Regardless, it is found that the oxidation will result in the increase of these coatings’ mass to a certain extent, according to the reactions (1)–(8) above. Regarding TaC, a portion of TaC is oxidized into Ta2 O5 . However, the majority of TaC (about 85%) is reserved. This phenomenon is closely related to the change in G of reactions (7) and (8). As shown in Table 3, the G values for reactions (7) and (8) are more positive when compared to those of the other reactions, which means that these two reactions proceed with more difficulty. The existence of TaC inhibits the invasion of O, further decreasing the oxidation mass gain to a certain extent. Moreover, the residual and dispersively distributed TaC make the oxide film denser and prevent oxygen farther away from diffusing toward inner oxygen, as well as playing a positive role in resisting oxidation. In conclusion, the addition of TaC is very important to protect these coatings from being significantly oxidized. Moreover, more addition of TaC contributes to the further improvement in oxidation resistance of the coatings. Based on the above results, the effect of the content of TaC on microstructure, high-temperature wear/oxidation resistance of the coatings can be summarized as follows: (1) microstructure of the coatings is significantly refined with the increase in content of TaC; (2) hardness of the coatings presents the increasing tendency with the increase in content of TaC; (3) high-temperature wear resistance of the coatings is improved with the increase in content of TaC; (4) oxidation resistance of the coatings is improved with the increase in content of TaC. 4. Conclusions The main conclusions of this study are as follows: (1) The laser-clad composite coatings were fabricated on Ti6Al4V substrate using mixed powders of Ni-based alloy with different contents of TaC (0, 5, 10, 15, 20, 30 and 40 wt.%). The composite coatings mainly consisted of TiNi/Ti2 Ni as the matrix and TiC/TiB2 /TiB as the reinforcements. Ta mainly existed in these compounds in solute form. (2) With increasing TaC content, the friction coefficient of these coatings became smoother and approximately showed a decreasing tendency due to the formation of Ta2 O5 . The wear rate of the coatings was decreased by almost 83–98% when compared to that of the substrate, and presented a decreasing tendency with increasing TaC content. The wear mechanism of the substrate (Ti6Al4V) was confirmed to be the combination

493

of oxidation, scratches produced on the substrate, and brittle debonding of the oxidation film. For the coatings, the oxidation and slight scratches produced on the coatings were predominant over the course of the entire wear process. The oxidation film more effectively shielded the coatings from destruction with increasing TaC content. (3) The mass gain data of the substrate and the coatings were in accord with the parabolic rule. Following the entire oxidation test, the mass gain of the coatings was significantly decreased, by 45.5-68.7%, when compared to that of the substrate. The oxidation rates of the substrate and the coatings at 1000 ◦ C were 12.170 (the substrate), 5.886 (0 wt.%), 4.937 (5 wt.%), 4.517 (10 wt.%), 4.394 (15 wt.%), 3.951 (20 wt.%), 4.239 (30 wt.%) and 3.530 (40 wt.%) mg2 cm−4 h−1 . Following the entire oxidation test, the mass gains of the substrate and these coatings were 83.27 (the substrate), 45.41 (0 wt.%), 37.99 (5 wt.%), 35.19 (10 wt.%), 33.53 (15 wt.%), 30.81 (20 wt.%), 31.17 (30 wt.%) and 26.10 (40 wt.%) mg cm−2 . When compared to the coating without TaC added, the addition of TaC can reduce the mass gain by nearly 16.3%–42.5%. (4) The addition of TaC can effectively improve the hightemperature wear resistance and oxidation resistance of titanium alloy (Ti6Al4V). Acknowledgments This work was financially supported by the National Natural Science Foundation of China (51471105), “Shu Guang” project of Shanghai Municipal Education Commission and Shanghai Education Development Foundation (12SG44) and “Graduate innovation” Project of Shanghai University of Engineering Science (15KY0504). References [1] S. Bruni, M. Martinesi, M. Stio, C. Treves, T. Bacci, F. Borgioli, Effects of surface treatment of Ti6Al4V titanium alloy on biocompatibility in cultured human umbilical vein endothelial cells, Acta Biomater. 1 (2005) 223–234. [2] S. Durdu, Ö.F. Deniz, I. Kutbay, M. Usta, Characterization and formation of hydroxyapatite on Ti6Al4V coated by plasma electrolytic oxidation, J. Alloys Comp. 551 (2013) 422–429. [3] H.M. Wang, Y.F. Liu, Microstructure and wear resistance of laser clad Ti5 Si3 /NiTi2 intermetallic composite coating on titanium alloy, Mater. Sci. Eng. A 338 (2002) 126–132. [4] H. Tian, Y.M. Wang, Y.F. Zhang, L.X. Guo, J.H. Ouyang, Y. Zhou, D.C. Jia, Oxidation resistance of AlPO4 bonded ceramic coating formed on titanium Alloy for high-temperature applications, Int. J. Appl. Ceram. Technol. 3 (2015) 614–624. [5] J.J. Dai, J.Y. Zhu, C.Z. Chen, F. Weng, High temperature oxidation behavior and research status of modifications on improving high temperature oxidation resistance of titanium alloys and titanium aluminides: A review, J. Alloys Comp. 685 (2016) 784–798. [6] X. Li, S.Q. Lu, K.L. Wang, High temperature deformation characteristic of as-cast Ti40 burn-resistant titanium alloy, Trans. Mater. Heat Treat. 33 (2011) 28–32. [7] Q.J. Wang, J.R. Liu, R. Yang, High temperature titanium alloys: status and perspective, J. Aeronaut. Mater. 34 (2014) 1–26. [8] J.M. Cai, Z.X. Li, J.M. Ma, X. Huang, C.X. Cao, Research and development of 600 ◦ C high temperature titanium alloys for aero-engine, Mater. Review. 19 (2005) 50–53. [9] N. Cinca, C. Lima, J. Guilemany, An over view of intermetallics research and application: status of thermal spray coatings, J. Mater. Res. Technol. 58 (2014) 75–86. [10] C. Martini, L. Ceschini, A comparative study of the tribological behaviour of PVD coatings on the Ti6Al4V alloy, Tribol. Int. 44 (2011) 297–308. [11] H. Kashani, M. Heydarzadeh Sohi, H. Kaypour, Microstructural and physical properties of titanium nitride coatings produced by CVD process, Mater. Sci. Eng. 286A (2000) 324–330. [12] N. Lin, X. Huang, X. Zhang, A. Fan, L. Qin, B. Tang, In vitro assessments on bacterial adhesion and corrosion performance of TiN coating on Ti6Al4V titanium alloy synthesized by multi-arc ion plating, Appl. Surf. Sci. 258 (2012) 7047–7051. [13] N. Lin, X. Huang, J. Zou, X. Zhang, L. Qin, A. Fan, B. Tang, Effects of plasma nitriding and multiple arc ion plating TiN coating on bacterial adhesion of commercial pure titanium via in vitro investigations, Surf. Coat. Technol. 209 (2012) 212–215.

494

Y.H. Lv et al. / Applied Surface Science 402 (2017) 478–494

Table 3 The changes in the standard Gibbs free energy (G) of Reactions (1)–(8) at 1200 K. Reaction

(1)

(2)

(3)

(4)

(5)

(6)

(7)

(8)

G value (/kJ• mol−1 )

−536.2

−605.8

−515.5

−475.4

−557.7

−603.5

−479.4

−452.1

[14] Y. Bai, L. Zhao, J.J. Tang, S.Q. Ma, C.H. Ding, J.F. Yang, L. Yu, Z.H. Han, Influence of original powders on the microstructure and properties of thermal barrier coatings deposited by supersonic atmospheric plasma spraying, part II: Properties, Ceram. Int. 39 (2013) 4437–4448. [15] L. Jin, G.F. Liu, P.Z. Li, H.B. Zhou, C. Wang, G.D. Zhou, Adhesion strength and thermal shock properties of nanostructured 5La3TiYSZ, 8LaYSZ and 8CeYSZ coatings prepared by atmospheric plasma spraying, Ceram. Int. 41 (2015) 12099–12106. [16] F. Wang, C. Chen, H. Yu, Research status of laser cladding on titanium and its alloys: a review, Mater. Des. 58 (2014) 412–425. [17] Y. Chen, H.M. Wang, Rapidly solidified MC carbide morphologies of a pulsed laser surface alloyed ␥-TiAl intermetallic with carbon, Scr. Mater. 50 (2004) 507–510. [18] B.S. Du, Z.D. Zou, X.H. Wang, S.Y. Qu, Laser cladding of in situ TiB2 /Fe composite coating on steel, Appl. Surf. Sci. 254 (2008) 6489–6494. [19] Q.Y. Wang, Y.F. Zhang, S.L. Bai, Z.D. Liu, Microstructures, mechanical properties and corrosion resistance of Hastelloy C22 coating produced by laser cladding, J. Alloys Comp. 553 (2013) 253–258. [20] F.C. Liu, Y.Q. Mao, X. Lin, B.S. Zhou, T. Qian, Microstructure and high temperature oxidation resistance of Ti-Ni gradient coating on TA2 titanium alloy fabricated by laser cladding, Opt. Laser Technol. 83 (2016) 140–147. [21] C. Huang, Y.Z. Zhang, J.Y. Shen, R. Vilar, Thermal stability and oxidation resistance of laser clad TiVCrAlSi high entropy alloy coatings on Ti6Al4V alloy, Surf. Coat. Technol. 206 (2011) 1389–1395. [22] H.X. Liu, X.W. Zhang, Y.H. Jiang, R. Zhou, Microstructure and high temperature oxidation resistance of in-situ synthesized TiN/Ti3 Al intermetallic composite coatings on Ti6Al4V alloy by laser cladding process, J. Alloys Comp. 670 (2016) 268–274. [23] X.B. Liu, X.J. Meng, H.Q. Liu, G.L. Shi, S.H. Wu, C.F. Sun, M.D. Wang, L.H. Qi, Development and characterization of laser clad high temperature self-lubricating wear resistant composite coatings on Ti6Al4V alloy, Mater. Des. 55 (2014) 404–409. [24] X.L. Lu, X.B. Liu, P.C. Yu, S.J. Qiao, Y.J. Zhai, M.D. Wang, Y. Chen, D. Xu, Synthesis and characterization of Ni60-hBN high temperature self-lubricating anti-wear composite coatings on Ti6Al4V alloy by laser cladding, Opt. Laser Technol. 78 (2016) 87–94. [25] C. Barbatti, J. Garcia, P. Brito, A.R. Pyzalla, Influence of WC replacement by TiC and (Ta, Nb)C on the oxidation resistance of Co-based cemented carbides, Int. J. Refract Met. Hard Met. 27 (2009) 768–776. [26] D.J. Yao, H.J. Li, Q.G. Fu, J. Tao, Y.J. Wang, Ta2 O5 -TaC/SiC multilayer coating for protecting C/C composites against oxidation and ablation, Mater. China 30 (2011) 1–5. [27] Y.N. Yang, H.F. Yin, B.K. Liu, H.D. Yuan, H. Jia, Effect of addition of TaC on properties of Ti3 SiC2 ceramics, J. Aeronaut. Mater. 33 (2013) 56–60.

[28] Z.F. Xiang, X.B. Liu, J. Ren, J. Luo, S.H. Shi, Y. Chen, G.L. Shi, S.H. Wu, Investigation of laser cladding high temperature anti-wear composite coatings on Ti6Al4V alloy with the addition of self-lubricant CaF2 , Appl. Surf. Sci. 313 (2014) 243–250. [29] H. Zhang, Y. Zou, Z.D. Zou, C.W. Shi, Effects of CeO2 on microstructure and corrosion resistance of TiC-VC reinforced Fe-based laser cladding layers, J.Rare Earth. 32 (2014) 1095–1100. [30] D.H. Buckley, Surface Effects in Adhesion, Friction, Wear and Lubrication, Elsevier, Amsterdam, 1981. [31] H. Goto, D.H. Buckley, 5th Eurotrip Conference Proceeding, Vol. 3, K. Holmberg, I. Nieminen (Finish Society for Tribology, Espoo), 1989 p. 10. [32] J. Takadoum, Z. Zsiga, Effect of wear vapour in air on friction and wear of Al2 O3 , Si3 N4 and partially stabilized zirconia sliding on various metals, J. Mater. Sci. Lett. 12 (1993) 791–793. [33] C. Xu, H. Tian, C.E. Reece, M.J. Kelley, Topographic power spectral density study of the effect of surface treatment processes on niobium for superconducting radio frequency accelerator cavities, Phys. Rev. ST Accel. Beams. 15 (2012) 1–18. [34] O.E. Lundberg, A. Nordborg, I.L. Arteaga, The influence of surface roughness on the contact stiffness and the contact filter effect in nonlinear wheel-track interaction, J. Sound Vib. 366 (2016) 429–446. [35] J. Hun, H.B. Xu, Friction and wear behavior analysis of the stainless steel surface fabricated by laser texturing underwater, Tribol. Int. 102 (2016) 371–377. [36] M.J. Chao, W.L. Wang, E.J. Liang, D.X. Ouyang, Microstructure and wear resistance of TaC reinforced Ni-based coating by laser cladding, Surf. Coat. Technol. 202 (2008) 1918–1922. [37] T. Yu, Q.L. Deng, G. Dong, J.G. Yang, Influence of Ta on microstructure and abrasive wear resistance of laser clad NiCrSiB coating, J. Wuhan Univ. Technol.-Mater. Sci. Ed. 28 (2013) 437–443. [38] Y. Enomoto, K. Mizuhara, Characterization of wear behavior of steel and ceramics in the VAMAS round robin tests, Wear 119 (1993) 162–164. [39] J.H. Liu, Micro-arc Oxidation Films on Titanium, Zirconium, Tantalum and Their Bio-tribological Behaviors, Jinan University, 2011. [40] D.L. Ye, J.H. Hu, Practical Inorganic Thermodynamics Manual, second edition, Metallurgical Industry Press, Beijing, 2002. [41] Y.X. Qin, W.J. Lv, D. Xu, D. Zhang, Oxidation kinetics of in situ synthesized TiC/Ti matrix composites at elevated temperature, Chin. J. Nonferr. Met. 15 (2005) 352–357.