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Effects of post-heat treatment on microstructure and properties of laser cladded composite coatings on titanium alloy substrate
Optics & Laser Technology 65 (2015) 66–75
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Effects of post-heat treatment on microstructure and properties of laser cladded composite coatings on titanium alloy substrate G.J. Li, J. Li n, X. Luo School of Materials Engineering, Shanghai University of Engineering Science, Shanghai 201620, China
art ic l e i nf o
a b s t r a c t
Article history: Received 12 November 2013 Received in revised form 11 June 2014 Accepted 2 July 2014
The composite coatings were produced on the Ti6Al4V alloy substrate by laser cladding. Subsequently, the coatings were heated at 500 1C for 1 h and 2 h and then cooled in air. Effects of post-heat treatment on microstructure, microhardness and fracture toughness of the coatings were investigated by scanning electron microscopy (SEM), X-ray diffraction (XRD), energy dispersive spectroscopy (EDS), optical microscopy (OM). Wear resistance of the coatings was evaluated under the dry sliding reciprocating friction condition at room temperature. The results indicated that the coatings mainly consist of a certain amount of coarse white equiaxed WC particles surrounded by the white-bright W2C, a great deal of fine dark spherical TiC particles and the matrix composed of the α(Ti), Ti2Ni and TiNi phases. Effects of the post-heat treatment on phase constituents and microstructure of the coatings were almost negligible due to the low temperature. However, the post-heat treatment could decrease the residual stress and increase fracture toughness of the coatings, and fracture toughness of the coatings was improved from 2.77 MPa m1/2 to 3.80 MPa m1/2 and 4.43 MPa m1/2 with the heat treatment for 1 h and 2 h, respectively. The mutual role would contribute to the reduction in cracking susceptibility. Accompanied with the increase in fracture toughness, microhardness of the coatings was reduced slightly. The dominant wear mechanism for all the coatings was abrasive wear, characterized by micro-cutting or micro-plowing. The heat treatment could significantly decrease the average friction coefficient and reduce the fluctuation of the friction coefficient with the change in sliding time. The appropriate heat treatment time (approximately 1 h) had a minimal effect on wear mass loss and volume loss. Moreover, the improvement in fracture toughness will also be beneficial to wear resistance of the coatings under the long service. & 2014 Elsevier Ltd. All rights reserved.
Keywords: Laser cladding Post-heat treatment Wear resistance
1. Introduction Titanium and its alloys are widely used in chemical industries, aerospace and medical applications owing to their excellent corrosion resistance, high toughness and exceptional strengthto-weight ratio. However, the low surface hardness and poor wear resistance of titanium alloys limited severely their further application in many industry fields [1,2]. Therefore, it is essential to improve the wear resistance of titanium alloys using surface modification techniques, for instances, physical or chemical vapor deposition and plasma spraying [3–5]. Nevertheless, these methods have their limitations, such as the poor bonding strength between coating and substrate. As a promising surface modification technology, laser cladding has unique properties on wear resistance, corrosion resistance, and metallurgical bonding between the coating and the substrate compared with these
n
Corresponding author. Tel.: þ 86 21 67791198; fax: þ 86 21 67791377. E-mail address: [email protected] (J. Li).
http://dx.doi.org/10.1016/j.optlastec.2014.07.003 0030-3992/& 2014 Elsevier Ltd. All rights reserved.
conventional surface modification technologies [6]. Over the past few years, laser cladding composite coating onto pure titanium or titanium alloy substrate has attracted extensive attention owing to the higher efficiency, excellent metallurgical bonding in the interface and easy-controlled process and the excellent wear properties of the coatings [7]. Kulka et al. [8] obtained composite boride layers composed of laser-borided re-melted zone (TiB, TiB2 and Tiα0 -phase), heat affected zone (Tiα0 -phase) and the substrate (Tiα-phase) using laser-boriding with boron on cylindrical surface of commercially pure titanium substrate, the composite layers present excellent wear resistance in comparison with commercially pure titanium. Savalani et al. [9] produced TiC reinforced Ti matrix composite layers on pure titanium substrate by laser cladding mixed powders of Ti and carbon-nanotube of different contents, and found that the coatings with higher carbon-nanotube content had better wear resistance. Li et al. [10] obtained in situ synthesized TiN and TiB particulatereinforced metal matrix composite coating with hardness of 800 HV–1200 HV by laser cladding with a Ti/h–BN powder mixture, and the composite coating present better wear
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resistance. Our research team [11–13] fabricated titanium-based composite coating reinforced by in situ synthesized TiB whiskers and TiC particles on Ti6Al4V by laser cladding. The above studies mainly focused on the Ti matrix composite coating, some works regarding Ni matrix composite coating on titanium substrate were also carried out. Guo et al. [14] produced NiCrBSi and NiCrBSi/WC–Ni composite coatings on pure Ti substrates by laser cladding and investigated the friction and wear behavior at elevated temperatures of 300 1C and 500 1C, and found the wear rate of the NiCrBSi/WC–Ni composite coating was approximately 24–47 times lower than that of the pure Ti substrate in the same case. A process of laser cladding NiCoCrAlY coating on Ti6Al4V substrate with pre-placed NiCoCrAlY powder was studied by Meng et al. [15], the average microhardness of the coating was 800 HV–1000 HV, which was two times higher than that of Ti6Al4V substrate. The above researches indicate that laser cladding can improve significantly hardness and wear resistance of titanium alloys. However, laser cladding composite coating also has some disadvantages, e.g. high cracking susceptibility, which critically restricts its extensive applications in industry. Cracks may occur in the coating during or after laser cladding, caused by high thermal stress and microstructural stress occurred during rapid heating and rapid solidification [16]. Once cracks have been generated, crack propagation will be difficult to deal with in the subsequent application process, even resulting in the sample's failing. Therefore, cracks have a great influence on the quality of the laser cladded coating. How to reduce the cracking susceptibility of the coating has attracted extensive and continuous attention [16–18]. The generation of cracks can be attributed to the presence of the residual stress in the cladded coating. Based on this, many researchers have proposed many creative methods to reduce cracking susceptibility of the coating, such as preheating the substrate [19], adding rare elements and alloying elements to the clad powder [16,20,21], optimizing the processing parameters [22], etc. Besides, considering many important components by preheating the substrate are often subjected to the damage, Zhou et al. [23] put forward laser induction hybrid rapid cladding (LIHRC), and obtained the composite coatings which were free of cracks and had a good metallurgical bonding with substrate. Based on the principle of crack generation and expansion during laser cladding, Wang et al. [24] proposed adding 316 stainless steel net with low yield strength and good plasticity in the coatings, and crack density was significantly reduced. The former researches were mostly focused on preventing from producing of cracks in laser cladding process. However, post-treatment after laser cladding, aims to adjust microstructure and relieve residual stress generated in laser cladding process to avoid cracking in the subsequent application process [25]. In previous reports, postheat treatment was widely used to improve the required properties of the bulk materials by refining crystal, relieving residual stresses and even changing the microstructure [26]. A few investigations reported the effect of the heat treatment on the laser cladded composite coating. Chen et al. [27] investigated the influence of post-cladding stress-relieving treatments on the residual stresses in laser clad AISI P20 tool steel on the prehardened wrought P20 substrate, and the results indicated that the heat treatment significantly influenced the magnitude and sign of the residual stress by changing the volume fraction of retained austenite in the coating. Liu et al. [28] investigated the effect of tempering treatment on the corrosion resistance and microhardness of the Ni60CuMoW composite coatings on 45 steel surfaces, and the results showed that the tempering treatment could improve microhardness and the maximum self-corrosion potential and reduce corrosion current density significantly. Previous studies mainly focused on the effects of the heat
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treatment on residual stress, microstructure and corrosion resistance. However, very few studies reported the combined influence of post-heat treatment on cracking susceptibility and wear resistance of the laser cladded coating. Hence, in this study, the authors attempted to explore and discuss the effects of the post-heat treatment on microstructure and properties of the laser cladding coating on titanium alloy substrate, especially focused on hardness and fracture toughness. On this basis, the effect of the post-heat treatment on wear resistance was further studied.
2. Experimental procedures Ti6Al4V cylindrical samples with 50 mm in diameter and 10 mm in height were used as the substrate. Before laser cladding, in order to ensure good surface finish, all the samples were ground with emery papers and ultrasonically cleaned in acetone. Commercial powder blends of pure Ni and pure WC in nominal chemical constituents (mass fraction, %) of 70Ni–30WC were selected as the cladding materials. The powder blends after fully mixed and drying treatment were pre-placed on the Ti6Al4V substrate with the thickness of approximately 1 mm and were compressed with a press force of 50000 N by a hydraulic press instrument in order to make the preplaced blends more dense. Then single track laser cladding was carried out by an HL-5000 type CO2 laser source with a power of 2.5 kW, a beam diameter of 4 mm and a scanning speed of 5 mm/s. Fig. 1 shows the schematic drawing of laser cladding. After laser cladding, the specimens were divided into three groups. One group was not disposed and remained the original status as cladding; others were heated respectively at 500 1C for 1 h and for 2 h in the muffle furnace, and cooled in air. Microstructure was observed by means of JSM6460 scanning electron microscope (SEM) equipped with EDAX GENESIS energy dispersive spectroscopy (EDS). Microhardness along the depth of the cross-section was measured by using an HXD-1000TM microhardness tester. The load used was 100 g and loading time was set at 15 s. Fracture toughness, representing the cracking susceptibility of the coating, was measured by the Vickers indentation method. Vickers indentations were prepared on the cross-sections of the coatings using an HV-120 Vickers-hardness with the load of 20 N. Morphologies of the indentations were observed immediately using the VHX-600K optical microscope (OM). Fig. 2 shows the schematic drawing of cracks around the Vickers indentation. Fracture toughness of the coatings was calculated by the following
Fig. 1. The schematic drawing of laser cladding.
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3. Results and discussion 3.1. Microstructure characterization Fig. 4 presents the XRD patterns of the coatings with and without the heat treatment. According to the indexed results of all diffraction peaks in terms of JCPDS, it can be seen that phase constituents of the coatings are almost unchanged before and after the heat treatment, which consist of α(Ti), Ti2Ni, TiNi, WC, W2C and TiC. Ti2Ni, TiNi, TiC and W2C as the new phases were in-situ synthesized during laser cladding. No Ti element was introduced into the preplaced layer, and the Ti element in the coatings was derived from the substrate. During laser cladding, both the surface layer of the substrate and the preplaced layer were melted rapidly due to their low melting points (approximately 1670 1C and 1455 1C), respectively, and prodigious amounts of Ti element came into the molten pool. According to the Ni–Ti binary phase diagram [32] (Fig. 5), when the temperature is below 765 1C, phase transition reactions cannot be occurred and no new precipitation phases are generated. Therefore the heat treatment has little effect on the phase
Fig. 2. Schematic drawing of cracks around the Vickers indentation.
Fig. 3. The schematic drawing of the friction test on the different views: (a) front view and (b) left view.
equation [29]: a P K IC ¼ 0:079 3=2 log 4:5 c a
ð1Þ
in which KIC is fracture toughness (MPa m1/2), P is the indentation load (N), a is half a length of the diagonal line of an indentation (m) and c is the sum of half indentation diagonal and crack length (m). In this study, dry sliding reciprocating friction tests of the Ti6Al4V substrate and the single-track coatings were carried out using the ball-on-disc mode on a CFT-I comprehensive friction and wear apparatus at room temperature. The aim of using the singletrack coatings, not the overlapping coatings, was to eliminate the effect of the overlapping. Microstructure may become coarse in the overlapping zone since the zone is subjected to the second laser radiation during laser cladding. That is to say, the overlapping can affect microstructure, further affect the mechanical properties of the coatings [30,31]. Hence, the single-track coatings were used in this study. Fig. 3 shows the schematic drawing of the friction test on the different views. The coatings' surface is rough and may affect the testing results. Therefore the surface parts of the coatings with a thickness of approximately 0.2 mm were removed by the spark machining method, and then were polished with 300grit SiC abrasive paper prior to the wear tests. Counterbody balls were made of Si3N4 of approximately 4 mm in diameter with a hardness of 1700 HV. The applied load was 20 N. The sliding speed was maintained at 5 m/min, and the sliding time was approximately 20 min. The friction coefficient was recorded by the computer connected to the tester. The wear volume loss and mass loss of the substrate and the coatings were measured by a surface mapping profiler coupled with the wear tester and a photoelectric balance (0.01 mg resolution), respectively. Besides, worn surface morphologies and wear debris of the coatings were investigated by SEM and EDS.
Fig. 4. XRD patterns of the coatings with and without the heat treatment.
Fig. 5. The Ni–Ti binary phase diagram.
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constituents of the coatings due to the low heat treatment temperature of 500 1C. In order to understand the formation mechanism of new phases in the coatings, thermodynamics calculation was applied. Based on the above XRD results, the reactions between the Ni–Ti, Ti–C and W–C systems were considered. All possible reactions occurring during laser cladding are as following:
Fig. 6. Relationship between standard Gibbs free energy and temperature for the possible reactions during laser cladding.
NiþTi ¼TiNi
(2)
Niþ2Ti ¼ Ti2Ni
(3)
TiþC ¼ TiC
(4)
WþC ¼W2C
(5)
Based on the relative data in Ref. [33], the standard Gibbs free energies of formation of those compounds at different temperatures are presented in Fig. 6. It is obvious that the standard Gibbs free energies of formation of those compounds are negative below 1100 K, which indicates that the formation reactions of those compounds are spontaneous, and the spontaneous trend of these compounds at the same temperature is as follows: TiC 4 Ti2Ni4TiNi 4W2C. Based on the Ni–Ti binary phase diagram (Fig. 5), it is impossible to obtain α(Ti), Ti2Ni and TiNi phases simultaneously when the content of Ti element is constant and uniformly Table 1 Chemical composition (at%) of the phases with different morphologies.
Fig. 7. Schematic drawing of Ti element distribution in the coatings.
Analysis zones
Ti
Ni
Al
W
C
V
A B C D
9.21 42.62 53.03 43.78
– 3.94 30.75 48.76
– 1.20 11.04 6.51
41.21 1.95 1.5 –
49.59 50.29 – –
– – 3.68 0.93
Fig. 8. Microstructure of the upper part of the coatings: (a) without the heat treatment, (b) with the heat treatment for 2 h, (c) morphological features of different phases: white equiaxed particle, dark spherical particle, gray phase and light gray phase are labeled A, B, C and D, respectively, and (d) morphology of the WC particles.
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distributed in the coatings. However, the XRD results show that the three phases are present simultaneously in the coatings, which indicates that chemical composition distribution is non-uniform in the coatings. As is well known, laser cladding is a transient process. Due to the quick cooling process of the molten pool, the diffusion and convection of Ti element from the substrate in the molten pool were not sufficient and presents the gradient distribution along the depth of the coatings, which affects the phase constituents and microstructure of different regions in the coatings. Fig. 7 shows the schematic drawing of Ti element distribution in the coatings. The content of Ti element in the upper part of the coating is relatively low, and when the relative atomic percentage
of Ti element ranges from 50% to 66.7%, as shown in A zone (Fig. 5 and Fig. 7), the reaction products are TiNi and Ti2Ni. On the contrary, the lower region of the coatings is rich in Ti element. When the relative atomic percentage of Ti element ranges from 66.7% to 100%, as shown in B zone (Fig. 5 and Fig. 7), the ultimate reaction products are Ti2Ni and α(Ti) in which a certain number of elements such as Al, Ni and W are dissolved. Fig. 8(a) and (b) clearly reveal the morphology of each phase in the coatings. It can be found that a certain number of white coarse equiaxed particles with the average size of approximately 6 μm and numerous dark fine spherical particles with the diameter of 1–4 μm are uniformly distributed in the matrix composed of the
Fig. 9. EDS analysis of different phases: (a) white equiaxed particle labeled A, (b) dark spherical particle labeled B, (c) gray matrix phase labeled C, and (d) light gray phase labeled D.
Fig. 10. Schematic drawing of the formation of W2C.
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gray phase and the light gray phase. Furthermore, a thin layer of bright white phase with a thickness of approximately 0.4 μm is observed around the white particles. Meanwhile, by contrast with the cross-sectional micrographs before and after heat treatment, as shown in Fig. 8(a) and (b), it shows that microstructure of the coatings has no obvious difference. That is to say, the effect of the heat treatment on microstructure of the coating is almost negligible. Fig. 8(c) shows the morphology of the coating with the heat treatment for 2 h at a higher magnification. The phases with different morphological features were analyzed by EDS. The results, as shown in Table 1 and Fig. 9, indicate that the white particles (marked A in Fig. 8(c)) are mainly composed of W and C and the calculated atomicity ratio of W to C is approximately 1:1. The phases should be the WC. Similarly, the fine spherical particles mainly consist of Ti and C and the atomicity ratio of Ti to C is also approximately 1:1, are determined as in-situ synthesized TiC (marked B in Fig. 8(c)). The gray phase (marked C in Fig. 8(c)) and the light gray phase (marked D in Fig. 8(c)) as the matrix contain a high concentration of Ti and Ni, and the ratio of Ti to Ni atoms are approximately 2:1 and 1:1, respectively, are identified as the Ti2Ni and TiNi intermetallic compounds. It is very difficult to determine accurately the chemical compositions of the thin layer around the WC particles by the spot analysis of EDS due to the restriction of the electron beam spot diameter. Fig. 8(d) shows the BSE image of the WC particle. It can be seen that some fine spherical TiC particles are close to the WC particle. More importantly, the thin layer along the edge of the WC particles can still be observed and are brighter and whiter than the WC particles to form a white-bright zone. Based on the imaging mechanism of back scattered electrons, the brightness of the BSE image is increased with the increase in average atomic number of the phase. It indicates that the white-bright zone has a higher atomic number than that of WC. According to Refs. [23,34], W2C phase was generated in the edge of original WC particles during laser
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cladding. Moreover, the XRD analysis result also shows that the W2C phase is present in the coatings. As a result, the white-bright phase can be confirmed as W2C. The formation mechanism of W2C will be discussed in the subsequent analyses. Fig. 10 shows the schematic drawing of formation of the W2C. During laser cladding, the protruding parts around some largesized WC particles might preferentially be melted due to its high melting temperature (2870 1C). The partial melting will result in the formation of the carbon-rich and tungsten-rich liquid zone around the WC particle. Compared with W atoms, carbon atoms have a higher diffusion coefficient due to their smaller atomic radius [35], which indicates that carbon atoms can diffuse rapidly outwards into the matrix. Moreover, the standard Gibbs free energy of formation of TiC is more negative (Fig. 6), which
Fig. 12. Microhardness profiles along depth direction of the coatings.
Fig. 11. SEM image and map distribution of Ti and Ni elements: (a) microstructure at the interface part of the coating without the heat treatment, (b) Ti elemental distribution, and (c) Ni elemental distribution.
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Fig. 13. Optical images of the indentations from the coatings: (a) without the heat treatment, (b) with the heat treatment at 500 1C for 1 h, and (c) with the heat treatment at 500 1C for 2 h.
Table 2 The wear mass loss and volume loss of the Ti6Al4V and the coatings. Samples
Wear mass loss/ mg
Wear volume loss/ mm3
The substrate The coating without the heat treatment The coating annealed for 1 h The coating annealed for 2 h
5.2 2.1
0.94 0.31
2.4 3.3
0.32 0.36
coating without the heat treatment were analyzed by SEM and EDS. It can be seen that the part mainly consists of dendrites and interdendritic cellular crystals, as shown in Fig. 11(a). The interdendritic structure is rich in Ti element while poor in Ni element, as shown in Fig. 11(b) and (c), thus was confirmed to be α(Ti). Fig. 14. The changes in friction coefficient of the coatings with sliding time.
Fig. 15. Surface profiles across the wear scars of the substrate and the coatings.
indicates that the titanium may preferentially react with carbon to form TiC. The formation of TiC can further promote the diffusion of C atoms into the matrix. As a result, the liquid zone around the WC particle is rich in tungsten while poor in carbon. W2C was formed in the subsequent cooling process. To confirm the existence of α(Ti) in the lower part of the coatings, microstructure and elemental distribution at the interface part of the
3.2. Microhardness and fracture toughness The microhardness profiles along the depth direction of the coatings are shown in Fig. 12. It can be clearly seen that the profile are divided into three zones, corresponding to the coating, the transition zone, and the substrate. Thickness of the coatings is approximately 1.1 mm. And the coating without the heat treatment has an average hardness of approximately 782 HV0.1, which is approximately 2 times that of the substrate (approximately 382 HV0.1). The high hardness of the coating is attributed to the presence of TiNi and Ti2Ni intermetallic compounds and dispersion hardening of some reinforcements, such as WC, W2C and TiC. Besides, the average microhardness values of the coatings heated for 1 h and 2 h are 710 HV0.1 and 687 HV0.1, respectively. It is worth noting that there is a hardness value drop of approximately 70– 100 HV0.1 after the heat treatment. Moreover, with the increase in heat treatment time, the average microhardness of the coatings can be further decreased. Fracture toughness of the coatings before and after the heat treatment was measured by the Vickers indentation method. When a certain load is applied to the surfaces of the samples, the corners of the indentations give rise to the stress concentration and cracks are easily generated and propagated. Optical images of Vickers indentations obtained from the cross-sections of the coatings are indicated in Fig. 13. Obviously, cracks length of the coatings with the heat treatment is less than that of the coating without the heat treatment. According to the formula (1), fracture toughness of the unheated and heated coatings at 500 1C for 1 h and 2 h is 2.77 MPa m1/2, 3.80 MPa m1/2 and 4.43 MPa m1/2, respectively.
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Comparatively speaking, the fracture toughness of the coatings after the heat treatment is obviously elevated with the heat treatment time prolonged. It indicates that the heat treatment can reduce the cracking susceptibility of the coatings to some extent. Residual stress can be introduced during laser cladding due to its characteristic of rapid heating and cooling, mainly consisting of thermal stress and structural stress. According to previous reports [19,36], the residual stress generated in the laser cladded coating belongs to the tensile stress, which improves the cracking susceptibility of the coating. As previously mentioned, effects of the heat treatment on phase constituents and microstructure of the coatings are almost negligible. However, in the annealed state, thermal motion of atoms could occur and cause the reduction in density of defects markedly. As a result, the residual stress was relieved, in accordance with Ref. [37], and then hardness and fracture toughness of the coatings were decreased. Based on the above analyses, the appropriate heat treatment time (1 h) is beneficial to the decrease of the cracking susceptibility on the premise of appropriate reduction of hardness. 3.3. Wear resistance The variation of friction coefficients of the coatings against the hardened Si3N4 ball with sliding time is shown in Fig. 14. It can be clearly observed that the friction coefficients of all the coatings are increased with the increase in sliding time. In addition, with the heat treatment time's increase, the average friction coefficients of the coatings present a decreasing trend and the friction coefficient curves are more stable, especially the coating heated at 500 1C for 2 h presents a lowest value of approximately 0.41, far less than the coating without heat treatment (approximately 0.58). The atmosphere may be related to the surface characteristics of the coatings. With the increase in fracture toughness of the coatings with the heat treatment, the hard particles are easily imbedded in the relative soft matrix under a certain stress during sliding, which
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contributes to the improvement of the surface micro-flatness and the reduction of the sliding resistance and hence, the value and the fluctuation of the friction coefficient are decreased. The surface profiles across the wear scars are shown in Fig. 15. Wear mass loss and volume loss of the coatings and the substrate are shown in Table 2. In comparison with the substrate, all the coatings with less wear mass loss and volume loss present the outstanding wear resistance. Moreover, with the increase in heat treatment time, both wear mass loss and volume loss present an increasing trend, which is mainly attributed to the hardness difference of the coatings (Fig. 12). Compared with the coating without the heat treatment, wear mass loss of the coating heated for 1 h presents a slight increase and the relative wear mass loss and wear volume loss are 14% and 3%, respectively. However, when the heat treatment time is increased to 2 h, the wear mass loss is markedly increased, and the relative wear mass loss and wear volume loss are increased to 57% and 16%, respectively. Fig. 16 shows SEM images of the worn surfaces of the substrate and the coatings. It is noted that the worn surface of the substrate shows a great number of plowing grooves coupled with severe plastic deformation along the edges of the wear scars. The surfaces of the coatings also present obvious features of micro-cutting and plowing grooves. However, the width and depth of the plowing grooves are decreased and no obvious signs of plastic deformation are visible. It can be concluded that the dominant wear mechanism of the coatings is almost unchanged and the difference of the worn surfaces is not significant whether or not the coatings were heated. Wear debris of the coatings heated for 2 h mainly consists of tiny powders and were pressed to form flake-like chips under repeated sliding, as shown in Fig. 17. The EDS analysis results indicate that wear debris is rich in Ti and Ni and consists of a small amount of W, C and Al. It can be confirmed that wear debris is mainly composed of the TiNi–Ti2Ni dual-phase intermetallic compounds and a few WC, W2C and TiC particles. No Si3N4 is found due to its high hardness of approximately 1700 HV when compared with that of the coatings (680–780 HV).
Fig. 16. SEM images of the worn surfaces of (a) the substrate, (b) the coatings without the heat treatment, (c) the coating with the heat treatment for 1 h, and (d) the coating with the heat treatment for 2 h.
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Fig. 17. SEM image and EDS result of wear debris of the coating heated for 2 h.
These observations reveal that the appropriate heat treatment (heated at 500 1C for 1 h) can significantly reduce the value and fluctuation of the friction coefficient on the premise that wear mass and volume are decreased slightly. Moreover, the improvement of the coating's fracture toughness makes a positive contribution to preventing the coating from micro-/macro-cracking or micro-/macro-fracturing [38]. Hence, the appropriate heat treatment (heated at 500 1C for 1 h) will help to improve the wear resistance of the coatings during the long-term service.
4. Conclusion (1) The laser cladded composite coatings with the thickness of approximately 1.1 mm were fabricated on the Ti6Al4V alloy substrate by using the optimized processing parameters, and microstructure mainly consisted of some equiaxed WC particles surrounded by the white-bright W2C, a great number of fine spherical TiC particles, and the matrix composed of the α (Ti), Ti2Ni, TiNi phases. Effects of the post-heat treatment on phase constituents and microstructure of the coatings were not significant. (2) The post-heat treatment decreased the residual stress and microhardness, and increased fracture toughness of the coatings. The average microhardness values of the coatings were reduced from 782 HV0.1 to 710 HV0.1 and 687 HV0.1 with the heat treatment for 1 h and 2 h, respectively. And the fracture toughness of the coatings were improved from 2.77 MPa m1/2 to 3.80 MPa m1/2 and 4.43 MPa m1/2 with the heat treatment for 1 h and 2 h, respectively. (3) The dominant wear mechanism for all the coatings was abrasive wear, characterized by micro-cutting and microplowing. The appropriate heat treatment (heated at 500 1C for 1 h) can significantly reduce the value and fluctuation of the friction coefficient, and had small effects on wear mass and volume loss. The improvement in fracture toughness will be helpful to improve the coating's resistance to micro-/macrocracking or micro-/macro-fracturing and improve the wear resistance of the coating during the long-term service.
Acknowledgments This work was financially supported by the National Natural Science Foundation of China (51002093) and “Shu Guang” project
of Shanghai Municipal Education Commission and Shanghai Education Development Foundation (12SG44).
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Contents lists available at ScienceDirect
Optics & Laser Technology journal homepage: www.elsevier.com/locate/optlastec
Effects of post-heat treatment on microstructure and properties of laser cladded composite coatings on titanium alloy substrate G.J. Li, J. Li n, X. Luo School of Materials Engineering, Shanghai University of Engineering Science, Shanghai 201620, China
art ic l e i nf o
a b s t r a c t
Article history: Received 12 November 2013 Received in revised form 11 June 2014 Accepted 2 July 2014
The composite coatings were produced on the Ti6Al4V alloy substrate by laser cladding. Subsequently, the coatings were heated at 500 1C for 1 h and 2 h and then cooled in air. Effects of post-heat treatment on microstructure, microhardness and fracture toughness of the coatings were investigated by scanning electron microscopy (SEM), X-ray diffraction (XRD), energy dispersive spectroscopy (EDS), optical microscopy (OM). Wear resistance of the coatings was evaluated under the dry sliding reciprocating friction condition at room temperature. The results indicated that the coatings mainly consist of a certain amount of coarse white equiaxed WC particles surrounded by the white-bright W2C, a great deal of fine dark spherical TiC particles and the matrix composed of the α(Ti), Ti2Ni and TiNi phases. Effects of the post-heat treatment on phase constituents and microstructure of the coatings were almost negligible due to the low temperature. However, the post-heat treatment could decrease the residual stress and increase fracture toughness of the coatings, and fracture toughness of the coatings was improved from 2.77 MPa m1/2 to 3.80 MPa m1/2 and 4.43 MPa m1/2 with the heat treatment for 1 h and 2 h, respectively. The mutual role would contribute to the reduction in cracking susceptibility. Accompanied with the increase in fracture toughness, microhardness of the coatings was reduced slightly. The dominant wear mechanism for all the coatings was abrasive wear, characterized by micro-cutting or micro-plowing. The heat treatment could significantly decrease the average friction coefficient and reduce the fluctuation of the friction coefficient with the change in sliding time. The appropriate heat treatment time (approximately 1 h) had a minimal effect on wear mass loss and volume loss. Moreover, the improvement in fracture toughness will also be beneficial to wear resistance of the coatings under the long service. & 2014 Elsevier Ltd. All rights reserved.
Keywords: Laser cladding Post-heat treatment Wear resistance
1. Introduction Titanium and its alloys are widely used in chemical industries, aerospace and medical applications owing to their excellent corrosion resistance, high toughness and exceptional strengthto-weight ratio. However, the low surface hardness and poor wear resistance of titanium alloys limited severely their further application in many industry fields [1,2]. Therefore, it is essential to improve the wear resistance of titanium alloys using surface modification techniques, for instances, physical or chemical vapor deposition and plasma spraying [3–5]. Nevertheless, these methods have their limitations, such as the poor bonding strength between coating and substrate. As a promising surface modification technology, laser cladding has unique properties on wear resistance, corrosion resistance, and metallurgical bonding between the coating and the substrate compared with these
n
Corresponding author. Tel.: þ 86 21 67791198; fax: þ 86 21 67791377. E-mail address: [email protected] (J. Li).
http://dx.doi.org/10.1016/j.optlastec.2014.07.003 0030-3992/& 2014 Elsevier Ltd. All rights reserved.
conventional surface modification technologies [6]. Over the past few years, laser cladding composite coating onto pure titanium or titanium alloy substrate has attracted extensive attention owing to the higher efficiency, excellent metallurgical bonding in the interface and easy-controlled process and the excellent wear properties of the coatings [7]. Kulka et al. [8] obtained composite boride layers composed of laser-borided re-melted zone (TiB, TiB2 and Tiα0 -phase), heat affected zone (Tiα0 -phase) and the substrate (Tiα-phase) using laser-boriding with boron on cylindrical surface of commercially pure titanium substrate, the composite layers present excellent wear resistance in comparison with commercially pure titanium. Savalani et al. [9] produced TiC reinforced Ti matrix composite layers on pure titanium substrate by laser cladding mixed powders of Ti and carbon-nanotube of different contents, and found that the coatings with higher carbon-nanotube content had better wear resistance. Li et al. [10] obtained in situ synthesized TiN and TiB particulatereinforced metal matrix composite coating with hardness of 800 HV–1200 HV by laser cladding with a Ti/h–BN powder mixture, and the composite coating present better wear
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resistance. Our research team [11–13] fabricated titanium-based composite coating reinforced by in situ synthesized TiB whiskers and TiC particles on Ti6Al4V by laser cladding. The above studies mainly focused on the Ti matrix composite coating, some works regarding Ni matrix composite coating on titanium substrate were also carried out. Guo et al. [14] produced NiCrBSi and NiCrBSi/WC–Ni composite coatings on pure Ti substrates by laser cladding and investigated the friction and wear behavior at elevated temperatures of 300 1C and 500 1C, and found the wear rate of the NiCrBSi/WC–Ni composite coating was approximately 24–47 times lower than that of the pure Ti substrate in the same case. A process of laser cladding NiCoCrAlY coating on Ti6Al4V substrate with pre-placed NiCoCrAlY powder was studied by Meng et al. [15], the average microhardness of the coating was 800 HV–1000 HV, which was two times higher than that of Ti6Al4V substrate. The above researches indicate that laser cladding can improve significantly hardness and wear resistance of titanium alloys. However, laser cladding composite coating also has some disadvantages, e.g. high cracking susceptibility, which critically restricts its extensive applications in industry. Cracks may occur in the coating during or after laser cladding, caused by high thermal stress and microstructural stress occurred during rapid heating and rapid solidification [16]. Once cracks have been generated, crack propagation will be difficult to deal with in the subsequent application process, even resulting in the sample's failing. Therefore, cracks have a great influence on the quality of the laser cladded coating. How to reduce the cracking susceptibility of the coating has attracted extensive and continuous attention [16–18]. The generation of cracks can be attributed to the presence of the residual stress in the cladded coating. Based on this, many researchers have proposed many creative methods to reduce cracking susceptibility of the coating, such as preheating the substrate [19], adding rare elements and alloying elements to the clad powder [16,20,21], optimizing the processing parameters [22], etc. Besides, considering many important components by preheating the substrate are often subjected to the damage, Zhou et al. [23] put forward laser induction hybrid rapid cladding (LIHRC), and obtained the composite coatings which were free of cracks and had a good metallurgical bonding with substrate. Based on the principle of crack generation and expansion during laser cladding, Wang et al. [24] proposed adding 316 stainless steel net with low yield strength and good plasticity in the coatings, and crack density was significantly reduced. The former researches were mostly focused on preventing from producing of cracks in laser cladding process. However, post-treatment after laser cladding, aims to adjust microstructure and relieve residual stress generated in laser cladding process to avoid cracking in the subsequent application process [25]. In previous reports, postheat treatment was widely used to improve the required properties of the bulk materials by refining crystal, relieving residual stresses and even changing the microstructure [26]. A few investigations reported the effect of the heat treatment on the laser cladded composite coating. Chen et al. [27] investigated the influence of post-cladding stress-relieving treatments on the residual stresses in laser clad AISI P20 tool steel on the prehardened wrought P20 substrate, and the results indicated that the heat treatment significantly influenced the magnitude and sign of the residual stress by changing the volume fraction of retained austenite in the coating. Liu et al. [28] investigated the effect of tempering treatment on the corrosion resistance and microhardness of the Ni60CuMoW composite coatings on 45 steel surfaces, and the results showed that the tempering treatment could improve microhardness and the maximum self-corrosion potential and reduce corrosion current density significantly. Previous studies mainly focused on the effects of the heat
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treatment on residual stress, microstructure and corrosion resistance. However, very few studies reported the combined influence of post-heat treatment on cracking susceptibility and wear resistance of the laser cladded coating. Hence, in this study, the authors attempted to explore and discuss the effects of the post-heat treatment on microstructure and properties of the laser cladding coating on titanium alloy substrate, especially focused on hardness and fracture toughness. On this basis, the effect of the post-heat treatment on wear resistance was further studied.
2. Experimental procedures Ti6Al4V cylindrical samples with 50 mm in diameter and 10 mm in height were used as the substrate. Before laser cladding, in order to ensure good surface finish, all the samples were ground with emery papers and ultrasonically cleaned in acetone. Commercial powder blends of pure Ni and pure WC in nominal chemical constituents (mass fraction, %) of 70Ni–30WC were selected as the cladding materials. The powder blends after fully mixed and drying treatment were pre-placed on the Ti6Al4V substrate with the thickness of approximately 1 mm and were compressed with a press force of 50000 N by a hydraulic press instrument in order to make the preplaced blends more dense. Then single track laser cladding was carried out by an HL-5000 type CO2 laser source with a power of 2.5 kW, a beam diameter of 4 mm and a scanning speed of 5 mm/s. Fig. 1 shows the schematic drawing of laser cladding. After laser cladding, the specimens were divided into three groups. One group was not disposed and remained the original status as cladding; others were heated respectively at 500 1C for 1 h and for 2 h in the muffle furnace, and cooled in air. Microstructure was observed by means of JSM6460 scanning electron microscope (SEM) equipped with EDAX GENESIS energy dispersive spectroscopy (EDS). Microhardness along the depth of the cross-section was measured by using an HXD-1000TM microhardness tester. The load used was 100 g and loading time was set at 15 s. Fracture toughness, representing the cracking susceptibility of the coating, was measured by the Vickers indentation method. Vickers indentations were prepared on the cross-sections of the coatings using an HV-120 Vickers-hardness with the load of 20 N. Morphologies of the indentations were observed immediately using the VHX-600K optical microscope (OM). Fig. 2 shows the schematic drawing of cracks around the Vickers indentation. Fracture toughness of the coatings was calculated by the following
Fig. 1. The schematic drawing of laser cladding.
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3. Results and discussion 3.1. Microstructure characterization Fig. 4 presents the XRD patterns of the coatings with and without the heat treatment. According to the indexed results of all diffraction peaks in terms of JCPDS, it can be seen that phase constituents of the coatings are almost unchanged before and after the heat treatment, which consist of α(Ti), Ti2Ni, TiNi, WC, W2C and TiC. Ti2Ni, TiNi, TiC and W2C as the new phases were in-situ synthesized during laser cladding. No Ti element was introduced into the preplaced layer, and the Ti element in the coatings was derived from the substrate. During laser cladding, both the surface layer of the substrate and the preplaced layer were melted rapidly due to their low melting points (approximately 1670 1C and 1455 1C), respectively, and prodigious amounts of Ti element came into the molten pool. According to the Ni–Ti binary phase diagram [32] (Fig. 5), when the temperature is below 765 1C, phase transition reactions cannot be occurred and no new precipitation phases are generated. Therefore the heat treatment has little effect on the phase
Fig. 2. Schematic drawing of cracks around the Vickers indentation.
Fig. 3. The schematic drawing of the friction test on the different views: (a) front view and (b) left view.
equation [29]: a P K IC ¼ 0:079 3=2 log 4:5 c a
ð1Þ
in which KIC is fracture toughness (MPa m1/2), P is the indentation load (N), a is half a length of the diagonal line of an indentation (m) and c is the sum of half indentation diagonal and crack length (m). In this study, dry sliding reciprocating friction tests of the Ti6Al4V substrate and the single-track coatings were carried out using the ball-on-disc mode on a CFT-I comprehensive friction and wear apparatus at room temperature. The aim of using the singletrack coatings, not the overlapping coatings, was to eliminate the effect of the overlapping. Microstructure may become coarse in the overlapping zone since the zone is subjected to the second laser radiation during laser cladding. That is to say, the overlapping can affect microstructure, further affect the mechanical properties of the coatings [30,31]. Hence, the single-track coatings were used in this study. Fig. 3 shows the schematic drawing of the friction test on the different views. The coatings' surface is rough and may affect the testing results. Therefore the surface parts of the coatings with a thickness of approximately 0.2 mm were removed by the spark machining method, and then were polished with 300grit SiC abrasive paper prior to the wear tests. Counterbody balls were made of Si3N4 of approximately 4 mm in diameter with a hardness of 1700 HV. The applied load was 20 N. The sliding speed was maintained at 5 m/min, and the sliding time was approximately 20 min. The friction coefficient was recorded by the computer connected to the tester. The wear volume loss and mass loss of the substrate and the coatings were measured by a surface mapping profiler coupled with the wear tester and a photoelectric balance (0.01 mg resolution), respectively. Besides, worn surface morphologies and wear debris of the coatings were investigated by SEM and EDS.
Fig. 4. XRD patterns of the coatings with and without the heat treatment.
Fig. 5. The Ni–Ti binary phase diagram.
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constituents of the coatings due to the low heat treatment temperature of 500 1C. In order to understand the formation mechanism of new phases in the coatings, thermodynamics calculation was applied. Based on the above XRD results, the reactions between the Ni–Ti, Ti–C and W–C systems were considered. All possible reactions occurring during laser cladding are as following:
Fig. 6. Relationship between standard Gibbs free energy and temperature for the possible reactions during laser cladding.
NiþTi ¼TiNi
(2)
Niþ2Ti ¼ Ti2Ni
(3)
TiþC ¼ TiC
(4)
WþC ¼W2C
(5)
Based on the relative data in Ref. [33], the standard Gibbs free energies of formation of those compounds at different temperatures are presented in Fig. 6. It is obvious that the standard Gibbs free energies of formation of those compounds are negative below 1100 K, which indicates that the formation reactions of those compounds are spontaneous, and the spontaneous trend of these compounds at the same temperature is as follows: TiC 4 Ti2Ni4TiNi 4W2C. Based on the Ni–Ti binary phase diagram (Fig. 5), it is impossible to obtain α(Ti), Ti2Ni and TiNi phases simultaneously when the content of Ti element is constant and uniformly Table 1 Chemical composition (at%) of the phases with different morphologies.
Fig. 7. Schematic drawing of Ti element distribution in the coatings.
Analysis zones
Ti
Ni
Al
W
C
V
A B C D
9.21 42.62 53.03 43.78
– 3.94 30.75 48.76
– 1.20 11.04 6.51
41.21 1.95 1.5 –
49.59 50.29 – –
– – 3.68 0.93
Fig. 8. Microstructure of the upper part of the coatings: (a) without the heat treatment, (b) with the heat treatment for 2 h, (c) morphological features of different phases: white equiaxed particle, dark spherical particle, gray phase and light gray phase are labeled A, B, C and D, respectively, and (d) morphology of the WC particles.
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distributed in the coatings. However, the XRD results show that the three phases are present simultaneously in the coatings, which indicates that chemical composition distribution is non-uniform in the coatings. As is well known, laser cladding is a transient process. Due to the quick cooling process of the molten pool, the diffusion and convection of Ti element from the substrate in the molten pool were not sufficient and presents the gradient distribution along the depth of the coatings, which affects the phase constituents and microstructure of different regions in the coatings. Fig. 7 shows the schematic drawing of Ti element distribution in the coatings. The content of Ti element in the upper part of the coating is relatively low, and when the relative atomic percentage
of Ti element ranges from 50% to 66.7%, as shown in A zone (Fig. 5 and Fig. 7), the reaction products are TiNi and Ti2Ni. On the contrary, the lower region of the coatings is rich in Ti element. When the relative atomic percentage of Ti element ranges from 66.7% to 100%, as shown in B zone (Fig. 5 and Fig. 7), the ultimate reaction products are Ti2Ni and α(Ti) in which a certain number of elements such as Al, Ni and W are dissolved. Fig. 8(a) and (b) clearly reveal the morphology of each phase in the coatings. It can be found that a certain number of white coarse equiaxed particles with the average size of approximately 6 μm and numerous dark fine spherical particles with the diameter of 1–4 μm are uniformly distributed in the matrix composed of the
Fig. 9. EDS analysis of different phases: (a) white equiaxed particle labeled A, (b) dark spherical particle labeled B, (c) gray matrix phase labeled C, and (d) light gray phase labeled D.
Fig. 10. Schematic drawing of the formation of W2C.
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gray phase and the light gray phase. Furthermore, a thin layer of bright white phase with a thickness of approximately 0.4 μm is observed around the white particles. Meanwhile, by contrast with the cross-sectional micrographs before and after heat treatment, as shown in Fig. 8(a) and (b), it shows that microstructure of the coatings has no obvious difference. That is to say, the effect of the heat treatment on microstructure of the coating is almost negligible. Fig. 8(c) shows the morphology of the coating with the heat treatment for 2 h at a higher magnification. The phases with different morphological features were analyzed by EDS. The results, as shown in Table 1 and Fig. 9, indicate that the white particles (marked A in Fig. 8(c)) are mainly composed of W and C and the calculated atomicity ratio of W to C is approximately 1:1. The phases should be the WC. Similarly, the fine spherical particles mainly consist of Ti and C and the atomicity ratio of Ti to C is also approximately 1:1, are determined as in-situ synthesized TiC (marked B in Fig. 8(c)). The gray phase (marked C in Fig. 8(c)) and the light gray phase (marked D in Fig. 8(c)) as the matrix contain a high concentration of Ti and Ni, and the ratio of Ti to Ni atoms are approximately 2:1 and 1:1, respectively, are identified as the Ti2Ni and TiNi intermetallic compounds. It is very difficult to determine accurately the chemical compositions of the thin layer around the WC particles by the spot analysis of EDS due to the restriction of the electron beam spot diameter. Fig. 8(d) shows the BSE image of the WC particle. It can be seen that some fine spherical TiC particles are close to the WC particle. More importantly, the thin layer along the edge of the WC particles can still be observed and are brighter and whiter than the WC particles to form a white-bright zone. Based on the imaging mechanism of back scattered electrons, the brightness of the BSE image is increased with the increase in average atomic number of the phase. It indicates that the white-bright zone has a higher atomic number than that of WC. According to Refs. [23,34], W2C phase was generated in the edge of original WC particles during laser
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cladding. Moreover, the XRD analysis result also shows that the W2C phase is present in the coatings. As a result, the white-bright phase can be confirmed as W2C. The formation mechanism of W2C will be discussed in the subsequent analyses. Fig. 10 shows the schematic drawing of formation of the W2C. During laser cladding, the protruding parts around some largesized WC particles might preferentially be melted due to its high melting temperature (2870 1C). The partial melting will result in the formation of the carbon-rich and tungsten-rich liquid zone around the WC particle. Compared with W atoms, carbon atoms have a higher diffusion coefficient due to their smaller atomic radius [35], which indicates that carbon atoms can diffuse rapidly outwards into the matrix. Moreover, the standard Gibbs free energy of formation of TiC is more negative (Fig. 6), which
Fig. 12. Microhardness profiles along depth direction of the coatings.
Fig. 11. SEM image and map distribution of Ti and Ni elements: (a) microstructure at the interface part of the coating without the heat treatment, (b) Ti elemental distribution, and (c) Ni elemental distribution.
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Fig. 13. Optical images of the indentations from the coatings: (a) without the heat treatment, (b) with the heat treatment at 500 1C for 1 h, and (c) with the heat treatment at 500 1C for 2 h.
Table 2 The wear mass loss and volume loss of the Ti6Al4V and the coatings. Samples
Wear mass loss/ mg
Wear volume loss/ mm3
The substrate The coating without the heat treatment The coating annealed for 1 h The coating annealed for 2 h
5.2 2.1
0.94 0.31
2.4 3.3
0.32 0.36
coating without the heat treatment were analyzed by SEM and EDS. It can be seen that the part mainly consists of dendrites and interdendritic cellular crystals, as shown in Fig. 11(a). The interdendritic structure is rich in Ti element while poor in Ni element, as shown in Fig. 11(b) and (c), thus was confirmed to be α(Ti). Fig. 14. The changes in friction coefficient of the coatings with sliding time.
Fig. 15. Surface profiles across the wear scars of the substrate and the coatings.
indicates that the titanium may preferentially react with carbon to form TiC. The formation of TiC can further promote the diffusion of C atoms into the matrix. As a result, the liquid zone around the WC particle is rich in tungsten while poor in carbon. W2C was formed in the subsequent cooling process. To confirm the existence of α(Ti) in the lower part of the coatings, microstructure and elemental distribution at the interface part of the
3.2. Microhardness and fracture toughness The microhardness profiles along the depth direction of the coatings are shown in Fig. 12. It can be clearly seen that the profile are divided into three zones, corresponding to the coating, the transition zone, and the substrate. Thickness of the coatings is approximately 1.1 mm. And the coating without the heat treatment has an average hardness of approximately 782 HV0.1, which is approximately 2 times that of the substrate (approximately 382 HV0.1). The high hardness of the coating is attributed to the presence of TiNi and Ti2Ni intermetallic compounds and dispersion hardening of some reinforcements, such as WC, W2C and TiC. Besides, the average microhardness values of the coatings heated for 1 h and 2 h are 710 HV0.1 and 687 HV0.1, respectively. It is worth noting that there is a hardness value drop of approximately 70– 100 HV0.1 after the heat treatment. Moreover, with the increase in heat treatment time, the average microhardness of the coatings can be further decreased. Fracture toughness of the coatings before and after the heat treatment was measured by the Vickers indentation method. When a certain load is applied to the surfaces of the samples, the corners of the indentations give rise to the stress concentration and cracks are easily generated and propagated. Optical images of Vickers indentations obtained from the cross-sections of the coatings are indicated in Fig. 13. Obviously, cracks length of the coatings with the heat treatment is less than that of the coating without the heat treatment. According to the formula (1), fracture toughness of the unheated and heated coatings at 500 1C for 1 h and 2 h is 2.77 MPa m1/2, 3.80 MPa m1/2 and 4.43 MPa m1/2, respectively.
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Comparatively speaking, the fracture toughness of the coatings after the heat treatment is obviously elevated with the heat treatment time prolonged. It indicates that the heat treatment can reduce the cracking susceptibility of the coatings to some extent. Residual stress can be introduced during laser cladding due to its characteristic of rapid heating and cooling, mainly consisting of thermal stress and structural stress. According to previous reports [19,36], the residual stress generated in the laser cladded coating belongs to the tensile stress, which improves the cracking susceptibility of the coating. As previously mentioned, effects of the heat treatment on phase constituents and microstructure of the coatings are almost negligible. However, in the annealed state, thermal motion of atoms could occur and cause the reduction in density of defects markedly. As a result, the residual stress was relieved, in accordance with Ref. [37], and then hardness and fracture toughness of the coatings were decreased. Based on the above analyses, the appropriate heat treatment time (1 h) is beneficial to the decrease of the cracking susceptibility on the premise of appropriate reduction of hardness. 3.3. Wear resistance The variation of friction coefficients of the coatings against the hardened Si3N4 ball with sliding time is shown in Fig. 14. It can be clearly observed that the friction coefficients of all the coatings are increased with the increase in sliding time. In addition, with the heat treatment time's increase, the average friction coefficients of the coatings present a decreasing trend and the friction coefficient curves are more stable, especially the coating heated at 500 1C for 2 h presents a lowest value of approximately 0.41, far less than the coating without heat treatment (approximately 0.58). The atmosphere may be related to the surface characteristics of the coatings. With the increase in fracture toughness of the coatings with the heat treatment, the hard particles are easily imbedded in the relative soft matrix under a certain stress during sliding, which
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contributes to the improvement of the surface micro-flatness and the reduction of the sliding resistance and hence, the value and the fluctuation of the friction coefficient are decreased. The surface profiles across the wear scars are shown in Fig. 15. Wear mass loss and volume loss of the coatings and the substrate are shown in Table 2. In comparison with the substrate, all the coatings with less wear mass loss and volume loss present the outstanding wear resistance. Moreover, with the increase in heat treatment time, both wear mass loss and volume loss present an increasing trend, which is mainly attributed to the hardness difference of the coatings (Fig. 12). Compared with the coating without the heat treatment, wear mass loss of the coating heated for 1 h presents a slight increase and the relative wear mass loss and wear volume loss are 14% and 3%, respectively. However, when the heat treatment time is increased to 2 h, the wear mass loss is markedly increased, and the relative wear mass loss and wear volume loss are increased to 57% and 16%, respectively. Fig. 16 shows SEM images of the worn surfaces of the substrate and the coatings. It is noted that the worn surface of the substrate shows a great number of plowing grooves coupled with severe plastic deformation along the edges of the wear scars. The surfaces of the coatings also present obvious features of micro-cutting and plowing grooves. However, the width and depth of the plowing grooves are decreased and no obvious signs of plastic deformation are visible. It can be concluded that the dominant wear mechanism of the coatings is almost unchanged and the difference of the worn surfaces is not significant whether or not the coatings were heated. Wear debris of the coatings heated for 2 h mainly consists of tiny powders and were pressed to form flake-like chips under repeated sliding, as shown in Fig. 17. The EDS analysis results indicate that wear debris is rich in Ti and Ni and consists of a small amount of W, C and Al. It can be confirmed that wear debris is mainly composed of the TiNi–Ti2Ni dual-phase intermetallic compounds and a few WC, W2C and TiC particles. No Si3N4 is found due to its high hardness of approximately 1700 HV when compared with that of the coatings (680–780 HV).
Fig. 16. SEM images of the worn surfaces of (a) the substrate, (b) the coatings without the heat treatment, (c) the coating with the heat treatment for 1 h, and (d) the coating with the heat treatment for 2 h.
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Fig. 17. SEM image and EDS result of wear debris of the coating heated for 2 h.
These observations reveal that the appropriate heat treatment (heated at 500 1C for 1 h) can significantly reduce the value and fluctuation of the friction coefficient on the premise that wear mass and volume are decreased slightly. Moreover, the improvement of the coating's fracture toughness makes a positive contribution to preventing the coating from micro-/macro-cracking or micro-/macro-fracturing [38]. Hence, the appropriate heat treatment (heated at 500 1C for 1 h) will help to improve the wear resistance of the coatings during the long-term service.
4. Conclusion (1) The laser cladded composite coatings with the thickness of approximately 1.1 mm were fabricated on the Ti6Al4V alloy substrate by using the optimized processing parameters, and microstructure mainly consisted of some equiaxed WC particles surrounded by the white-bright W2C, a great number of fine spherical TiC particles, and the matrix composed of the α (Ti), Ti2Ni, TiNi phases. Effects of the post-heat treatment on phase constituents and microstructure of the coatings were not significant. (2) The post-heat treatment decreased the residual stress and microhardness, and increased fracture toughness of the coatings. The average microhardness values of the coatings were reduced from 782 HV0.1 to 710 HV0.1 and 687 HV0.1 with the heat treatment for 1 h and 2 h, respectively. And the fracture toughness of the coatings were improved from 2.77 MPa m1/2 to 3.80 MPa m1/2 and 4.43 MPa m1/2 with the heat treatment for 1 h and 2 h, respectively. (3) The dominant wear mechanism for all the coatings was abrasive wear, characterized by micro-cutting and microplowing. The appropriate heat treatment (heated at 500 1C for 1 h) can significantly reduce the value and fluctuation of the friction coefficient, and had small effects on wear mass and volume loss. The improvement in fracture toughness will be helpful to improve the coating's resistance to micro-/macrocracking or micro-/macro-fracturing and improve the wear resistance of the coating during the long-term service.
Acknowledgments This work was financially supported by the National Natural Science Foundation of China (51002093) and “Shu Guang” project
of Shanghai Municipal Education Commission and Shanghai Education Development Foundation (12SG44).
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