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Phase composition and tribological properties of Ti–Al coatings produced on pure Ti by laser cladding
Applied Surface Science 253 (2007) 9301–9310 www.elsevier.com/locate/apsusc
Phase composition and tribological properties of Ti–Al coatings produced on pure Ti by laser cladding Baogang Guo a,b, Jiansong Zhou a, Shitang Zhang a,b, Huidi Zhou a, Yuping Pu c, Jianmin Chen a,* a
State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou 730000, China b Graduate School, Chinese Academy of Sciences, Beijing 100039, China c Central Iron and Steel Research Institute, Beijing 100081, China Received 1 May 2007; received in revised form 23 May 2007; accepted 23 May 2007 Available online 2 June 2007
Abstract Ti–Al coatings with 14.7, 18.1, 25.2 and 29.7 at.% Al contents were fabricated on pure Ti substrate by laser cladding. The laser cladding Ti– Al coatings were analyzed with X-ray diffraction (XRD), scanning electron microscope (SEM) and X-ray energy dispersive spectroscopy (EDS). It was found that with the increase of Al content, the diffraction peaks shifted gradually to higher 2u values. The laser cladding Ti–Al coatings with 14.7 and 18.1 at.% Al were composed of a-Ti and a2-Ti3Al phases, while those with 25.2 and 29.7 at.% Al were composed of a2-Ti3Al phase. With the increase of Al content, the cross-sectional hardness increased, while the fracture toughness decreased. For the laser cladding Ti–Al coatings, when the Al content was 18.1 at.%, the wear mechanism was adhesive wear and abrasive wear; while when the Al content 25.2 at.%, the wear mechanism was adhesive wear, abrasive wear and microfracture. With the increase of Al content, the wear rate of laser cladding Ti–Al coatings decreased under 1 N normal load, while the wear rate firstly decreased and then increased under a normal load of 3 N. Due to its optimized combination of high hardness and high fracture toughness, the laser cladding Ti–Al coating with 18.1 at.% Al showed the best anti-wear properties at higher normal load. # 2007 Elsevier B.V. All rights reserved. Keywords: Laser cladding; Ti–Al coating; Phase composition; Tribological properties
1. Introduction Titanium and its alloys are extensively used in aeronautical, marine and chemical industries owing to their specific properties, such as low density, high strength, and excellent corrosion resistance [1,2]. However, the low hardness and poor tribological properties of titanium and titanium alloys may become a critical factor when wear phenomena are involved [2]. Therefore, various surface modification technologies have been developed to improve their wear resistance. Among them, laser processing is a promising technique in enlarging their application scope [2,3]. To improve the wear resistance of titanium and its alloys, laser surface coating of a metal matrix composite (MMC) layer
* Corresponding author. Tel.: +86 931 4968018. E-mail address: [email protected] (J. Chen). 0169-4332/$ – see front matter # 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2007.05.056
on the matrix surface have been widely used. It was found that the wear resistance of titanium and its alloy was noticeably enhanced when a TiN/Ti MMC coating was in situ synthesized on their surface by laser gas nitriding in nitrogen atmosphere [4,5]. However, cracking is a common problem associated with laser gas nitriding [5,6], which can be ameliorated by preheating the substrate [5] or using a mixture of argon and nitrogen as the treatment gas at the expense of hardness [6]. Courant et al. [7] in situ synthesized a TiC/C/Ti MMC coating on the surface of Ti–6Al–4V alloy with graphite powder by laser alloying. It was found that the TiC/C/Ti MMC coating showed higher microhardness (TiC as hard phase) and lower friction coefficient (numerous graphite inclusions acting as solid lubricant) compared with the as received sample, and that the wear resistance of the laser treated sample was obviously improved. Tian et al. in situ synthesized a series of TiC/TiB/Ti [8], TiC/Ti5Si3/Ti [9], TiB/TiB2/Ti/NiTi [10] MMC coatings on the surface of titanium and its alloys with C + B, C + Si and
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B + Ni mixed powders respectively by laser alloying. The MMC coatings all showed high hardness and excellent wear resistance due to the formation of TiC, TiB, TiB2, or Ti5Si3 hard ceramic compounds [8–10]. Studies on WC/W2C/TiC/W/Ti MMC coating, fabricated by laser melt injection of WCp particles into Ti–6Al–4V alloy, also showed that the MMC coating demonstrated substantially improved wear resistance in comparison with the substrate, and ‘the most wear resistant’ area was the hard surroundings of the WC particle [11]. Other MMC coatings can also be fabricated by laser melt injection of SiCp [12], TiB2 [13] into titanium alloy, and a substantial improvement in sliding wear resistance under boundary lubrication conditions was observed for TiB2/Ti MMC coatings in comparison with substrate alloys [13]. In addition to MMC coatings, intermetallic matrix composite (IMC) coatings have also been used to improve the wear resistance of titanium alloy. For example, Ti5Si3/NiTi2 [14], Ti2Ni3Si/NiTi [15] IMC coatings were in situ synthesized on the surface of titanium alloy with Ni + Ti + Si mixed powders by laser cladding. Wear tests revealed that the Ti5Si3/ NiTi2 IMC coating exhibited excellent wear resistance under dry sliding wear condition due to the combination of high hardness and strong intermetallic atomic bonding of Ti5Si3 phase, hard and ductile NiTi2 intermetallic matrix and the fine microstructure [14]. The Ti2Ni3Si/NiTi IMC coating also showed excellent abrasive and adhesive wear resistance under dry sliding wear conditions due to the high hardness, strong intermetallic atomic bond, and anomalous hardness–temperature relations of Ti2Ni3Si primary dendrite and the highly ductile NiTi intermetallic phase [15]. In the past studies, Ti–Al coatings with 17–36 at.% Al were in situ synthesized on pure Ti substrate by laser cladding [16–18], and the microstructures of these Ti–Al coatings were systemically investigated by Abboud and West [16]. Unfortunately, there has been scarce studies on the effect of Al content on the friction and wear properties of Ti–Al coatings. In the present study, Ti–Al coatings with 14.7–29.7 at.% Al were in situ synthesized on the surface of pure Ti substrate with Ti + Al mixed powders by laser cladding. And the effects of Al content on the phase composition and tribological properties of these laser cladding Ti–Al coatings were investigated.
(particle size, 74–150 mm, 99% purity) elemental powders were selected as the starting precursor materials for fabricating the Ti–Al coatings by laser cladding. Three homogeneously mixed powders in composition (at.%) of 80Ti–20Al, 70Ti– 30Al and 60Ti–40Al were prepared and then pre-placed on the substrates by an organic binding material with a thickness of approximately 1.0 mm. All the specimens were placed in a furnace at 120 8C for 2 h to vaporize the water moisture in the pre-placed coatings. The powder-bedded specimens were preheated in a furnace at approximately 220 8C before laser cladding treatment and were cooled down slowly to avoid the formation of thermal crack. The laser cladding was conducted on a 10-kW transverse-flow continuous-wave CO2 laser material processing system equipped with a 4-axis computer numerical controlled (CNC) laser material processing machine tool. The laser cladding process was conducted in argon protective atmosphere with a laser beam size of 10 mm 1 mm. The laser processing parameters, the sample labels in this study were shown in Table 1. After laser cladding, the Ti–Al coatings were sectioned and polished for the following tests. Metallographic samples were prepared using standard mechanical polishing procedures and etched in a solution of HF, HNO3 and H2O in volume ratio of 5:5:90 at room temperature for approximately 30–50 s. The coating thickness was defined by JSM-5600LV scanning electron microscopy (SEM) and shown in Table 1. The Al content of each sample was determined by EDS analysis tool attached to SEM, in which six replicable tests were conducted and the averaged value was given in Table 1. The crosssectional hardness and fracture toughness of the laser cladding Ti–Al coatings was measured on an MH-5 Vickers microhardness tester with a dwell time of 5 s. For the microhardness and fracture toughness test, the load was set at 0.49 and 9.8 N, respectively. Fracture toughness KIC of the laser cladding Ti–Al coatings was calculated from Eq. (1), where E is the Young’s modulus, H is the Vickers hardness, P is the indentation load, and c is the crack length [19]. The averages of 20 toughness values were given and the test deviation of toughness was 30%. Prior to the direct indentation tests, all laser cladding Ti–Al coatings were polished to a thickness of 340 mm.
2. Experimental
K IC ¼ 0:016ðE=HÞ1=2 PðcÞ3=2
Polished pure Ti discs (31 mm in diameter, 10 mm in thickness) were used as the substrates. The commercial pure titanium (particle size, 50–74 mm, 99% purity), pure aluminum
The phase presented in the laser cladding Ti–Al coatings was identified using an automatic Philips X’ Pert-MRD X-ray diffractometer (40 kV, 30 mA, Cu Ka radiation) by scanning in
(1)
Table 1 Laser processing parameters, Al contents and thickness of the laser cladding Ti-Al coatings Samples number
1 2 3 4
Powder composition
80Ti–20Al 70Ti–30Al 60Ti–40Al 60Ti–40Al
Laser processing parameters Laser power (kW)
Scanning speed (mm/s)
5.0 5.0 5.0 4.0
4.0 4.0 4.0 3.3
Al content in the coating (at.%)
Thickness (mm)
14.7 1.2 18.1 1.7 25.2 2.1 29.7 2.4
421 672 595 622
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the 2u = 15–908 range. After the XRD tests, Sample 1 and Sample 2 were annealed in a furnace at 650 8C for 100 h in Argon protection [20], and then cooled slowly in the furnace. After the heat treatment, the phase composition of Sample 1 and Sample 2 were then determined by XRD. The aim of heat treatment was to determine the phase composition of Sample 1 and Sample 2. The dry sliding wear tests at room temperature were performed on a reciprocating ball-on-disc UMT-2MT tribometer (Center for Tribology, Inc., California, USA) in air. The counterpart was GCr15 steel balls with hardness of 700 HV and diameter of 3 mm. The wear tests were carried out under a normal load 1 and 3 N at a sliding speed of 0.1 m/s over a period of 30 min. The worn surfaces were analyzed by SEM. The wear volume losses were measured by MicroXAM 3D non-contact surface mapping profiler (ADE Corporation, Massachusetts, USA). Three replicated tests were conducted for each test condition and the averaged value was given. 3. Results and discussions 3.1. Microstructure and composition of laser cladding Ti– Al coatings Table 1 shows the Al contents of the laser cladding Ti–Al coatings by EDS analysis. It can be seen that the Al content was 14.7 1.2, 18.1 1.7, 25.2 2.1 and 29.7 2.4 at.% to Sample 1, Sample 2, Sample 3 and Sample 4, respectively. Fig. 1 shows the surface morphologies of laser cladding Ti– Al coatings. It is clear that many small particles and balls (the balls mainly composed of Ti (Al 6.1 at.%), some nearly composed of Ti) distributed on the surfaces of the laser cladding
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Ti–Al coatings. However, on the surfaces of Sample 3 and Sample 4, the balls were smaller and lesser than those on the surfaces of Sample 1 and Sample 2, which was attributed to the lower Ti content in the Ti + Al mixed powders (see Table 1). The formation of these balls were attributed to the surface tension driven caused by the capillary instability effect [21], and was likely to occur when the molten metal did not wet the underlying substrate due to an oxide layer being present on the substrate and on the surface of the melt [21]. In addition, some long ridges appeared on the surfaces of Sample 3 and Sample 4, which was quite different from those of Sample 1 and Sample 2. Fig. 2 shows the interface morphologies of the laser cladding Ti–Al coatings. As can be seen from Fig. 2, the bonding between the coating and the substrate was of high-quality metallurgical fusion bonding. For Sample 1, its microstructure near the interface was made up of randomly distributed lamellae phase (13.8 at.% Al and 86.2 at.% Ti), while for Sample 2, its microstructure near the interface was composed of lamellae phase (14.8 at.% Al and 85.2 at.% Ti) and equiaxed phase (18.2 at.% Al and 81.8 at.% Ti). For Sample 3 and Sample 4, they had similar microstructure near the interface with honeycomb-like phase (24.6 at.% Al and 75.4 at.% Ti for Sample 3, 28.8 at.% Al and 71.2 at.% Ti for Sample 4). Figs. 3–7 show the XRD patterns of pure Ti and the laser cladding Ti–Al coatings before and after heat treatment. As can be seen from Fig. 3, no pure Al peaks was found in the laser cladding Ti–Al coatings. The absence of diffraction lines of pure aluminum indicated that all Al atoms substituted Ti atoms, which resulted in the decrease of Ti lattice parameters [22]. Therefore, the diffraction peaks shifted gradually to higher 2u values with increasing Al content from pure Ti to Sample 4 (see Fig. 3). As can also be seen from Fig. 3, Sample 3 (25.2 at.%
Fig. 1. Surface morphologies of laser cladding Ti-Al coatings: (a) Sample 1, (b) Sample 2, (c) Sample 3, and (d) Sample 4.
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Fig. 2. Interface morphologies of the laser cladding Ti–Al coatings: (a) Sample 1, (b) Sample 2, (c) Sample 3, and (d) Sample 4.
Al) and Sample 4 (29.7 at.% Al) were composed of a2-Ti3Al phases. However, the phase composition of Sample 1 (14.7 at.% Al) and Sample 2 (18.1 at.% Al) could not be assigned as a-Ti or a2-Ti3Al phase, because their d values had a deviation compared with those of a-Ti or a2-Ti3Al phase (see the ‘‘d values’’ of the strongest peaks marked in Fig. 3). The phases in Sample 1 and Sample 2 were not determined directly by Ti–Al binary phase diagram presented by Murray [23] because of the rapid solidification in the laser cladding process. In order to determine the phase composition of Sample 1 and Sample 2, heat treatment was taken. For Sample 1, before heat treatment, there was only one a2-Ti3Al peak in the lower 2u values (due to the low content of a2-Ti3Al phase in Sample 1,
see Fig. 4), while two new a2-Ti3Al peaks appeared after heat treatment (see Fig. 6). In the case of Sample 2, after heat treatment, the intensity of a2-Ti3Al peaks enhanced (see Fig. 7), which all showed the precipitation of a2-Ti3Al phase during the heat treatment and were similar to the heat treatment behavior (the precipitation of a2-Ti3Al) of Ti–Al alloys with 16.0 at.% Al and 19.2 at.% Al in reference [20]. Furthermore, the XRD patterns of Sample 1 and Sample 2 had no change after heat treatment (no new peak or shoulder peak at higher 2u values appeared after the precipitation of a2-Ti3Al, see Fig. 5) compared with those before heat treatment (see Fig. 3) except the slight increase of d values (see the ‘‘d values’’ of the strongest peaks marked in Figs. 5 and 3) and the enhancing
Fig. 3. XRD patterns of pure Ti and the laser cladding Ti–Al coatings. T: a-Ti; TA: a2-Ti3Al.
Fig. 4. Magnified XRD patterns of pure Ti and the laser cladding Ti–Al coatings in Fig. 3. TA: a2-Ti3Al.
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Fig. 5. XRD patterns of Sample 1 and Sample 2 after heat treatment. T: a-Ti; TA: a2-Ti3Al.
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Fig. 7. XRD patterns of Sample 2 at lower 2u values before and after heat treatment. TA: a2-Ti3Al.
intensities of the peaks (see Fig. 7). Considering the above, conclusions can be drawn that Sample 1 and Sample 2 were composed of mixed a-Ti and a2-Ti3Al phases according to the Ti–Al binary phase diagram [20,23]. As can also be seen from Fig. 4, the intensity of a2-Ti3Al peaks of Sample 2 at lower 2u values was stronger than that of Sample 1, which indicated that the content of a2-Ti3Al phase in Sample 2 was higher than that in Sample 1. Fig. 8 shows the microhardness profile along the depth direction of the laser cladding Ti–Al coatings. It is clear that the laser cladding Ti–Al coatings had a high and wavy hardness distribution within the main coatings except in the coating/ substrate bonding zone where the hardness showed a gradient decrease to the substrate. As can also be seen from Fig. 8, the hardness of the laser cladding Ti–Al coatings increased with the increase of Al content. For the two phase (a-Ti + a2-Ti3Al) Ti– Al coatings (Sample 1 and Sample 2), because of the effective Al solid solution strengthening and the precipitation of the hard a2-Ti3Al phase, their hardness was higher than that of pure Ti.
Moreover, the more Al content, the more effective Al solid solution strengthening and more precipitation of hard a2-Ti3Al phase, therefore, the higher hardness of Sample 2 than that of Sample 1. Sample 3 and Sample 4 were both composed of the hard a2-Ti3Al phase, so their hardness was higher than that of Sample 1 and Sample 2, which was similar to the results presented in reference [24]. Table 2 shows the fracture toughness of laser cladding Ti–Al coatings measured by the direct indentation method. The studies carried out by Evans indicated that the a2-Ti3Al precipitation could lead to the lowering of ductility of titanium alloy [25], that is to say, the brittle a2-Ti3Al phase had a detrimental effect on the ductility of titanium alloy. Therefore, it is reasonable to postulate that the ductility of Sample 2 was lower than that of Sample 1 due to its higher content of brittle a2-Ti3Al phase than that in Sample 1. Moreover, according to the results presented in reference [26], Sample 4 (29.7 at.% Al) was more brittle than Sample 3 (25.2 at.% Al). In summary, the ductility between the laser cladding Ti–Al coatings was in this
Fig. 6. Magnified XRD patterns of Sample 1 and Sample 2 in Fig. 5. TA: a2Ti3Al.
Fig. 8. Microhardness profile across the laser cladding Ti–Al coatings.
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Table 2 The average length of crack and fracture toughness of the laser cladding Ti-Al coatings in the direct indentation method Materials
Average length of crack (mm)
Fracture toughness (MPam1/2)
Pure Ti Sample Sample Sample Sample
No crack 5.6 9.7 22.0 32.1
– 167.8 72.3 21.0 11.7
1 2 3 4
order: Sample 1 > Sample 2 > Sample 3 > Sample 4. It is known that the ductility had a critical effect on the fracture toughness of materials [27]. In ductile materials, the material near the tip of the flaw can deform, causing the tip of any crack to become blunt, reducing the stress intensity factor, and preventing growth of the crack [27]. So, in the direct indentation tests, the more brittle the Ti–Al coatings, the easier crack formation or growth under the same applied load, therefore, the fracture toughness of Ti–Al coatings decreased with increasing Al contents due to their decreasing ductility (see Table 2). In another study, the fracture toughness of a2Ti3Al based alloy was measured using pre-cracked three-point bend methods in accordance with ASTM-E399 [28]. Unfortunately, this method is not appropriate for these thin laser cladding Ti–Al coatings. The direct indentation method may be not appropriate for measuring the fracture toughness of the laser cladding Ti–Al coatings, but it can give the order of fracture toughness of the laser cladding Ti–Al coatings. 3.2. Tribological properties of laser cladding Ti–Al coatings 3.2.1. Variation of friction coefficient Fig. 9 shows the friction coefficient of pure Ti and the laser cladding Ti–Al coatings sliding against GCr15 ball under 1 and 3 N normal loads. It can be seen that the friction coefficient of the laser cladding Ti–Al coatings varied with increasing Al content. As a whole, the friction coefficient of
Fig. 9. Friction coefficient of pure Ti and the laser cladding Ti–Al coatings sliding against GCr15 steel ball.
laser cladding Ti–Al coatings under 1 N normal load was higher than that of pure Ti, while those under 3 N normal load was lower than that of pure Ti except Sample 4. Under the present experimental conditions, the lowest friction coefficient (0.33) could be obtained under 3 N normal load for Sample 3. From the above, it can be seen that the laser cladding Ti–Al coatings exhibited better friction behavior under higher normal load. 3.2.2. Wear mechanism of laser cladding Ti–Al coatings Fig. 10 shows the worn surfaces of pure Ti and laser cladding Ti–Al coatings sliding against GCr15 steel ball. Under the sliding wear conditions, since the hardness of pure Ti (205 HV) was much lower than that of the GCr15 steel ball (700 HV), it was easy for the hard asperities on the surface of the GCr15 steel ball to penetrate into the contact surface of pure Ti, which resulted in effective micro-cutting (see Fig. 10a and b). Moreover, because of the extremely high affinity of titanium to metallic adhesion, serious adhesion occurred between the contact surface (see Fig. 10a and b). Thus, the wear behavior of pure Ti was featured as abrasive and adhesive wear evidenced by the grooves and adhesive craters on its worn surface. In addition, severe plastic deformation was also visible on its worn surface. The worn surface of Sample 1 was similar to that of pure Ti due to its lower Al content (less Al solid strengthening effect and lower hardness than that of Sample 2). However, because of its higher hardness than pure Ti, the worn surface of Sample 1 was characterized by the presence of shallower grooves and less adhesive craters compared with that on the worn surface of pure Ti (see Fig. 10c and d). For Sample 2, only slight adhesive features were observed on its worn surface due to its higher hardness than that of Sample 1. Besides the adhesive wear, the abrasive wear features were also observed evidenced by the grooves on its worn surface (see Fig. 10e and f). For Sample 3 and Sample 4, under the action of repeating sliding cycles, cracks were firstly formed on the worn surfaces, and then peeled off at places where the crack propagated due to their drastically decreased fracture toughness compared with that of Sample 1 and Sample 2, therefore, microfracutre were observed on their worn surfaces besides adhesive wear and abrasive wear (see Fig. 10g–j). Fig. 11 shows the worn surfaces of GCr15 steel ball sliding against pure Ti, Sample 4 under 3 N normal load and the corresponding distributions of Fe and Ti elements. As can be seen from Fig. 11, during the sliding wear test, pure Ti and the laser cladding Ti–Al coating transferred to the GCr15 steel ball surface and formed the protuberances with high content of Ti element because of adhesiveness, and the adhesive behavior was more severe against pure Ti than against Sample 4. As can also be seen from Fig. 11, the worn surfaces of GCr15 steel ball sliding against pure Ti and Sample 4 also showed abrasive wear features evidenced by slight scratches. In a word, the surface morphologies of the worn steel balls corresponded well to those of the worn surfaces of pure Ti and Sample 4, which was in accordance with the wear mechanism of pure Ti and the laser cladding Ti–Al coatings.
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Fig. 10. Typical worn surfaces of (a) pure Ti; (c) Sample 1; (e) Sample 2; (g) Sample 3; (i) Sample 4 at 1 N load, and (b) pure Ti; (d) Sample 1; (f) Sample 2; (h) Sample 3; (j) Sample 4 at 3 N load (arrows shows the cracks).
3.2.3. Wear volumes of laser cladding Ti–Al coatings Fig. 12 shows the wear volume of pure Ti and laser cladding Ti–Al coatings sliding against GCr15 steel ball under 1 and 3 N normal loads. It is clearly seen that the wear
volumes of all the laser cladding Ti–Al coatings were lower than that of pure Ti, that is to say, the wear resistance of Ti was improved by forming laser cladding Ti–Al coatings on its surface.
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Fig. 11. Worn surfaces of GCr15 steel ball sliding against (a) pure Ti; (b) Sample 4 under 3 N normal load and the corresponding map distributions of Fe and Ti elements, 330.
The variation of wear volume of pure Ti and the laser cladding Ti–Al coatings can be explained by Archard Eq. (2) [29]. V¼
kWS 3H
(2)
where V is the sliding wear volume; W, the normal load; S, the total sliding distance (S); H, the hardness of the wearing
surface; and k is a probability factor that a given area contact will fracture within the weaker material rather than at the original interface. According to Eq. (2), the hardness of laser cladding Ti–Al coatings was higher than that of pure Ti, so the wear volume of the laser cladding Ti–Al coatings was lower than that of pure Ti. And the wear volume at the higher normal load (3 N) was
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program of the light in China’s western region (2006)’’ for financial support. References
Fig. 12. Wear volume of pure Ti and the laser cladding Ti–Al coatings sliding against GCr15 steel ball.
higher than that at lower normal load (1 N) due to the higher load in Eq. (2). Under a lower normal load (1 N), the wear volume of laser cladding Ti–Al coatings decreased with the increase of Al content, which was mainly attributed to the increasing hardness. However, under a higher normal load (3 N), the wear rate firstly decreased due to the increasing hardness, and then increased due to the severe brittle fracture caused by the decreased fracture toughness (the increasing k value in Eq. (2)). Sample 2 had the optimized combination of high hardness and high fracture toughness, so it showed the best anti-wear properties at the higher normal load. 4. Conclusions 1. With the increase of Al content, the diffraction peaks of the laser cladding Ti–Al coatings shifted gradually to higher 2u values. The laser cladding Ti–Al coatings with 14.7 and 18.1 at.% Al were composed of mixed a-Ti and a2-Ti3Al phases, while the laser cladding Ti–Al coatings with 25.2 and 29.7 at.% Al were composed of a2-Ti3Al phases. 2. With the increase of Al content, the cross-sectional hardness increased, while the fracture toughness decreased. 3. For Sample 1 and Sample 2, the wear mechanism was adhesive wear and abrasive wear, while the wear mechanisms shifted to adhesive wear, abrasive wear and microfracture for Sample 3 and Sample 4. 4. With the increase of Al content, the wear rate of laser cladding Ti–Al coatings decreased under 1 N load, while the wear rate firstly decreased and then increased under 3 N normal load. Sample 2 had the optimized combination of high hardness and high fracture toughness, so it had the best anti-wear properties under the higher normal load. Acknowledgements The authors are grateful to the National Natured Science Foundation of China (Grant No.5057521), the Innovative Group Foundation from NSFC (Grant No.50421502), and ‘‘The
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[27] D.R. Askeland, P.P. Phule´, Mechanical properties and behavior, The Science and Engineering of Materials, fourth edition, Brooks/Cole, 2002. [28] S.V. Kamat, A.K. Gogia, D. Banerjee, Effect of alloying elements and heat treatment on the fracture toughness of Ti–Al–Nb alloys, Acta. Mater. 46 (1) (1998) 239–251. [29] P.J. Blau, Glossary of terms, Friction, Lubrication, and Wear Technology, vol. 18, ASM Handbook, ASM International, 1992.
Phase composition and tribological properties of Ti–Al coatings produced on pure Ti by laser cladding Baogang Guo a,b, Jiansong Zhou a, Shitang Zhang a,b, Huidi Zhou a, Yuping Pu c, Jianmin Chen a,* a
State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou 730000, China b Graduate School, Chinese Academy of Sciences, Beijing 100039, China c Central Iron and Steel Research Institute, Beijing 100081, China Received 1 May 2007; received in revised form 23 May 2007; accepted 23 May 2007 Available online 2 June 2007
Abstract Ti–Al coatings with 14.7, 18.1, 25.2 and 29.7 at.% Al contents were fabricated on pure Ti substrate by laser cladding. The laser cladding Ti– Al coatings were analyzed with X-ray diffraction (XRD), scanning electron microscope (SEM) and X-ray energy dispersive spectroscopy (EDS). It was found that with the increase of Al content, the diffraction peaks shifted gradually to higher 2u values. The laser cladding Ti–Al coatings with 14.7 and 18.1 at.% Al were composed of a-Ti and a2-Ti3Al phases, while those with 25.2 and 29.7 at.% Al were composed of a2-Ti3Al phase. With the increase of Al content, the cross-sectional hardness increased, while the fracture toughness decreased. For the laser cladding Ti–Al coatings, when the Al content was 18.1 at.%, the wear mechanism was adhesive wear and abrasive wear; while when the Al content 25.2 at.%, the wear mechanism was adhesive wear, abrasive wear and microfracture. With the increase of Al content, the wear rate of laser cladding Ti–Al coatings decreased under 1 N normal load, while the wear rate firstly decreased and then increased under a normal load of 3 N. Due to its optimized combination of high hardness and high fracture toughness, the laser cladding Ti–Al coating with 18.1 at.% Al showed the best anti-wear properties at higher normal load. # 2007 Elsevier B.V. All rights reserved. Keywords: Laser cladding; Ti–Al coating; Phase composition; Tribological properties
1. Introduction Titanium and its alloys are extensively used in aeronautical, marine and chemical industries owing to their specific properties, such as low density, high strength, and excellent corrosion resistance [1,2]. However, the low hardness and poor tribological properties of titanium and titanium alloys may become a critical factor when wear phenomena are involved [2]. Therefore, various surface modification technologies have been developed to improve their wear resistance. Among them, laser processing is a promising technique in enlarging their application scope [2,3]. To improve the wear resistance of titanium and its alloys, laser surface coating of a metal matrix composite (MMC) layer
* Corresponding author. Tel.: +86 931 4968018. E-mail address: [email protected] (J. Chen). 0169-4332/$ – see front matter # 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2007.05.056
on the matrix surface have been widely used. It was found that the wear resistance of titanium and its alloy was noticeably enhanced when a TiN/Ti MMC coating was in situ synthesized on their surface by laser gas nitriding in nitrogen atmosphere [4,5]. However, cracking is a common problem associated with laser gas nitriding [5,6], which can be ameliorated by preheating the substrate [5] or using a mixture of argon and nitrogen as the treatment gas at the expense of hardness [6]. Courant et al. [7] in situ synthesized a TiC/C/Ti MMC coating on the surface of Ti–6Al–4V alloy with graphite powder by laser alloying. It was found that the TiC/C/Ti MMC coating showed higher microhardness (TiC as hard phase) and lower friction coefficient (numerous graphite inclusions acting as solid lubricant) compared with the as received sample, and that the wear resistance of the laser treated sample was obviously improved. Tian et al. in situ synthesized a series of TiC/TiB/Ti [8], TiC/Ti5Si3/Ti [9], TiB/TiB2/Ti/NiTi [10] MMC coatings on the surface of titanium and its alloys with C + B, C + Si and
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B + Ni mixed powders respectively by laser alloying. The MMC coatings all showed high hardness and excellent wear resistance due to the formation of TiC, TiB, TiB2, or Ti5Si3 hard ceramic compounds [8–10]. Studies on WC/W2C/TiC/W/Ti MMC coating, fabricated by laser melt injection of WCp particles into Ti–6Al–4V alloy, also showed that the MMC coating demonstrated substantially improved wear resistance in comparison with the substrate, and ‘the most wear resistant’ area was the hard surroundings of the WC particle [11]. Other MMC coatings can also be fabricated by laser melt injection of SiCp [12], TiB2 [13] into titanium alloy, and a substantial improvement in sliding wear resistance under boundary lubrication conditions was observed for TiB2/Ti MMC coatings in comparison with substrate alloys [13]. In addition to MMC coatings, intermetallic matrix composite (IMC) coatings have also been used to improve the wear resistance of titanium alloy. For example, Ti5Si3/NiTi2 [14], Ti2Ni3Si/NiTi [15] IMC coatings were in situ synthesized on the surface of titanium alloy with Ni + Ti + Si mixed powders by laser cladding. Wear tests revealed that the Ti5Si3/ NiTi2 IMC coating exhibited excellent wear resistance under dry sliding wear condition due to the combination of high hardness and strong intermetallic atomic bonding of Ti5Si3 phase, hard and ductile NiTi2 intermetallic matrix and the fine microstructure [14]. The Ti2Ni3Si/NiTi IMC coating also showed excellent abrasive and adhesive wear resistance under dry sliding wear conditions due to the high hardness, strong intermetallic atomic bond, and anomalous hardness–temperature relations of Ti2Ni3Si primary dendrite and the highly ductile NiTi intermetallic phase [15]. In the past studies, Ti–Al coatings with 17–36 at.% Al were in situ synthesized on pure Ti substrate by laser cladding [16–18], and the microstructures of these Ti–Al coatings were systemically investigated by Abboud and West [16]. Unfortunately, there has been scarce studies on the effect of Al content on the friction and wear properties of Ti–Al coatings. In the present study, Ti–Al coatings with 14.7–29.7 at.% Al were in situ synthesized on the surface of pure Ti substrate with Ti + Al mixed powders by laser cladding. And the effects of Al content on the phase composition and tribological properties of these laser cladding Ti–Al coatings were investigated.
(particle size, 74–150 mm, 99% purity) elemental powders were selected as the starting precursor materials for fabricating the Ti–Al coatings by laser cladding. Three homogeneously mixed powders in composition (at.%) of 80Ti–20Al, 70Ti– 30Al and 60Ti–40Al were prepared and then pre-placed on the substrates by an organic binding material with a thickness of approximately 1.0 mm. All the specimens were placed in a furnace at 120 8C for 2 h to vaporize the water moisture in the pre-placed coatings. The powder-bedded specimens were preheated in a furnace at approximately 220 8C before laser cladding treatment and were cooled down slowly to avoid the formation of thermal crack. The laser cladding was conducted on a 10-kW transverse-flow continuous-wave CO2 laser material processing system equipped with a 4-axis computer numerical controlled (CNC) laser material processing machine tool. The laser cladding process was conducted in argon protective atmosphere with a laser beam size of 10 mm 1 mm. The laser processing parameters, the sample labels in this study were shown in Table 1. After laser cladding, the Ti–Al coatings were sectioned and polished for the following tests. Metallographic samples were prepared using standard mechanical polishing procedures and etched in a solution of HF, HNO3 and H2O in volume ratio of 5:5:90 at room temperature for approximately 30–50 s. The coating thickness was defined by JSM-5600LV scanning electron microscopy (SEM) and shown in Table 1. The Al content of each sample was determined by EDS analysis tool attached to SEM, in which six replicable tests were conducted and the averaged value was given in Table 1. The crosssectional hardness and fracture toughness of the laser cladding Ti–Al coatings was measured on an MH-5 Vickers microhardness tester with a dwell time of 5 s. For the microhardness and fracture toughness test, the load was set at 0.49 and 9.8 N, respectively. Fracture toughness KIC of the laser cladding Ti–Al coatings was calculated from Eq. (1), where E is the Young’s modulus, H is the Vickers hardness, P is the indentation load, and c is the crack length [19]. The averages of 20 toughness values were given and the test deviation of toughness was 30%. Prior to the direct indentation tests, all laser cladding Ti–Al coatings were polished to a thickness of 340 mm.
2. Experimental
K IC ¼ 0:016ðE=HÞ1=2 PðcÞ3=2
Polished pure Ti discs (31 mm in diameter, 10 mm in thickness) were used as the substrates. The commercial pure titanium (particle size, 50–74 mm, 99% purity), pure aluminum
The phase presented in the laser cladding Ti–Al coatings was identified using an automatic Philips X’ Pert-MRD X-ray diffractometer (40 kV, 30 mA, Cu Ka radiation) by scanning in
(1)
Table 1 Laser processing parameters, Al contents and thickness of the laser cladding Ti-Al coatings Samples number
1 2 3 4
Powder composition
80Ti–20Al 70Ti–30Al 60Ti–40Al 60Ti–40Al
Laser processing parameters Laser power (kW)
Scanning speed (mm/s)
5.0 5.0 5.0 4.0
4.0 4.0 4.0 3.3
Al content in the coating (at.%)
Thickness (mm)
14.7 1.2 18.1 1.7 25.2 2.1 29.7 2.4
421 672 595 622
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the 2u = 15–908 range. After the XRD tests, Sample 1 and Sample 2 were annealed in a furnace at 650 8C for 100 h in Argon protection [20], and then cooled slowly in the furnace. After the heat treatment, the phase composition of Sample 1 and Sample 2 were then determined by XRD. The aim of heat treatment was to determine the phase composition of Sample 1 and Sample 2. The dry sliding wear tests at room temperature were performed on a reciprocating ball-on-disc UMT-2MT tribometer (Center for Tribology, Inc., California, USA) in air. The counterpart was GCr15 steel balls with hardness of 700 HV and diameter of 3 mm. The wear tests were carried out under a normal load 1 and 3 N at a sliding speed of 0.1 m/s over a period of 30 min. The worn surfaces were analyzed by SEM. The wear volume losses were measured by MicroXAM 3D non-contact surface mapping profiler (ADE Corporation, Massachusetts, USA). Three replicated tests were conducted for each test condition and the averaged value was given. 3. Results and discussions 3.1. Microstructure and composition of laser cladding Ti– Al coatings Table 1 shows the Al contents of the laser cladding Ti–Al coatings by EDS analysis. It can be seen that the Al content was 14.7 1.2, 18.1 1.7, 25.2 2.1 and 29.7 2.4 at.% to Sample 1, Sample 2, Sample 3 and Sample 4, respectively. Fig. 1 shows the surface morphologies of laser cladding Ti– Al coatings. It is clear that many small particles and balls (the balls mainly composed of Ti (Al 6.1 at.%), some nearly composed of Ti) distributed on the surfaces of the laser cladding
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Ti–Al coatings. However, on the surfaces of Sample 3 and Sample 4, the balls were smaller and lesser than those on the surfaces of Sample 1 and Sample 2, which was attributed to the lower Ti content in the Ti + Al mixed powders (see Table 1). The formation of these balls were attributed to the surface tension driven caused by the capillary instability effect [21], and was likely to occur when the molten metal did not wet the underlying substrate due to an oxide layer being present on the substrate and on the surface of the melt [21]. In addition, some long ridges appeared on the surfaces of Sample 3 and Sample 4, which was quite different from those of Sample 1 and Sample 2. Fig. 2 shows the interface morphologies of the laser cladding Ti–Al coatings. As can be seen from Fig. 2, the bonding between the coating and the substrate was of high-quality metallurgical fusion bonding. For Sample 1, its microstructure near the interface was made up of randomly distributed lamellae phase (13.8 at.% Al and 86.2 at.% Ti), while for Sample 2, its microstructure near the interface was composed of lamellae phase (14.8 at.% Al and 85.2 at.% Ti) and equiaxed phase (18.2 at.% Al and 81.8 at.% Ti). For Sample 3 and Sample 4, they had similar microstructure near the interface with honeycomb-like phase (24.6 at.% Al and 75.4 at.% Ti for Sample 3, 28.8 at.% Al and 71.2 at.% Ti for Sample 4). Figs. 3–7 show the XRD patterns of pure Ti and the laser cladding Ti–Al coatings before and after heat treatment. As can be seen from Fig. 3, no pure Al peaks was found in the laser cladding Ti–Al coatings. The absence of diffraction lines of pure aluminum indicated that all Al atoms substituted Ti atoms, which resulted in the decrease of Ti lattice parameters [22]. Therefore, the diffraction peaks shifted gradually to higher 2u values with increasing Al content from pure Ti to Sample 4 (see Fig. 3). As can also be seen from Fig. 3, Sample 3 (25.2 at.%
Fig. 1. Surface morphologies of laser cladding Ti-Al coatings: (a) Sample 1, (b) Sample 2, (c) Sample 3, and (d) Sample 4.
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Fig. 2. Interface morphologies of the laser cladding Ti–Al coatings: (a) Sample 1, (b) Sample 2, (c) Sample 3, and (d) Sample 4.
Al) and Sample 4 (29.7 at.% Al) were composed of a2-Ti3Al phases. However, the phase composition of Sample 1 (14.7 at.% Al) and Sample 2 (18.1 at.% Al) could not be assigned as a-Ti or a2-Ti3Al phase, because their d values had a deviation compared with those of a-Ti or a2-Ti3Al phase (see the ‘‘d values’’ of the strongest peaks marked in Fig. 3). The phases in Sample 1 and Sample 2 were not determined directly by Ti–Al binary phase diagram presented by Murray [23] because of the rapid solidification in the laser cladding process. In order to determine the phase composition of Sample 1 and Sample 2, heat treatment was taken. For Sample 1, before heat treatment, there was only one a2-Ti3Al peak in the lower 2u values (due to the low content of a2-Ti3Al phase in Sample 1,
see Fig. 4), while two new a2-Ti3Al peaks appeared after heat treatment (see Fig. 6). In the case of Sample 2, after heat treatment, the intensity of a2-Ti3Al peaks enhanced (see Fig. 7), which all showed the precipitation of a2-Ti3Al phase during the heat treatment and were similar to the heat treatment behavior (the precipitation of a2-Ti3Al) of Ti–Al alloys with 16.0 at.% Al and 19.2 at.% Al in reference [20]. Furthermore, the XRD patterns of Sample 1 and Sample 2 had no change after heat treatment (no new peak or shoulder peak at higher 2u values appeared after the precipitation of a2-Ti3Al, see Fig. 5) compared with those before heat treatment (see Fig. 3) except the slight increase of d values (see the ‘‘d values’’ of the strongest peaks marked in Figs. 5 and 3) and the enhancing
Fig. 3. XRD patterns of pure Ti and the laser cladding Ti–Al coatings. T: a-Ti; TA: a2-Ti3Al.
Fig. 4. Magnified XRD patterns of pure Ti and the laser cladding Ti–Al coatings in Fig. 3. TA: a2-Ti3Al.
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Fig. 5. XRD patterns of Sample 1 and Sample 2 after heat treatment. T: a-Ti; TA: a2-Ti3Al.
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Fig. 7. XRD patterns of Sample 2 at lower 2u values before and after heat treatment. TA: a2-Ti3Al.
intensities of the peaks (see Fig. 7). Considering the above, conclusions can be drawn that Sample 1 and Sample 2 were composed of mixed a-Ti and a2-Ti3Al phases according to the Ti–Al binary phase diagram [20,23]. As can also be seen from Fig. 4, the intensity of a2-Ti3Al peaks of Sample 2 at lower 2u values was stronger than that of Sample 1, which indicated that the content of a2-Ti3Al phase in Sample 2 was higher than that in Sample 1. Fig. 8 shows the microhardness profile along the depth direction of the laser cladding Ti–Al coatings. It is clear that the laser cladding Ti–Al coatings had a high and wavy hardness distribution within the main coatings except in the coating/ substrate bonding zone where the hardness showed a gradient decrease to the substrate. As can also be seen from Fig. 8, the hardness of the laser cladding Ti–Al coatings increased with the increase of Al content. For the two phase (a-Ti + a2-Ti3Al) Ti– Al coatings (Sample 1 and Sample 2), because of the effective Al solid solution strengthening and the precipitation of the hard a2-Ti3Al phase, their hardness was higher than that of pure Ti.
Moreover, the more Al content, the more effective Al solid solution strengthening and more precipitation of hard a2-Ti3Al phase, therefore, the higher hardness of Sample 2 than that of Sample 1. Sample 3 and Sample 4 were both composed of the hard a2-Ti3Al phase, so their hardness was higher than that of Sample 1 and Sample 2, which was similar to the results presented in reference [24]. Table 2 shows the fracture toughness of laser cladding Ti–Al coatings measured by the direct indentation method. The studies carried out by Evans indicated that the a2-Ti3Al precipitation could lead to the lowering of ductility of titanium alloy [25], that is to say, the brittle a2-Ti3Al phase had a detrimental effect on the ductility of titanium alloy. Therefore, it is reasonable to postulate that the ductility of Sample 2 was lower than that of Sample 1 due to its higher content of brittle a2-Ti3Al phase than that in Sample 1. Moreover, according to the results presented in reference [26], Sample 4 (29.7 at.% Al) was more brittle than Sample 3 (25.2 at.% Al). In summary, the ductility between the laser cladding Ti–Al coatings was in this
Fig. 6. Magnified XRD patterns of Sample 1 and Sample 2 in Fig. 5. TA: a2Ti3Al.
Fig. 8. Microhardness profile across the laser cladding Ti–Al coatings.
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Table 2 The average length of crack and fracture toughness of the laser cladding Ti-Al coatings in the direct indentation method Materials
Average length of crack (mm)
Fracture toughness (MPam1/2)
Pure Ti Sample Sample Sample Sample
No crack 5.6 9.7 22.0 32.1
– 167.8 72.3 21.0 11.7
1 2 3 4
order: Sample 1 > Sample 2 > Sample 3 > Sample 4. It is known that the ductility had a critical effect on the fracture toughness of materials [27]. In ductile materials, the material near the tip of the flaw can deform, causing the tip of any crack to become blunt, reducing the stress intensity factor, and preventing growth of the crack [27]. So, in the direct indentation tests, the more brittle the Ti–Al coatings, the easier crack formation or growth under the same applied load, therefore, the fracture toughness of Ti–Al coatings decreased with increasing Al contents due to their decreasing ductility (see Table 2). In another study, the fracture toughness of a2Ti3Al based alloy was measured using pre-cracked three-point bend methods in accordance with ASTM-E399 [28]. Unfortunately, this method is not appropriate for these thin laser cladding Ti–Al coatings. The direct indentation method may be not appropriate for measuring the fracture toughness of the laser cladding Ti–Al coatings, but it can give the order of fracture toughness of the laser cladding Ti–Al coatings. 3.2. Tribological properties of laser cladding Ti–Al coatings 3.2.1. Variation of friction coefficient Fig. 9 shows the friction coefficient of pure Ti and the laser cladding Ti–Al coatings sliding against GCr15 ball under 1 and 3 N normal loads. It can be seen that the friction coefficient of the laser cladding Ti–Al coatings varied with increasing Al content. As a whole, the friction coefficient of
Fig. 9. Friction coefficient of pure Ti and the laser cladding Ti–Al coatings sliding against GCr15 steel ball.
laser cladding Ti–Al coatings under 1 N normal load was higher than that of pure Ti, while those under 3 N normal load was lower than that of pure Ti except Sample 4. Under the present experimental conditions, the lowest friction coefficient (0.33) could be obtained under 3 N normal load for Sample 3. From the above, it can be seen that the laser cladding Ti–Al coatings exhibited better friction behavior under higher normal load. 3.2.2. Wear mechanism of laser cladding Ti–Al coatings Fig. 10 shows the worn surfaces of pure Ti and laser cladding Ti–Al coatings sliding against GCr15 steel ball. Under the sliding wear conditions, since the hardness of pure Ti (205 HV) was much lower than that of the GCr15 steel ball (700 HV), it was easy for the hard asperities on the surface of the GCr15 steel ball to penetrate into the contact surface of pure Ti, which resulted in effective micro-cutting (see Fig. 10a and b). Moreover, because of the extremely high affinity of titanium to metallic adhesion, serious adhesion occurred between the contact surface (see Fig. 10a and b). Thus, the wear behavior of pure Ti was featured as abrasive and adhesive wear evidenced by the grooves and adhesive craters on its worn surface. In addition, severe plastic deformation was also visible on its worn surface. The worn surface of Sample 1 was similar to that of pure Ti due to its lower Al content (less Al solid strengthening effect and lower hardness than that of Sample 2). However, because of its higher hardness than pure Ti, the worn surface of Sample 1 was characterized by the presence of shallower grooves and less adhesive craters compared with that on the worn surface of pure Ti (see Fig. 10c and d). For Sample 2, only slight adhesive features were observed on its worn surface due to its higher hardness than that of Sample 1. Besides the adhesive wear, the abrasive wear features were also observed evidenced by the grooves on its worn surface (see Fig. 10e and f). For Sample 3 and Sample 4, under the action of repeating sliding cycles, cracks were firstly formed on the worn surfaces, and then peeled off at places where the crack propagated due to their drastically decreased fracture toughness compared with that of Sample 1 and Sample 2, therefore, microfracutre were observed on their worn surfaces besides adhesive wear and abrasive wear (see Fig. 10g–j). Fig. 11 shows the worn surfaces of GCr15 steel ball sliding against pure Ti, Sample 4 under 3 N normal load and the corresponding distributions of Fe and Ti elements. As can be seen from Fig. 11, during the sliding wear test, pure Ti and the laser cladding Ti–Al coating transferred to the GCr15 steel ball surface and formed the protuberances with high content of Ti element because of adhesiveness, and the adhesive behavior was more severe against pure Ti than against Sample 4. As can also be seen from Fig. 11, the worn surfaces of GCr15 steel ball sliding against pure Ti and Sample 4 also showed abrasive wear features evidenced by slight scratches. In a word, the surface morphologies of the worn steel balls corresponded well to those of the worn surfaces of pure Ti and Sample 4, which was in accordance with the wear mechanism of pure Ti and the laser cladding Ti–Al coatings.
B. Guo et al. / Applied Surface Science 253 (2007) 9301–9310
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Fig. 10. Typical worn surfaces of (a) pure Ti; (c) Sample 1; (e) Sample 2; (g) Sample 3; (i) Sample 4 at 1 N load, and (b) pure Ti; (d) Sample 1; (f) Sample 2; (h) Sample 3; (j) Sample 4 at 3 N load (arrows shows the cracks).
3.2.3. Wear volumes of laser cladding Ti–Al coatings Fig. 12 shows the wear volume of pure Ti and laser cladding Ti–Al coatings sliding against GCr15 steel ball under 1 and 3 N normal loads. It is clearly seen that the wear
volumes of all the laser cladding Ti–Al coatings were lower than that of pure Ti, that is to say, the wear resistance of Ti was improved by forming laser cladding Ti–Al coatings on its surface.
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Fig. 11. Worn surfaces of GCr15 steel ball sliding against (a) pure Ti; (b) Sample 4 under 3 N normal load and the corresponding map distributions of Fe and Ti elements, 330.
The variation of wear volume of pure Ti and the laser cladding Ti–Al coatings can be explained by Archard Eq. (2) [29]. V¼
kWS 3H
(2)
where V is the sliding wear volume; W, the normal load; S, the total sliding distance (S); H, the hardness of the wearing
surface; and k is a probability factor that a given area contact will fracture within the weaker material rather than at the original interface. According to Eq. (2), the hardness of laser cladding Ti–Al coatings was higher than that of pure Ti, so the wear volume of the laser cladding Ti–Al coatings was lower than that of pure Ti. And the wear volume at the higher normal load (3 N) was
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program of the light in China’s western region (2006)’’ for financial support. References
Fig. 12. Wear volume of pure Ti and the laser cladding Ti–Al coatings sliding against GCr15 steel ball.
higher than that at lower normal load (1 N) due to the higher load in Eq. (2). Under a lower normal load (1 N), the wear volume of laser cladding Ti–Al coatings decreased with the increase of Al content, which was mainly attributed to the increasing hardness. However, under a higher normal load (3 N), the wear rate firstly decreased due to the increasing hardness, and then increased due to the severe brittle fracture caused by the decreased fracture toughness (the increasing k value in Eq. (2)). Sample 2 had the optimized combination of high hardness and high fracture toughness, so it showed the best anti-wear properties at the higher normal load. 4. Conclusions 1. With the increase of Al content, the diffraction peaks of the laser cladding Ti–Al coatings shifted gradually to higher 2u values. The laser cladding Ti–Al coatings with 14.7 and 18.1 at.% Al were composed of mixed a-Ti and a2-Ti3Al phases, while the laser cladding Ti–Al coatings with 25.2 and 29.7 at.% Al were composed of a2-Ti3Al phases. 2. With the increase of Al content, the cross-sectional hardness increased, while the fracture toughness decreased. 3. For Sample 1 and Sample 2, the wear mechanism was adhesive wear and abrasive wear, while the wear mechanisms shifted to adhesive wear, abrasive wear and microfracture for Sample 3 and Sample 4. 4. With the increase of Al content, the wear rate of laser cladding Ti–Al coatings decreased under 1 N load, while the wear rate firstly decreased and then increased under 3 N normal load. Sample 2 had the optimized combination of high hardness and high fracture toughness, so it had the best anti-wear properties under the higher normal load. Acknowledgements The authors are grateful to the National Natured Science Foundation of China (Grant No.5057521), the Innovative Group Foundation from NSFC (Grant No.50421502), and ‘‘The
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