Effect of Laser Cladding Stellite 6-Cr3C2-WS2 Self-Lubricating Composite Coating on Wear Resistance and Microstructure of H13

In order to prevent the wear failure of the hot-working die, the composite coatings of Stellite 6-Cr3C2-WS2 was fabricated on H13 hot-working die steel by laser cladding. The composite coating was prepared through the in-situ generation technology, that can give H13 the ability of self-lubricating at the working temperature (about 200 °C). The effect of the various WS2 percentages on the properties of the coating was studied by X-ray diffraction (XRD), scanning electron microscopy (SEM), energy dispersive spectroscopy (EDS), microhardness test, friction and wear test. In addition, the phase constitutions, microstructures and wear properties were also investigated systematically. The obtained hardness of the cladding coating is approximately 2.5 times higher than the substrate because of the constituents of γ-(Fe, Co)/Cr7C3 eutectic colony, (Cr, W)C carbide and dendritic crystals in the coating. Furthermore, the friction coefficient decreases to 70% of the substrate due to the CrS self-lubricating phase. The analyses results suggest that an 85% Stellite 6-10% Cr3C2-5% WS2 composite coating has excellent material properties.


Introduction
Due to the excellent high-temperature resistance, high strength and oxidation resistance, H13 steel has been widely used in hot-working die, casting die and extrusion die [1][2][3]. Unfortunately, the intense friction in the deformation of sheet metal can result in cracks and fatigue wear on H13 steel in the hot stamping process [4][5][6]. In order to improve the surface properties, composite coating has been adopted on the die surface with poor wear resistance [7,8]. Adding alloy powder with the similar chemical composition into the base material can ensure an excellent metallurgical bonding interface. As a competitive and effective surface strengthening technique, laser cladding can produce microstructure with fine and even crystals, thereby the mechanical properties of the coating are improved [9][10][11][12][13]. Compared with the traditional technology of improving surface properties, the laser cladding exhibits many outstanding advantages, such as concentrated heat input, short processing time, fast coating cooling, low dilution ratio of the cladding layer, good metallurgical combination of substrate and cladding coatings, wide selective range of laser cladding powder, and so on [14][15][16][17]. In addition, the in-situ self-lubricating method is one of the methods used to generate the lubricating phase [18][19][20]. As compared to the directly added lubricating phase, the advantages of the in-situ self-lubricating are obvious, such as avoiding poor bonding performance and cracking between the directly added lubricating phase and the coatings. Thus, many researchers have conducted studies on laser cladding using in-situ self-lubricating technology to improve the surface properties of materials [21,22]. Currently,

Materials and Methods
In this paper, the H13 steel which is often used in high temperature conditions, was chosen as the substrate and its chemical compositions are shown in the third row of Table 1. The dimension size of the substrate was 50 mm × 20 mm × 10 mm, covered by a laser cladding surface with a size of 50 mm × 20 mm. According to the order of gravel size, the sandpaper was used to clean the substrate surface from dust and oxide film. Then alcohol was used to remove the oil on the surface of the substrate. In order to obtain the high hardness and great wear resistance of the cladding coatings, Co-based powder, such as Stellite 6 (Kennametal Stellite Co., Ltd., Shanghai, China), was adopted for the coatings (see the chemical compositions of Stellite 6 in the fourth row of Table 1). In addition, three kinds of composite powders with different mass percentages of Stellite 6, Cr 3 C 2 and WS 2 were prepared and used in the laser cladding. The purity of Cr 3 C 2 (Aidun Co., Ltd., Heibei, China) and WS 2 (Shanghai Onway Co., Ltd., Shanghai, China) are 99.9%. Then the corresponding three specimens were made from the substrates and coatings, as shown in Table 2. The proportioned powder as shown in Table 2 was ground by a speed of 300 r/min in a QM-3SP2 planetary mill (Nanjing Chi Shun Technology Development Co., Ltd., Nanjing, China) for 2 hours and then dried at a temperature of 200 • C for 1 h before laser cladding. The powders were fed to the substrate surface by side axis method and then melted by a LDM3000-60 semiconductor laser (Laserline GmbH, Mülheim-Kärlich, Germany). During the laser cladding test, these processing parameters were selected: laser power 1600 W, scanning speed 400 mm/min, powder feeding amount 12 L/min, spot diameter 4 mm, wavelength 900 nm and protective gas nitrogen. Thus, good quality coatings were achieved with those parameters. To avoid cracks on the surface during laser cladding, the substrate H13 was preheated.
After the laser cladding process, the specimens were crosscut and metallographic specimens were prepared following standard mechanical polishing procedures. The phase constituents of cladding coatings were analyzed by a D8-Advance X-ray diffractometer (Bruker AXS GmbH, Karlsruhe, Germany) at a scanning speed of 5 • /min ranging from 20 • to 100 • with Cu target Kα radiation, and accelerating voltage of 40 kV and a current of 40 mA; the microstructure and element of cladding coatings were identified by a S-3400N scanning electron microscope equipped with an X-ray spectrometer (JEOL, Tokyo, Japan); the microhardness of cladding coatings was also measured using an HV-1000 hardness tester (Suzhou Jingli Co., Ltd., Suzhou, China) with a test load of 300 g and a dwell time of 15 s. Moreover, the wear resistance of cladding coatings was measured by a MFT-5000 multi-functional friction and wear machine (Rtec GmbH, San Jose, CA, USA). The coupled ball is GCr15 steel (ø9.5 mm) and the corresponding rotating speed, loading force, wearing time, and working temperature of hot die were 25 r/min, 50 N, 60 min, and 25 • C/200 • C, respectively. The weight loss of the specimen was measured by an electronic balance whose precision was 0.1 mg (Bond West Instrument Technology Co., Ltd., Shanghai, China) and the worn morphology was observed by an optical microscope (Olympus GmbH, Tokyo, Japan).

Constituent Phase and Microstructure
The cross section of the Co-based composite coatings 1-3 are shown in Figure 1, in which the bonding line between the substrate and coating could be observed. The cladding coatings are uniform, smooth and free of defects such as pores and impurities. The good metallurgical combination is obtained between the cladding coating and the substrate, because the bonding line is stable and of low dilution ratio. The cross section of the Co-based composite coatings 1-3 are shown in Figure 1, in which the bonding line between the substrate and coating could be observed. The cladding coatings are uniform, smooth and free of defects such as pores and impurities. The good metallurgical combination is obtained between the cladding coating and the substrate, because the bonding line is stable and of low dilution ratio. The XRD diffraction patterns of coatings from specimens 1, 2 and 3 are shown in Figure 2. The phases of specimen 1 mainly consist of γ-(Fe, Co), Cr7C3, Cr3C2 and other carbides. The main chemical reactions are described in Equations (1)-(4). WS2 and Cr3C2 decompose under high temperature [30,31]. The S and C elements react with Cr to form CrS and Cr7C3, respectively. The hard phases γ-(Fe, Co) and Cr7C3 mainly contribute to improve the hardness of the cladding coatings. Besides this, no oxides are generated in the cladding coatings because of the good self-fusibility of the Co-based alloy powder. Compared to specimen 1, specimens 2 and 3 generate the extra phases of CrS, (Cr, W)C and WS2 due to the addition of WS2 in the cladding powder. The (Cr, W)C carbides and CrS self-lubricating phase is generated due to the decomposition of WS2 at 450 °C and the reaction. The lubricating property and wear resistance are improved because of the CrS and residual WS2.  The XRD diffraction patterns of coatings from specimens 1, 2 and 3 are shown in Figure 2. The phases of specimen 1 mainly consist of γ-(Fe, Co), Cr 7 C 3 , Cr 3 C 2 and other carbides. The main chemical reactions are described in Equations (1)-(4). WS 2 and Cr 3 C 2 decompose under high temperature [30,31]. The S and C elements react with Cr to form CrS and Cr 7 C 3 , respectively. The hard phases γ-(Fe, Co) and Cr 7 C 3 mainly contribute to improve the hardness of the cladding coatings. Besides this, no oxides are generated in the cladding coatings because of the good self-fusibility of the Co-based alloy powder. Compared to specimen 1, specimens 2 and 3 generate the extra phases of CrS, (Cr, W)C and WS 2 due to the addition of WS 2 in the cladding powder. The (Cr, W)C carbides and CrS self-lubricating phase is generated due to the decomposition of WS 2 at 450 • C and the reaction. The lubricating property and wear resistance are improved because of the CrS and residual WS 2 .

Constituent Phase and Microstructure
The cross section of the Co-based composite coatings 1-3 are shown in Figure 1, in which the bonding line between the substrate and coating could be observed. The cladding coatings are uniform, smooth and free of defects such as pores and impurities. The good metallurgical combination is obtained between the cladding coating and the substrate, because the bonding line is stable and of low dilution ratio. The XRD diffraction patterns of coatings from specimens 1, 2 and 3 are shown in Figure 2. The phases of specimen 1 mainly consist of γ-(Fe, Co), Cr7C3, Cr3C2 and other carbides. The main chemical reactions are described in Equations (1)-(4). WS2 and Cr3C2 decompose under high temperature [30,31]. The S and C elements react with Cr to form CrS and Cr7C3, respectively. The hard phases γ-(Fe, Co) and Cr7C3 mainly contribute to improve the hardness of the cladding coatings. Besides this, no oxides are generated in the cladding coatings because of the good self-fusibility of the Co-based alloy powder. Compared to specimen 1, specimens 2 and 3 generate the extra phases of CrS, (Cr, W)C and WS2 due to the addition of WS2 in the cladding powder. The (Cr, W)C carbides and CrS self-lubricating phase is generated due to the decomposition of WS2 at 450 °C and the reaction. The lubricating property and wear resistance are improved because of the CrS and residual WS2.   always goes forward with the liquid-solid interface in the whole cooling crystallization process. The initial crystallization takes place at the solid-phase substrate where the temperature gradient (temperature/solidification rate) of the coating near the substrate is relatively large [32]. The dendrite epitaxial growth has a strong directionality and the plane crystal grows vertically to the substrate. 3 7 Cr 3C r 7 C C   (4) Figure 3 shows the microstructures in the cladding coating of specimens 1-3. The grains at the bottom region in Figure 3a,d,g are mainly plane crystals which are fine and uniform, and the grains closest to the substrate are denser and smaller. According to the solidification theory, the solidification always goes forward with the liquid-solid interface in the whole cooling crystallization process. The initial crystallization takes place at the solid-phase substrate where the temperature gradient (temperature/solidification rate) of the coating near the substrate is relatively large [32]. The dendrite epitaxial growth has a strong directionality and the plane crystal grows vertically to the substrate. As shown in Figure 3b,e,h the plane crystals turn into dendrite crystals in the middle region because of the gradual increase of the grain growth rate as a result of the increase in temperature and decrease in solidification speed. The crystals initially grow vertically to the As shown in Figure 3b,e,h the plane crystals turn into dendrite crystals in the middle region because of the gradual increase of the grain growth rate as a result of the increase in temperature and decrease in solidification speed. The crystals initially grow vertically to the interface between coating and substrate and then grow into dendrites in the middle of the coating due to the non-directional heat dissipation. The top of the cladding coating is shown in Figure 3c,f,i. It is full of cellular grains because the heat dissipation at the top region of the coating under high temperature is slow and then nucleation has sufficient time to grow cellular crystals.
For comparison, the microstructures in the middle region of specimens 1-3 are shown in Figure 4. During the laser cladding, the crystals precipitated in the welding pool and then grew into dendrites. The dendrite and intergranular network eutectic structure are generated due to the eutectic reaction in the transformation of the intergranular metal solution to the solid-state. The microstructures at the middle regions of cladding coatings of specimens 2 and 3 are more compact and uniform than that of specimen 1 because of the high thermal conductivity of the cladding coatings with WS 2 . The grains could refine when the growth rate of grains is less than the nucleating rate. Moreover, the high absorption capacity of the laser beam on the W element results in better energy conversion and then improves the combination between coating and substrate. beam on the W element results in better energy conversion and then improves the combination between coating and substrate.
In addition, white point in zone A and black point in zone B are found in Figure 4b. The element compositions of the white point in zone A and black point in zone B are analyzed by EDS analysis and indicated in Table 3. The white point is hard phase carbide, which consists of Cr7C3 and Cr3C2, and the black point is the lubricating phase that is CrS, which are consistent with the results from XRD analysis.

Microhardness
The microhardness variations along different coating depths in the three specimens are plotted in Figure 5. The microhardness in the three specimens are approximately 2.5 times that  Table 3. The white point is hard phase carbide, which consists of Cr 7 C 3 and Cr 3 C 2 , and the black point is the lubricating phase that is CrS, which are consistent with the results from XRD analysis.

Microhardness
The microhardness variations along different coating depths in the three specimens are plotted in Figure 5. The microhardness in the three specimens are approximately 2.5 times that of the substrate and exhibit three distinguishable zone: zones I, II, and III. The microhardness in zone I decreases slightly with increasing depths and seems stable due to the evenly distributed materials in the laser cladding region. While in the narrow bonding region II, the microhardness increases with a big jump due to the following two reasons: at first, the grains in the bonding zone are refined and dendritic because of rapid temperature fluctuation in the laser cladding; secondly, the γ-(Fe, Co) is generated by the diffusion of Fe element from the substrate to the bonding zone and the amount of hard carbide (Cr, W)C in the region. In zone III, the microhardness decreases rapidly because of more plane crystals, less dendrites and much less carbide in the deeper region. The average microhardness of specimens 1, 2 and 3 are 556 ± 10 HV, 527 ± 12 HV and 516 ± 11 HV, respectively, which are all around 2.5 times that of the substrate (215 HV ± 15 HV). This is caused by the solution and dispersion strengthening resulting from the dissolution of Cr and C atoms in the Co-phase. The microhardness of specimen 1 is higher than that of specimens 2 and 3 due to the amount of Cr 7 C 3 in the coatings. Ⅲ , the microhardness decreases rapidly because of more plane crystals, less dendrites and much less carbide in the deeper region. The average microhardness of specimens 1, 2 and 3 are 556 ± 10 HV, 527 ± 12 HV and 516 ± 11 HV, respectively, which are all around 2.5 times that of the substrate (215 HV ± 15 HV). This is caused by the solution and dispersion strengthening resulting from the dissolution of Cr and C atoms in the Co-phase. The microhardness of specimen 1 is higher than that of specimens 2 and 3 due to the amount of Cr7C3 in the coatings.

Friction Coefficient and Weight Loss
The evolutions of friction coefficients obtained from the three specimens and the substrate at a room temperature of 25 °C are plotted in Figure 6a. Overall, the friction coefficients of the specimens reduce by 20-25%. This is because of the presence of Cr7C3 in specimen 1 and the continuous and uniform lubricating film generated by WS2 powder in specimens 2 and 3. The friction coefficients in all of the four cases initially increase rapidly and then tend to be stable after 200 s, and the coefficient of the substrate fluctuated between 0.35 and 0.4 due to the insufficient hardness and coarse surface. We also notice that the friction coefficient of specimen 2 is smaller and more stable than that of specimen 1 and specimen 3 because of more self-lubricating phases and hard phases generated.

Friction Coefficient and Weight Loss
The evolutions of friction coefficients obtained from the three specimens and the substrate at a room temperature of 25 • C are plotted in Figure 6a. Overall, the friction coefficients of the specimens reduce by 20-25%. This is because of the presence of Cr 7 C 3 in specimen 1 and the continuous and uniform lubricating film generated by WS 2 powder in specimens 2 and 3. The friction coefficients in all of the four cases initially increase rapidly and then tend to be stable after 200 s, and the coefficient of the substrate fluctuated between 0.35 and 0.4 due to the insufficient hardness and coarse surface. We also notice that the friction coefficient of specimen 2 is smaller and more stable than that of specimen 1 and specimen 3 because of more self-lubricating phases and hard phases generated. The evolutions of corresponding friction coefficients in all the four cases at 200 °C are also plotted, as shown in Figure 6b. The friction coefficients are smaller and more stable than those at 25 °C in Figure 6a due to the fact that more CrS separated at high temperature can act as the lubricant. In addition, the friction coefficient of specimen 2 is approximately 0.26-0.28, just 70% of that of the substrate, and its evolution curve is more stable than that of specimen 3 because of the excessive S element from the decomposition of WS2 under laser irradiation.
The wear weight loss of the three specimens and the substrate at 25 °C and 200 °C are both demonstrated in Figure 7. The loss of the three specimens are significantly less than that of the substrate due to the solid solution strengthening effect of γ-(Fe, Co) and the improved wear resistance of hard phase Cr7C3. In addition, the wear weight loss of specimens 2 and 3 are smaller than that of specimen 1, suggesting that WS2 and CrS play the critical role of The evolutions of corresponding friction coefficients in all the four cases at 200 • C are also plotted, as shown in Figure 6b. The friction coefficients are smaller and more stable than those at 25 • C in Figure 6a due to the fact that more CrS separated at high temperature can act as the lubricant. In addition, the friction coefficient of specimen 2 is approximately 0.26-0.28, just 70% of that of the substrate, and its evolution curve is more stable than that of specimen 3 because of the excessive S element from the decomposition of WS 2 under laser irradiation. The wear weight loss of the three specimens and the substrate at 25 • C and 200 • C are both demonstrated in Figure 7. The loss of the three specimens are significantly less than that of the substrate due to the solid solution strengthening effect of γ-(Fe, Co) and the improved wear resistance of hard phase Cr 7 C 3 . In addition, the wear weight loss of specimens 2 and 3 are smaller than that of specimen 1, suggesting that WS 2 and CrS play the critical role of producing lubrication film and reducing wear, especially at 200 • C.
The evolutions of corresponding friction coefficients in all the four cases at 200 °C are also plotted, as shown in Figure 6b. The friction coefficients are smaller and more stable than those at 25 °C in Figure 6a due to the fact that more CrS separated at high temperature can act as the lubricant. In addition, the friction coefficient of specimen 2 is approximately 0.26-0.28, just 70% of that of the substrate, and its evolution curve is more stable than that of specimen 3 because of the excessive S element from the decomposition of WS2 under laser irradiation.
The wear weight loss of the three specimens and the substrate at 25 °C and 200 °C are both demonstrated in Figure 7. The loss of the three specimens are significantly less than that of the substrate due to the solid solution strengthening effect of γ-(Fe, Co) and the improved wear resistance of hard phase Cr7C3. In addition, the wear weight loss of specimens 2 and 3 are smaller than that of specimen 1, suggesting that WS2 and CrS play the critical role of producing lubrication film and reducing wear, especially at 200 °C .

Wear Mechanism
Wear mechanism plays a significant role of changing the friction coefficient and wear weight loss, and was studied deeply in the three specimens. Figure 8 shows the wear surface

Wear Mechanism
Wear mechanism plays a significant role of changing the friction coefficient and wear weight loss, and was studied deeply in the three specimens. Figure 8 shows the wear surface morphology of the three specimens and the substrate at 25 • C. In Figure 8a, large numbers of block shedding and deep furrow caused by low hardness of the substrate and absence of the hard phase led to serious abrasive wear and adhesive wear on the substrate surface. In contrast, the coating surfaces in the three specimens in Figure 8b-d observed with shallow scratching and small plastic deformation are almost smooth and flat as a result of the mild abrasive and adhesive wear. In the wear process, the soft phase structure was removed at the initial stage. Since the contact stress between the grinding ball and the hard phase is beyond the yield limit of the hard phase, the hard phase particles drop off and act as abrasive particles in the process. The remaining hard phase effectively resists wear, deformation and scratches. In specimen 3, the white areas pointed to by arrows, shown in Figure 8d, mainly include elements of C, Cr and W observed by EDS analysis, and consist of (Cr, W)C carbides according to the XRD analysis. Therefore, the material surface, effectively protected by the (Cr, W)C carbides is hard to be worn.
The wear morphologies of specimens 1-3 and the substrate at 200 • C are also demonstrated in Figure 9a-d, respectively. The wear surface of the substrate and specimen 1 was seriously damaged without the protection of lubrication film at 200 • C. The surfaces of specimens 2 and 3 are almost smooth and with few slight scratches. Besides this, many black regions on the wear surface include large numbers of self-lubricating phases of WS 2 /CrS obtained by energy spectrum analysis and XRD analysis. Hence, the self-lubricating phase CrS that separated at a high temperature acts as the lubricating film at the contact surfaces to reduce the friction coefficient. It results in mild abrasive and adhesive wear on the contact surfaces. In addition, there is less plastic deformation and furrow on the surface than those in the case of 25 • C. initial stage. Since the contact stress between the grinding ball and the hard phase is beyond the yield limit of the hard phase, the hard phase particles drop off and act as abrasive particles in the process. The remaining hard phase effectively resists wear, deformation and scratches. In specimen 3, the white areas pointed to by arrows, shown in Figure 8d, mainly include elements of C, Cr and W observed by EDS analysis, and consist of (Cr, W)C carbides according to the XRD analysis. Therefore, the material surface, effectively protected by the (Cr, W)C carbides is hard to be worn. The wear morphologies of specimens 1-3 and the substrate at 200 °C are also demonstrated in Figure 9a-d, respectively. The wear surface of the substrate and specimen 1 was seriously damaged without the protection of lubrication film at 200 °C . The surfaces of specimens 2 and 3 are almost smooth and with few slight scratches. Besides this, many black regions on the wear surface include large numbers of self-lubricating phases of WS2/CrS obtained by energy spectrum analysis and XRD analysis. Hence, the self-lubricating phase CrS that separated at a high temperature acts as the lubricating film at the contact surfaces to reduce the friction coefficient. It results in mild abrasive and adhesive wear on the contact surfaces. In addition, there is less plastic deformation and furrow on the surface than those in the case of 25 °C .

Conclusion
To improve the hardness and wear resistance of H13, different proportions of composite powder were covered on the surface of H13 steel by laser cladding. The microstructure, hardness, and wear resistance of different laser cladding coatings on H13 were investigated.