International Journal of Minerals, Metallurgy and Materials Volume 23, Number 2, February 2016, Page 184 DOI: 10.1007/s12613-016-1226-z

Microstructure characteristics of Ni/WC composite cladding coatings Gui-rong Yang1), Chao-peng Huang1), Wen-ming Song1,2), Jian Li3), Jin-jun Lu4), Ying Ma1), and Yuan Hao1) 1) State Key Laboratory of Advanced Processing and Recycling of Nonferrous Metals, Lanzhou University of Technology, Lanzhou 730050, China 2) Lanzhou Petroleum Machinery Institute, Lanzhou 730000, China 3) Wuhan Research Institute of Materials Protection, Wuhan 430030, China 4) College of Chemistry and Materials Science, Northwest University, Xi’an 710069, China (Received: 11 June 2015; revised: 27 July 2015; accepted: 3 August 2015)

Abstract: A multilayer tungsten carbide particle (WCp)-reinforced Ni-based alloy coating was fabricated on a steel substrate using vacuum cladding technology. The morphology, microstructure, and formation mechanism of the coating were studied and discussed in different zones. The microstructure morphology and phase composition were investigated by scanning electron microscopy, optical microscopy, X-ray diffraction, and energy-dispersive X-ray spectroscopy. In the results, the coating presents a dense and homogeneous microstructure with few pores and is free from cracks. The whole coating shows a multilayer structure, including composite, transition, fusion, and diffusion-affected layers. Metallurgical bonding was achieved between the coating and substrate because of the formation of the fusion and diffusion-affected layers. The Ni-based alloy is mainly composed of -Ni solid solution with finely dispersed Cr7C3/Cr23C6, CrB, and Ni+Ni3Si. WC particles in the composite layer distribute evenly in areas among initial Ni-based alloying particles, forming a special three-dimensional reticular microstructure. The macrohardness of the coating is HRC 55, which is remarkably improved compared to that of the substrate. The microhardness increases gradually from the substrate to the composite zone, whereas the microhardness remains almost unchanged in the transition and composite zones. Keywords: cladding; composite coatings; microstructure characteristics; formation mechanisms; hardness

1. Introduction Ni-based alloy coatings are widely used in engineering fields because of their excellent resistance to wear and corrosion [1–2]. Ni-based alloy powders exhibit a self-fluxing characteristic by adding B and Si, resulting in the formation of low-melting Ni–B eutectic phase. This self-fluxing characteristic enables Ni-based alloy powders to melt onto and fuse well with steel substrates. The addition of Si further improves the self-fluxing property of Ni-based alloy powders [3–4]. In addition, elements Si and B can serve as deoxidizers during the coating formation. Cr and Ni can form a Ni-based solid solution, and the Cr remainder can form carbides and borides that improve the hardness and tribological properties of the coating through solution strengthening and hard-phase strengthening. To provide high abraCorresponding author: Gui-rong Yang

high abrasion or wear resistance, tungsten carbide (WC) particles are often used as reinforcement in Ni-based alloys because of the excellent properties of WC, which include relatively high hardness, high heat stability, and good wettability to molten metals; these combined characteristics result in the composite materials with excellent mechanical properties [5–6]. Previous studies have investigated the microstructures of Ni-based alloy coatings fabricated using different surface treatment technologies. The coatings deposited by plasma transferred arc welding are composed of -Ni primary dendritic phase with Ni+Ni3B or Ni+Ni3Si eutectic phase and Cr-based compounds (CrB, Cr3C2, and Cr7C3) situated in the interdendritic regions [7]. In the case of the coatings fabricated using laser cladding, the coatings are mainly composed of -Ni dendrites, M23C6, and CrB [8]. In the case of the coatings fabricated using high-velocity oxyfuel (HVOF)

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© University of Science and Technology Beijing and Springer-Verlag Berlin Heidelberg 2016

G.R. Yang et al., Microstructure characteristics of Ni/WC composite cladding coatings

spraying, the main phase is -Ni solid solution with small fractions of Cr7C3 and Ni3B [9]. Ni-based alloys fused under an argon atmosphere are demonstrated to be constituted by the equiaxial grains of Ni-based solid solution surrounded with a B-rich intergranular net, in addition to CrB and finely dispersed chromium carbides (like Cr2C, Cr23C6, and Cr7C3, as inferred by X-ray diffraction patterns) [3]. Ni-based coatings deposed by vacuum melting are reported to contain the Ni solid–solution matrix, Ni3B, chromium carbides (Cr7C3 and Cr23C6), and chromium borides (mainly CrB) [4,10]. Thus, the microstructure characteristics and composition of the coatings differ among those fabricated by different technologies. In addition to affecting the microstructure of the coatings, the fabrication method also influences the morphology. Problems such as cracking, agglomeration of WC or hard phases, low interfacial binding strength, and the formation of tiny pores can occur in the coatings by laser cladding, thermal spraying, HVOF spraying, and other surface processing technologies [11–12]. However, these problems do not tend to occur in the coatings deposited using vacuum melting. The properties are profoundly affected by the size and distribution of ceramic particles in the coatings and the bonding strength between the coating and substrate [13–14]. The composite coatings have been previously deposited directly onto the substrate surface. However, with this approach, inevitably, the addition of WC with high mass frac-

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tion to obtain good abrasion resistance would negatively affect the binding strength between the substrate and coating [15]. In the present work, to avoid such problems, the coatings with a transition layer between the composite coating (Ni/WC) and the substrate were fabricated by vacuum cladding technology. A coating with a special structure was obtained through temperature and heating time control in a vacuum furnace. The microstructure morphology and phase composition were investigated, and the formation mechanism was also discussed. Furthermore, the macrohardness and microhardness in the different areas of the coating were evaluated.

2. Experimental In this study, a mixture of commercially available Ni-based powder with a nominal composition of 16.00wt% Cr, 3.50wt% B, 4.20wt% Si, 5.00wt% Fe, and 0.90wt% C and WC powder with a small particle size were used as raw materials for fabricating a composite coating with a transition layer between the coating and substrate. The particle size of the Ni-based alloy powder ranged from 48 to 106 μm, and that of WC was 18 μm or less. The powder appearance is shown in Fig. 1. The Ni-based alloy powder is characteristic of near-perfect spherical particles. In contrast, the WC powder is characteristic of irregular polyhedral particles, as shown in Fig. 1(b).

Fig. 1. Appearance of the Ni-based alloy particles (a) and WC particles (b).

The substrate was ordinary Q235 (Fe360) low-carbon steel, and it was machined to dimensions of 50 mm × 50 mm × 10 mm. The substrate surface was ground by 200 to 1000 grit sandpaper and cleaned by acetone and distilled water before the preformed layer was prepared. Two kinds of powders were prepared for fabricating the coating. A mixed powder with 40wt% WC and 60wt% Ni-based alloy powder was first prepared; Ni-based alloy powder was also used to form a transition layer. The preformed layer was

prepared by overlaying the mixture of particles and homemade binder (sodium silicate) on the substrate surface. The preformed layer included two layers: Ni-based alloy layer (transition layer) and mixed powder layer (composite layer, Ni-based alloy powder combined with WC particles). The preformed transition layer was first deposited onto the substrate surface, and then the composite layer was deposited onto the transition layer. The thickness of the preformed layer was approximately 2 mm, including a 1 mm transition

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layer and a 1 mm composite layer. The preformed layer on the substrate was dried at low temperature. After drying, the substrates were placed into a vacuum carbon tube furnace under a vacuum of 1 Pa or better. They were heated to 1050°C over a period of 20 min, maintained at this temperature for 7 min, and then cooled inside the furnace to 150°C. All specimens were cut transversely from the coating surface to the substrate, mechanically ground, polished with diamond paste, and then etched by aqua regia (75 mL hydrochloric + 25 mL nitric acid), FeCl3 solution (75 mL hydrochloric + 25 mL nitric acid + 1 g FeCl3), and nitric acid solution (4 mL nitric acid + 96 ml ethanol) to obtain an optimum microstructure appearance. The microstructure morphology was observed by optical microscopy (OM) and scanning electron microscopy (SEM). Energy-dispersive X-ray spectroscopy (EDS) attached to SEM was employed to qualitatively evaluate the elemental distribution by point and line scannings. Phase identification was performed by X-ray diffraction (XRD).

3. Results and discussion 3.1. Microstructure of the whole coating The cross-section morphology of the whole coating is shown in Fig. 2(a). The whole composite coating includes four regions with different microstructure characteristics;

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these regions are the transition zone (TZ) (marked as A zone), composite zone (CZ) (marked as B zone), fusion zone (FZ) (marked as C zone), and diffusion-affected zone (DZ) (marked as D zone). The region marked as a red-dotted line is the interface between the substrate and vacuum cladding coating. The left and right areas are the substrate and the vacuum cladding coating, respectively. The whole cross-sectional morphology of the coating indicates that the coating is free from inclusions and microcracks. Fig. 2(b) presents the backscattered-electron image (BSEI) of FZ and DZ, which provides a clearer comparison to Fig. 2(a). The FZ is a white, bright belt whose width is as large as approximately 50 μm, as shown in Fig. 2(b); the main elements in this zone are Fe and Ni, which originate from the substrate and the coating, respectively. More dark phases are evident in DZ than in the area more distant from the interface. Dark phases in this area are pearlite, and the thickness of DZ is approximately 280 μm, as confirmed in Fig. 2(b). Fig. 2(c) shows a magnified image of the TZ morphology. The microstructure is uniform, and the extensive blocky reinforcement is dispersed throughout the Ni matrix. The thickness of TZ is approximately 2 mm, which could provide excellent metallurgical fusion between the composite coating and the substrate. The composite zone is strengthened by WC particles or hard phases, as shown in Fig. 2(d). It also shows that the distribution of WC particles

Fig. 2. Microstructure of the whole coating: (a) cross-section of the whole coating; (b) fusion zone (FZ) and diffusion-affected zone (DZ); (c) transition zone (TZ); (d) composite zone (CZ).

G.R. Yang et al., Microstructure characteristics of Ni/WC composite cladding coatings

among the composite layer forms a reticulate structure and that not every WC particle is directly connected to another particle. The thickness of CZ is approximately 1 mm. Fig. 3 shows the XRD patterns of the Ni/WC vacuum cladding composite coating. The analysis results indicate that the main phases in the composite coating are WC, Cr7C3, Cr23C6, Ni3Si, Ni3B, CrB, and a Ni-based solid solution. The hard phase WC originates from raw composite powder materials. Three types of phases form in the Ni-based alloy coating fabricated through any coating fabricating technology as long as sufficient heat is supplied to melt the Ni-based powder, as discussed in the introduction [3–4,7–10]. The three types of phases are -Ni solid solution, eutectic phase, and Cr–B or Cr–C compound. These three types of phases are also observed in the Ni/WC composite coating fabricated through vacuum cladding technology.

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area between the Ni-based particles in the preformed layer or in the connected area between the Ni-based particles, as marked in the image. A more detailed SEM micrograph of the lamellar eutectic structure is shown in Fig. 4(c). The eutectic structure can be divided into different areas, and the orientation of the laminar structure is the same within a given area. The orientation of the laminar structure differs among the different areas. As evident in Fig. 4(c), the dispersed hard-phase particles are also present in the eutectic region. The composition of different phases marked with arrows in Fig. 4 was analyzed through EDS and XRD; the results are presented in Table 1 and Fig. 5.

3.2. Microstructure of the transition layer Fig. 4(a) shows the BSEI of TZ, the dark-gray phase marked with arrow 3 and the black phase marked with arrow 2 are distributed evenly among the white-gray matrix marked with arrow 1. The shape of the dark-gray phase marked with arrow 3 is hexagonal, rectangular, or other shapes, as shown in Fig. 4(a). Fig. 4(b) shows the microstructure morphology of TZ, as observed by OM. A lamellar eutectic structure is distributed throughout this zone. The lamellar eutectic structure is located mainly in the capillary

Fig. 3. XRD pattern of the whole coating.

In the area labeled by arrow 1 in Fig. 4(a), the main elements are Ni, Cr, Fe, and a small amount of Si; the results

Fig. 4. Microstructures of the transition layer: (a) BSEI of TZ observed by OM; (b) deeply etched appearance of TZ observed by OM; (c) magnified view of eutectic area observed by SEM.

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Table 1. Main element composition of different phases marked with arrows in Fig. 4 at% Phases Arrow 1

Cr

Ni

Fe

3.90 67.20 19.70

Si 9.10

Arrow 2

94.00

1.40

4.30

―

Arrow 3

68.40

6.60 24.90

―

Arrow 4

4.10 59.10

4.70 19.20

Fig. 5. XRD pattern of the transition layer.

indicate that the main phase in this area is likely -Ni solid solution. The main elements of phases labeled as “3” are Cr, Ni, and Fe. As is known, the complete crystal of Cr7C3 is a hexagonal prism whose cross-sectional morphology can be hexagonal, rectangular, or another shape. These cross-section shapes agree well with those observed for many dark-gray particles labeled as “3” in Fig. 4(a). Cr atoms in Cr7C3 are easily replaced by Fe atoms or Ni atoms during the formation of Cr7C3 crystal because Cr7C3 is an interstitial compound. Therefore, a portion of Cr atoms are replaced by Fe or Ni atoms during the formation of Cr7C3, resulting in an interstitial solid solution. This explanation is consistent with elemental analysis results shown in Table 1. In the BSEI of a sample, a higher average atomic number of a zone results in a stronger backscattered-electron signal, which, in turn, results in a brighter appearance of the zone; conversely, a smaller average atomic number will result in a darker appearance [16]. The average atomic number of Cr23C6 is very similar to that of M7C3. Therefore, Cr23C6 phases could not be accurately distinguished from M7C3 phases. It is ascertained that Cr23C6 phases and M7C3 phases appear as dark-gray phases. The main element of phases labeled as “2” in Fig. 4(a) is Cr, as shown in Table 1. The darker appearance of phases labeled as “2” compared to phases labeled as “3” in Fig. 4(a) indicates that the average atomic number in phases labeled as “2” is lower than that in phases labeled as “3”. The peak intensity of CrB is the

smallest among all hard phases detected by XRD in the transition layer, as shown in Fig. 5. Analysis of the elemental content, backscattered electron signal and XRD results indicates that black phases labeled as “2” may be CrB. The XRD analysis results show that the TZ is mainly composed of Ni-based solid solution, Ni3Si, Ni3B, CrB, Cr7C3, and Cr23C6. The dark-gray phases are Cr7C3 and Cr23C6. The Ni3Si and Ni3B easily form eutectic phases with -Ni, and the microstructure of the eutectic phase is very fine [17]. The main elements in the area labeled by arrow 4 in Fig. 4(c) are Ni and small amounts of Si or other elements, as shown in Table 1. The lamellar eutectic phases are composed of -Ni, Ni3Si, or Ni3B. With increasing temperature, the contact interfaces between the Ni-based alloy particles melt first during the vacuum cladding process. The sintering necks then form gradually over time in the areas between Ni-based particles, and grow with increasing temperature and heat retention. The porosity of the preformed layer decreases gradually as the sintering necks grow. The solid solubility of Si and B in the -Ni phase is low [7], especially under the condition of non-equilibrium crystallization. Si and B atoms tend to precipitate in regions such as crystal boundaries or interdendritic spaces, and form a eutectic structure with -Ni in the sintering-necks area. Most of the Si and B atoms diffuse to the capillary area between the Ni-based particles during densification and crystallization, which results in a regional and discontinuous eutectic phase distribution. Cr exhibits a strong tendency to form carbides and borides and to form a solid solution with Ni. The uniform distribution of Cr in the Ni-based alloy leads to the uniform distribution of Cr–B and Cr–C, as shown in Fig. 4. 3.3. Microstructure of the composite layer As shown in Fig. 6(a), no demarcation line is evident between TZ and CZ. The difference between TZ and CZ is judged on the basis of the position of WC particles. Fig. 6(b) shows a magnified image of CZ. The matrix of the Ni/WC composite cladding coating is composed of Ni-based alloy. Numerous fine ceramic particles are distributed at the interface area between the spherical components. The position of the spherical components is the initial position of Ni-based alloy particles in the preformed layer. The radius of spherical area is approximately 20–100 μm, which is smaller than that of the raw particles, compared with Fig. 2(d) and Fig. 6(b). A more detailed microstructure morphology of the Ni/WC composite cladding coating is illustrated in Fig. 6(c), in which the eutectic structure and hard phases are observed to be distributed in the spherical area. The hard-phase particles

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Fig. 6. Microstructures of the composite layer: (a) interface microstructure of TZ and CZ; (b) morphology of CZ; (c) partial morphology of the spherical area in CZ.

and WC particles do not contact each other. The size of most of the WC particles is smaller than 10 μm, which indicates that the wettability between the Ni-based alloy melt and WC particles is excellent and that the fusion between the Ni-based alloy and WC occurs to some extent. The main phases of CZ are WC, Cr7C3, Cr23C6, Ni3Si, Ni3B, CrB, and Ni-based solid solution, as shown in Fig. 3. The eutectic structure is close to the edge of the spherical area. When the heating temperature exceeded the melting point of the Ni-based alloy powder during the vacuum cladding process, molten metal located near the surface of Ni-based alloy particles filled the interstice between particles; the wetting process between molten Ni-based alloy liquid and WC particles was accomplished simultaneously. WC particles in the molten metal liquid were approximately in suspension and moved slightly under the action of metal liquid surface tension. The WC particles adjusted their positions at the interface area by a certain width that was larger than the width of the interstice among the preformed layer under the control of capillary pressure that formed in the micropores between WC particles. The melting point of the Ni-based alloy powder was approximately 1000°C, and it increased slightly because of the addition of WC particles. The viscosity of the molten liquid metal decreased with increasing temperature. This viscosity was relatively high at experimental temperature which was only 50°C higher than the melting point of the Ni-based alloy powder. Therefore, no molten metal liquid was flowing in the range of the whole

composite layer, as inferred from Fig. 2(d). The viscous flow of the molten metal liquid occurred only at the area between adjacent spherical particles during the densification process, which allowed Ni-based alloy particles to retain their spherical morphology after solidification. The three-dimensional reticulation framework of WC particles played an important role during the formation of the composite coating. Due to the high surface energy and high lattice distortion energy of powdery materials [18], the decrease in total surface free energy of Ni-based particles could provide driving force during the vacuum cladding process. The Ni-based alloy provided sufficient metal liquid to completely surround WC particles with molten Ni-based alloy liquid during the vacuum cladding process. Diffusion occurred between the edge of WC particles and the Ni-based alloy under vacuum when the temperature of the Ni-based alloy exceeded 1000°C. Excellent wettability between the Ni-based molten liquid and WC under experimental conditions ensured that no inclusions were present at the interface between WC and the Ni-based alloy. No aggregation was observed within the composite coating layer because of the high viscosity of the Ni-based alloy liquid under experimental conditions. All of these factors enabled the composite coating to exhibit good tribological properties. The Ni-based alloy crystallized primarily during the cooling stage. This crystallization process of the Ni-based alloy is fundamentally similar to the crystallization process of the transition layer.

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3.4. Microstructure of the fusion zone and diffusion-affected zone The bonding strength between the coating and substrate is an important property that affects the service life of the composite coating. Fig. 7 is the microscopic analysis of the interface between the substrate and coating. A white belt is marked with an arrow adjacent to the substrate in Fig. 7(a). The white belt exhibits the characteristics of metallurgical fusion; therefore, it is named as the fusion zone (FZ). Short, rod-like grains are formed near the substrate in FZ; these grains are originated from the coating side and oriented toward the substrate, as confirmed in Fig. 7(b). The elemental distribution in FZ is shown in Fig. 7(c). The Fe content increases gradually from the coating to the substrate, as indicated by the blue curve in Fig. 7(c), and the Ni content decreases gradually from the coating to the substrate, as indicated by the cyan curve in Fig. 7(c). The fusion zone includes area E and area F, as marked in Fig. 7(b). The main phase of area E is the Ni-based solid solution, and that of

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area F is the Fe-based solid solution, as confirmed by the distribution of elements. Area F exhibits a microstructure consisting of parallel, short, rod-like grains. According to previously reported results, a small amount of B and C atoms diffuse from the coating to the substrate and Fe atoms diffuse from the substrate to the coating during the coating fabricating process [4]. The diffusion of B and C atoms from the coating to the substrate would lower the melting point of the substrate to a certain extent, and the diffusion of Fe from the substrate to the coating would accordingly increase the melting point of the coating materials [4]. The main phase is the solid solution, and no hard phases are observed in FZ, as indicated by EDS analysis, because neither Cr nor other elements (except Ni and Fe) were present in this zone, as shown in Fig. 7(c). Cr forms mainly carbides and borides, and these hard phases cannot diffuse under experimental conditions. The typical microstructure of DZ is shown in Fig. 7(d), which is composed entirely of the pearlite that is not the original ferrite and the pearlite in the substrate.

Fig. 7. Microscopic analysis results of FZ and DZ: (a) BSEI image; (b) OM image; (c) elemental distribution; (d) SEM microstructure.

A schematic showing the formation of FZ is presented in Fig. 8. The melting point of the white belt area increased because of B and C diffusion from the coating to the substrate and Fe diffusion from the substrate to the coating. The Ni-based solid solution formed during crystallization at the white belt area. The negative temperature gradient formed in

front of the solid/liquid interface, which resulted in a certain undercooling. This condition satisfied requirements for short rod-like grain growth. The melting point of materials in the area that was near the substrate increased because of the diffusion of elements, which led to a decrease of undercooling degree. The new nuclei could not form under this condi-

G.R. Yang et al., Microstructure characteristics of Ni/WC composite cladding coatings

tion, and the crystallization process was the continuous growth of grains along the direction of undercooling degree. The main element in the liquid near the substrate was Fe. Therefore, the columnar crystals were composed primarily of the Fe-based solid solution, as shown in Figs. 7(b) and (d). The diffusion of B from the coating to the substrate resulted in a leftward shift of the eutectic point in the Fe–C equilibrium diagram [19]. In addition, as previously mentioned, C diffused from the coating to the substrate. The diffusion of C increased the C content at DZ close to the eutectoid steel. Therefore, a 280-μm wide zone composed entirely of pearlite formed. The pearlite content decreased with increasing distance from FZ until a normal carbon steel structure was formed, as shown in Fig. 2(b). This area was referred to as the diffusion-affected zone, DZ.

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improvement of microhardness in TZ. In CZ, the maximum microhardness is Hv 950 and the average microhardness is Hv 850. The comprehensive reinforcement by WC and hard phases further improves the microhardness in CZ.

Fig. 9. Microhardness of the whole composite coating from the substrate to the coating surface.

4. Conclusions

Fig. 8. Schematic showing the formation of fusion zone (FZ).

3.5. Hardness The average macrohardness of the surface composite coating was as high as HRC 55, and the macrohardness of the substrate was approximately HRC 18. The macrohardness of the composite coating was substantially greater than that of the substrate because of the strengthening action of hard phases, such as WC, Cr7C3, Cr23C6, Ni3Si, Ni3B, and CrB. Fig. 9 shows the microhardness of the whole coating from the substrate to the coating surface. The microhardness is approximately Hv 150 at the substrate area, as shown in Fig. 9, and increases gradually to Hv 720. These areas are the DZ and FZ areas, where the main phases are solid solutions, as previously analyzed. The microhardness curve became smooth, which indicates that the microhardness remains almost unchanged in TZ. The average microhardness is approximately Hv 750 at this area (TZ). Some metallic compounds, such as Cr7C3, Cr23C6, Ni3Si, Ni3B, and CrB, are present in TZ. The strengthening action of hard phases in the Ni-based solid solution is primarily responsible for the

(1) A Ni/WC composite was fabricated by vacuum cladding technology on a low-carbon steel substrate. The microstructure of the coating is uniform, and a metallurgical fusion layer forms between the substrate and coating. The whole coating consists of four zones with different microstructural characteristics: composite layer, transition layer, fusion layer, and diffusion-affected layer. (2) The distribution of WC is reticular in 1-mm thick CZ. The spherical area is separated by discontinuous WC. The spherical area is mainly composed of Ni-based solid solution, a eutectic phase, and hard phases, such as Cr7C3, Cr23C6, and CrB. In TZ, the main phases are hard phases (Cr7C3, Cr23C6, CrB) and the eutectic phase, which are distributed evenly in Ni-based solid solution along approximately 2-mm width. The 80-μm wide fusion zone includes two areas: Ni-based solid solution area adjacent to the coating and Fe-based solid solution area adjacent to the substrate. The phase in the diffusion-affected zone is mainly pearlite, and the width of this zone was approximately 280 μm. (3) The macrohardness of the composite coating is approximately HRC 55, and that of the substrate is HRC 18. The microhardness of the substrate is approximately Hv 150, and it increases gradually from the substrate to TZ. The average microhardness is Hv 750 in TZ, whereas it is approximately Hv 850 in CZ. The maximum value of microhardness in CZ is Hv 950. The hardness of the coating is dramatically improved compared to that of the substrate.

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Acknowledgements This work was financially supported by the National Natural Science Foundation of China (No. 51205178), the Natural Science Foundation of Gansu Province, China (No. 1208RJZA189), and the Doctor Fund Project of Lanzhou University of Technology.

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