Metal-matrix composite fabricated with gas tungsten arc melt injection and precoated with NiCrBSi alloy to increase the volume fraction of WC particles

1 Introduction

The wear resistance of metal-matrix composites (MMCs) with remarkable content of hard phases with very high hardness, such as carbides, is well known. The investigation of Liyanage et al. [1] on plasma transferred arc welded Ni-WC overlays shows that the MMC overlays are from two to five times more wear resistant than the matrix alloys without WC particles.

Many welding and cladding methods have been used to produce WC particle-reinforced MMCs on low-cost substrates. However, there are still some problems to be solved. The first problem is the degree of dissolution of WC particles. In the process of welding or cladding, WC particles are heated by a heat source and dissolve into the molten pool partially or even completely. Katsich and Badisch [2] have accessed the effect of carbide degradation in a WC/W₂C-reinforced Ni-based hard facing, and the results show significant carbide degradation with increasing welding current, resulting in a significant reduced primary carbide content and carbide diameter. Reduced carbide content indicated a significant wear rate increase under pure three-body abrasion conditions. Dissolution results in lowered wear resistance, as there are less WC particles that remained to provide wear protection. The contents of carbon and tungsten in the matrix increase and the toughness of the matrix decreases, so the risk of cracking increases. The extra carbon and tungsten precipitate as M₆C-type carbides, which are less wear resistant than WC particles.

The second problem is the volume fraction of WC particles in the matrix. Jankauskas et al. [3] have investigated the effect of WC grain size and content on low stress abrasive wear. They have found that the wear rate of hard facing decreases with the increase in WC content, and a factor of 9 has been achieved when WC content is 42–43 wt%. A further increase in WC content is generally considered difficult because of the problem of WC particle dissolution.

The third problem is the distribution of WC particles in the matrix. Fernández et al. [4] have studied the tribological improvement of NiCrBSi laser cladding coating reinforced with different weight percentages of WC particles. They have found that WC particles tend to precipitate at the bottom of the melt coating; hence, the percentage of carbides increases with the depth of the coating. The density of WC particles is much higher than that of the matrix. They are prone to sink down at the bottom of the molten pool. The tendency of segregation of WC particles is related to the time they spend in the molten pool, which is determined by the procedure and process parameters. With a certain procedure and fixed parameters, little can be done to control the distribution of WC particles in the matrix. The segregation of WC particles decreases the wear resistance of the top layer of the MMC and increases the cracking tendency of the MMC/substrate interface.

To solve the problem of dissolution and segregation of WC particles, Vreeling et al. [5] proposed a laser melt injection process. In the laser melt injection process, a laser beam is used to melt the substrate to form a molten pool. With the movement of the laser beam, a tail is formed behind the molten pool, and the carbide particles are injected into the tail of the molten pool, avoiding direct interaction with the laser beam. With this method, they have strengthened titanium alloy with WC particles [5] and aluminum alloy with SiC particles [6]. Zhao et al. [7] use a plasma arc to melt and inject WC particles into low carbon steel substrates. Other carbide particles, such as Cr₃C₂-NiCr [8] and SiC [9], have also been successfully incorporated into overlays with little dissolution by the melt and injection processes.

However, the volume fraction of the incorporated carbide particles is still limited. It is found that only the particles with velocity higher than a critical velocity v_c can overcome the surface tension and enter the molten pool [6]. To increase the volume fraction of the carbide particles in the MMCs, the velocity of the injected particles must be increased or the surface tension of the molten pool must be decreased. NiCrBSi alloys, which have low melting point, low surface tension when melted, and excellent wear resistance, are often sprayed or cladded on industrial components for wear protection [10]. If they are precoated on the substrate, the surface tension of the molten pool will be greatly decreased. Therefore, the volume fraction of WC particles in the MMCs can be increased.

In this paper, WC particle-reinforced MMCs were fabricated on Q235 low carbon steel substrates with the gas tungsten arc melt injection (GTAMI) process. NiCrBSi alloy was deposited on the substrate before melt and injection, to lower the surface tension, to allow more WC particles to be incorporated into the MMCs. The microstructures and wear resistance of the MMCs were reported.

2 Materials and methods

The substrate was of Q235 low carbon steel (similar to ASTM A570 Gr.A; Anshan Steel Co., Ltd., Anshan, China). The chemical composition and mechanical properties of the substrate are shown in Table 1. The dimensions of the substrate were 250×50×5 mm³. The NiCrBSi alloy powder used for precoating was manufactured by VEHA Co., Ltd. (Keorekovsk, Russia), a Russian company that manufactures thermal spray equipment and materials. The morphology of the NiCrBSi powder particles is shown in Figure 1. The nominal diameters of the particles were 5–45 μm. The composition of the NiCrBSi powder is shown in Table 2. The injected particles were cast and crushed WC-8% Co particles (Harbin Welding Institute, Harbin, China). The size of the WC-8% Co particles was 350–700 μm.

Figure 1:

Morphology of NiCrBSi powder particles.

The substrate was degreased, dried, and grit blasted. The NiCrBSi powder was flame sprayed onto the substrate. The coating was 1 mm thick.

The schematic diagram of the GTAMI system is shown in Figure 2. The coated substrate was fixed on a movable platform. The moving velocity of the platform was adjustable. A gas tungsten arc torch was set above the platform with an angle α to the normal of the platform. The substrate with coating was heated by the arc, and a molten pool formed under the arc. The platform moved horizontally with a velocity of 4 mm/s. The molten pool left a tail behind the arc. The injection nozzle was fixed together with the torch with an angle β to the normal of the platform. WC particles were stored in a hopper. The height of the hopper was adjustable. WC particles were driven out of the hopper by a roller and were accelerated down through the pipe and out the injection nozzle by gravity. WC particles left the nozzle with a certain velocity, entered the tail of the molten pool, and were caught in the matrix after the solidification of the molten pool. The parameters of the GTAMI process are shown in Table 3. Single-pass specimens with different hopper heights and multipass specimens were prepared. The overlap of the multipass specimen was 3 mm.

Figure 2:

Schematic diagram of the GTAMI system.

The specimens were sectioned, mounted, ground, polished, and etched. The etchant was FeCl₃ solution (5 g FeCl₃, 25 ml HCl, 25 ml ethanol). The microstructures of the MMCs were investigated with a Hitachi model S-570 scanning electron microscopy (SEM). The compositions of different phases were analyzed with a Tracor Northern model TN-5502 energy-dispersive X-ray spectroscopy (EDS). The phases in the MMCs were analyzed with a model D/max-rB X-ray diffraction (XRD) analyzer (Rigaku, Japan).

The wear resistance of the MMCs was tested with a M-200 wear tester (Jinan Yihua Testing Equipment Co., Ltd., Jinan, China). The specimen was sectioned to 20×6×5 mm³ and ground to obtain a flat surface for testing. The counterpart was a Al₂O₃ ring with a diameter of 500 mm. The rotation rate of the ring was 200 rpm. The load on the ring was 490 N. The wear loss of the specimen was weighed every 20 min in the test with a CHI604C electric balance with precision of 0.1 mg (Shanghai Chenhua Co., Ltd., Shanghai, China). The substrate was also tested for comparison. The wear loss of the substrate was weighed every minute.

3 Results and discussion

Figure 3 shows the cross-sections of specimens prepared with different hopper height H. Both the precoated NiCrBSi alloy and the substrate were melted and mixed together, forming the matrix of the MMCs. Many irregular-shaped WC particles were found embedded in the matrix. The edges of WC particles were sharp. The distribution of WC particles in the matrix of the specimen with H=100 mm (Figure 3A) was more uniform than those with larger H (Figure 3B–D). When H=400 mm, almost all WC particles were in the lower part of the MMCs. All MMCs were crack free.

Figure 3:

Cross-sections of the specimens prepared with different hopper heights: (A) H=100 mm, (B) H=200 mm, (C) H=300 mm, and (D) H=400 mm.

The volume fraction of WC particles shown in Figure 3 was estimated with Image-Pro Plus and is shown in Figure 4. Figure 4 shows that the volume fraction of WC particles increased with the decrease in H in the tested range. Compared with the MMCs produced with plasma arc melt injection process [11], the volume fraction of WC particles of the MMCs produced with the GTAMI process and precoated NiCrBSi alloy was greatly increased.

Figure 4:

Influence of hopper height H on the volume fraction of WC particles in the MMCs.

The increase in the volume fraction of WC particles in the MMCs was caused by precoating the substrate with NiCrBSi alloy. WC particles were driven out of the hopper by the roller at height H and were accelerated down through the pipe by gravity. The velocities of these particles were within a certain range when they arrived at the surface of the molten pool. The mean velocity and distribution of velocities of the particles depended on H and the injection system. When a certain WC particle arrived at the surface of the molten pool, whether it could enter the molten pool depended on its velocity. Only when its velocity was higher than a critical velocity v_c could it overcome the surface tension to enter the molten pool and was embedded in the matrix after the solidification of the molten pool. The critical velocity v_c is related to the radius and density of the carbide particles and the surface tensions between phases. It is determined by Vreeling et al. [6] as

where R is the radius of the particle (m), ρ is the density of the particle (kg/m³), σ_lv is the surface tension between the liquid and the vapor (N/m), σ_lp is the surface tension between the liquid and the particle (N/m), and σ_pv is the surface tension between the particle and the vapor (N/m).

From Equation (1), it was easy to know that the decrease in σ_lv and σ_lp would result in a lower critical velocity v_c, which meant that the velocities of more WC particles would exceed v_c and would enter the molten pool. NiCrBSi is a self-fluxing alloy powder used for cladding and hard facing. The surface tension of its melt was much lower than that of the liquid steel. Therefore, precoating the Q235 substrate with NiCrBSi alloy powder resulted in a low surface tension of the molten pool, and more WC particles entered the molten pool.

When the hopper was moved higher, the distance from the tip of the nozzle to the substrate increased. WC particles leaving the nozzle scattered to a larger area, and less WC particles arrived at the molten pool. Therefore, the volume fraction of WC particles in the MMCs decreased with increasing hopper distance (Figure 4). The mean velocity of WC particles increased when the hopper was moved higher, so WC particles entered deeper into the molten pool. With the increase in hopper height H, the segregation of WC particles was achieved. When H=400 mm, the segregation of WC particles became very evident, as shown in Figure 3D. Limited by the structure of the system, H could not be adjusted to <100 mm. Results with shorter H could not be obtained, but it was reasonable to guess a shallower distribution of WC particles in the MMCs.

Figure 5 shows the XRD patterns of the single-pass MMCs. Besides WC particles, two other phases, Fe-Ni solid solution and Fe₃W₃C carbide, were found in the MMCs. The precoated NiCrBSi alloy was melted and mixed with the substrate, forming the Fe-Ni solid solution phase. This was the matrix of the MMCs. The XRD patterns showed the Fe₃W₃C phase in the MMCs, which meant dissolution of WC particles. However, SEM images showed that the dissolution of WC particles was not serious. Most of the WC particles retained their original shapes (Figure 3). Only some surface of WC particles dissolved into the molten pool and precipitated as Fe₃W₃C carbides.

Figure 5:

XRD patterns of the single-pass MMCs.

Figure 6 shows the microstructures of the single-pass specimen. Besides WC particles, three other kinds of microstructures were found in the matrix, marked as phases A–C, respectively. Phase B looked a little darker than phase A in the SEM image. Phase C was a white herringbone structure. Both phases A and C were located at the grain boundary of phase B.

Figure 6:

Microstructures of the single-pass specimen in the vicinity of a WC particle at the (A) top, (B) middle, and (C) bottom parts of the MMCs.

The EDS results of these three microstructures are shown in Table 4. The EDS analysis results showed that both phases A and B were Fe-Ni solid solutions. Although there were more alloying elements (tungsten, silicon, and chromium) in phase A (comparing the amount of the alloying elements at A1 and B1, A2 and B2, and A3 and B3, as shown in Table 4), there were less alloying elements in phase B, and its melting point must be higher. In the solidifying process, phase B solidified first. With less chromium content, its corrosion resistance was worse than phase A. Therefore, it looked darker in the SEM image after etching. There were more alloying elements in phase A, and it solidified after phase B in the solidifying process at the grain boundary of phase B and looked a little whiter in the SEM image. After the solidification of the Fe-Ni solid solution, more alloying elements were left in the remaining liquid metal, and the liquid metal solidified as eutectic herringbone structure.

For the same kind of Fe-Ni solid solution, such as phase A, the contents of nickel, tungsten, silicon, and chromium decreased from the top part to the bottom part of the matrix (comparing the amount of the alloying elements at A1, A2, and A3, as shown in Table 4). The precoated NiCrBSi coating and the substrate were melted and mixed when heated. Nickel, chromium, boron, and silicon were mixed with the iron from the substrate to form the matrix (boron could not be detected with the EDS analysis). The NiCrBSi coating was on the top of the substrate, so the contents of nickel, chromium, and silicon decreased with distance from the surface of the MMCs. The distribution of tungsten was the same as that of other elements, but the mechanism was different. When a WC particle entered the molten pool, it was heated by the liquid metal and began to dissolve. The dissolution process continued until the particle was caught by the solidifying matrix metal. The particle found at the bottom part must get through the top part of the molten pool. This meant that the dissolution of the particles at the bottom part of the matrix had a contribution to the tungsten content at the top part, but the dissolution of the particles at the top part of the matrix had no contribution to the tungsten content at the bottom part. Therefore, the tungsten content decreased with distance from the surface of the MMCs. This kind of distribution of the alloying elements was beneficial to the mechanical property of the MMCs. The ratio of phase C in Figure 6A–C to the matrix (phase A+phase B) was calculated with Image-Pro Plus as 25.5%, 22.1%, and 14.7%, respectively. This trend was the same as the distribution of elemental tungsten.

The microstructures of the multipass specimen are shown in Figure 7. Figure 7A is an overview of the overlapping area of the first and second passes. To investigate the influence of the second pass, the boundary between the first and second passed must be determined. The local enlargement of area 1 in Figure 7A is shown in Figure 7B. The white belt consisted of herringbone structure. A columnar crystal was found growing from the white belt towards the second pass. This suggested that the right edge of the white belt was the boundary between the first and second passes, as indicated by the dashed line in Figure 7A and B. The white belt was located at the heat-affected zone of the second pass, and its shape fit the fusion boundary of the second pass. This meant that the herringbone structure was formed in the melting process of the second pass.

Figure 7:

Microstructures of the multipass specimen: (A) overview of the overlapping area of the first and second passes, (B) local enlargement of area 1 in (A), (C) local enlargement of area 2 in (A), and (D) local enlargement of area 3 in (A).

The dashed line in Figure 7A crossed the particle marked P. This suggested that particle P was injected into the molten pool of the first pass in the melt and injection processes of the first pass, and its top right corner was reheated by the arc in the melt and injection processes of the second pass. The local enlargement of area 2 in Figure 7A is shown in Figure 7C. For comparison, the local enlargement of the bottom left corner of particle P, area 3 in Figure 7A, is shown in Figure 7D. The bottom left corner of particle P was not reheated by the arc, and the microstructures in its vicinity were the same as the ones in the single-pass specimen. Although the top right corner of particle P was reheated by the arc, more WC particles dissolved into the liquid metal. The EDS results of different microstructures in Figure 7C are shown in Table 5. Comparing Table 5 with Table 4, the tungsten content at point A4 was much higher than that at points A1, A2, and A3, and it was also higher for point B4. This was the result of more WC dissolution. The difference of the tungsten content between points A4 and B4 was less than that between points A1 and B1. The same was true for other elements. The remelting provided more time for the elements to distribute uniformly. A herringbone structure was not fully developed, so the alloying element contents at point C4 were very close to those at point A4. Around the WC particle, some carbide blocks with high tungsten content appeared. Judged by its appearance and tungsten content, it was believed to be Fe₃W₃C. The remelting of WC particles provided the matrix with more tungsten and carbon, and when the tungsten content was high, Fe₃W₃C was more likely to precipitate other than the herringbone structure.

The wear loss of the specimen produced with H=100 mm and the substrate versus time are shown in Figure 8. The wear test of the substrate could not be carried on after 4 min because of severe wear. The wear loss of the substrate after 4 min was 524.9 mg. The wear loss of the MMCs after 40 min test was only 6.9 mg. It could be concluded that the wear resistance of the substrate was greatly improved with WC particle-reinforced MMCs produced by the GTAMI process.

Figure 8:

Wear loss of the specimen produced with H=100 mm and the substrate versus time.

4 Conclusions

MMCs were fabricated using the GTAMI process, with WC particles injected into low carbon steel Q235 substrate precoated with NiCrBSi alloy. The volume fraction of WC particles in the MMCs was found to be much higher than that in the MMCs produced with plasma arc injection process, without precoated coatings. The volume fraction increased with decreasing hopper height and was as high as 44% when the hopper height was 100 mm. The dissolution of WC particles was very little, and WC particles kept their sharp edges. The dissolved WC particles precipitated as herringbone structure. The contents of the alloying elements in the matrix decreased from the top to the bottom of the matrix. WC particles in the overlapping area of the successive passes were reheated by the following melting and dissolved more. Fe₃W₃C carbide blocks could be found in the overlapping area. More herringbone structure appeared in the previous pass, at the heat-affected zone of the following pass. The wear loss of the MMCs after 40 min was 6.9 mg, 76 times less than that of the substrate after the 4 min test.