Effects of B content on wear and corrosion resistance of laser-cladded Co-based alloy coatings

TC4 alloy was selected as the substrate; its main chemical composition is shown in Table 1. The surface of TC4 alloy was polished with sandpaper to ensure a smooth, flat surface. Secondly, ultrasonic cleaning was carried out with acetone solution to remove oil stains on the surface, and the alloy was dried. The laser cladding materials selected in this experiment were Co-based alloy powders with different B contents (Co₃₄Fe₈Cr₂₉Ni₈Si₇)_100–xB_x (x = 0,4,6,8,12,16), and the particle size of the powder was about 200–400 mesh. The marked numbers and nominal components of the samples are shown in Table 2. The morphology of the B4 powder is shown in Figures 1a, and 1b illustrates its structure. Nd: YAG solid-state pulsed laser (RH-700 type Nd: YAG laser, China) was used in this laser cladding experiment. The processing parameters of the laser cladding included an output current of 380A, a scanning speed of 100 mm/min, a frequency of 4 HZ, a pulse width of 8 ms, and a preset powder thickness of 200 μm. The preinstalled powder and substrate were placed in a chamber made of quartz glass, and argon with a purity of 99.99% was sprayed continuously during cladding to prevent oxidation of the coatings during preparation.

Fig. 1.

To study the microstructure of the alloy coatings, a wire-cutting machine was used to make the alloyed coatings into a sample with a size of 10 ×10 × 10 mm. The cross-section and surface of the samples were successively polished with a different mesh of metallographic sandpaper, and mechanical polishing was carried out. Aqua regia was used to etch the coating (HNO₃: HCl = 1:3) for 90 s. The microstructure of the samples was characterized by scanning electron microscope (SEM, Zeiss-MERLIN Compact, Germany) equipped with an energy dispersion spectrometer. The phase analysis was performed by Cu-Ka radiation on an X-ray diffractometer (XRD, D/max 2500, Japan). The XRD spectra were measured in 2θ range from 10° to 90° with a step size of 0.02° and a scanning speed of 5°/min. Transmission electron microscopy (TEM, JEOL-2100, Japan) was used to verify some phases in the coatings.

The microhardness was measured with a microhardness tester (TH701, China) under an applied load of 200 g and a residence time of 10 s. According to the standard of ASTM D1044-18, the friction and wear tests were carried out on a high-temperature abrasive testing machine (HT-1000, China) at a temperature of 20°C, a load of 1150 g, a rotation speed of 1120 r/min, and a wear time of 40 min. ZrO₂ balls with a diameter of 2.5 mm were selected as grinding materials. The electrochemical corrosion test was carried out in the electrochemical workstation (Zahner, Germany) with 3.5 wt% NaCl electrolyte. A three-electrode device was used in the experiment, in which the assistant electrode was a platinum plate, and the reference electrode was a saturated calomel electrode. Before the measurement of the polarization curves, the samples were immersed in the solution for 20 min, followed by the measurement of the open-circuit potentials (OCP) in the solution for 1 h. Potentiodynamic polarization curves were tested at a scanning rate of 1 mV/s in the range of -0.9 V to 0.3 V.

Figure 2 shows the XRD spectrum of the designed alloys, indicating that the coatings mainly consisted of TiCr, Fe_0.1Ti_0.18V_0.72, CoTi, Ti₂Ni. A phase of TiB would appear with the gradual increase of B content. The reason was that the hard phase TiB was formed in the process of laser cladding due to the decrease of the Co-based alloy element content and the increase of B content. TiB-reinforced titanium matrix composites were favorable for improving the coatings in this study because of their excellent properties [28]. XRD patterns showed that a wide diffraction peak appeared in the 2θ region, ranging from 44° to 45°, which was attributed to the existence of the amorphous phase in the coatings. At the same time, a strong crystal diffraction peak was superimposed on the diffuse peak, indicating the coexistence of amorphous and crystalline phases in the coatings. The volume percentage of amorphous phase in the coatings was calculated by Pseudo-Voigt function fitting with the method of Verdon with the calculation results shown in Table 3. It could be seen that with the increase of B content, the amorphous content in the coatings increased first and then decreased. When the B content was 6%, the volume fraction of amorphous phase in the coating was highest at around 65.86%.

Fig. 2.

Figure 3 shows the microstructure of three different regions of a typical laser coating crosssection. Figure 3 panels b, c, and d are the enlarged SEM images of the B, C, and D regions in Figure 3a, respectively. The coating exhibited excellent metallurgical bonding with the substrate with no pores. Nevertheless, because of the different thermal expansion coefficients of the substrate and coating materials, as well as the high heating and cooling rates of LC, a small number of cracks appeared in the coating [29] as shown in Figure 3a. Numerous dendrites were found in the middle and bottom regions of the coating, while the top region of the coating was mainly composed of equiaxed crystals. According to the theory of metallurgical solidification, the growth of coating microstructure is closely related to the temperature gradient and solidification rate [30]. The components near the heat-affected zone were not undercooling because of the large temperature gradient and low solidification rate. As a result, the interaction between laser and cladding powder could facilitate the flow and agitation of the molten pool; and the microstructure of the coating exhibited columnar morphology, honeycomb morphology, or dendritic morphology. As a result of the high undercooling rate, the microstructure at the top of the coating was mainly equiaxed crystal.

Fig. 3.

Figure 4 shows the microstructure of Co-based coatings with different B contents. As shown in Figure 4a, the coating without B element exhibited a typical equiaxed crystal structure, and there was no amorphous region. As shown in Figure 4b, with the addition of B element, the dendrite growth in the coating was improved, forming a microstructure in which the amorphous phase and the crystalline phase were nested and wrapped. The B6 coating was very bright and smooth with no existence of dendrite, as shown in Figure 4c. Therefore, it could be concluded that the coating was mainly composed of amorphous phase, or the corrosion resistance of the coating was higher than that of other coatings. From Figure 4d, e, and f, with the further increasing of B contents, the microstructure of the coatings showed dendrite structure again, and the crystallinity was significantly improved. It could be seen that a change of B contents would affect the microstructure of Co-based alloy coatings, which might exert different effects on the coating properties. Figure 5 shows the SEM image and EDS elemental distribution of the B6 coating, indicating that the elements of B6 coating were uniformly distributed in the amorphous region, and there was basically no large-scale segregation of Ti, Co, Fe, Cr, Ni, or Si metal elements. There was a slight enrichment of B element, which further proved that the addition of an appropriate amount of B element could effectively inhibit the segregation and crystallization of metal elements and improve the formation ability of Co-based amorphous phases.

Fig. 4.

Fig. 5.

In order to confirm that the coating contained amorphous phase, the B6 coating was studied in detail using TEM. Figure 6a was the bright field TEM image from the upper region of the B6 coating, and Figure 6b gives detailed information about the border between the amorphous phase and the crystalline phase in Figure 6a. The white bright spots in Figure 6a were identified as amorphous. The atomic arrangement in the amorphous phase was disordered, so the bright field region was determined as the amorphous region in Figure 6b. In contrast, the atoms in the dark field were arranged in an orderly fashion, which was determined to be the crystalline area. The existence of both crystalline and amorphous phases in the coating was further confirmed by TEM.

Fig. 6.

The microhardness distribution of (Co₃₄Fe₈Cr₂₉Ni₈Si₇)_100–x B_x alloy coatings presented a typical step-like pattern that was about 200 μm as shown in Figure 7. The hardness of the Co-based coatings first increased and then decreased with the increase of B content. When the B content was 6%, the hardness of the B6 coating was up to 1210 HV_0.2 which was 3.4 times that of the substrate (350 HV_0.2). The microhardness of the coatings was significantly improved due to the combination of grain refinement strengthening, solid solution strengthening, and second-phase strengthening by the hard compound TiB. Certainly, the different content of amorphous phase was also one of the reasons why the microhardness of the coatings was unequal. The coating deviated from the nominal composition, and the microhardness decreased significantly because of the mutual diffusion of elements between the coating and the substrate during LC.

Fig. 7.

The friction coefficient of the (Co₃₄Fe₈Cr₂₉Ni₈Si₇)_100–xB_x alloy coatings first increased and then decreased along with the increase of B contents as shown in Figure 8. Because of the peeling and wear of the coating surfaces during the friction process, the friction coefficient curve fluctuated slightly and finally stabilized at about 0.18. The friction coefficient of B6 coating was the highest, and the friction curve had gone through a gradually declining process. However, the wear loss was the decisive parameter to evaluate wear performance. Under the same conditions, the smaller the wear loss, the higher the wear resistance. The wear loss of B6 coating was as low as 5.8 mg which was one-fourth of that of substrate as shown by the wear loss of (Co₃₄Fe₈Cr₂₉Ni₈Si₇)_100–xB_x alloy coatings and substrate presented in Figure 9. To further analyze the wear mechanism, the wear surface was observed using SEM.

Fig. 8.

Fig. 9.

Figure 10 shows the wear morphology of the (Co₃₄Fe₈Cr₂₉Ni₈Si₇)_100–xB_x alloy coatings. As shown in Figure 10a, severe plastic deformation, spalling, and cracks along the wear direction were observed on the wear surface of B0 coating, indicating that serious adhesive wear and fatigue wear occurred during the wear test. According to EDS analysis, the oxygen content of the B0 coating was higher than that of the B6 coating, consequently the B0 coating was also accompanied by oxidative wear as shown in Table 4. The brittle Fe0.1Ti0.18V0.72 intermetallic compounds with dendrite structure in the coating would give birth to stress concentration, which in turn led to the initiation of fatigue cracks on the coating surface and finally extended to the surface along the direction of friction, resulting in spalling of the coating. At the same time, under the rolling of the grinding ball, the formation of abrasive particles in the spalling pit intensified the wear, resulting in the increase of the friction coefficient of the coating and relatively high wear loss [31]. Unlike with B6 coating, cracks were not found on the wear surface of B4 coating as shown in Figure 10b. Shallow furrow tracks along the wear direction were observed, however, indicating that the wear mechanism of B4 coating was abrasive wear. Figure 10c shows the wear surface of the B6 coating with few grooves, which was much smoother as compared with the other coatings. As seen in Figure 10 (d, e, and f), the wear morphology of each sample was similar, but the deformation was smaller than that of B0 coating. The curved grooves, fine debris, and scratches on the whole wear surface were observed, indicating that the wear mechanism was both abrasive wear and slight adhesive wear. Due to the addition of excessive B element, the content of amorphous phase in the coatings began to decrease, the density of defects such as dislocations increased, and brittle intermetallics and borides began to form again. The coatings with less hardness would produce higher wear during use, but the wear resistance of the coatings containing boride and a small amount of amorphous phase was still higher than that of the B0 coating.

Fig. 10.

Based on the above test results, the wear morphology of the coatings was closely related to the content of internal amorphous phase. The addition of B element was theoretically beneficial to the formation of amorphous phase and the reduction of micro-defects of Co-based coatings. The coatings exhibited different wear mechanisms which gradually transitioned from severe fatigue spalling and oxidative wear to slight abrasive wear. When the content of B reached 6%, the coating showed the best wear resistance, and the wear mechanism was mainly slight abrasive wear.

Figure 11 shows the polarization curve of (Co₃₄Fe₈Cr₂₉Ni₈Si₇)_100–xB_x alloy coatings in 3.5 wt% NaCl solution at room temperature together with the corresponding electrochemical parameters as shown in Table 5. According to the electrochemical corrosion theory, the free corrosion current density (I_corr) of the alloy is the standard for measuring the corrosion rate, and smaller I_corr implied stronger corrosion resistance [32]. A comparison of the corrosion resistance of (Co₃₄Fe₈Cr₂₉Ni₈Si₇)_100–xB_x alloy coatings demonstrated that the I_corr value first decreased and then increased with the increase of B contents, indicating that the addition of B could improve the corrosion resistance of Co-based alloy coatings within a certain range. Under the same corrosion conditions, the I_corr of B4 coating was the lowest at 5.846 × 10^–5 A/cm², indicating that the B6 coating had provided excellent corrosion protection effect on the substrate in comparison to other coatings.

Fig. 11.

Amorphous metal alloys, or so-called metal glass (MG), were recognized to have excellent properties due to their unique atomic structure, such as high strength, high hardness, and strong corrosion resistance [33–37]. Generally, the difficulty of forming amorphous metal alloys was usually represented by glass-formation ability (GFA) [38]. According to Gibbs free energy formula [39], coating materials with higher mixing entropy and lower mixing enthalpy showed better GFA. The study had also shown that high condensation rate and low dilution rate could significantly improve the GFA of the coatings [40]. Therefore, the Co-based coatings with high mixing entropy and low mixing enthalpy prepared by LC technology in this paper completely conformed to the conditions of amorphous formation. Under the same compressive stress, unlike the crystalline materials with dislocation slip and plastic deformation on the macro and micro scales, the coatings with high amorphous content exhibited less deformation and high hardness due to the absence of defects such as dislocations [41]. This was very consistent with the results shown in Figure 7. Secondly, under the influence of matrix melting and dilution, the chemical composition at the bottom of the molten pool clearly deviated from the deep eutectic composition, which reduced GFA and increased the critical cooling rate required for the formation of amorphous alloy. However, the chemical composition in the upper part of the coatings still maintained a relatively pure original composition due to low dilution, so the amorphous phase was mainly concentrated in the region. Overwhelmingly higher amorphous content had given birth to better wear resistance, yet the crack resistance deteriorated, which made the deeper corrosion of Ti alloy substrate and resulted in its relatively lower corrosion resistance.