Objective To further improve the performance of laser cladding CoCrFeNiW0.6 high entropy alloy coating by annealing. Methods CoCrFeNiW0.6 high entropy alloy coating was prepared on the surface of 45 steel by RFL-C1000 fiber laser. The high entropy alloy coating was annealed at different temperatures (600, 800, 1 000 ℃) in SXL-1200 tubular resistance furnace. The holding time was 2 h and the cooling method was furnace cooling. The microstructure, microhardness and friction and wear properties of the cladding layer were analyzed and tested by X-ray diffractometer (XRD), scanning electron microscope (SEM), energy dispersive spectrometer (EDS), microhardness tester, friction and wear testing machine. Results The CoCrFeNiW0.6 high entropy alloy coating was composed of FCC phase and μ phase (Fe7W6). After annealing at different temperatures, no new phase was precipitated in the coating, and the intensity of the μ phase diffraction peak showed a trend of first decreasing and then increasing. The microstructure of the coating changed significantly after high temperature annealing (800 ℃, 1 000 ℃, 2 h). After annealing at 800 ℃/2 h ℃, a large amount of μ phase precipitated between dendrites. After annealing at 1 000 ℃/2 h ℃, the grain boundaries began to fracture and decompose, and a large amount of W-rich particle phase (μ phase) precipitated inside the grains and at the grain boundaries. After annealing at 1 000 /2 h ℃, the cladding layer has a higher average microhardness of 475.68HV0.3, which is about 45% higher than that of the cladding layer without annealing. After annealing at 600 /2 h ℃, the average friction coefficient of the coating is the lowest, about 0.226, and the wear loss is the smallest. Compared with the coating without annealing, the wear loss is reduced by about 28%. The increase in annealing temperature did not significantly change the wear mechanism, which was mainly abrasive wear. Conclusion High temperature annealing can promote the formation of μ phase; after annealing, the hardness of CoCrFeNiW0.6 high entropy alloy coating is significantly improved, and the friction and wear properties of the coating are improved. The strengthening mechanism is solid solution strengthening and second phase strengthening.
Traditional alloys generally use one or two elements as the main elements, and a small amount of other functional elements is added to improve the performance of certain aspects of the alloy. Due to the limitation of alloy entropy, traditional alloys can no longer meet people’s increasing demands. Therefore, many scholars at home and abroad have gradually gone beyond the research field of traditional alloys and carried out extensive research on multi-principal alloys, and proposed a series of new alloy design methods. In 2004, Professor Ye Junwei et al. [1] proposed the concept of “high entropy alloy”, which is composed of 5 to 13 principal elements. In order to achieve high entropy value, the content of each principal element (expressed in atomic number fraction) is stipulated to be between 5% and 35%. High entropy alloys break the design principles of traditional alloys and create a precedent for the research of new alloys. Because high entropy alloys have unique high entropy effects, hysteresis diffusion effects, lattice distortion effects, and “cocktail” effects [2-5], they have good strength, hardness, wear resistance, corrosion resistance and thermal stability [6-7], and have gradually become a current research hotspot.
The atomic radius and electronegativity of elements such as Co, Cr, Fe, and Ni are similar, and the mixing enthalpy between atoms of different elements is small. The alloy composed of these elements easily forms a single face-centered cubic (FCC) solid solution structure and has good plasticity and toughness, but its hardness, strength, wear resistance and other aspects are average [8-9]. Therefore, other functional elements are often added to CoCrFeNi by alloying to improve the comprehensive performance of high entropy alloys. Liu et al. [10] prepared CoCrFeNiMox high entropy alloy blocks by vacuum arc melting technology. The study found that with the increase of Mo element content, the precipitation of intermetallic compounds continued to increase, the microhardness and yield stress of the alloy continued to increase, and the plasticity decreased; when x=1, the alloy has good comprehensive performance.
The failure of mechanical parts often occurs first on the surface of the material. Therefore, preparing a high entropy alloy coating on the surface of the part can not only ensure that the part has excellent performance, but also improve its economy. Ding et al. [11] prepared CoCrFeNiTiNbBx high entropy alloy coating. After adding B, TiB phase was synthesized in situ in the coating, and more short rod-shaped dendrites and equiaxed crystals appeared in the structure. The structure was finer and more uniform. After 60 min of friction and wear experiment, it was found that when x = 1.25, the wear amount of the coating was only about 0.45 times that of the substrate, and the wear mechanism changed from abrasive wear and adhesive wear (x < 1) to single abrasive wear (x ≥ 1). Gao et al. [12] prepared FeCoNiCrAl0.5Ti0.5 high entropy alloy coating on the surface of AISI 1045 steel and found that the coating had a chrysanthemum-shaped eutectic structure and had good wear resistance. Its wear rate was only one-third of that of the substrate. In the preparation process of high entropy alloy coating, composition segregation and various lattice structure defects are prone to occur, which will affect the structure and performance of high entropy alloy coating. The use of appropriate heat treatment process can cause the alloy to undergo phase transformation without changing the alloy element composition, thereby improving the alloy’s microstructure, reducing defects inside the alloy, reducing stress, and improving the alloy’s performance [13]. Xiong et al. [14] subjected FeMnCrNiCo+TiC (TiC mass fraction is 20%) high entropy alloy coating to 600/750/℃ ℃900 ℃ heat treatment and kept it warm for 75 h. The results showed that heat treatment made the TiC particles more evenly distributed and promoted the decomposition of TiC, which enhanced the solid solution strengthening effect of the coating. At 900 ℃, new M23C6 carbides were generated, which improved the microhardness, crack resistance and wear resistance of the coating.
Based on the above research and findings, this paper will use laser cladding technology to prepare CoCrFeNiW0.6 high entropy alloy coating, and through annealing treatment, explore the effects of different annealing temperatures (600, 800, 1000 ℃) on the coating structure and performance.
1 Experiment
The substrate is made of 45 steel, which is cut into 50 mm×50 mm×10 mm substrates. The 50 mm×50 mm surface of the substrate is selected as the cladding surface. The cladding surface is roughly ground with an angle grinder to remove the oxide scale, and the cladding surface is polished flat with sandpaper of 200 to 800 meshes for standby use. JA2003 electronic precision balance is used to weigh different masses of W single substance powder (purity of 99.9%, particle size of 300 mesh) and CoCrFeNi alloy powder (particle size of 45~105 μm). The atomic number fraction of each element in the powder is shown in Table 1. The alloy powders were mixed evenly using a MSK–SFM–1 horizontal planetary ball mill. The speed of the ball mill was 300 r/min and the mixing time was 120 min. The grinding balls were made of cemented carbide with a diameter of 3~5 mm and a ball-to-material ratio of 3:1. The mixed alloy powders were placed in a drying oven and dried at 80 °C (12 h) to ensure the dryness and fluidity of the powders.
The single-layer laser cladding experiment was carried out using an RFL–C1000 laser. The pre-coating method was selected as the powder supply method. The thickness of the pre-coating was 1 mm. Argon was introduced during the experiment to prevent the material from oxidizing at high temperatures. The argon flow rate was 10 L/min. The laser cladding process parameters are shown in Table 2. The high entropy alloy coating was annealed at different temperatures using an SXL–1200 tubular resistance furnace. The treatment temperatures were 600, 800, and 1 000 °C, respectively. The holding time was 2 h, and the cooling method was furnace cooling. For simplicity, the initial samples without annealing and the samples annealed at 600, 800, and 1 000 ℃ are represented by W0.6, W0.6/600 ℃, W0.6/800 ℃, and W0.6/1 000 ℃, respectively.
The clad samples were cut into samples with sizes of 10 mm×10 mm×10 mm and 10 mm×5 mm×10 mm using an EDM wire cutting machine. The coating surface was polished with 240-1000 mesh sandpaper, and the phase composition of multiple coatings was detected using a D8-X-ray diffractometer. The sample cross section was polished with 240-1500 mesh sandpaper and polished. After being corroded by saturated FeCl3 hydrochloric acid solution, the microstructure and element distribution of the cladding layer were observed and tested by GeminiSEM 500 thermal field emission scanning electron microscope (FSEM) and its configured characteristic X-ray energy spectrometer (EDS). The microhardness of the single-pass cladding layer was tested by HV1000Z automatic turret microhardness tester. The applied load was 3 N, the loading time was 10 s, and one point was punched every 0.1 mm along the cross section of the cladding layer. Three points were tested at the same depth and the average value was taken. The coating surface was continuously polished in the same direction using 800-grit sandpaper to ensure that the coating surface used in the experiment was flat and at the same roughness. The friction and wear test of the cladding layer was carried out by a ring-block friction and wear tester (M-2000 type). The friction ring was GCr15 steel after quenching. The experimental parameters are shown in Table 3.
2 Analysis and discussion
2.1 Microstructure analysis
The XRD patterns of the coatings at different annealing temperatures are shown in Figure 1. Without annealing, W0.6 showed FCC phase and μ phase diffraction peaks. Compared with the standard PDF card, the main component of the μ phase was Fe7W6. After annealing at 600, 800, and 1 000 ℃, the intensity of the μ phase diffraction peak decreased first and then increased. The intensity of the FCC phase diffraction peak was significantly lower than that of the initial sample, and at 1 000 ℃, the μ phase diffraction peak became more. Compared with the standard PDF card, the μ phase at this time was mainly composed of Fe7W6 and Co7W6. After calculation, the lattice constants of W0.6, W0.6/600℃, W0.6/800℃, and W0.6/1000℃ on the (200) crystal plane are 3.624, 3.628, 3.620, and 3.592 Å (1 Å=0.1 nm), respectively. The lattice constants increase first and then decrease. This is because laser cladding has a high heating rate and cooling rate. The prepared coating is prone to component segregation and various lattice structure defects without heat treatment, and the internal stress is large [13]. Without heat treatment, the precipitated μ phase reduces the content of W element in the solid solution, and the coating structure is unbalanced. After annealing at 600/2 h℃, the residual stress is released, the lattice structure defects are alleviated, and the atomic diffusion capacity is enhanced. The W atoms in the precipitated μ phase enter the solid solution and occupy the position in the original lattice lattice. Due to the large radius of W atoms, the lattice is distorted, the lattice constant increases, the volume fraction of μ phase decreases, and the intensity of μ phase diffraction peak decreases. In traditional high-temperature alloys, μ phase usually precipitates at 700~1000℃[15], which serves as a reference for this experiment. After annealing at 800 and 1000℃ for 2h, W atoms in the solid solution precipitate in the form of μ phase compounds, and the content of W element in the solid solution decreases, which makes the lattice constant smaller, the volume fraction of μ phase increases, and the intensity of μ phase diffraction peak increases.
The microstructure of the coating and the EDS point scanning results at different positions at different annealing temperatures are shown in Figure 2 and Table 4. Without annealing, the microstructure of the coating is a dendritic structure, and a small amount of precipitation is precipitated between the dendrites. EDS analysis shows that the precipitation is enriched with W element, and the dendrite width is about 4~5 μm. After annealing at 600/2 h ℃, the dendrites continued to grow with a width of about 4 μm, and the W content of the precipitated elements between the dendrites decreased. After annealing at 800/2 h ℃, the dendrite width was about 2~3 μm, and a large amount of W-rich precipitates were precipitated in the microstructure, and the Ni content was low. After annealing at 1000/2 h ℃, the grain morphology of the microstructure changed significantly, the grain boundaries began to fracture and decompose, and a large number of fine particles appeared inside the grains and at the grain boundaries. The width was less than 0.5 μm and tended to be evenly distributed. The dendrite width was about 3~4 μm, the W content of the fractured grain boundaries was high, and the W content in the matrix was relatively poor. Due to the high energy density of laser cladding, the substrate and the alloy powder were melted together to form a convection molten pool, resulting in the Fe content measured in the cladding layer being higher than the theoretical value [16]. The local EDS surface scanning image of the W0.6/1 000 ℃ high entropy alloy cladding layer is shown in Figure 3. It can be seen that the W element is obviously segregated in the bright area (precipitated particles and fracture grain boundaries), and the W content in the matrix is low.
Combining XRD images with the solidification theory of metal materials [17-18], under solidification conditions, the solute at the interface front will be enriched into a boundary layer, resulting in a difference between the liquid phase solidification temperature and the actual temperature at the interface front, thereby causing composition supercooling. At this time, the interface becomes unstable during the growth process and continues to grow deep into the liquid to form dendrites. The diffusion rate of the solute is much smaller than the solidification rate. Laser cladding has a faster heating and cooling rate, and it is very difficult to achieve equilibrium solidification. Therefore, when annealing is not performed, the faster cooling rate hinders the diffusion of W atoms, and some W atoms enter the solid solution lattice to form a substitutional solid solution. Another part of the W atoms combines with Fe atoms to form a new compound precipitate (Fe7W6), namely the μ phase, as shown in Figure 2a. After annealing at 600/2℃h, the residual stress is reduced, the lattice structure defects are alleviated, and the increase in temperature provides energy for atomic diffusion, allowing the W atoms to further diffuse into the solid solution lattice. At this time, the volume fraction of μ phase hard precipitates decreases, as shown in Figure 2b. After 800/2 ℃ h annealing, high temperature annealing causes W atoms inside the solid solution to precipitate from the solid solution[19], and aggregate with Fe atoms between dendrites, promoting the formation and growth of μ phase precipitation, as shown in Figure 2c. After 1000/2 h ℃ annealing, W atoms inside the solid solution further diffuse and precipitate from the solid solution, forming a large number of W-rich nano-scale particle phases, as shown in Figure 2d, which is very similar to the precipitation behavior of TCP phase after annealing of high-temperature alloys[20-21]. It can be seen in the XRD diagram that the μ phase diffraction peaks increase and are stronger at this time, so the particle phase is judged to be μ phase, proving that 1000/2 ℃ h annealing can play a role in refining the μ phase. In previous studies, many scholars have also reached similar conclusions[19,22].
2.2 Microhardness analysis
As can be seen from Figure 4, the increase in annealing temperature gradually increases the hardness of the cladding layer. After calculation, the average hardness of the W0.6, W0.6/600 ℃, W0.6/800 ℃, and W0.6 /1 000 ℃ cladding layers are (327.75± 15.35) HV0.3, (380.27±32.27) HV0.3, (419.15±34.15) HV0.3, and (475.68±32.92) HV0.3, respectively. After 1 000 /2 h ℃ annealing, the hardness of the cladding layer reaches the maximum value, and its hardness is increased by about 45% compared with the cladding layer without annealing. The improvement of microhardness is mainly due to the following reasons: W atoms have a large atomic radius. When they enter the lattice to form a substitutional solid solution, they will produce a lattice distortion effect, resulting in a more obvious solid solution strengthening effect; W elements promote the formation of μ phase, which is a hard phase [23] and can effectively inhibit dislocation movement, increase the resistance of dislocation slip, and achieve second phase strengthening; high temperature annealing treatment (800/2 h ℃, 1000/2 h ℃) promotes the precipitation of μ phase, which further improves the hardness of the alloy. This also shows that the CoCrFeNiW0.6 high entropy alloy coating has the phenomenon of annealing hardening. Before annealing treatment, the hardness of the heat affected zone is high and varies greatly. After annealing treatment, the hardness of the heat affected zone decreases and is close to the hardness of the substrate, with little change. Combined with the analysis of the heat affected zone microstructure (Figure 5), the heat affected zone structure has changed significantly before and after annealing, from the quenched structure before annealing to the pearlite structure after annealing, so that the hardness of this area decreases and approaches the hardness of the substrate.
2.3 Friction and wear performance
The friction coefficient curves of the coatings at different annealing temperatures are shown in Figure 6. Due to the running-in stage in the early stage of the experiment, the friction coefficient of the coating is generally unstable. As the contact surface between the friction pair and the coating increases, the friction coefficient will gradually stabilize [24]. The average friction coefficient and wear loss of the coating are shown in Figure 7. The average friction coefficients of the W0.6, W0.6/600 ℃, W0.6/800 ℃, and W0.6/1 000 ℃ coatings in the stable wear stage (5~20 min) are 0.311, 0.226, 0.288, and 0.291, respectively, and the wear loss of the coatings are 2.9, 2.1, 2.6, and 2.8 mg, respectively. After annealing, the friction coefficient and wear volume of the coating decreased. After annealing at 600/2 ℃ h, the friction coefficient and wear volume of the coating reached the lowest, and the wear volume was reduced by about 28% compared with the coating without annealing. After annealing at 800/2 h ℃ and 1 000/2 h ℃, the friction coefficient and wear volume of the coating increased to varying degrees, but were lower than those of the coating without annealing. Combined with the microstructure analysis, after annealing at 600/2 h ℃, the lattice structure defects were alleviated, and the embedding of a small amount of W-rich precipitation greatly improved the stability of the microscopic interface; the W element tended to diffuse into the solid solution, aggravating the lattice distortion and significantly enhancing the solid solution strengthening. After annealing at 800/2 h ℃, W atoms precipitate from the solid solution and combine with Fe atoms between dendrites to form a large number of μ phase hard precipitations, which increase the resistance of dislocation slip and improve the hardness of the coating. Excessive μ phase hard precipitation causes the coating to produce more wear debris with higher hardness during the friction and wear process. As abrasive particles, it continuously micro-cuts the coating surface, aggravating the wear of the coating. After annealing at 1000/2 h ℃, the precipitated μ phase is finer and denser, which further increases the hard precipitation wear debris produced during the friction and wear process. At the same time, the fracture decomposition of the grain boundary also reduces the stability of the micro interface, making the friction coefficient of the coating slightly higher than W0.6/800℃. It can be seen that the μ phase plays a “double-edged sword” role in room temperature friction and wear. It can not only improve the hardness of the coating, increase the resistance of dislocation sliding, and improve the friction and wear performance, but also produce abrasive particles with higher hardness in the friction and wear process, reducing the friction and wear performance. The final effect of the μ phase on the room temperature friction and wear performance is the combined result of the two situations.
2.4 Discussion on wear mechanism
The wear morphology of the coating is shown in Figure 8. Before annealing, the wear surface of the cladding layer showed wider and deeper plowing-shaped wear marks, and a large amount of wear debris accumulated in the plowing grooves. The main wear mechanism is abrasive wear. This is because before annealing, the hardness of the cladding layer is low, and the hard precipitated particles fall off during the friction and wear process. Under the “bonding” effect of the FCC solid solution with good plasticity [25], large-particle wear debris is formed. Under the action of the tangential force of the grinding ring, it moves relative to the cladding layer, forming wide and deep plowing-shaped wear marks. After annealing, the plow-shaped wear marks on the wear surface become shallower, and a large number of flaky fine wear debris appear. The wear mechanism is mainly abrasive wear. This is because after annealing at 600/2h℃, W atoms tend to diffuse into the solid solution lattice, which aggravates the lattice distortion, significantly enhances the solid solution strengthening effect, increases the hardness, effectively hinders dislocation slip, and reduces the μ phase hard particles, making the wear surface flatter and the wear marks shallower. After annealing at 800/2h℃ and 1000/2h℃, the large amount of μ phase hard precipitation significantly improves the hardness of the cladding layer, but it will also cause more hard particles to fall off during the friction and wear process, forming abrasive particles. Under the action of the tangential force of the grinding ring, the micro-cutting effect on the cladding layer is intensified, and elongated wear marks are formed.
3 Conclusions
1) The CoCrFeNiW0.6 high entropy alloy coating is composed of FCC phase and μ phase. After annealing at different temperatures, no new phase is precipitated in the coating, and the intensity of the μ phase diffraction peak shows a trend of first decreasing and then increasing. The microstructure of the coating is mainly composed of dendrites and a small amount of μ phase precipitation before annealing and after 600/2 h ℃ annealing. After 800/2 h ℃ annealing, a large amount of μ phase precipitation is precipitated between dendrites. After 1000/2 ℃ h annealing, the grain boundaries begin to fracture and decompose, and a large amount of W-rich particle phase (μ phase) is precipitated in the organization, which is evenly distributed inside the grains and at the grain boundaries, and the μ phase is refined.
2) After annealing at 1 000 /2 h ℃, the microhardness of the cladding layer reached the maximum value (475.68HV0.3), which was about 45% higher than that of the cladding layer without annealing. After annealing at 600 /2 h ℃, the average friction coefficient of the coating in the stable wear stage was the lowest (0.226), and the wear loss was the smallest (2.1 mg). Compared with the coating without annealing, the wear loss was reduced by about 28%, and the friction and wear performance of the coating was the best. The strengthening mechanism was mainly solid solution strengthening and second phase (μ phase) strengthening. After annealing, the wear mechanism did not change significantly, mainly abrasive wear, the wear surface became flat, and the wear mark became shallow.
