Microstructural evolution and corrosion resistance of CoCuNiTiX0.6 (X = Mn, Al) HEA coatings on the 45-steel substrate

3.1. Phase composition

Fig.1 shows XRD patterns of CoCuNiTiX_0.6 (X= Mn, Al) HEACs. The $\mathrm{s}\mathrm{i}{\mathrm{n}}^2\theta$ ratios were calculated for the diffraction peaks of the two alloys, marked as FCC and BCC, as shown in Table 1. According to the extinction law of the lattice [26], when the $\mathrm{s}\mathrm{i}{\mathrm{n}}^2\theta$ ratio of the diffraction peak angle is 3:4:8:11:12:16: ... and 2:4:6:8:10:12: ..., the crystal structures are face-centered cubic and body-centered cubic, respectively. Based on the Bragg equation and the lattice structure theory [26]:

where $d$ is the interplanar spacing, $\theta$ is the angle between the incident X-ray and the corresponding crystal plane, and $\lambda$ is the wavelength of the target material emitting X-rays, which is the wavelength of the $K_{\alpha }$ characteristic radiation from a Cu target, approximately 1.5418 Å. $a$ is the lattice constant, and $h$, $k$, and $l$ are Miller indices. The Miller indices corresponding to the diffraction peak of each alloy phase were calculated using Eqs. (1-3), as shown in Table 1.

Fig. 1XRD patterns of CoCuNiTiMn0.6 (top) and CoCuNiTiAl0.6 (bottom) HEACs

Compared with the Joint Committee on Powder Diffraction Standards (JCPDS) cards, the shape and number of diffraction peaks of the two alloys are similar to the superimposed peaks of Ni (JCPDS No. 65-2865) with FCC structure and Ti (JCPDS No. 65-5970) with BCC structure. Based on the strongest diffraction peak of Ni (44.493°) of the standard card, the strongest peaks of FCC labeling in CoCuNiTiMn_0.6 and CoCuNiTiAl_0.6 shifted toward lower angles by 0.970° and 0.540°, respectively. Based on the strongest peak angle of Ti (38.415°) diffraction peak of the standard card, the strongest peaks of BCC labeling in CoCuNiTiMn_0.6 and CoCuNiTiAl_0.6 shifted to the small angle direction by 6.206° and 6.464°, respectively. Table 2 shows the characteristic parameters of each alloying element in CoCuNiTiX_0.6 (X = Mn. Al) HEACs. As shown in Table 2, the atomic radii of Ni (0.125 nm) and Ti (0.146 nm) are the minimum and maximum values of all elements in the two alloys, respectively. When other alloying elements enter and occupy the lattice sites of Ni with an FCC structure, this inevitably leads to cell expansion. Occupying the lattice sites of Ti with BCC inevitably causes cell shrinkage. From Eq. (1), cell expansion and shrinkage cause the diffraction peak angle to shift to lower and higher angles, respectively. Therefore, the diffraction peak angles of the two alloys were shifted compared to the data of the standard diffraction card.

Fig. 2WPFRR analysis of XRD data of CoCuNiTiX0.6 (X = Mn, Al) HEACs

a) CoCuNiTiMn_0.6

b) CoCuNiTiAl_0.6

Fig. 2 shows the whole pattern fitting and Rietveld refinement (WPFRR) analysis of the X-ray diffraction data of CoCuNiTiX_0.6 (X = Mn, Al) HEACs. The refinement weight factors ($R$) of the CoCuNiTiMn_0.6 and CoCuNiTiAl_0.6 alloys are 4.66 and 3.31 %, respectively. The refinement results are presented in Table 3. As shown in Table 3, compared with the CoCuNiTiMn_0.6 alloy, the lattice constants and unit cell volumes of both BCC and FCC phases in the CoCuNiTiAl_0.6 alloy are relatively larger, which is due to the fact that the atomic radius of Al is significantly larger than that of Mn. The densities of both phases in CoCuNiTiAl_0.6 alloy are lower, which is because the density of Al (2.702 g/cm³) is significantly smaller than that of Mn (7.300 g/cm³, see Table 2), that is, the “cocktail effect” of high entropy alloy. Table 3 also shows that the CoCuNiTiMn_0.6 alloy consists of FCC main phase and BCC phase, while the CoCuNiTiAl_0.6 alloy consists of BCC main phase and FCC phase. This is in line with reports in the literature that an increase in the proportion of Al leads to an increase in the BCC phase content in high-entropy alloys [27-29].

The formation of FCC and BCC dual-phase structures in CoCuNiTiX_0.6 (X = Mn, Al) alloys originate from the high configurational entropy effect. According to the Gibbs’ free energy ($\mathrm{\Delta }{G}_{mix}$), entropy of mixing ($\mathrm{\Delta }{S}_{mix}$), enthalpy of mixing ($\mathrm{\Delta }{H}_{mix}$), atomic radius difference ($\delta$), VEC, phase formation parameters ($\mathrm{\Omega }$) and $\mathrm{\Lambda }$ [31-35]:

where $T$ is the absolute temperature (298.15 K at room temperature). $R$ is the gas constant (8.314 J·mol^-1·K^-1). $N$ represents the number of components. $c_i$ is the molar fraction of the component $i$, and $\sum _{i=1}^N{c}_i=1$. ${\mathrm{\Omega }}_{ij}$ is the interaction parameter between components $i$ and $j$ of the regular solution. $\mathrm{\Delta }{H}_{AB}^{mix}$ is the A-B binary mixing enthalpy calculated using the Miedema model. The mixing enthalpy values between any two alloying elements in the CoCuNiTiX_0.6 (X = Mn, Al) alloy in Table 4 were obtained from the literature [30]. The atomic radius difference in Table 4 was calculated using Equation (7). The $\mathrm{\Delta }{S}_{mix}$, $\mathrm{\Delta }{H}_{mix}$, $\mathrm{\Delta }{G}_{mix}$, $\delta$, VEC, $\mathrm{\Omega }$, and $\mathrm{\Lambda }$ for CoCuNiTiX_0.6 (X = Mn, Al) HEACs in Table 5 were calculated using Eqs. (4-11). Guo et al. [31,32] pointed out that when –22 kJ/mol ≤ $\mathrm{\Delta }{H}_{mix}$ ≤ 7 kJ/mol, 0 ≤ $\delta$ ≤ 8.5, and 11 J/K/mol ≤ $\mathrm{\Delta }{S}_{mix}$ ≤19.5 J/K/mol, it is easy to form HEA phase. The enthalpy of mixing, atomic radius difference, and entropy of mixing of CoCuNiTiMn_0.6 and CoCuNiTiAl_0.6 alloys are –13.65 kJ/mol, 6.40, 13.24 J/K/mol and –19.89 kJ/mol, 6.87, 13.24 J/K/mol, respectively, indicating that both alloys readily form HEA phases. Yang et al. [33] proposed the $\mathrm{\Omega }$ criterion for multi-principal element alloys. When $\mathrm{\Omega }\ge$ 1.1 and $\delta \le$ 6.6 %, the alloy tends to form a simple SSP. The $\mathrm{\Omega }$ and $\delta$ values of the CoCuNiTiMn_0.6 alloy are 1.61 and 6.40 %, respectively. This indicates that the CoCuNiTiMn_0.6 HEAC easily forms a simple SSP structure. Zhang et al. [34] extended the theory of $\mathrm{\Omega }$ parameter criterion, and pointed out that when $\mathrm{\Omega }$ > 1.0, the alloy is easy to form a solid solution. Intermetallic compounds are easily formed when $\mathrm{\Omega }$ < 1.0. From Table 5, the $\mathrm{\Omega }$ parameter of CoCuNiTiAl_0.6 (1.03) alloy is within the solid solution formation range. Guo et al. [35] also pointed out that, when 6.87 ≤ VEC ≤ 8.0, a dual-phase structure readily forms FCC and BCC, and the larger the VEC value, the more conducive to the formation of FCC phase. As shown in Table 5, the VEC of the CoCuNiTiAl_0.6 alloy is 7.78, which easily formed FCC and BCC dual-phase structures. The VEC of the CoCuNiTiMn_0.6 alloy is 8.30, indicating that the FCC phase in the CoCuNiTiMn_0.6 alloy is easier to form than that in the CoCuNiTiAl_0.6 alloy. This is in line with XRD analysis results obtained in this study. Singh et al. [32] proposed the geometric parameter $\mathrm{\Lambda }$ when studying the formation law of solid solution phase in multi-principal alloy. When 0.24 < $\mathrm{\Lambda }$ < 0.96, the multi-principal alloy is easy to form a dual-phase structure. When $\mathrm{\Lambda }$ > 0.96, a single SSP structure is easily formed. From Table 5, the $\mathrm{\Lambda }$ values of the CoCuNiTiMn_0.6 and CoCuNiTiAl_0.6 alloys are 0.32 and 0.28, respectively, which further proves that these two alloys easily form a dual-phase structure.

3.2. Microstructure analysis

Fig. 3 and Fig. 4 show the OM photos and SEM images of CoCuNiTiX_0.6 (X = Mn, Al) HEACs, respectively. As shown in Figs. 3 and 4, the microstructure of CoCuNiTiX_0.6 (X = Mn, Al) exhibits a typical dendritic structure. As shown in Fig.3, there are thin and approximately equal-width small zones on the side of the cladding layer of the two alloys close to the substrate. This is because the cooling rate and temperature gradient near the substrate are the largest, and the fine grain zone is rapidly formed on the surface of the substrate [36, 37]. The columnar crystal zone grows into the liquid metal next to the fine grain zone on the surface. Owing to the maximum temperature gradient perpendicular to the surface of the substrate, columnar crystals with unfavorable orientations are eliminated by other solid phases during the growth process. During the competition for preferential growth, columnar crystals rapidly grow in the direction opposite to the temperature gradient, resulting in a microstructure nearly perpendicular to the fusion line. The length and width of columnar crystal growth of CoCuNiTiMn_0.6 are significantly larger than those of CoCuNiTiAl_0.6. This is because the growth of the columnar crystal zone ends with the formation of an internal equiaxed crystal region. According to Table 2, the melting point of all elements in the CoCuNiTiMn_0.6 alloy is between 1660 °C and 1083 °C, and the temperature range is less than 600 °C. During the preparation of alloy coatings by laser cladding, as the heat source leaves, the smaller volume of the molten pool rapidly solidifies, resulting in high viscosity, poor fluidity, slow atomic diffusion, and long diffusion distance of the CoCuNiTiMn_0.6 alloy melt. It takes a long time to achieve heterogeneous nucleation and generate internal equiaxed crystals through “composition undercooling” inside the melt, thereby forming a coarse columnar crystal structure. Owing to the high viscosity and poor fluidity, the volume shrinkage caused by thermal expansion and cold contraction at the end of solidification is difficult to be fed, so it is easy to form defects such as pores (see Fig. 3(a)). In the CoCuNiTiAl_0.6 alloy, Al with the lowest melting point can significantly reduce the viscosity inside the melt and enhance the fluidity of the liquid metal due to its relatively late precipitation time, which is conducive to improving the atomic diffusion rate inside the melt and reducing the diffusion time and resistance. Simultaneously, this is also beneficial for the forced convection of atomic clusters with different compositions and structures (such as density flow) to reduce the temperature gradient inside the melt, which is conducive to the growth of equiaxed dendrites, and the welding defects caused by volume shrinkage. This comprehensive effect promotes the formation of a dendritic structure composed of a narrow columnar crystal zone and a wide equiaxed crystal zone in the CoCuNiTiAl_0.6 alloy.

Fig. 3OM photos of CoCuNiTiX0.6 (X = Mn, Al) HEACs

a) CoCuNiTiMn_0.6

b) CoCuNiTiAl_0.6

Fig. 4SEM images of CoCuNiTiX0.6 (X= Mn, Al) HEACs

a) CoCuNiTiMn_0.6

b) CoCuNiTiAl_0.6

As shown in Fig. 4, CoCuNiTiX_0.6 (X = Mn, Al) HEACs are composed of dendrite (DR) regions embedded with primary phases (PP) and interdendrite (ID) regions. To better understand the distribution of various alloying elements in different regions, the segregation factor $K$ and the average segregation factor $K_a$ were introduced after deducting the Fe element introduced by the “excavation effect” of the substrate during the laser cladding process [38]:

where $\xi$ represents the directional coefficient. When the alloying element is enriched in the dendrite region, $\xi =$ 1; when the alloying element is enriched in the interdendrite region, $\xi =$ –1. $C_{ID}$ and $C_{DR}$ represent the concentrations of the same alloying element in ID and DR regions, respectively. $K_i$ represents the element segregation coefficient of the $i$-th component. The larger the absolute value of $K$, the larger the difference of the same alloying element in DR and ID regions. $K\ge$0 and $K<$0 indicate the enrichment of alloying elements in DR and ID regions, respectively. The larger the $K_a$ value, the greater the overall segregation degree of each alloying element in the alloy system. Table 6 shows the EDS results for CoCuNiTiX_0.6 (X = Mn, Al) HEACs. The $K$ value in Table 6 was calculated using Eq. (12). Using Eq. (13), the average segregation coefficients of CoCuNiTiMn_0.6 and CoCuNiTiAl_0.6 alloys are 23.39 % and 10.64 %, respectively, indicating that the average segregation degree of various alloying elements in Al-containing alloys is lower than that of Mn-containing alloys. This is related to the previous analysis that Al addition can improve the alloy fluidity, reduce the melt viscosity, and decrease the diffusion distance and resistance of the alloy. In addition, as shown in Table 6, the Ti content in the primary phase of the two alloys is the highest, mainly because Ti has the highest melting point (see Table 2). During the cooling of the melt, Ti preferentially reaches and meets the required undercooling degree for precipitation, thus forming a Ti-rich primary phase. As the melt temperature further decreases, other alloying elements also gradually reach the precipitation temperature. All other alloying elements exhibit the maximum negative mixing enthalpy with Ti, so the atomic clusters enriched with these elements in the melt tightly bind with Ti. To reduce nucleation resistance, later-precipitated solid phases adhere to the rough surface of Ti-rich primary phase for non-uniform nucleation, thereby forming DR regions inlaid with primary phases. Consequently, Ti concentrations are relatively high in DR regions. This is also consistent with the results in Table 6, which show that both alloys have positive segregation coefficients for Ti. According to Table 6, both alloys exhibit negative segregation coefficients for Cu, indicating its tendency to enrich in ID regions. This is mainly because Cu has the smallest negative mixing enthalpy or the largest positive mixing enthalpy with other alloying elements and is easily squeezed into the edge region of the atomic cluster in the melt. In addition, owing to its low melting point and subsequent precipitation, it is easy to enrich in the interdendrite region during the latest solidification.

3.3. Corrosion performance

Fig.5 shows the potentiodynamic polarization curves of the CoCuNiTiX_0.6 (X = Mn, Al) HEACs in 3.5 wt% NaCl solution. From Fig.5, CoCuNiTiMn_0.6 and CoCuNiTiAl_0.6 HEACs showed an obvious passivation phenomenon. During the polarization process, the formation of a passivation film helps to resist the corrosion of the NaCl solution. The self-corrosion current density ($i_{corr}$) and potential ($E_{corr}$), anode (${\beta }_a$) and cathode (${\beta }_c$) Tafel slope of both alloys can be obtained by extrapolation, and the fitting results are listed in Table 7. The polarization resistance $R_p$ and annual corrosion rate $K_{corr}$ in Table 7 can be expressed as [39]:

where $q$ is approximately 3272 mm/(A·cm·year), EW is the equivalent weight of the sample ($g$), $\rho$ is the sample density (g·cm^-3), $f_i$ is the mass fraction of component $i$, $n_i$ is the exchange electron number of component $i$ (mol^-1), and $A_i$ is the atomic weight of component $i$ (g/mol). According to Table 7, compared with CoCuNiTiAl_0.6 HEAC, CoCuNiTiMn_0.6 HEAC has a higher corrosion potential and a lower self-corrosion current density, indicating a lower corrosion tendency of the alloy. From Eq. (15), the annual corrosion rate is positively correlated with $i_{corr}$. Therefore, the annual corrosion rate of CoCuNiTiMn_0.6 HEAC is also relatively low. However, a large polarization amplitude in the polarization test may destroy the passivation film, resulting in the distortion of $i_{corr}$ measured using the polarization curve. Slightly disturbed electrochemical impedance spectroscopy (EIS) better reflects the true protective performance of the film layer. This phenomenon is consistent with literature [40, 41].

Fig. 5Potentiodynamic polarization curves of CoCuNiTiX0.6 (X = Mn, Al) HEACs in 3.5 wt% NaCl solution

Fig. 6 shows the electrochemical impedance profiles of CoCuNiTiX_0.6 (X = Mn, Al) HEACs in 3.5 wt% NaCl solution. From Fig. 6(a), the Nyquist plots of the CoCuNiTiMn_0.6 and CoCuNiTiAl_0.6 HEACs show a flattened semicircular arc. This indicates that charge transfer occurs between the coating surface and the solution, and the electric field distribution of the electric double layer formed by this charge separation is not uniform. The capacitive arc radius of Al-containing alloy is obviously larger than that of Mn-containing alloy, indicating that Al-containing alloys have stronger corrosion resistance. According to Fig. 6(b), the maximum phase angles of CoCuNiTiMn_0.6 and CoCuNiTiAl_0.6 alloys are –54.95° and –56.91°, respectively. Under similar corrosion systems and equivalent circuit models, the larger the absolute value of the maximum phase angle, the closer the electrode performance to an ideal and dense capacitance interface. The constant phase angle element CPE1, phase shift $n$, and charge transfer resistance $R_{ct}$ in Table 8 were fitted using the equivalent circuit illustrated in Fig. 6(a). The fitting results are marked in Fig. 6 using a combination of dotted lines and hollow symbols. The data obtained from the experiments are marked by a combination of solid lines and symbols. The fitting curve (FC) was close to the experimental curve (EC). The impedance $Z_{CPE}$ can be expressed as [42]:

where $Y_0$ is the scaling factor, $j$ is the imaginary unit, and $\omega$ is the angular frequency, and $m$ represents the degree of capacitance dispersion. The smaller the value of $\left|m-1\right|$, the less influence the porous structure exerts [43]. The $m$ values of CoCuNiTiMn_0.6 and CoCuNiTiAl_0.6 HEACs are 0.7318 and 0.7544, respectively. The $m$ value of CoCuNiTiAl_0.6 is closer to 1, indicating that its passivation film is denser and the surface uniformity is better than that of the Mn-containing alloy, which is in line with SEM and EDS results. The formation of pores increases the surface area exposed to corrosion and distorts the double layer, providing preferential diffusion pathways for Cl⁻ ions, which causes the CPE behavior to deviate significantly from the ideal capacitance behavior [39, 44]. The $R_{ct}$ of CoCuNiTiAl_0.6 is 2028.00 Ω·cm², which is approximately 2.49 times that of CoCuNiTiMn_0.6 (815.70 Ω·cm²), indicating that its charge transfer resistance is greater. According to the theory of the impedance modulus of a passivation alloy or coating system [45, 46], at sufficiently low frequencies (such as 0.1 Hz), the response of electrochemical impedance spectroscopy mainly reflects the charge transfer process occurring at the corrosion interface or the capacitive behavior of the passivation film, and its corresponding values are related to the charge transfer resistance or the passivation film impedance. In the passivation system, the impedance modulus $\left|Z\right|$ at 0.1 Hz can be used as an approximate index to evaluate the corrosion resistance of the coatings. The $\left|Z\right|$ of CoCuNiTiAl_0.6 HEAC at 0.1 Hz is 1698.35 Ω·cm², which is approximately 2.53 times that of the Mn-containing alloy (670.34 Ω·cm²) and similar to the $R_{ct}$ ratio. This further confirms that porosity defects significantly reduce the overall protective performance and CoCuNiTiAl_0.6 HEAC exhibits better corrosion resistance. The corrosion resistance of CoCuNiTiAl_0.6 HEAC is relatively good owing to its dense passive film, few defects, and relatively uniform composition distribution.

Fig. 6Electrochemical impedance profiles of CoCuNiTiX0.6 (X = Mn, Al) HEACs in 3.5 wt% NaCl solution

a) Nyquist plots

b) Bode plots