3.1. Microstructural Characterization
Figure 3 presents the XRD diffraction patterns of the coatings fabricated at different scanning speeds. The main peaks of Coating 1 could clearly be observed at 2θ = 36.123°, 2θ = 39.087°, 2θ = 41.560°, 2θ = 42.460°, 2θ = 44.482°, 2θ = 61.141°, and 2θ = 77.745°. Comparison of the d values of these peaks with those in the corresponding JCPDS cards confirmed that TiNi and Ti
2Ni were formed as the matrix phase, while TiC and TiB
2 were formed as the reinforcement phase during laser cladding. The four compounds are rich in Ti, which indicates that this element plays a dominant role in the resultant phases of the coating. The coating obtained at a scanning speed of 11 mm∙s
−1 revealed an XRD diffraction pattern showing significant changes compared with that obtained for Coating 1; this finding implies that the phase constituents of these coatings have obvious differences. The indexed result shows that the matrix phases of Coating 2 are γ(Ni) and Ni
3Ti, which are significantly different from those formed in Coating 1. TiB
2 and TiC were still observed, along with four new peaks related to Cr
7C
3 at 2θ = 41.886°, 2θ = 44.242°, 2θ = 49.872°, and 2θ = 81.099°. Ni and Cr as the main elements in Coating 2 greatly influenced the phase constituents of the coating ad weakened the role of Ti. When the scanning speed was increased to 17 mm∙s
−1, the XRD diffraction pattern obtained demonstrated slight changes relative to that of Coating 2. Specifically, an increase in the intensity of the main peaks associated with γ(Ni) and Cr
7C
3 and the appearance of some new peaks associated with CrB were found. This result implies that the cladding material plays a significant role in the phase constituents of the coating at high scanning speeds. The changes observed in the XRD patterns may be mainly attributed to increases in scanning speed, which causes great reductions in the dilution ratio (D) of the coatings. The D of all coatings could be calculated by the following formula:
where
refers to the penetration depth of the melt pool below the substrate and
represents the thickness of the melt pool above the original substrate level.
The calculated D values show a downward trend with increasing scanning speed (Coating 1, 64.7%; Coating 2, 53.9%; Coating 3, 37.3%). Rahman Rashid et al. [
45] investigated the effects of specific energy (specific energy = laser power/(scanning speed × beam diameter)) on D. The results indicated that the D of the laser-clad coatings is closely related to the specific energy; specifically, the average D of the coatings gradually increased with decreasing specific energy. This finding is contrary to the results obtained in the present research. This difference in results may be attributed to differences in the powder feeding method between the two studies. In the present study, powder layers with a constant thickness were placed on Ti6Al4V prior to laser cladding, which means the amount of molten powder was constant in all zones during laser cladding. That is to say, the specific energy could be regarded as the only essential factor responsible for D. However, the study reported in [
46] delivered the powder by a coaxial nozzle during laser cladding, which means the amount of molten powder may vary at different specific energies. In addition to the specific energy, the molten powder is another factor strongly affecting D. Under the given feeding rate, beam diameter, and laser power, increases in scanning speed cause a reduction in specific energy and, in turn, a reduction in D because of the reduction in
. However, the amount of molten powder correspondingly decreased, thereby causing an increase in D owing to the reduction in
. When the role of the latter suppresses that of the former, D may present an upward tendency with increasing scanning speed, as reported in Ref. [
46], which contradicts the findings in the present study.
The volume fractions of different phases in the three coatings were quantitatively analyzed by Highscore Plus software (
Table 1). The change in phase constituents of the coatings with the scanning speed may be expected to result in the corresponding changes in mechanical properties and corrosion resistance, as discussed below.
The microstructures of the three coatings were observed under low magnification (
Figure 4). As depicted in
Figure 4a, some coarse black blocky particles with an average size of approximately 8 μm × 10 μm are uniformly distributed in the gray matrix of Coating 1. The gray matrix is composed of two distinct phases. Clear inspection reveals that a small number of fine dendrites adhere to the surfaces of blocky particles. As the scanning speed increased to 11 mm∙s
−1, the microstructure of Coating 2 (
Figure 4b) presented a slight change relative to that of Coating 1; here, a large number of light-gray honeycomb-like structures, as well as some separated dark-gray phases, were observed. The microstructure of Coating 3 (
Figure 4c) demonstrated significant differences compared with the first two coatings with further increases in the scanning speed. In this case, the blocky particles presented a downward trend in number as well as the synthesis of strip-like particles.
High-magnification SEM images and the corresponding EDS analyses were obtained to identify the categories and morphologies of the different phases of the coatings (
Figure 5,
Table 2). Coating 1 shows four phases (
Figure 5(a1,a2)), which correspond to the protrusion phase (marked A), the depression phase (marked B), coarse black blocky particles (marked C), and fine dendrites (marked D). The EDS analysis results indicate that Zones A and B are rich in Ti and Ni; small amounts of Al, Si, and Cr could also be detected. Combining these findings with the XRD results, the two phases may be concluded to be two secondary solid solutions with Ti-N compounds as the solvent and Al, Si, and Cr as the solute. Cr and Al are closer to Ti and Si is closer to Ni in terms of atomic radius and electronegativity (atomic radius: 1.85, 1.82, and 2.00 Å for Cr, Al, and Ti, respectively; 1.46 and 1.62 Å for Si and Ni, respectively; electronegativity: 1.66, 1.61, and 1.54 for Cr, Al, and Ti, respectively; 1.90 and 1.92 for Si and Ni, respectively). Therefore, Cr and Al atoms could substitute Ti atoms, while Si atoms could displace Ni atoms in the lattices of the Ti-N compounds. The calculation results demonstrate that the atomic ratios of Ti + Cr + Al and Ni + Si in Zones A and B are 2.1:1 and 0.9:1, respectively. The two phases could be identified as Ti
2Ni and TiNi secondary solid solutions. Zone C is rich in Ti and B, and Zone D mainly consists of Ti and C. The black blocky phase and fine dendrites are likely TiB
2 and TiC. These results correspond well with those obtained from the XRD analyses. In addition to blocky TiB
2 particles and fine TiC dendrites, a new coarse strip-like phase (marked G) was synthesized in Coating 2 (
Figure 5(b1,b2)). The reinforcement is rich in C and Cr and was determined to be Cr
7C
3 via the XRD results. In addition to the equiaxed phase (marked E), a honeycomb-like structure (marked F) was also synthesized in the matrix; this structure was not observed in Coating 1. A magnified backscattered electron (BSE) image of Zone F indicates that the honeycomb-like structure is mainly composed of light-gray reticular phases, among which a small number of dark-gray phases are located. The dark-gray and equiaxed phases have approximately the same atomic number contrast, which implies that they belong to a single phase. Zone E contains a large amount of Ni of approximately 60 at.%, which is slightly higher than that in the honeycomb-like structure (approximately 54 at.%). Combining the XRD results with the Ti-Ni binary phase diagram, the equiaxed phase may be concluded to be primary γ(Ni), and the honeycomb-like structure can be identified as the eutecticum of γ(Ni) + Ni
3Ti. When the scanning speed is increased to 17 mm∙s
−1, the volume fraction of γ(Ni) gradually increases, and the volume fraction of the eutecticum correspondingly decreases. This change may be mainly associated with the increase in Ni content in the coatings due to weaker dilution at higher scanning speeds. Other fine strip-like particles (marked H) rich in Cr and B appeared to be in situ synthesized in Coating 3 (
Figure 5(c1,c2)); these particles were identified as CrB by XRD. The area fractions of these phases are shown in
Table 3. It can be seen that the results are approximately similar to those shown in
Table 1. The slight changes should be attributed to the differences in the used unit; namely, the used units in
Table 1 and
Table 3 are the volume fraction and the area fraction, respectively.
3.3. Corrosion Performance
Figure 11 shows the potentiodynamic polarization curves of the coatings and substrate in 3.5 wt.% NaCl solution. Some parameters (
Table 5), such as corrosion potential (E
corr), corrosion current density (i
corr), and current density (i
s) in the active and comparatively stable corrosion state can be acquired from
Figure 11.
Among the coatings obtained, Coating 2 possesses the most positive E
corr. E
corr refers to the potential at which the material reaches a stable corrosion state; this parameter can be applied to characterize the tendency of a material to corrode in a corrosion system. A material with low E
corr is more inclined to act as a cathode in a galvanic cell without an applied voltage and suffer from serious corrosion by losing electrons when compared with that with a high E
corr. Therefore, in terms of E
corr, Coating 2 demonstrates the best corrosion resistance among the three coatings. Coating 2 also demonstrates the lowest i
corr among the three coatings. i
corr, which refers to the corrosion current density at E
corr, is used to characterize the corrosion rate of a material. Therefore, in terms of i
corr, the corrosion resistance of Coating 2 is also superior to those of Coatings 1 and 3. Changes in E
corr and i
corr may be attributed to differences in the composition and phases of the samples. When the applied potential exceeds E
corr, the electrode surface enters the active state because i
s rapidly increases as the potential increases. These phenomena promote oxidation reactions on the electrode surface, which releases metals in their ionic form to the electrolyte. As the potential further increases, i
s increases slowly and finally reaches a comparatively stable value, thus indicating that the electrode enters a comparatively stable corrosion state. The transformation from active state to comparatively stable corrosion state is closely related to the formation of a thin and dense oxidation film formed on the electrode surface, which isolates the material from the electrode and greatly retards the release rate of metal ions. The formation preference of the passive film can be evaluated by assessing differences between E
corr and the critical potential (E
a-s) as the electrode transforms from the active state to the comparatively stable corrosion state. As shown in
Table 5, the E
a-s − E
corr values of the coatings show the trend Coating 3 > Coating 1 > Coating 2, thus indicating that the oxidation film is more easily formed on the surface of Coating 2 than on the surfaces of the two other coatings. The corrosion rate in the comparatively stable corrosion state can be described by i
s, which is the average value of all current densities obtained in Platform 1, as shown in
Figure 11. The results of i
s demonstrated the following order: Coating 2 (3.664 × 10
−7 A∙cm
−2) ≈ Coating 1 (3.500 × 10
−7 A∙cm
−2) < Coating 3 (7.345 × 10
−5 A∙cm
−2). These results indicate that Coating 2 possesses the lowest corrosion rate in the comparatively stable corrosion state. When the potential exceeds another critical potential (E
s-a), i
s is sharply enhanced with increasing potential, thereby indicating that the electrode enters the active state once more because a portion of the oxidation film suffers from destruction. Coating 2 clearly demonstrates a higher E
s-a (0.313 V) compared with those of Coating 1 (0.050 V) and Coating 3 (0.290 V), which means that the oxidation film formed on Coating 2 is more stable than those formed on the other coatings. Further increases in potential clearly identify a third critical potential (E’
s-a), beyond which the electrode enters a second comparatively stable corrosion state because of the repair of the damaged oxidation film. The i
s of the three coatings in this state is similar (approximately 1.093 × 10
−2 A∙cm
−2), which corresponds to the average value in Platform 2. Thus, according to the above analyses, Coating 2 exhibits greater corrosion resistance than Coating 1 or Coating 3.
XPS was applied to detect the elemental composition and corresponding valences of the oxidation film formed on Coating 2 to reveal the passive mechanism of the electrode, as shown in
Figure 12.
Figure 12a shows the XPS survey spectrum of the oxidation film formed on Coating 2. The oxidation film is composed of Ni, Al, Ti, Cr, Si, and O. Then, the high-resolution narrow spectra of these metal elements were obtained to detect their valence states. Four strong peaks at 855.57, 873.21, 861.09, and 879.72 eV were clearly detected in the Ni
2p spectrum (
Figure 12b), which fit the standard peaks of Ni in NiO well. The strong peak at 74.49 eV in the Al
2p spectrum (
Figure 12c) confirms the existence of Al
2O
3. Two strong peaks at 458.43 eV and 464.63 eV in the Ti
2p spectrum (
Figure 12d) prove the presence of TiO
2 in the oxidation film. Two obvious peaks at 576.57 and 585.93 eV, which are highly consistent with the binding energies of Cr in Cr
2O
3, could be detected in the Cr
2p spectrum (
Figure 12e). The clear peak with a binding energy of 101.95 eV in
Figure 12f confirms the existence of SiO
2. According to the above analyses, the oxidation film may be concluded to be mainly composed of NiO, Al
2O
3, TiO
2, Cr
2O
3, and SiO
2. The molar fraction of oxides in the coatings can be calculated from the XPS results (
Table 6). As the scanning speed increased, the contents of TiO
2 and Al
2O
3 demonstrated a gradual downward trend, whereas the contents of NiO, Cr
2O
3, and SiO
2 revealed an upward trend due to reductions in the dilution rate of the coatings.
Differences in the formation preference of the oxidation film among the samples may be reasonably explained by thermodynamic calculations. According to the XPS results, the following reactions occur during oxidation:
The change in the standard Gibbs free energy (ΔGθ) at room temperature (approximately 298 K) can be calculated as −889.53 KJ∙mol−1 for Reaction (4), −3163.88 KJ∙mol−1 for Reaction (5), −424.90 KJ∙mol−1 for Reaction (6), −2096.35 KJ∙mol−1 for Reaction (7), and −856.50 KJ∙mol−1 for Reaction (8). As the ΔGθ of all reactions is negative, these reactions spontaneously occur. The reaction order in terms of thermodynamics is Reaction (5) > Reaction (7) > Reaction (4) > Reaction (8) > Reaction (6). Given the calculated results of ΔGθ, determining which sample the oxidation film preferentially forms on is very difficult.
The compactness of the formed oxidation film is also responsible for the formation preference of the oxidation film, which can be characterized via the Pilling–Bedworth ratio (PBR). The PBR reflects the change in volume of a given metal element (
X) subject to oxidation (
AX +
BO2 =
CXmOn) and could be calculated by the following formula [
45]:
where
A and
C refer to the molar numbers of the reactant (
X) and product (
XmOn), respectively,
and
signify the atomic weight of the reactant (
X) and the molecular weight of the product (
XmOn), respectively, and
and
represent the density of the reactant (
X) and product (
XmOn), respectively.
The PBR of the above reactions could be calculated as 1.77 for Reaction (4), 1.46 for Reaction (5), 1.66 for Reaction (6), 2.02 for Reaction (7), and 2.27 for Reaction (8). The PBRs of the above reactions are all greater than 1, which indicates that the metal subject to oxidation undergoes volume expansion, which could improve the compactness of the oxidation film and result in better blocking effects between the sample and electrolyte. However, when the PBR is higher than 2, high levels of stress may be produced and cause cracking of the oxidation film. The electrolyte may still be in contact with the sample via the fracture areas and cause the active dissolution of these areas, thus retarding the change process from the active state to the comparatively stable corrosion state. Coatings 1 and 2 can enter into the comparatively stable corrosion state more rapidly than Coating 3 because of their low Ea-s − Ecorr values, which could be attributed to their suitable contents of Cr and Si with a high PBR. As the contents of Cr and Si are further enhanced in Coating 3, some cracks may be generated and delay the change process from the active state to the comparatively stable corrosion state.