3.1. Subsection Microstructural Characterization
Figure 3 presents the metallographic morphology along the cross-section of HSC specimens. In the unmodified coating, the long-range dendritic solidification microstructure growing along certain angles can be clearly observed, as well as several larger defect points, as shown in
Figure 3a,b [
27]. However, long-term dendritic solidifications are reduced and equiaxed grains are enlarged in the specimen modified with 0.4 wt.% YNPs, while the concentration of internal defects is significantly reduced, as shown in
Figure 3c,d. The RE element can improve the undercooling point during the solidification of alloy system, thereby reducing the metastable eutectic transition temperature [
22,
28]. During the rapid heating process, dispersed Y-ions partially dissolve from part of YNPs with low solubility, segregating at the front of primary dendrites with low electronegativity, reducing the undercooling degree of the steel system, and providing an enhanced nucleation rate [
29]. In addition, YNPs improve the fluidity of melt in the molten pool, resulting in multi-directional heat flow at all locations in the molten pool. The dendrite phase is reduced as dendrites dissociated into smaller species by undercooling effect, resulting in grain refinement and multi-directional crystallization. Consequently, long-term dendritic solidifications are transformed into equiaxed grains and short-term columnar crystals, resulting in a dense microstructure and less defects within the modified coating. However, large pores or defects are formed in coating with 0.8 wt.% YNPs, compromising the coating performance, as shown in
Figure 3e,f. In the PTA process, due to the technical characteristics of rapid heating and cooling, large columnar crystals grow vertically from the bottom of the molten pool to the center of the coating, while the addition of YNPs will induce the dismember of dendritic crystals, so as to refine the grains and enhance the properties of the coating [
23].
The XRD patterns of alloy powder and coating sample are shown in
Figure 4.
Figure 4a shows that the main phases of all alloy powders are α-Fe phase (PDF#34-0529), Fe-Cr phase (PDF#85-1410), and Taenite phase (PDF#47-1417). The decrease in γ-Fe peak was due to the rapid cooling process during solidification [
20,
30].
Figure 4b shows that the main phases of all coatings are α-Fe phase (PDF#34-0529), Fe-Cr phase (PDF#85-1410), and Cr phase (PDF#88-2323), corresponding to the diffraction peaks at 44.5°, 64.5°, and 43.5°, respectively. Notably, for Cr, it exhibits an obvious enrichment in the HSC0 coating samples [
31]. However, the Cr phase in HSC02, HSC04, HSC06, and HSC08 samples decreased significantly with the addition of Y
2O
3. The addition of Y
2O
3 nanoparticles provides sufficient undercooling [
29] for PTA coating in the rapid solidification process, which promotes the diffusion rate of Cr and reduces non-equilibrium segregations [
22,
31]. Several weak XRD peaks are shown in HSC08 specimens, corresponding to the Y
2O
3 phase (PDF#65-3178) with a crystal plane of (222), (125), and (440), which peaked at 29.2°, 46.9, and 48.5°. It suggested that the extra Y
2O
3 nanoparticles will exist in the steel matrix to form inclusion, and its melting point is higher than that of the steel matrix.
To further understand the influence of the Y element on 14CrSiMnV alloy coating, the EDS maps are measured to analyze the elemental composition of the secondary phase particles. Sphere-like secondary phase particles in HSC0 sample (
Figure 5a) mainly consist of Si and O element. However,
Figure 5b,c shows the secondary phase particles in the specimens modified with 0.4 wt.% and 0.8 wt.% YNPs, confirming the presence of Y, Mn, Si, and O. Hence, Y element segregated into Mn-oxides, followed by Y element playing a role in purifying grains and grain boundaries, which is consistent with XRD data. Furthermore, with the addition of YNPs, 18.9–23.0 wt.% of Y, 2.4–5.4 wt.% of Mn, and 18.2–24.5 wt.% of O, EDS was detected in the secondary phase particles, as shown in
Table 3. This is due to the active chemical properties of the Y element, which has a strong affinity with O, Si, and Mn. The ICP mass spectrometer and OND determinate are measured to analyze the Y and O content of the 14CrSiMnV alloy coating, and it was found that the Y and O contents in the coating increased with the increase of YNP content (
Table 4).
In the conventional Y
2O
3 addition mechanism, the compounds formed with dissolved Y ions as a slag float on the melt surface of PTA coating, playing the role of slag removal, tissue purification, porosity, and crack elimination [
11,
12,
32]. On the other hand, abundant Y
2O
3 can provide conditions for the formation of Y oxides and impurity compounds to aggregate and grow up, leading to a decrease in the number of effective crystal nuclei, resulting in coarsening of the tissue and an increase in the number of inclusions. However, instead of the slag in modified coating, dissolved Y forms compound with Mn, Si, and O as secondary phase particles, while Y (0.180 nm), Mn (0.127 nm), and Si (0.146 nm) are suited to generate a substitutional solid solution during the rapid solidification process, which is relatively similar to the solid solution of Fe phase (0.172 nm), and provides abundant nucleation sites [
33,
34].
3.2. Mechanical Characterization
Figure 6 exhibits the tensile performance of HSC specimens. It is revealed that all test specimens display linear elastic strain curves. The tensile strength of the alloy coating specimen is 667 MPa (HSC0), 810 MPa (HSC02), 1281 MPa (HSC04), 1066 MPa (HSC06), and 946 MPa (HSC08), respectively. The elongation of the alloy coating specimens is 2.1% (HSC0), 2.8% (HSC02), 3.2% (HSC04), 3.0% (HSC06), and 2.9% (HSC08), respectively, as shown in
Table 5. The tensile strength of the modified specimen (HSC04) increased by 92.0%, and the toughness of the coating increased by 55.6%. The YNP-modified alloy coating possesses higher tensile strength and elongation than the unmodified coating. One should note that grain refinement improves the tensile strength and ductility of the alloy coatings. During the PTA forming process, the Y element is combined with Mn, Si, and O elements to form new Y-Mn-Si-O compounds. These compounds are dispersed in the coating and become hetero-nucleation sites for grain nucleation, promoting grain nucleation and resulting in increased concentration equiaxed grains and short-term columnar crystals. The dispersion of the secondary phase particles in the coating results in a stress field, which hinders the deformation and enhances the mechanical properties of the modified coating.
Figure 7 shows the fracture morphology of the tensile specimens.
Figure 7a displays a long-range river-like pattern and fan-sliding morphology in the cross-section of HSC0 fracture sample. On the other hand, a mass of cleavage planes and step features, with secondary micro-cracks and pores, are observed on the surface of HSC0 sample (
Figure 7b), confirming the occurrence of cleavage and brittle fracture. It is revealed that the cleavage fracture occurred along the growth direction of long-range dendritic solidification microstructures, because cleavage and slip easily occur in microstructurally uneven regions of BCC-phase steel under loading, leading to linear elastic strain curves and cleavage fracture with low tensile properties [
22].
Moreover, short-range river patterns, small cleavage steps, and dimples in the sliding direction were observed in the modified coating, as shown in the fracture morphology of HSC04 (
Figure 7c,d) and HSC08 (
Figure 7e,f) specimens, confirming the failure mechanism was based on hybrid-type with quasi-cleavage, brittle fracture, and granular ductile fracture [
35]. On the one hand, extension sliding cracks were disassembled and blocked by short-range dendritic solidification and enlarged equiaxed grain zone. It is relatively easy to hinder the extension of cleavage steps during the tensile process due to the difference between secondary phase particles and coating in the elastic–plastic region and bonding ability, resulting in small cleavage steps under the influence of tensile stress. As a result, the brittle capacity of modified specimens is significantly improved. The addition of YNPs improves the internal grain structure of the coating, which plays a role in fine grain strengthening and second phase strengthening, while the dispersion distribution of the second phase particles hinders the expansion of slip dislocation and improves the strength of the coating [
36]. However, excessive amounts of YNPs increase O content within the Fe phase (
Table 4), resulting in large pores and micro-cracks (
Figure 7d), and thereby decreasing the tensile strength of HSC08.
Figure 8 shows the longitudinal section morphology of the tensile specimens. The crystallographic facets with the size of a few microns at longitudinal section of the tensile specimens can be observed, as shown in
Figure 8b–f. During the tensile process, the crack expands along the grain boundary and forms the crystallographic facets on the fracture surfaces [
29]. Thus, the larger the grain size, the longer the crystallographic facets. The addition of YNPs will refine the grains, hinder the slip dislocation, and enhance the tensile property of the coating [
23]. As a result, the crystallographic facets of HSC04 and HSC08 is smaller than that of HSC0.
Figure 9 shows the microhardness along the cross-section of coatings. The average microhardness of the coating specimen is 627 HV0.1(HSC0), 658 HV0.1(HSC02), 698 HV0.1(HSC04), 669 HV0.1(HSC06), and 642 HV0.1(HSC08). Moreover, the microhardness of the modified coating increased with the YNPs content. When 0.4 wt.% YNPs are added, the modified coating exhibited the highest microhardness distribution, which is 11.3% higher than the unmodified coating. During the solidification process, the Y-Mn-Si-O compound disperses in the grains of coating and produces a pinning effect, inhibiting the grain growth and playing a role in grain strengthening and secondary phase strengthening. However, with the increase of YNPs content, excessive Y increases the viscosity of molten metal in the cladding pool and deteriorates the fluidity of molten metal, forming a large number of inclusions at grain boundaries, weakening the binding force between grains, thereby reducing the mechanical strength of the coating, including Vickers microhardness.
Figure 10 shows reciprocating sliding wear surface morphology of the coatings under a low-speed dry friction sliding. It displays a fish-scale pattern on the wear surface of all coatings, indicating the plastic deformation and adhesive wear characteristics due to delamination [
37,
38]. As shown in
Figure 10a, both compact particles and deep grooves can be observed on uneven worn surfaces of the HSC0 specimen, whereas micro-plowing is the main feature of abrasive wear [
37] due to third-body interactions. Owing to the dendritic solidification along the
z-axis during the fast solidification and cooling process, a fine well-aligned structure is formed with a high-temperature gradient, leading to a large friction coefficient gap between the build direction (
z-axis) and cladding direction (
x-axis) of the PTA coating [
3]. During the initial stage of sliding wear process, abundant debris is produced along the
x-axis of coating due to brittle fracture and micro-cutting, promoting the adhesion and abrasive wear friction of unmodified coatings. On the other hand, HSC02 (
Figure 10b), HSC04 (
Figure 10c), and HSC06 (
Figure 10d) exhibited abrasive wear, originating from the uniform microstructure and low amounts of debris. It is worth noting that the dendritic crystals inside the coating transform into equiaxed crystals after the addition of an optimal amount of YNPs, resulting in a lower friction coefficient and similar sliding wear resistance in different directions and less debris [
39,
40]. This is due to the addition of YNPs, improving the grain structure, hindering the growth of columnar grains, making the inner structure of the coating close, and improving the wear resistance of the coating. However, excessive YNPs will lead to superfluous O content and inclusion enrichment at grain boundaries in the coating, and the uneven internal composition will lead to reduce wear resistance. The EDS maps were measured to analyze the elemental composition of the reciprocating wear surface, as shown in
Table 6. The results display that no new element (such as N) was introduced into the coating surface during the reciprocating friction process, while the wear resistance of the coating is improved by added small amount of YNPs.
Figure 11 plots the weight loss–wear cycle curves and relative wear resistance coefficient of different coatings. During the initial stage of friction process, amounts of abrasive and debris particles are peeled off by brittle fracture from the coating surface, resulting in cracks or scars. Then, such particles are brought into the grooves to promote micro-cutting between surrounding matrix and eutectic structure during third-body interactions, resulting in a plastic deformation phenomenon and wear loss. The wear weight loss of specimens is found to be 0.1345 g (Q235), 0.0755 g (HSC0), 0.0752 g (HSC02), 0.0587 g (HSC04), 0.0693 g (HSC06), and 0.0732 (HSC08). The wear loss of HSC04 coating is the lowest, which is 56.3% and 22.2% less than the Q235 and unmodified coating, respectively.
Meanwhile, the relative wear resistance coefficient (RWSc) of Q235 steel was set at 1, while the RWSc of modified coating was found to be 1.7788 (HSC0), 1.7859 (HSC02), 2.2827 (HSC04), 1.9380 (HSC06), and 1.8347 (HSC08). The RWSc of the HSC04 specimen is increased by 28.3%.
Figure 12a displays a worn surface with deep and coarse scratches, as well as large irregular spalling pits for unmodified coatings. In contrast, shallower friction scratches and smoother slender furrows were exhibited in the modified coating due to lesser amounts of brittle fracture and debris. It can be ascribed to a decrease of dendritic grains and enlarged equiaxed grains in the modified coatings, resulting in the transformation of the wear mechanism from plastic deformation to slight peeling.