3.1. Microstructural Characterization
Figure 1 illustrates the microstructural characteristics of the non-inoculated and TiC/TiB
2-inoculated samples.
Figure 1a–f represent the micrograph and EDS maps of the non-inoculated sample, revealing Cr-rich δ-ferrite in lathy and vermicular forms. This retention at room temperature is due to the high solidification rate of the WAAM process and the high Cr content (~13 wt.%) in the alloy, which acts as a ferrite stabilizer [
4]. The TEM image shown in
Figure 1b depicts a lath martensitic matrix with retained austenite, resulting from the presence of approximately 8 wt.% Ni, which acts as an austenite stabilizer and inhibits the complete transformation from austenite to martensite during cooling. Additionally, the EDS map in
Figure 1e identifies Al-rich oxides, likely originating from the shielding gas, ambient air, or residual moisture. For more comprehensive information on the microstructure of WAAM-fabricated PH 13-8Mo parts, refer to the authors’ recent publications [
3,
5].
Figure 1g–l present SEM micrographs and EDS maps of the TiC-reinforced sample, revealing islands of residual δ-ferrite with a different morphology compared to the lathy and vermicular forms observed in the non-inoculated sample. This difference is likely due to the presence of TiC particles acting as nucleation sites, resulting in the formation of equiaxed primary δ-ferrite grains. The EDS maps, shown in
Figure 1i–l, confirm the presence of TiC phases that are evenly distributed throughout the TiC-inoculated sample. These TiC phases can be identified as either intact preplaced TiC nano-powders or larger in situ TiC particles formed on Al
2O
3 inclusions during the solidification process. The formation of larger in situ TiC particles can be justified by partial dissolution of the preplaced powders during the deposition of each layer. Zhai et al. [
6] have also observed the partial dissolution of micron-sized TiC particles added to 316L powders during selective laser melting, leading to the formation of in situ TiC particles in the final printed component. Bahramizadeh et al. [
7] suggested that the lower Gibbs free energy of formation for Al
2O
3 particles promotes their early formation during solidification, enabling them to act as favorable sites for the heterogeneous nucleation of cubic TiC particles. Similarly, Sharifitabar et al. [
8] reported the heterogeneous nucleation of TiC on Al
2O
3 particles during gas tungsten arc cladding of 1045 steel, using a mixture of TiO
2, Al, C, and Fe powders.
According to
Figure 1m–r, intergranular residual δ-ferrite and Ti-rich particles were also formed in the TiB
2-inoculated sample, following a mechanism similar to that occurring in the TiC-inoculated sample. Additionally, M
3B
2 phases were found in a skeleton-like morphology. During the deposition process, TiB
2 nano-powders partially dissolved, resulting in the formation of in situ TiC phases due to a rapid Ti-C reaction. The remaining boron segregated into the liquid phase, forming Cr-rich M
3B
2-type borides with a skeleton morphology. Sigolo et al. [
9] similarly observed the formation of skeleton-shaped M
3B
2-type borides during plasma transferred arc (PTA) cladding utilizing boron-modified stainless steels.
3.2. Crystallographic Orientation Characterization
To comprehensively characterize the crystallographic orientation of the samples, the inverse pole figure (IPF) map, phase map, and grain boundaries misorientation maps of each sample at high magnifications are shown in
Figure 2. In the non-inoculated sample, the microstructure predominantly exhibited a lath martensitic structure with a dominant body-centered tetragonal (BCT) structure, with only ~2% of retained austenite (FCC), which increased to ~12% and ~7% in the TiC-inoculated and TiB
2-inoculated samples. The higher fraction of retained austenite in the TiC-inoculated sample can be attributed to the introduction of carbon as a strong austenite stabilizer into the microstructure due to the partial dissolution of preplaced TiC nanoparticles. Chen et al. [
10] also reported that the incorporation of TiC particles into stainless steels during additive manufacturing processes increases the fraction of retained austenite. On the other hand, the segregation of additional boron in the TiB
2-inoculated sample at the grain boundaries of primary austenite could reduce its grain size and increase the phase stability of austenite at room temperature [
11].
According to the grain boundary maps, the inoculated samples demonstrated a significantly high content of low-angle grain boundaries (LAGBs) (
Figure 2h,i), surpassing that observed in the non-inoculated counterpart (
Figure 2g). This can be attributed to the effective role of TiC/TiB
2 nanoparticles serving as heterogeneous nucleation sites for grain formation. Upon introduction into the molten metal, these particles facilitate grain nucleation at multiple locations simultaneously. This controlled nucleation process promotes a more uniform distribution of grains, which may favour the formation of low-angle grain boundaries. The specific interactions between the growing grains and the inoculants could encourage alignment along specific axes, thereby creating conditions conducive to the formation of low-angle boundaries. Also, the presence of inoculants can influence the solidification process, affecting the cooling rate and thermal gradients established during solidification, which, in turn, can impact how grains nucleate and grow relative to inoculants, favoring the formation of low-angle boundaries over high-angle ones.
3.3. Corrosion Properties
Electrochemical impedance spectroscopy (EIS) results, illustrated through Nyquist and Bode plots (
Figure 3a,b), provide a detailed insight into the electrochemical performance and stability of the passive films on the non-inoculated and inoculated samples. The EIS data highlight the direct influence of microstructure, particularly the distribution of grain boundaries on corrosion resistance. According to
Figure 3a, the Nyquist plots of the inoculated samples showed a relatively larger capacitive loop radius compared to the non-inoculated sample, indicating the presence of a more robust and stable passive film, which can be attributed to the high-volume fraction of LAGBs in their microstructure. LAGBs are known for their low energy levels, which makes them less susceptible to corrosion attacks [
12]. Certain microstructural features, such as retained austenite content, can also enhance the electrochemical stability of stainless steels due to the higher solubility of substitutional alloying elements, such as Ni and Mo in the austenite phase, with an FCC crystal structure compared to BCC-Fe [
13]. As observed in the Bode plots, the inoculated samples exhibited higher impedance at low frequencies, indicative of better passive film barrier properties than the non-inoculated sample. The EIS data were fitted using an equivalent circuit (EC) commonly applied to stainless steels to quantify the EIS measurements [
14]. This selected EC, R
s(CPE
P[R
P(R
ct CPE
dl)]) [
1], represents an electrode coated by a porous layer, which aligns with the microporous passive films on the surface of PH 13-8Mo MSS. Based on the fitted data, the overall corrosion resistance of the passive layers (R
Total) was measured to be 5.99 × 10
6 Ω·cm
2, 5.02 × 10
6 Ω·cm
2, and 3.03 × 10
6 Ω·cm
2 for the non-inoculated, TiC-inoculated, and TiB
2-inoculated samples, confirming the slightly higher corrosion resistance of the inoculated samples.
The CPP results (
Figure 3c) show that all samples maintained a robust passive behaviour, with a wide passivation range and consistent pitting potentials (E
pit), indicating minimal degradation of the passive layer. It is evident that the corrosion potentials of inoculated samples were higher compared to the non-inoculated sample. The corrosion resistance did not decrease due to the unchanged chemical composition, especially Cr content, while the observed improvement in electrochemical stability could be attributed to two main microstructural alterations induced by the inoculation process, including (i) higher fraction of LAGBs and (ii) higher content of retained austenite in the inoculated samples. In corrosion science, the intricate relationship between microstructure and passivity breakdown, particularly pitting behaviour, is a complex phenomenon. Several microstructural factors, including grain boundary misorientation angles, the presence of precipitates, phase structures, and residual stresses, significantly influence the passivity behaviour of stainless steels.
In the present study, the relative fractions of low-angle and high-angle grain boundaries could be one of the contributing factors to the corrosion performance of the non-inoculated and inoculated samples. LAGBs, characterized by dislocations that form regular arrays across the boundary plane, have lower energy levels compared to high-angle grain boundaries (HAGBs). This difference is primarily due to the fewer atomic misorientations involved in LAGBs [
12]. As a result, LAGBs tend to be more stable and less prone to corrosion-induced degradation. On the other hand, the point defect model (PDM) suggests that passivity breakdown is primarily driven by cation vacancies’ condensation at the metal–barrier layer interface [
13]. When the flux of cation vacancies surpasses their annihilation rate, accumulation of vacancies leads to the detachment of the barrier layer, ultimately triggering passivity breakdown [
15]. The elevated concentration of Mo in the reversed austenite phase contributes to improved pitting resistance, as Mo reduces cation vacancy flux, thereby increasing the breakdown potential and delaying the onset of pitting [
16]. Furthermore, the high Ni content in reversed austenite also enhances pitting resistance by promoting the crystallization of the passive film and strengthening the barrier layer’s resistance against breakdown [
17]. It has also been reported that increased Ni content in the Cr
2O
3 passive layer fortifies the barrier layer, thereby improving the resistance of the material to pitting corrosion [
18]. These factors collectively explain the observed improvement in pitting resistance in the inoculated samples.
However, it is notable that the Epit of the TiB2-inoculated sample did not improve as significantly as that of the TiC-inoculated sample. This indicates a weaker protective passive film and less stable electrochemical behaviour for the TiB2-inoculated sample. This observation aligns with the microstructural heterogeneities posed by the presence of high-chromium M3B2 phases in the TiB2-inoculated sample, which enhances micro-galvanic coupling effects between the chromium-enriched area (M3B2) and the martensitic matrix. This interaction leads to localized corrosion sites, thereby deteriorating the overall corrosion resistance.
These findings emphasize the critical role of microstructural characteristics, particularly the type and distribution of grain boundaries and the presence of different phases, in influencing the corrosion resistance of stainless steels. The superior performance of the TiC-inoculated samples underscores the benefits of using inoculation techniques that promote a higher fraction of LAGBs and higher content of retained austenite. The challenges observed with the TiB2-inoculated samples highlight the detrimental effects of unfavorable chromium-enriched borides (M3B2), which result in micro-galvanic interactions within a heterogeneous microstructure.
Confocal microscopy (
Figure 3d) was employed to further analyze the corrosion effects on the samples following cyclic potentiodynamic polarization tests. The TiC-inoculated samples exhibited minimal pitting, with pit depths significantly shallower than those found in other samples, indicating superior resistance to localized corrosion. This finding aligns with their enhanced electrochemical stability demonstrated in previous tests.
The presence of large pits observed in
Figure 3d is attributed to the unfavorable area ratio that develops between the corroding pit (small anode) and the passivated surface (large cathode). The 7-day passivation process results in the formation of a stable passive film on the surface. Consequently, when pitting initiates, the establishment of an unfavorable area ratio between the corroding pit and the surrounding surface becomes inevitable. This phenomenon is comparable to the large pits observed on the surface of galvanostatically passivated 316 stainless steel [
19]. Furthermore, the absence of fluctuations in the anodic branch of the CPP curves (
Figure 3c), which would indicate metastable pitting, further supports this observation.
There is a well-founded concern regarding the potential impact of strengthening mechanisms on the corrosion resistance of materials. However, the corrosion analysis conducted in this study has demonstrated that applying nano-inoculation strengthening not only did not compromise the corrosion resistance but also resulted in a slight improvement. This enhancement in corrosion resistance is a significant advantage for the in-service performance of materials in industrial applications. This finding underscores the potential of nano-inoculation strengthening as a valuable technique for enhancing both the mechanical and corrosion properties of materials used in challenging industrial applications.