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Dela Cruz et al. Microstructures 2023;3:2023012 https://dx.doi.org/10.20517/microstructures.2022.33 Page 9 of 25

generated prior austenite columnar grains that grow in their <001> direction parallel to the building
direction. Meanwhile, Figure 4B, E, H, and K display the randomly orientated HCP martensite phase within
the austenite grains.

The X-ray spectra of the Fe-30Mn-6Si reference alloy and the LPBF alloy made from different processing
parameters were gathered and quantified using the Rietveld refinement method. The results are then shown
in Figure 5 and Table 2. The major phases identified in the LPBF alloy are γ-austenite and ε-martensite
because of their intense XRD peaks and composition that is ≥ 19 wt.%, as seen in Table 2. A dual-phased
[59]
microstructure is expected in the Fe-30Mn-6Si alloy that underwent post-process treatment , while the
homogenised alloy may be single-phase austenite , and such is observed in Figure 5. The existence of the γ
[60]
and ε phases in the LPBF alloy is due to the far-from-equilibrium process conditions of the technique.
Table 2 also reveals three other phases in the LPBF alloy; α-FeMn, α-FeSi, and FeO. The presence and
composition of these phases are observed to vary in the LPBF alloy prepared for different parameters. For
example, FeO was identified at 0.25 J/mm and 0.44 J/mm but not at 0.29 J/mm and 0.88 J/mm. Upon close
inspection at 54.3° 2θ in the 0.29 J/mm, its 10 peak is visible. Several trials were made to include the low-
intensity peaks from those three phases for a detailed analysis, but the quality of the resulting Rietveld
refinement was unsatisfactory. A more detailed XRD scan is therefore necessary for a comprehensive
analysis of those three phases.

Key microstructural features associated with LPBF processing, such as the types and volume fraction of
phases present, solidified grain size, morphology, and texture of the processed samples, were strongly
influenced by the laser power, scan speed, and re-scan strategy. This shows that the desired microstructure
is tailored by controlling laser power and scan speed to change the LED. The information on the thermal
history of the resultant product is, however, necessary to completely understand the development of the
microstructure.

Melt pool of single laser track scan
A polished surface of the reference alloy was subjected to single track laser scans at various LEDs. This
resulted in the melting and subsequent solidification along the laser tracks, which generated a certain melt
pool morphology for a given LED, when viewing a cross section perpendicular to the laser track. The effect
of LED on the cross section of melt pool morphology is shown in Figure 6. Figure 6C and D show that a
high LED creates both a deep and wide melt pool that penetrates at least 120 μm below the polished surface.
In contrast, a low LED generates a relatively shallow melt pool of 50 μm deep [Figure 6A]. In Figure 6B, the
melt pool became wide and deep when the LED was slightly increased from 0.25 J/mm to 0.29 J/mm by
increasing the laser power from 100 W to 175 W and scan speed from 400 mm/s to 600 mm/s. Overall, there
is sufficient lateral overlap of the melt pool tracks because the width of the melt pool is wider than the
0.45 μm distance of the parallel laser tracks.

The melting mode at low LED (0.25 J/mm), as defined by Tenbrock et al., is conduction mode, and the rest,
0.29 J/mm to 0.88 J/mm, are in keyhole mode . In the authors’ single laser track investigation on 316 L
[61]
stainless steel, the group used the melt pool depth-to-width ratio threshold of less than 0.8 as the
conduction mode; above 0.8, the keyhole mode of melting transpired. Conduction mode of melting was
observed at low LED, where the underlying regions are heated through the energy conducted from the
surface . In the keyhole mode of melting, the high LED evaporated the metal and left a vapor cavity in the
[62]
[63]
melt pool that enhanced laser absorption and enabled a deeper melt pool than in conduction mode .