Page 42 - ESAM-1-2

Engineering Science in
Additive Manufacturing Multi-material additive manufacturing of metals

propagated from the SS316L interface region and expanded metallurgical bonding. Considering the CTE of SS316L,
into the IN625, as confirmed by scanning electron IN625, and IN718, it is evident that IN718 has a lower
microscopy imaging. Such patterned cracks formed due CTE compared to IN625. In SS316L/IN625 structures,
to the implemented scanning strategy, which induced although the actual temperature increases from SS316L
residual stresses and mismatches in CTEs between the sub- to IN625 due to the heat source, a gradual decrease in
layers (Figure 4A and B). the liquid temperature within the transition zone inhibits
Similarly, Chen et al. fabricated multiple layers the protrusion of SS316L into this region, resulting in
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of SS316L/IN625 using MM-LDED to understand the a smoother interface. As reported by Bodnar et al.,
microstructural characteristics at two types of interfaces: it is evident that the process parameters significantly
type I (SS316L as a substrate) and type II (IN625 as a influence the formation of zigzag cracks at the SS316L/
substrate) (Figure 4C and D). This section focuses on type I IN625 interface. In contrast, for SS316L/IN718 bimetallic
interfaces, while type II will be discussed in Subsection 3.3. structures, porosity and cracks were clearly observed in
The type I interface exhibits a flat transition region due to MM-LPBF rather than MM-LDED. This behavior was
less penetration compared to type II (Figure 4E and F). attributed to the formation of brittle intermetallic phases
The shallower penetration in type I was attributed to the during solidification, coupled with the lower CTE of IN718
lower energy absorptivity and thermal conductivity of compared to SS316L, which exacerbates thermal stress at
IN625 relative to SS316L. Fractography images revealed no the interface.
cracks or pores at the interface due to the lower density Besides the influence of brittle intermetallic phase
and viscosity of SS316L compared to IN625. Beyond formation and CTE, the authors speculate that the shielding
surface observation, the inverse pole figure map of the gas flow rate differences between SS316 and IN718 could
type I interface exhibited fine columnar grains in IN625 potentially have an impact on the interfacial porosity
and equiaxed grains on SS316L. Similarly, in a hybrid-AM formation, attributed to the deposition of spattering and
process where SS316L was fabricated using LPBF and welding plume. The lack of data on shielding gas flow rate
IN625 using LDED, no cracks or porosity were visible on at the interface provides no evidence of the influence of gas
the surface, which validates the earlier findings. In all flow rate but could be found on monolithic materials of
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cases, the interface exhibited a good transition zone with SS316 and IN718. 182
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smooth, defect-free bonding due to comparable melting Stainless steel/Cu alloy combinations are utilized in
points and CTE values of the two alloys. fields such as energy and automotive, leveraging the
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The SS316L/IN718 interface has been similarly analyzed mechanical strength of SS and the high thermal conductivity
as the SS316L/IN625-LPBF interfaces. Yusuf et al. of Cu. Figure 5 illustrates the interfacial microstructural
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observed that the transition zone exhibits equiaxed grain characteristics of dissimilar SS/Cu materials with various
growth with a transition zone thickness of approximately Cu alloys.
100 μm and was characterized by columnar grains in A bimetallic SS316L/C18400 structure fabricated
the individual element-rich region (Figure 4G and H). through MM-LPBF exhibited sporadic cracking (Figure 5B)
Despite the presence of irregular-shaped lack-of-fusion on the SS side of the interface, with a transition width of
pores (highlighted in red in Figures 4I and J), no cracks 750 μm (Figure 5A). These cracks were attributed to the
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were detected. Similar observations were made by Mei thermal mismatch between SS and Cu, whereby Cu diffuses
et al. and Wei et al., who attributed the formation into austenitic grain boundaries of SS316L, exerting
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of pores at the interface to energy-material mismatch pressure and causing cracking. Optical microscopic images
and uneven powder deposition during the MM-LPBF revealed SS-rich regions free of porosity or defects, while
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fabrication. In contrast, Ghanavati et al. reported a a considerable number of pores were observed on the
nearly defect-free interface between SS316L and IN718 Cu-rich region due to the high reflectivity and thermal
(Figure 4K and L), which was attributed to the presence conductivity of Cu (Figure 5C and D). EDS analysis
of a strong local temperature gradient at the interface. The revealed a good metallurgical bonding between SS316L
microstructural solidification at the interface proceeded and C18400. With the understanding of the interfacial
through a narrow planar solidification mode, driven by characteristics of SS316L/C18400, the bimetallic structure
the high local temperature gradient. An EDS analysis of SS316L/C25400 could also be examined. Similarly, an
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confirmed a smooth transition from the Fe- to Ni-rich SS316L/C25400/SS316L MMAM structure manufactured
region (Figure 4M). using MM-LPBF exhibited characteristics akin to the
Across various AM processes, the SS/Ni bimetallic SS316L/C18400. Bai et al. observed the formation
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structure with SS as substrate generally exhibited good of a few cracks at the transition region extending to the

Volume 1 Issue 2 (2025) 10 doi: 10.36922/ESAM025180010