Selective laser melting of bulk immiscible alloy with enhanced strength: Heterogeneous microstructure and deformation mechanisms

doi:10.1016/j.jmst.2021.06.062

Abstract

Abstract:

To overcome the dimension limits of immiscible alloys produced by traditional techniques and enhance their mechanical properties, bulk Cu-Fe-based immiscible alloy with abundant nanotwins and stacking faults was successfully produced by selective laser melting (SLM). The SLM-produced bulk immiscible alloy displays a heterogeneous microstructure characterized by micro-scaled γ-Fe particles dispersed in fine ε-Cu matrix with a high fraction (~92%) of high-angle grain boundaries. Interestingly, abundant nanotwins and stacking faults are generated in the interior of nano-scaled γ-Fe particles embedded within ε-Cu matrix. The heterogeneous interface of soft domains (ε-Cu) and hard domains (γ-Fe) not only induces the geometrically necessary dislocations (GNDs) but also affects the dislocation propagation during plastic deformation. Therefore, the bimodal heterogeneous interface, and the resistance of nanotwins and stacking faults to the propagation of partial dislocation make the bulk immiscible alloy exhibit an enhanced strength of ~590 MPa and a good ductility of ~8.9%.

Key words: Immiscible alloy, Selective laser melting, Heterogeneous microstructure, Nanotwins, Stacking faults

Shengfeng Zhou, Min Xie, Changyi Wu, Yanliang Yi, Dongchu Chen, Lai-Chang Zhang. Selective laser melting of bulk immiscible alloy with enhanced strength: Heterogeneous microstructure and deformation mechanisms[J]. J. Mater. Sci. Technol., 2022, 104: 81-87.

Figures/Tables 6

Table 1 The chemical composition (wt.%) of 316L stainless steel powder used during SLM.

Cr	Ni	Mo	C	Si	Mn	S	P	Fe
16.42	11.24	2.12	0.03	0.37	1.42	0.011	0.04	Bal.

Table 2 The optimized parameters adopted to produce the Cu-Fe-based immiscible alloys.

Laser power (W)	Laser scanning speed (mm/s)	Spot size (μm)	Layer thickness (μm)
195	287	100	50

Fig. 1. (a) XRD patterns of SLM-produced Cu-Fe-based immiscible alloy and Fe-based alloy powder, (b) the cross-section microstructure of Cu-Fe-based immiscible alloy and the magnified image presented in the right inset, (c) representative EBSD image of Cu-rich matrix, (d) distribution of grain boundary misorientation angle for Cu-rich matrix, (e) bright-field TEM image of Cu-rich matrix and the SAED pattern of grain A corresponding to the fcc structure along [011] zone axis presented in the inset on left bottom. The inset on the right top shows a magnified TEM image of Cu grain, (f, g) bight-field TEM image of Fe-rich particle embedded within Cu-rich matrix, the SAED pattern and magnified image of stacking faults and twin boundaries presented in the inset.

Fig. 2. (a) Tensile engineering stress-strain curves of SLM-produced Cu-Fe-based immiscible alloy and traditional pure copper, (b) the engineering stress and engineering strain of the SLM-produced Cu-Fe-based immiscible alloy and other Cu-Fe-based immiscible alloy produced by traditional techniques, (c) the work hardening rates of SLM-produced Cu-Fe-based immiscible alloy and traditional pure copper.

Fig. 3. The fracture morphologies after tensile tests for (a) traditional pure Cu and the magnified image presented in the right inset, (b) the SLM-produced Cu-Fe-based immiscible alloy, (c) cleavage fracture with river pattern in the γ-Fe particles, and (d) equiaxed dimples in the ε-Cu matrix.

Fig. 4. Bright-field TEM images of dislocation distributions in the SLM-produced Cu-Fe-based immiscible alloy after tensile tests: (a) interactions of ε-Cu matrix and dislocations, and the magnified image of dislocation presented in the left inset, (b, c) interactions of stacking faults and dislocations, and the circled regions corresponding to the SAED of stacking faults and the magnified image of dislocation presented in the inset, (d, e) interactions of γ-Fe particles and dislocations.

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