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Article

Interfacial Characteristics and Mechanical Performance of IN718/CuSn10 Fabricated by Laser Powder Bed Fusion

by
Xiao Yang
1,†,
Guangsai Zou
1,2,†,
Zheng Wang
1,
Xinze He
1,
Mina Zhang
1,* and
Jingyu Xu
1,*
1
Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, Ningbo 315201, China
2
College of Mechanical Engineering, Zhejiang University of Technology, Hangzhou 310023, China
*
Authors to whom correspondence should be addressed.
†
These authors contributed equally to this work.
Crystals 2025, 15(4), 344; https://doi.org/10.3390/cryst15040344
Submission received: 19 March 2025 / Revised: 2 April 2025 / Accepted: 3 April 2025 / Published: 6 April 2025
(This article belongs to the Special Issue Advances of High Entropy Alloys (2nd Edition))
Browse Figures
Versions Notes

Abstract

:
To address the critical applications of heterogeneous structures involving nickel-based superalloys (IN718) and copper alloys (CuSn10) under extreme operating conditions, and to address the limitations of traditional joining techniques in terms of interfacial brittleness and geometric constraints, this study employs Laser Powder Bed Fusion (LPBF) technology, specifically multi-material LPBF (MM-LPBF). By precisely melting IN718 and CuSn10 powders layer by layer, the study directly fabricates multi-material IN718/CuSn10 joint specimens, thereby simplifying the complexity of traditional joining processes. The research systematically investigates the interfacial microstructure and mechanical property evolution laws and underlying mechanisms. It reveals that sufficient element diffusion and hardness gradients are present at the IN718/CuSn10 interface, indicating good metallurgical bonding. However, due to significant differences in thermophysical properties, cracks inevitably appear at the interface. Mechanical property tests indicate that the strength of the IN718/CuSn10 joint specimens falls between that of IN718 and CuSn10, but with lower elongation, and fractures primarily occur at the interface. This research provides theoretical support for establishing a process database for LPBF formed of nickel–copper heterogeneous materials, advancing the manufacturing technology of aerospace multi-material components.

1. Introduction

IN718 is renowned for its excellent high-temperature strength (creep resistance at 540–1000 °C), high-temperature stability (yield strength > 700 MPa at 650 °C), and oxidation resistance, while CuSn10 imparts components with exceptional thermal conductivity (≥60 W/m·K) and fatigue resistance [1,2]. The IN718/CuSn10 multi-material system integrates the exceptional high-temperature strength and fatigue resistance of an IN718 nickel-based superalloy with the outstanding thermal conductivity and electrical properties of CuSn10 bronze, creating a synergistic solution for advanced thermal-structural applications. This hybrid architecture demonstrates substantial potential in aerospace engineering, particularly for high-value components like rocket combustion liners and regeneratively cooled nozzles. However, conventional dissimilar material bonding techniques, such as physical vapor deposition (PVD) and brazing, face inherent limitations in geometric flexibility [3,4]. PVD facilitates surface-level bonding through thin-film deposition, while brazing enables bulk bonding via high-temperature filler-metal diffusion. Nevertheless, both methods struggle to fabricate intricate three-dimensional architectures, restricting their applicability in advanced aerospace components.
Multi-material Addictive manufacturing (MM-AM) technology, particularly multi-material LPBF (MM-LPBF), not only offers the capability to form complex geometries when fabricating multi-material structures but also enables compositional gradient transitions. Furthermore, owing to its ultrahigh cooling rate and complex thermal cycle, MM-LPBF achieves fine microstructural control, thereby enhancing the mechanical properties of interfaces, demonstrating great potential in the field of multi-material fabrication [5,6,7]. Currently, the research in the field predominantly focuses on steel/copper [8,9], steel/nickel [10], and titanium/nickel systems [11]. Interface integrity remains a critical challenge in heterogeneous material systems due to intermetallic compound formation and thermophysical property disparities, which can induce interfacial cracking and performance degradation [12,13]. Despite these advancements, the nickel–copper material system remains understudied. Onuike et al. [14] reported macroscopic cracking at IN718/GRCop-84 interfaces caused by significant CTE mismatch. Grandhi et al. [15] achieved metallurgical bonding in DED-fabricated IN718/CuSn10 interfaces but observed dimensional inaccuracies and microcracks. These findings highlight persistent challenges in defect mitigation and geometric precision for nickel/copper multi-material systems. Compared to DED, LPBF offers superior resolution (20–100 μm layer thickness) and densification (>99.5%), positioning it as a promising alternative for high-quality IN718/CuSn10 interfaces.
This study systematically investigates the interfacial microstructure evolution and mechanical property gradients in MM-LPBF fabricated IN718/CuSn10 heterostructures, elucidating geometric bonding mechanisms and fracture behavior at heterogeneous interfaces. The aim is to provide theoretical support for establishing a process database for nickel/copper heterogeneous materials and to advance the manufacturing technology of aerospace multi-material components.

2. Experimental Procedure

2.1. Materials

The experimental materials consisted of IN718 and CuSn10 alloy powders supplied by Asia New Materials. Figure 1a,a1 illustrate the morphology and particle size distribution of IN718 powder, while Figure 1b,b1 depict the morphology and particle size distribution of CuSn10 powder. The particle size distribution of both IN718 and CuSn10 powders ranges from 10 to 70 μm, with average particle sizes (Dv(50)) of 31.88 μm and 30.78 μm, respectively. The powders exhibit a predominantly spherical morphology, with no significant agglomeration or adhesion observed. In conclusion, the IN718 and CuSn10 powders are suitable for MM-LPBF fabrication.

2.2. Laser Powder Bed Fusion Process

The experiment utilized a DiMetal-100Pro LPBF machine (manufactured by Guangdong Leijia Additive Manufacturing Technology Co., Ltd., Zhuhai, China) to fabricate samples with dimensions of 10 mm × 22 mm × 22 mm (details in Table 1). The machine is equipped with a multi-powder canister system, enabling seamless switching between IN718 and CuSn10 powders.
During the experiment, IN718 powder was first used to form the substrate. The powder was delivered to the scraper via a conveying system, evenly spread, and then selectively melted by a laser beam following a preset path, layer by layer, until the substrate was fully formed. Subsequently, the system switched to CuSn10 powder, which was deposited layer by layer on the surface of the IN718 substrate, ultimately producing a CuSn10/IN718 multi-material specimen. An interlayer rotation scanning strategy was employed during the forming process, with a rotation angle of 67° between adjacent layers, as illustrated in Figure 2a. The specific process parameters for IN718 and CuSn10 were provided by the equipment supplier (Guangdong Leijia Additive Manufacturing Technology Co., Ltd., Zhuhai, China) and are detailed in Table 2. To ensure the quality of the formed parts, the entire printing process was conducted under an inert gas atmosphere, with the oxygen concentration in the chamber strictly controlled below 0.3%. After forming, tensile specimens were cut from the bulk using wire cutting technology (Figure 2b). Considering the characteristics of additive manufacturing and the size limitations of the equipment, non-standard sized specimens were used for mechanical property testing in this study. Relevant research on the application of this equipment has been reported in the literature [16].

2.3. Microstructural Characterization

The bulk specimens were sequentially ground using 200#, 550#, and 1200# sandpapers, followed by polishing with 9 μm, 3 μm, and 1 μm diamond suspensions and a 0.25 μm silica suspension to achieve a smooth surface. To clearly observe the interfacial microstructure, the specimens were chemically etched using an etching solution (HNO3:HCl:H2O = 1:1:1) for approximately 40 s and then ultrasonically cleaned with alcohol to remove residues. Prior to tensile testing, the tensile specimens were polished with 550# and 1200# sandpapers to eliminate surface defects and avoid stress concentration that could affect the accuracy of the experimental results. The interfacial microstructure and surface morphology of the specimens were observed using a super-depth microscope (VHX-7000, Keyence, Osaka, Japan). For further microstructural analysis, a field-emission scanning electron microscope (FEI QUANTA 250 FEG, Thermo Fisher Scientific, Hillsboro, OR, USA) equipped with an energy-dispersive X-ray spectroscopy (D8 DISCOVER, Bruker, Karlsruhe, Germany) detector was employed to analyze the compositional distribution of the specimens. Additionally, electron backscatter diffraction (EBSD) (Oxford Instruments, High Wycombe, UK) was utilized to characterize the grain structure, orientation, and stress distribution of the specimens, with a scanning step size set to 0.6 μm. The experimental data were processed and analyzed using AztecCrystal 2.0 software.

2.4. Mechanical Characterization

Microhardness testing was conducted using a fully automatic Vickers hardness tester (VH3100, Buehler, Lake Bluff, IL, USA). A 3 mm section along the interface of the specimen in the build direction was selected, with four measurement points taken at intervals of 0.2 mm, totaling 64 points. The average value of the measurement points was calculated. A load of 500 g was applied, with a dwell time of 10 s. Tensile testing was performed using a tensile stage (DDS4, Kammrath & Weiss GmbH, Schönefeld, Germany) at a constant strain rate of 2 μm/s.

3. Results and Discussion

3.1. Microstructure

As shown in Figure 3a, the macroscopic view of the IN718/CuSn10 fabricated specimen reveals a distinct boundary between IN718 and CuSn10, with a clean and intact appearance and well-controlled dimensions, demonstrating the excellent forming capability of LPBF technology for IN718/CuSn10. No warping, bulging, or macroscopic interfacial cracks were observed during the printing process, indicating good metallurgical bonding between IN718 and CuSn10. Figure 3b presents an optical micrograph of the IN718/CuSn10 interface, where localized cracks along the printing direction are observed at the boundary. The primary reasons for these cracks are as follows: the significant difference in melting points between IN718 (1608 °C) and CuSn10 (1273 °C) leads to poor stability of the molten pool at the interface [17]; the difference in thermal expansion coefficients between nickel (13 × 10−6/°C) and copper (17 × 10−6/°C) introduces excessive stress concentration, making the molten pool prone to cracking [18]; and the difference in laser absorption rates between IN718 (35% at 1064 nm) and CuSn10 (12% at 1064 nm) results in uneven energy input, causing insufficient melting of the powder [19,20]. These thermophysical property differences between the two materials are the main causes of crack formation at the interface. Additionally, pores are observed on the CuSn10 side, which are attributed to insufficient energy input caused by the high thermal conductivity and reflectivity of copper, leading to incomplete melting of the powder [21]. Additionally, heterogeneous pores are observed at the interface, which are attributed to the shrinkage caused by the solidification of the liquid copper alloy. Simultaneously, the pre-solidified nickel alloy hinders the timely replenishment of the surrounding liquid copper alloy, leading to the formation of these pores [22].
Figure 3c displays the interfacial corrosion metallographic image of the IN718/CuSn10 sample. The boundaries of the molten pool are clearly visible, accompanied by the presence of cracks, indicating that significant thermal stress existed in the molten pool of the fusion zone during the solidification process. A distinct transition zone is observed at the interface, with its color and micro-morphology differing from those of the two base materials. The width of the transition zone is approximately 750 μm, suggesting sufficient diffusion and metallurgical reactions between the two materials. This zone is defined as extending from the fine-grained CuSn10 boundary to the IN718-side crack front. The fusion zone width is determined by Ni-Cu interfacial diffusion, which is process-dependent and controlled by LPBF parameters [18]. Both the copper alloy and the nickel alloy exhibit clear fish-scale patterns at the molten pool boundaries. Additionally, distinct columnar crystals are observed on the copper alloy side, as well as cracks in the transition zone near the IN718 region. In Figure 3d, different regions formed by the cyclic flow within the molten pool can be clearly observed at the interface, primarily due to the Marangoni effect. During the solidification of the molten pool, solute redistribution leads to the phenomenon of compositional segregation [23,24]. Compared to Ni, Cu has poorer corrosion resistance, which is why it appears gray-black after corrosion. Based on this, it can be inferred that the silvery-white regions are Ni-rich phases, while the gray-black regions are Cu-rich phases. Under the influence of the Marangoni effect, the Ni-rich phases from the nickel matrix diffuse into the copper matrix, and the Cu-rich phases from the copper matrix diffuse into the nickel matrix, as evidenced by the elemental segregation and compositional variations shown in Figure 4. Due to the limited solubility of Ni and Cu in the respective matrices, these two heterogeneous phases exist as separate phases in the matrices after solidification [25,26].

3.2. Elemental Distribution

Figure 4a shows the elemental distribution at the interface. From the elemental distribution map, it can be clearly observed that Cu is primarily distributed in the CuSn10 region, while Ni is mainly concentrated in the IN718 region. In the transition zone between the two materials, Ni exhibits more uniform diffusion compared to Cu, whereas Cu shows segregation within the molten pool. Along the building direction from IN718 to CuSn10, the high-energy laser input penetrates the current deposited layer upon melting CuSn10, leading to remelting of the adjacent IN718 interfacial layer and triggering Ni-Cu molten pool convection. As layer-by-layer deposition advances, Marangoni-driven convection in the molten pool promotes the upward transport of alloying elements from underlying layers. Consequently, a characteristic cup-shaped Cu-rich phase distribution forms near the CuSn10 of the interfacial region. Under the influence of the Marangoni effect, the aggregation of Ni-rich and Cu-rich phases is formed. Figure 4b displays the variation in elemental concentration across the interface. The results indicate a gradual transition of Ni and Cu elements; however, due to the aggregation of Ni-rich and Cu-rich phases, significant fluctuations in concentration are observed. Figure 4c presents the composition at three specific points, showing that the Ni and Cu elements do not exhibit a linear change along the build direction. This further confirms the presence of sufficient elemental mixing at the interface. In summary, elemental diffusion and segregation at the interface are key factors enabling the metallurgical bonding of IN718 and CuSn10.

3.3. Crystallography

Through EBSD characterization analysis, the grain structure and stress distribution at the IN718/CuSn10 interface were further investigated. Figure 5a,b presents the Kernel Average Misorientation (KAM) map and Inverse Pole Figure (IPF) map of the IN718/CuSn10 interface. The grains of IN718 exhibit a columnar structure, with grain size decreasing sharply near the interface and then gradually transitioning to CuSn10 columnar grains with a smaller aspect ratio. A significant temperature gradient along the deposition direction, which is the optimal growth direction for grains, promotes the formation of columnar grains in both IN718 and CuSn10 [27,28]. During laser remelting, IN718 facilitates the epitaxial growth of grains from the previously solidified layer, resulting in large columnar grains. Due to the high thermal conductivity and reflectivity of the CuSn10 alloy, the energy available to drive its epitaxial growth is insufficient, resulting in the formation of columnar grain structures with a relatively small aspect ratio in the CuSn10 alloy.
The IN718 region exhibits significant stress concentration, while the CuSn10 region shows relatively lower stress, attributed to the good plasticity and larger grain size of the copper alloy, which effectively disperses and alleviates stress concentration. Stress in IN718 is primarily concentrated at grain boundaries, where cracks tend to initiate and propagate toward the IN718 side. When the stress exceeds the material’s strength limit, cracks serve as channels for stress release. As shown in Figure 5b, cracks mainly propagate along the growth direction of columnar grains, manifesting as intergranular cracks parallel to the deposition direction. Cracks typically propagate along the direction of weakest atomic bonding, and the weaker adhesion at columnar grain boundaries results in poor resistance to crack propagation, promoting longitudinal crack extension [29].
In the fine-grained zone at the interface (as shown in Figure 5c), the differences in physical properties between IN718 and CuSn10 alter the thermodynamics of the molten pool, leading to the formation of equiaxed grains [30]. The grain morphology is determined by the ratio of temperature gradient (G) to solidification rate (R), with lower G/R values favoring the formation of equiaxed grains [31,32]. The addition of the copper alloy may further reduce the G/R value, promoting the formation of equiaxed grains. Additionally, heterogeneous nucleation and recrystallization of IN718 and CuSn10 may contribute to significant grain refinement. The fine-grained zone is observed in regions with higher copper content, where nickel particles exist as nanoscale and submicron spherical phases [33]. Due to the higher melting point of nickel, it solidifies earlier, hindering the growth of copper-phase grains while providing nucleation sites for their growth. Furthermore, during the layer-by-layer laser melting process, overlapping regions generate residual stress due to the large temperature gradient and rapid cooling inherent to the LPBF process. When the laser melts the next layer, energy is transferred to the already solidified layer, inducing recrystallization in the previously solidified regions [34,35]. The combined effects of these factors enhance the grain nucleation rate, leading to grain refinement. The average grain size in the fine-grained zone is approximately 0.191 μm, as shown in Figure 5e. The equiaxial grain with smaller size increases the proportion of grain boundaries, effectively dispersing stress concentration. Consequently, the KAM map shows a more uniform stress distribution, preventing crack formation, as illustrated in Figure 5d.

3.4. Mechanical Properties

Figure 6 illustrates the microhardness distribution across the multi-material IN718/CuSn10 interface. The microhardness on the IN718 side is measured at 328.4 ± 6.7 HV, indicating high strength and hardness, whereas the CuSn10 side exhibits a hardness of 153.1 ± 4.1 HV. The intermediate transition zone demonstrates a gradual variation in hardness, lying between that of IN718 and CuSn10. The hardness values are uniformly distributed in the pure nickel alloy and pure copper alloy regions, consistent with the values reported for nickel alloys [36] and copper alloys [37] in the literature. A distinct gradient in hardness is observed from the IN718 region to the CuSn10 region, with values decreasing from higher to lower levels. At the interface, diffusion and mixing between the two materials have resulted in a gradient in composition and microstructure. Near the IN718 side of the transition zone, the hardness values initially exhibit significant fluctuations, suggesting a complex microstructure in this area, which is associated with microstructural inhomogeneities such as compositional segregation and microcracks, as observed in microstructural characterization. As the transition zone approaches the CuSn10 side, the hardness values stabilize, primarily due to the formation of a fine-grained region free from crack defects.
Figure 7 presents the stress–strain curves of the tensile specimens. The curves exhibit excellent repeatability, indicating stable tensile properties of the specimens. The CuSn10 material demonstrates relatively low tensile strength (approximately 200–300 MPa) but high elongation (exceeding 35%), reflecting its excellent ductility. Fracture occurs in the middle of the specimen, and the fracture surface shows significant necking, further confirming its superior plasticity. Table 3 summarizes the tensile properties of the specimens. The IN718 material exhibits high tensile strength (approximately 800–1000 MPa) but relatively low elongation (around 25–30%), indicating high strength with moderate ductility. In the fractured middle section of the IN718 specimen, the degree of necking is less pronounced compared to that of CuSn10. The tensile strength of the IN718/CuSn10 sample lies between that of CuSn10 and IN718, at approximately 400 MPa, while its elongation is notably low, at around 3%. This suggests that the IN718/CuSn10 interface or transition region may exhibit limited ductility, potentially due to the complex microstructure and stress concentration at the interface. Further analysis of the fracture morphology and microstructure in the transition zone would provide deeper insights into the mechanical behavior of the multi-material system.
Figure 8a,b presents the fracture morphology of the multi-material IN718/CuSn10 joint. The fracture surface is relatively smooth and flat, with river-like patterns observed at the interface (Figure 8b), indicating distinct cleavage characteristics and predominantly brittle fracture. Cracks are visible on the fracture surface (Figure 8a), primarily located on the IN718 side of the interface, suggesting that the fracture occurred at the interface and was biased toward the IN718 side. Figure 8c,d shows the morphology and elemental distribution of the tensile fracture surfaces of the two materials. The fracture surface appears rough, with evident tearing marks, indicating some degree of plastic deformation during the fracture process. Irregular protrusions and depressions are observed at the fracture edges, likely resulting from necking and fracture during tensile loading. The elemental distribution map reveals that the fracture primarily occurred at or near the interface of the two materials, which is associated with stress concentration caused by crack defects at the interface. Although the fracture occurred at the interface, the elemental distribution map shows some diffusion and mixing of the two elements at the interface, suggesting that metallurgical reactions took place. Future research may employ green lasers to address unmelted defects and porosity caused by the high reflectivity of copper alloys. Alternatively, remelting treatment could be implemented to enhance interfacial metallurgical bonding; furthermore, the gradient interlayer design could be optimized to suppress crack initiation and improve processability, thereby significantly enhancing interfacial bonding strength [38,39].

4. Conclusions

In this study, IN718/CuSn10 multi-material samples were successfully fabricated using MM-Laser Powder Bed Fusion (MM-LPBF) technology. The research findings are as follows:
(1)
Good metallurgical bonding was achieved at the IN718/CuSn10 interface, with a transition zone width of approximately 750 μm, indicating sufficient element diffusion and metallurgical reactions between the two materials.
(2)
Due to the significant differences in thermophysical properties between IN718 and CuSn10, microcracks occurred at the interface. These microcracks significantly adversely affected the hardness and tensile properties, impairing the interfacial bonding performance.
(3)
The grain size at the interface was significantly decreased, forming a fine-grained zone that effectively dispersed stress concentration.
(4)
Microcracks, elemental segregation, and microstructural inhomogeneity were identified as the main factors affecting the mechanical properties. Further improvements in the overall performance of the material could be achieved by promoting the formation of equiaxed grains to disperse stress and eliminate defects.

Author Contributions

Conceptualization, X.Y. and G.Z.; methodology, X.Y.; validation, Z.W. and X.H.; formal analysis, G.Z.; investigation, X.Y.; resources, M.Z.; data curation, X.Y.; writing—original draft preparation, X.Y. and G.Z.; writing—review and editing, X.Y. and G.Z.; supervision, J.X.; project administration, M.Z.; funding acquisition, M.Z. All authors have read and agreed to the published version of the manuscript.

Funding

This work was supported by the “Ningbo Natural Science Foundation” (Grant No. 2023J327), National Key R&D Program (Grant No. 2023YFB4607000), and National Natural Science Foundation of China under Grant 62201373.

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding author.

Conflicts of Interest

The authors declare no conflicts of interest.

Abbreviations

The following abbreviations are used in this manuscript:
LPBFLaser Powder Bed Fusion
AMAddictive manufacturing
MM-AMMulti-material Addictive manufacturing
MM-LPBFMulti-material LPBF

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Crystals 15 00344 g001
Figure 1. (a,b) SEM images of the IN718 and CuSn10 powders; (a1,b1) the frequency distribution diagrams of particle size.
Figure 1. (a,b) SEM images of the IN718 and CuSn10 powders; (a1,b1) the frequency distribution diagrams of particle size.
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Figure 2. (a) Schematic diagram of laser scanning strategy formation for IN718/CuSn10; (b) schematic diagram of tensile specimen sampling and dimensions for multi-material IN718/CuSn10, IN718, and CuSn10.
Figure 2. (a) Schematic diagram of laser scanning strategy formation for IN718/CuSn10; (b) schematic diagram of tensile specimen sampling and dimensions for multi-material IN718/CuSn10, IN718, and CuSn10.
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Figure 3. (a) Macroscopic morphology of the multi-material IN718/CuSn10 specimen; (b) corresponding optical micrograph of the interface; (c) optical micrograph of the sample interface after etching; and (d) corresponding magnified view.
Figure 3. (a) Macroscopic morphology of the multi-material IN718/CuSn10 specimen; (b) corresponding optical micrograph of the interface; (c) optical micrograph of the sample interface after etching; and (d) corresponding magnified view.
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Figure 4. (a) Band contrast image; (a1,a2) corresponding EDS elemental maps illustrating the spatial distribution of Ni and Cu; (b) EDS line scan demonstrating the concentration gradients of Ni and Cu along the scanning path; and (c) magnified view of (b), combined with EDS point analysis, revealing the Ni and Cu content and their distribution characteristics at the interface region.
Figure 4. (a) Band contrast image; (a1,a2) corresponding EDS elemental maps illustrating the spatial distribution of Ni and Cu; (b) EDS line scan demonstrating the concentration gradients of Ni and Cu along the scanning path; and (c) magnified view of (b), combined with EDS point analysis, revealing the Ni and Cu content and their distribution characteristics at the interface region.
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Figure 5. (a,b) The Kernel Average Misorientation (KAM) and Inverse Pole Figure-Z (IPF-Z) maps of the multi-material IN718/CuSn10 interface, respectively; (c,d) the corresponding KAM and IPF-Z maps of the fine-grained region highlighted in (b); and (e) the grain size distribution histogram for the region depicted in (c).
Figure 5. (a,b) The Kernel Average Misorientation (KAM) and Inverse Pole Figure-Z (IPF-Z) maps of the multi-material IN718/CuSn10 interface, respectively; (c,d) the corresponding KAM and IPF-Z maps of the fine-grained region highlighted in (b); and (e) the grain size distribution histogram for the region depicted in (c).
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Figure 6. Microhardness distribution profile across the multi-material IN718/CuSn10 interface.
Figure 6. Microhardness distribution profile across the multi-material IN718/CuSn10 interface.
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Figure 7. Tensile specimen stress–strain curves: (a) CuSn10; (b) IN718; (c) multi-material IN718/CuSn10.
Figure 7. Tensile specimen stress–strain curves: (a) CuSn10; (b) IN718; (c) multi-material IN718/CuSn10.
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Figure 8. (a,b) Fracture surface morphology of the tensile specimens; (c,d) morphology of the fracture location in the tensile specimens, along with the distribution maps of Ni and Cu elements.
Figure 8. (a,b) Fracture surface morphology of the tensile specimens; (c,d) morphology of the fracture location in the tensile specimens, along with the distribution maps of Ni and Cu elements.
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Table 1. Technical specifications of the DiMetal-100Pro Laser Powder Bed Fusion system.
Table 1. Technical specifications of the DiMetal-100Pro Laser Powder Bed Fusion system.
ItemParameter
Laser typeIPG single laser, 500 W
Build size150 × 150 × 150 mm
Layer thickness0.02–0.1 mm
Scanning speed≤7 m/s
Table 2. SLM process parameters for IN718 and CuSn10.
Table 2. SLM process parameters for IN718 and CuSn10.
ItemIN718CuSn10
Laser power (W)275300
Laser scanning speed (mm/s)1050700
Scanning space (mm)0.080.08
Layer thickness (mm)0.030.03
Table 3. Tensile properties of IN718, CuSn10, and multi-material IN718/CuSn10.
Table 3. Tensile properties of IN718, CuSn10, and multi-material IN718/CuSn10.
SampleYS (MPa)UTS (MPa)EL (%)
CuSn10359.6 ± 5.2487.7 ± 534.6 ± 1.2
IN718688.9 ± 8.3911.7 ± 1229.6 ± 3.1
IN718/CuSn10360.5 ± 3415.2 ± 15.33.0 ± 0.6
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MDPI and ACS Style

Yang, X.; Zou, G.; Wang, Z.; He, X.; Zhang, M.; Xu, J. Interfacial Characteristics and Mechanical Performance of IN718/CuSn10 Fabricated by Laser Powder Bed Fusion. Crystals 2025, 15, 344. https://doi.org/10.3390/cryst15040344

AMA Style

Yang X, Zou G, Wang Z, He X, Zhang M, Xu J. Interfacial Characteristics and Mechanical Performance of IN718/CuSn10 Fabricated by Laser Powder Bed Fusion. Crystals. 2025; 15(4):344. https://doi.org/10.3390/cryst15040344

Chicago/Turabian Style

Yang, Xiao, Guangsai Zou, Zheng Wang, Xinze He, Mina Zhang, and Jingyu Xu. 2025. "Interfacial Characteristics and Mechanical Performance of IN718/CuSn10 Fabricated by Laser Powder Bed Fusion" Crystals 15, no. 4: 344. https://doi.org/10.3390/cryst15040344

APA Style

Yang, X., Zou, G., Wang, Z., He, X., Zhang, M., & Xu, J. (2025). Interfacial Characteristics and Mechanical Performance of IN718/CuSn10 Fabricated by Laser Powder Bed Fusion. Crystals, 15(4), 344. https://doi.org/10.3390/cryst15040344

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