Graphical abstract

1. Introduction

The AlCoCrFeNi alloy and its derivatives within the high-entropy alloy (HEA) class have emerged as a focal point in recent alloy design efforts due to their compositional complexity and tunable properties [1, 2]. Among the principal elements, aluminum has been shown to exert the most pronounced effect on both the microstructure and mechanical behavior of the alloy. In the Al_xCoCrFeNi system, progressive addition of Al induces a transformation in the primary crystal structure—from a face-centered cubic (fcc) lattice at lower Al concentrations to a body-centered cubic (bcc) lattice at higher concentrations. Specifically, the fcc structure dominates for x ≤ 0.4, a dual-phase (fcc + bcc) region appears within 0.5 ≤ x ≤ 0.9, and a single bcc phase becomes stable for x ≥ 0.9 [3]. Building upon earlier studies focused on pure single-phase matrices, recent alloy development strategies have increasingly emphasized the deliberate formation of multi-phase microstructures and/or intermetallic precipitates to enhance performance [4, 5].

Among the fcc-structured Al_xCoCrFeNi high-entropy alloys, the Al_0.1 composition has emerged as a particularly attractive candidate due to its advantageous combination of mechanical resilience and microstructural stability, especially under cryogenic and moderately elevated temperature regimes [6, 7]. The stability of the single-phase FCC matrix facilitates deformation through dislocation glide and mechanical twinning, mechanisms that synergistically enhance strain hardening capacity and confer exceptional resistance to brittle fracture, even at cryogenic temperatures [6]. Additionally, its demonstrated hot workability and resistance to thermomechanical cracking underscore its suitability for industrial forming operations such as hot forging and rolling [7]. Despite these favorable mechanical attributes, a notable limitation lies in its inherently modest thermal conductivity, which may hinder its deployment in applications where efficient heat dissipation is critical—such as heat exchangers in power plants, thermal shielding structures, or load-bearing supports interfacing with superconducting components [8]. A promising strategy to overcome this constraint is inspired by the recent study conducted by Mirzababaei et al. [9], who achieved a 2.5- to 6.6-fold enhancement in thermal conductivity of 316L stainless steel through the incorporation of 60 vol% Cu via laser powder bed fusion (LPBF). This substantial increase highlights the potential of Cu-rich alloying in promoting thermal transport, suggesting that similar compositional modification of the Al_0.1CoCrFeNi alloy may enable the development of multifunctional materials that simultaneously meet structural and thermal performance requirements.

Traditionally, the fabrication of HEAs including Al_0.1CoCrFeNi alloy has been primarily achieved through arc melting and mechanical alloying techniques [10]. In the arc melting route, constituent elements are first weighed according to the desired atomic ratios and then melted under an inert atmosphere, typically argon, using a high-temperature electric arc. Alternatively, mechanical alloying enables the synthesis of HEA powders via high-energy ball milling, wherein repeated fracturing and cold welding of elemental or pre-alloyed powders result in intimate mixing and alloy formation. These powders can then be consolidated through spark plasma sintering, hot isostatic pressing, or conventional sintering. Despite the success of these techniques in establishing the foundational understanding of HEAs, both routes face challenges related to compositional homogeneity, limited microstructural tunability, and labor-intensive post-processing requirements. Recent efforts have thus explored Cu additions to Al_0.1CoCrFeNi alloys via arc melting, revealing that Cu induces the formation of coherent Cu-rich nanoprecipitates within the FCC matrix, which effectively enhance yield strength through precipitation strengthening [11]. However, increasing Cu content can reduce ductility due to strain localization and microcrack formation at precipitate–matrix interfaces [11]. These limitations have motivated the exploration of advanced processing methods—particularly additive manufacturing—as a means to produce compositionally complex alloys with refined microstructures and geometrically flexible architectures [6].

In this study, Cu-added Al_0.1CoCrFeNi composites were fabricated using the LPBF) process to establish an optimized processing window across a range of laser powers (100–250 W) and scanning speeds (200–800 mm/s). The objective was to evaluate the feasibility of producing high-performance, Cu-enriched Al_0.1CoCrFeNi-based alloys through in situ laser processing of blended elemental powders, thereby exploring the potential of compositional tuning for enhanced multifunctional properties. A Cu content of 21.8 wt%—corresponding to approximately 19.6 at%—was selected not only to enable sufficient Cu enrichment for Cu-rich phase formation, but also to shift the overall composition toward a near-equiatomic Al_0.1CoCrFeNiCu configuration.

2. Experimental

2.1. Preparation of Mixed Al_0.1CoCrFeNi + Cu composite powders

The Al_0.1CoCrFeNi powder was synthesized via a high-purity gas atomization process at Kongju National University (Cheonan, South Korea). High-purity elemental metals were weighed according to the stoichiometric composition and induction-melted at approximately 1600 °C to produce a homogeneous molten alloy. The melt was then discharged through a boron-nitride nozzle (8 mm diameter) into a gas atomization chamber, where high-pressure argon gas (1.3 MPa) was employed to atomize the stream into fine droplets. These droplets rapidly solidified into spherical powders, which were collected at the base of the chamber. The atomized powders exhibited a smooth, spherical morphology with a particle size distribution characterized by D₁₀, D₅₀, and D₉₀ values of 16.9 μm, 34.0 μm, and 58.5 μm, respectively.

For the preparation of Al_0.1CoCrFeNi + Cu composite powders, commercially available copper powders (99.7% purity, −325 mesh) were procured from Royal Metal Powders Inc. (Maryville, TN, USA) and blended with Al_0.1CoCrFeNi powders at a weight ratio of 21.8 wt% Cu to 78.2 wt% Al_0.1CoCrFeNi. This composition was selected to develop a HEA system in which no single element dominates the overall chemistry. Powder mixing was conducted using a Retsch PM 100 planetary ball mill (Haan, Germany). To prevent fragmentation or mechanical deformation of particles, no grinding media were introduced into the milling jar. A total of 500 g of blended powder was loaded, and the mixing was performed at 100 rpm for 1 hour. The SEM image and chemical composition maps of mixed Al0.1CoCrFeNi + Cu powders are given in Fig. 1. This procedure ensured uniform dispersion of the Cu particles within the Al_0.1CoCrFeNi matrix, resulting in a homogeneously blended precursor powder suitable for subsequent laser processing.

2.2. LPBF process of Al_0.1CoCrFeNiCu Composite Mixture

The fabrication of Al_0.1CoCrFeNi + Cu composite structures was carried out using a 2ONELAB LPBF system (Coherent, Germany). The machine is equipped with an infrared fiber laser operating at a wavelength of 1070 nm and a focused beam diameter of approximately 40 μm. For the purpose of laser process parameter optimization, sixteen cubic samples—each measuring 7 mm × 7 mm × 7 mm—were fabricated simultaneously in a single build. These samples were spatially separated by approximately 10 mm to minimize thermal interactions during fabrication.

The LPBF process employed a layer thickness of 40 μm and a hatch spacing of 40 μm. To systematically investigate the influence of processing conditions on part geometry and melt behavior, four laser power levels (100 W, 150 W, 200 W, and 250 W) and four scanning speeds (200 mm/s, 400 mm/s, 600 mm/s, and 800 mm/s) were selected. These combinations enabled a comprehensive evaluation of laser-material interaction under varying volumetric energy density (VED) conditions. The corresponding VED values for each parameter set are summarized in Table 1.

2.3. Microstructure Characterization

Following LPBF fabrication, the specimens were detached from the build platform using a Sodick VL600Q electrical discharge machining (EDM) system (Yokohama, Japan). To expose the internal features for microstructural evaluation, cross-sectional samples were sectioned along the central plane of each 7 mm cube in the build direction using the same EDM system. The extracted sections were then mounted in conductive carbon-filled phenolic resin using a hot compression mounting press to ensure mechanical stability during subsequent metallographic processing.

The mounted specimens underwent sequential grinding with silicon carbide (SiC) abrasive papers, beginning with coarse 120 grit and advancing to 1200 grit to achieve a flat and uniform surface. This was followed by mechanical polishing using 1 μm diamond suspension, and final surface finishing was completed with 0.05 μm colloidal silica to obtain a mirror-like finish suitable for high-resolution microstructural analysis.

Optical microscopy was performed using a Zeiss optical microscope (Oberkochen, Germany) to assess part relative density and surface integrity. For statistical reliability, a minimum of 10 cross-sectional images were acquired at 50× magnification from each sample to account for repeatability and measurement uncertainty.

To further investigate elemental distribution within the printed structures, scanning electron microscopy (SEM) was conducted using a Hitachi TM4000 tabletop SEM (Tokyo, Japan) equipped with a Bruker energy-dispersive spectroscopy (EDS) detector (Billerica, MA). This allowed qualitative mapping of elemental homogeneity and potential segregation at the microstructural scale.

3. Results and Discussion

3.1 Cross-Sectional Analysis of LPBF-Fabricated Al_0.1CoCrFeNiCu Composite Structures

The cross-sectional micrographs of the LPBF-fabricated Al_0.1CoCrFeNi + Cu composites (Fig. 2) revealed a strong dependency of defect morphology on the applied VED. At lower VEDs—particularly those corresponding to laser powers of 100–200 W combined with scanning speeds of 400–800 mm/s—pronounced lack-of-fusion (LOF) defects were observed. These defects arise from insufficient energy input, leading to incomplete melting and poor interlayer bonding [12]. The resulting melt pools are typically shallow and narrow, which inhibits capillary-driven consolidation and results in unbonded regions between tracks and layers. Conversely, processing conditions with higher VEDs, such as 250 W at scanning speeds of 200–400 mm/s, induced deeper, more unstable melt pools characteristic of keyhole-mode melting. In this regime, excessive energy input causes vapor-induced recoil pressure and dynamic melt pool instabilities, promoting the formation of gas entrapment and spherical porosity—commonly referred to as keyholes [13]. These observations are consistent with previous findings in HEA systems [14], where porosity was shown to increase significantly at both insufficient and excessive energy input levels due to LOF and keyhole phenomena, respectively. Importantly, none of the fabricated samples achieved a fully dense structure, as solidification cracking persisted across the processing window, suggesting that in addition to energy input, factors such as elemental segregation [15], thermal gradients, and residual stress accumulation [16] play critical roles in governing microstructural integrity during LPBF.

The consistent presence of solidification cracking observed in the LPBF-fabricated Al_0.1CoCrFeNi + Cu composites, even under conditions optimized to mitigate LOF and keyhole porosity, underscores the critical role of alloy chemistry and solidification dynamics in crack formation. Solidification cracking in HEAs is often attributed to a combination of low melting-point eutectic films, wide solidification temperature ranges, and high thermal gradients inherent to laser-based additive manufacturing processes [14]. In multicomponent systems such as Al_0.1CoCrFeNi + Cu, the differential solidification behavior between constituent elements can lead to compositional segregation at grain boundaries, particularly of Cu, which has a relatively low melting point and poor solubility in several transition metals. Guo et al. [15] demonstrated that microsegregation of Si and N in an FeCoCrNiMn-based HEA induced localized eutectic formation and weakened intergranular cohesion during the terminal stages of solidification. A similar mechanism is likely at play in the current alloy system, where segregation of Cu to interdendritic regions reduces the local solidus temperature and promotes the formation of thin liquid films along grain boundaries. These liquid films are susceptible to tensile stress-induced rupture during cooling, especially under the high thermal gradients and rapid solidification conditions characteristic of LPBF. Furthermore, the insufficient backfilling of liquid into the solidification front due to poor fluidity of Cu-rich melt may exacerbate crack propagation.

To gain a comprehensive understanding of how laser processing parameters affect the structural integrity of LPBF-fabricated Al_0.1CoCrFeNi + Cu composites, two data visualizations were constructed: a relative density plot as a function of VED (Fig. 3a) and a laser power–scan speed process map (Fig. 3b). As shown in Fig. 3a, the relative density of the printed parts generally increases with VED, reaching a plateau at approximately 200 J/mm³. This behavior suggests that energy input beyond this threshold does not significantly enhance densification, likely due to saturation in melt pool penetration and consolidation efficiency.

However, to disentangle the individual contributions of laser power and scan speed—factors that are otherwise convolved in the VED expression—the relative density data were further plotted as a two-variable process map in Fig. 3b. The results reveal a distinct trend: relative density decreases with increasing scan speed and decreasing laser power. Notably, at lower laser powers (100 W and 150 W), the relative density is highly sensitive to scan speed variations, indicating insufficient melting and reduced energy coupling at high scan velocities. In contrast, at higher laser powers (e.g., 250 W), the effect of scan speed on relative density becomes less pronounced, approaching an apparent insensitivity within the examined range.

This phenomenon can be partially explained by the laser absorptivity behavior of metallic powders under varying power inputs. As demonstrated by Trapp et al. [17], the laser absorptivity of stainless steel 316L powders increases exponentially from 0.3 to 0.8 as laser power increases from 50 W to 220 W. A similar effect is likely present in the Al_0.1CoCrFeNi + Cu system, where higher laser powers enhance powder-bed absorptivity, promoting more efficient energy coupling and stable melt pool formation. Consequently, this increased absorptivity at higher powers diminishes the dependence of melt pool behavior—and thus relative density—on scan speed. It also reduces the likelihood of LOF defects, which are more prevalent under conditions of poor absorptivity and shallow melting.

Despite these improvements, the maximum relative density achieved in this study was approximately 0.96, which remains below the industrially acceptable threshold (> 0.99) for critical structural applications. This shortfall is attributed not only to suboptimal melting dynamics but also to inherent issues in powder morphology, including incomplete packing and potential moisture content. Prior studies [18, 19] have shown that residual moisture in composite powders can lead to spattering and irregular melt pool behavior, ultimately increasing porosity. These results collectively emphasize the need for optimized powder handling and pre-processing protocols, in addition to refined laser parameters, to achieve full densification in LPBF-fabricated compositionally complex alloys.

3.2. Elemental Distribution Observation of Low and High VED Conditions

While the overall relative density of the LPBF-fabricated Al_0.1CoCrFeNi + Cu composites was found to be strongly influenced by VED and associated laser absorptivity behaviors, densification alone does not fully account for the nature and severity of microstructural defects. As previously discussed, LOF and keyhole-induced porosity were prevalent under low and high energy input conditions, respectively. However, beyond void morphology, the presence of these processing-induced artifacts may also be linked to elemental segregation, particularly in compositionally complex systems where multi-element interactions govern phase stability and local melting behavior. To further elucidate this relationship, EDS elemental mapping was performed on samples fabricated under the lowest and highest VED conditions. This comparative analysis aims to correlate the distribution of major alloying elements with the observed defect structures, providing deeper insight into the role of localized chemical inhomogeneity in the formation of process-induced artifacts.

To further understand the influence of laser energy input on the distribution of alloying elements and potential defect mechanisms, EDS elemental mapping was performed on specimens fabricated under the lowest and highest VED conditions. At the lowest VED setting (100 W, 800 mm/s), where the largest LOF defects and the lowest overall relative density were observed, significant Cu segregation was identified (Fig. 4). In contrast to the uniform distribution of the primary HEA-forming elements (Al, Co, Cr, Fe, Ni), Cu appeared concentrated in localized regions, often separated from the surrounding matrix. Partial co-segregation between Cu and Ni was also noted, which is thermodynamically plausible given the favorable mixing enthalpy between these two elements, as reported by Mirzababaei et al. [9]. Although X-ray diffraction analysis of the Cu-enriched phases within the Al_0.1CoCrFeNi matrix presents inherent challenges due to their limited volume fraction and potential overlap with matrix peaks, EDS clearly identified the presence of Cu-rich regions, consistent with prior observations reported by Lee et al. [20]. In that study, an increased Cu feed rate during the LPBF process led to a higher volume fraction of Cu-enriched islands within a stainless steel matrix, which was correlated with enhanced thermal conductivity. By analogy, the Cu-rich domains observed in Fig. 4 of the present work may similarly contribute to improved thermal transport, owing to the intrinsically high thermal conductivity of Cu and its localized enrichment within the composite microstructure.

The observed Cu segregation under low VED conditions is attributed primarily to insufficient melt pool residence time and reduced thermal enthalpy, both of which hinder the full dissolution and mixing of Cu into the HEA matrix. Trapp et al. [16] demonstrated that laser absorptivity—and therefore energy absorption—of metallic powders increases significantly with laser power. At lower power levels, limited absorptivity results in shallow and short-lived melt pools, which do not provide the thermal or temporal conditions required for effective diffusion and homogenization of dissimilar powder constituents. Additionally, as described by Afrasiabi et al. [11], inadequate energy input reduces melt pool convection and Marangoni flow, further suppressing elemental intermixing.

In contrast, under the highest VED condition (250 W, 200 mm/s), elemental mapping revealed a much more homogeneous distribution of Cu throughout the matrix (Fig. 5), suggesting effective melting, mixing, and solute incorporation due to deeper and more thermally stable melt pools. Importantly, while Cu segregation was evident at low VED, no direct correlation was observed between this segregation and the occurrence of solidification cracking. Cracks were detected even in regions where elemental distributions were homogenous, implying that cracking in the current Al_0.1CoCrFeNi + Cu system arises from mechanisms other than segregation-induced liquation.

One plausible explanation is the development of high thermal gradients and solidification-induced stresses intrinsic to LPBF processing of compositionally complex alloys. As highlighted by Guo et al. [15], cracking susceptibility in HEAs may also be governed by a combination of solidification shrinkage and grain boundary cohesion, independent of visible segregation. Platl et al. [21] similarly reported that intergranular cracking in FeCoMo alloys could occur due to intrinsic brittleness and residual stress development, even in the absence of any significant impurity phases or compositional heterogeneities. Additionally, as noted by Khodashenas and Mirzadeh [14], post-solidification tensile stresses generated during rapid thermal cycling in LPBF are sufficient to nucleate and propagate cracks in regions with unfavorable grain orientations or insufficient ductility at high cooling rates.

Now, Fig. 6a and 6b present EDS elemental maps highlighting Cu-enriched regions (defined as areas containing >50 wt% Cu) under two representative LPBF conditions: a high VED condition (200 W, 200 mm/s) and a low VED condition (150 W, 800 mm/s), respectively. In the low VED condition (Fig. 6a), extensive LOF cracking is observed in close proximity to Cu-segregated zones, suggesting that the presence of unmixed Cu phases contributes to localized mechanical incompatibility and crack initiation. In contrast, under the high VED condition (Fig. 6b), where cracking is narrower and less prevalent, there is neither a high density of Cu-segregated regions nor spatial correlation between Cu enrichment and crack location. This indicates that in the high VED regime, cracking is more likely driven by rapid solidification-induced embrittlement rather than compositional heterogeneity. A schematic summarizing the distinct cracking mechanisms under both processing conditions is provided in Fig. 6c.

Therefore, in the present alloy system, solidification cracking under high VED conditions is more likely attributed to the combined effects of rapid cooling, thermal stress accumulation, and insufficient strain accommodation at the solid–liquid interface, rather than elemental segregation. In contrast, at lower VED, pronounced Cu segregation is observed and appears to exacerbate crack formation, both in severity and frequency, by promoting localized mechanical and thermal mismatches during solidification. These results highlight the critical importance of optimizing thermal gradients and promoting grain refinement to mitigate cracking in LPBF-processed, compositionally complex alloys. Potential strategies to mitigate cracking frequency in LPBF-fabricated alloys include preheating the build substrate to reduce thermal gradients and suppress thermal shock during processing [22], modifying laser parameters—such as employing pulsed rather than continuous wave lasers to control energy input and cooling rates [23]—and incorporating strengthening dispersoids, such as oxide particles, to enhance crack resistance through microstructural stabilization [24].

5. Conclusion

In this study, Al_0.1CoCrFeNi + Cu composite were fabricated via LPBF using systematically varied laser powers (100–250 W) and scan speeds (200–800 mm/s) to explore the process-structure relationships governing densification, elemental distribution, and defect formation. The key findings and conclusions are summarized as follows:

1. Relative density improved with increasing VED, reaching a plateau near 200 J/mm³, while low VED led to pronounced LOF and Cu segregation due to insufficient melt pool residence time. In contrast, high-VED conditions promoted uniform Cu dissolution and suppressed segregation.

2. Cracking persisted even under homogeneous compositions at high VED, driven mainly by thermal gradients, shrinkage, and stress buildup. At low VED, Cu segregation further promoted cracking via local thermal and mechanical mismatch.

3. Suppressing defects in LPBF-fabricated HEA composites requires not only appropriate energy input but also strategies to refine grains, reduce thermal stress, and tailor alloying elements to minimize cracking susceptibility.

Article information

Funding

None.

Conflict of Interest

The authors have no conflicts of interest to declare.

Data Availability Statement

The data in this research will be available upon request to corresponding author.

Author Information and Contribution

First Author: PhD candidate; conceptualization, writing–original draft

Second Author: Professor; supervision, writing–review & editing

Third Author: Professor; writing–original draft, funding acquisition, supervision

Acknowledgements

The corresponding author gratefully acknowledges the invaluable research support provided by Prof. Brian K. Paul and Prof. Somayeh Pasebani. Their guidance and provision of access to the LPBF facilities were instrumental in the successful execution of this work. Their contributions significantly enhanced both the experimental design and the technical depth of this study.

Fig. 1.

Scanning electron micrograph with energy-dispersive X-ray spectroscopy elemental mapping of Al_0.1CoCrFeNiCu composite powder feedstock for laser powder bed fusion.

Fig. 2.

Cross-sectional view observation of fabricated composites at laser conditions of 100 – 250 W and 200 – 800 mm/s using optical microscopy, revealing artifacts such as lack of fusions, solidification cracks, and keyholes. The building direction is vertical to the image.

Fig. 3.

Relative density variation in laser powder bed fusion-printed Al_0.1CoCrFeNi + Cu composite with (a) volumetric energy density profile (b) a processing window of laser power and scan speed. The maximum relative density of the printed part is ≈ 0.96 where indicated with dashed lines.

Fig. 4.

Scanning electron microscopy–energy-dispersive X-ray spectroscopy mapping result for the lowest volumetric energy density condition (100 W, 800 mm/s), showing Cu segregation with some Ni concentration separated from other Al_0.1CoCrFeNi-driven elements.

Fig. 5.

Scanning electron microscopy–energy-dispersive X-ray spectroscopy mapping result for the highest volumetric energy density condition (250 W, 200 mm/s), showing the absence of Cu segregation.

Fig. 6.

Correlation between Cu segregation and solidification cracking behavior under different volumetric energy density (VED) conditions in laser powder bed fusion-processed Al_0.1CoCrFeNi + Cu composites. (a) EDS elemental map under a low VED condition (150 W, 800 mm/s). (b) EDS mapping under a high VED condition (200 W, 200 mm/s) (c) Schematic illustration comparing the dominant cracking mechanisms under low and high VED conditions. Red regions in (a) and (b) are where Cu concentration is higher than 50 wt%.

References

• 1. C.-J. Tong, Y.-L. Chen, J.-W. Yeh, S.-J. Lin, S.-K. Chen, T.-T. Shun, C.-H. Tsau and S.-Y. Chang: Metall. Mater. Trans. A, 36 (2005) 881.ArticlePDF
• 2. Y.-F. Kao, T.-J. Chen, S.-K. Chen and J.-W. Yeh: J. Alloys Compd., 488 (2009) 57.Article
• 3. W.-R. Wang, W.-L. Wang, S.-C. Wang, Y.-C. Tsai, C.-H. Lai and J.-W. Yeh: Intermetallics, 26 (2012) 44.Article
• 4. E. P. George, D. Raabe and R. O. Ritchie: Nat. Rev. Mater., 4 (2019) 515.ArticlePDF
• 5. Z. Li, A. Ludwig, A. Savan, H. Springer and D. Raabe: J. Mater. Res., 33 (2018) 3156.Article
• 6. W.-R. Wang, W.-L. Chen, S.-C. Wang, Y.-C. Tsai, C.-H. Lai and J.-W. Yeh: Intermetallics, 19 (2012) 698.
• 7. A. Sourav, S. Yebaji and S. Thangaraju: Mater. Sci. Eng. A, 793 (2020) 139877.Article
• 8. Z. Sun, C. Shi, L. Gao, S. Lin and W. Li: J. Alloys Compd., 901 (2022) 163554.Article
• 9. S. Mirzababaei, V. V. K. Doddapaneni, K. Lee, G. E. Paul, H. Pirgazi, K.-S. Tan, O. Ertorer, C.-H. Chang, B. K. Paul and S. Pasebani: Addit. Manuf., 70 (2023) 103576.Article
• 10. C. Li, M. Ferry, J. J. Kruzic and X. Li: J. Mater. Sci., 57 (2022) 9903.ArticlePDF
• 11. Q. K. Wang, L. F. Tang, Y. L. Bian, M. Majeed, H. W. Tang, Y. Cai, N. B. Zhang, L. Lu and S. N. Luo: Mater. Sci. Eng. A, 934 (2025) 148280.Article
• 12. M. Afrasiabi, D. Keller, C. Lüthi, M. Bambach and K. Wegener: Procedia CIRP, 113 (2022) 378.Article
• 13. W. E. King, H. D. Barth, V. M. Castillo, G. F. Gallegos, J. W. Gibbs, D. E. Hahn, C. Kamath and A. M. Rubenchik: J. Mater. Process. Technol., 214 (2014) 2915.Article
• 14. H. Khodashenas and H. Mirzadeh: J. Mater. Res. Technol., 21 (2022) 3795.Article
• 15. L. Guo, J. Gu, B. Gan, S. Ni, Z. Bi, Z. Wang and M. Song: J. Alloys Compd., 865 (2021) 158892.Article
• 16. Y. S. Lee and W. Zhang: Addit. Manuf., 12 (2016) 178.Article
• 17. J. Trapp, A. M. Rubenchik, G. Guss and M. J. Matthews: Appl. Mater. Today, 9 (2017) 341.Article
• 18. N. R. Comins: Philos. Mag., 25 (1972) 817.Article
• 19. S. Ly, S. A. Khairallah, A. M. Rubenchik, G. Guss and M. J. Matthews: Sci. Rep., 7 (2017) 4085.Article
• 20. K. Lee, V. V. K. Doddapaneni, S. Mirzababaei, S. Pasebani, C.-H. Chang and B. K. Paul: Addit. Manuf., 86 (2024) 104202.Article
• 21. J. Platl, D. Rainer, H. Leitner, C. Turk, F. Galbusera, A. G. Demir, B. Previtali and R. Schnitzer: Berg. Huettenmaenn. Monatsh., 167 (2022) 325.ArticlePDF
• 22. A. V. Müller, G. Schlick, R. Neu, C. Anstätt, T. Klimkait, J. Lee, B. Pascher, M. Schmitt and C. Seidel: Nucl. Mater. Ener., 19 (2019) 184.Article
• 23. H. Huang, B. Nie, P. Wan, L.-M. Yang, S. Bai and J. Liu: Proc. SPIE, 9353 (2015) 93530A.Article
• 24. Z. Hu, Y. Zhao, K. Guan, Z. Wang and Z. Ma: Addit. Manuf., 36 (2020) 101579.Article

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