Influence of pre-weld heat treatment temperatures and AlSi10Mg-Er-Zr filler powder on microstructure and mechanical properties of welded joints produced using laser metal deposition for L-PBF AlSi10Mg alloys

Influence of pre-weld heat treatment temperatures and AlSi10Mg-Er-Zr filler powder on microstructure and mechanical properties of welded joints produced using laser metal deposition for L-PBF AlSi10Mg alloys | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Research Article Influence of pre-weld heat treatment temperatures and AlSi10Mg-Er-Zr filler powder on microstructure and mechanical properties of welded joints produced using laser metal deposition for L-PBF AlSi10Mg alloys Yingying Liu, Jingchuan Li, Zhaotong Li, Li Cui, Dingyong He, and 1 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-7438037/v1 This work is licensed under a CC BY 4.0 License Status: Posted Version 1 posted You are reading this latest preprint version Abstract This study investigates the influence of pre-weld heat treatment (PWHT) temperatures (250 ~ 400 ℃) on laser metal deposition (LMD) welding of laser powder bed fusion (L-PBF) AlSi10Mg alloys using novel AlSi10Mg-Er-Zr filler powders, with particular emphasis on porosity characteristics, mechanical performance, and microstructural evolution of the welded joints. Results show that elevated PWHT temperatures effectively mitigated hydrogen porosity, reducing both pore density and maximum diameter. The synergistic combination of low PWHT temperatures at 250 ~ 280 ℃ using AlSi10Mg-Er-Zr filler powder demonstrated superior mechanical enhancement, achieving 4.3 ~ 7.2% improvement in ultimate tensile strength (UTS) and 3.8 ~ 94.6% increase in elongation at fracture (EF) compared to the conventional AlSi10Mg filler powder. Welded joints produced using AlSi10Mg-Er-Zr filler powder at 280 ℃ PWHT yielded the optimal mechanical performance, achieving a remarkable balance between strength and ductility with an UTS of 259.0 MPa and EF reaching 10.9%. The synergistic combination of AlSi10Mg-Er-Zr filler powder and PWHT not only enhanced microstructure improvement through substantial grain refinement, increased proportions of high-angle grain boundaries (HAGBs), Σ3 grain boundaries, and hard-oriented grains, but also reduced hydrogen porosity in the welds, collectively contributing to the superior mechanical performance of L-PBF AlSi10Mg welded joints. L-PBF AlSi10Mg alloys Laser metal deposition (LMD) welding Pre-weld heat treatment (PWHT) AlSi10Mg-Er-Zr filler powder hydrogen porosity microstructure improvement Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Figure 6 Figure 7 Figure 8 Figure 9 Figure 10 Figure 11 Figure 12 Figure 13 Figure 14 1. Introduction AlSi10Mg alloys fabricated via the laser powder bed fusion (L-PBF) process have gained significant attention in the aerospace, aeronautics, and automotive industries owing to their excellent printability, lightweight properties, and relatively low production costs [ 1 , 2 ]. Nevertheless, the production of complex and large-scale components using the L-PBF process remains impractical due to limitations in build chamber size [ 3 , 4 ]. To overcome the size limitations of L-PBF components, welding techniques have been increasingly employed to join additively manufactured parts with other L-PBF components or with conventionally fabricated structures [ 5 – 7 ]. This approach effectively circumvents the build volume constraints inherent to the L-PBF process while maintaining its advantages. However, the weldability of L-PBF AlSi10Mg alloys using fusion welding techniques has been limitedly reported in literature. It was found that fusion welding of L-PBF AlSi10Mg alloys remains particularly challenging [ 8 , 9 ], primarily due to the significantly higher hydrogen pore susceptibility generated in the welds compared to casting materials of the same alloy. The extensive porosity was directly attributable to the melting of the base metal (BM) during the welding, releasing very high content of hydrogen gas and entering the welding pool [ 10 – 12 ], which can be the hydrogen source for the nucleation of hydrogen pores. As a result, lowering the weld dilution reduces hydrogen pickup from the BM, thereby mitigating porosity. In this regard, some special welding processes, such as high-pressure or high-vacuum laser welding [ 13 , 14 ], laser metal deposition (LMD) process [ 15 – 17 ] were recently introduced to reduce the hydrogen porosity when welding of L-PBF AlSi10Mg alloys. However, the hydrogen porosity of the welds in L-PBF AlSi10Mg alloys is still higher than compared to conventional Al-based alloys, and it is unlikely that the porosity can be entirely avoided by optimization of the welding processes only. Pre-weld heat treatment (PWHT) is a widely adopted process during welding of metal and its alloys, serving to relieve residual stresses, mitigate porosity and crack formation, and consequently enhance microstructural and mechanical properties. However, research on the influence of PWHT on weldability of L-PBF AlSi10Mg alloys remains limited. Mäkikangas et al. implemented stress relief annealing at 300℃ before welding [ 8 ]. Although the study highlighted significant reduction of pores in the laser-welded AlSi10Mg alloys, the PWHT resulted in a strength reduction of approximately 19 MPa compared to non-heat-treated conditions. Chen et al. investigated the influence of solution heat treatment (SHT) at 520 ℃ on pore distribution, microstructural evolution, and mechanical properties of laser-welded joints [ 18 ]. Their findings revealed that vacuum SHT effectively reduced weld porosity to 0.14%. However, the heat treatment resulted in a substantial reduction in tensile strength, decreasing to 143 MPa, which was significantly lower than that observed in non-heat-treated joints. This strength deterioration primarily stemmed from microstructural changes, as elevated annealing temperatures and extended processing times promoted coarsening of the eutectic Si-rich phases while simultaneously decreasing their number density [ 19 ]. The aforementioned studies demonstrate that pre-weld vacuum SHT effectively minimizes porosity in laser welded L-PBF AlSi10Mg alloys, whereas the PWHT at 520 ℃ substantially degrades joint strength. Notably, a desired microstructure can be achieved using lower temperatures and shorter durations than those prescribed by ASTM T6 standards [ 20 ], suggesting that microstructural optimization in welded L-PBF AlSi10Mg alloys is feasible through controlled PWHT at lower temperatures. To bridge this knowledge gap, the present study systematically explores PWHT at temperatures below 520 ℃, with the dual objectives of significantly suppressing hydrogen porosity while preserving the tensile strength of the welded joints. The weld composition in welded joints is typically designed to match the base metal (BM), but may be intentionally modified to achieve compatible properties. Recent advances have shown that, the principle of grain refinement through inoculate addition with scandium (Sc) and zirconium (Zr) element, well-established in aluminum alloy casting [ 21 ], has recently been extended to additive manufacturing research. While the Sc and Zr combination has demonstrated significant grain-refining effects [ 22 , 23 ], erbium (Er) has emerged as a cost-effective rare earth alternative, exhibiting microalloying effects comparable to Sc. Numerous studies have confirmed Er's strengthening potential in aluminum alloys [ 24 ]. Our prior research revealed that minor additions of Er and Zr substantially enhance the mechanical performance of both L-PBF AlSi10Mg alloys and welded L-PBF AlSi10Mg alloys [ 25 , 26 ]. This improvement stems from the formation of nanoscale coherent Al₃Er and Al₃(Zr,Er) precipitates, achieving an exceptional strength-ductility balance. However, the synergistic effects of Er and Zr microalloying and PWHT on microstructural evolution and mechanical properties in welded L-PBF AlSi10Mg remain insufficiently understood. This study systematically investigates the effects of low PWHT temperatures (250–400 ℃) and AlSi10Mg-Er-Zr filler powder on the LMD welded L-PBF AlSi10Mg alloys. The research methodology involved conducting vacuum PWHT on L-PBF AlSi10Mg sheets at four specific temperatures (250 ℃, 280 ℃, 300 ℃, and 400 ℃) followed by LMD welding in a butt joint configuration using both conventional AlSi10Mg and the prepared AlSi10Mg-Er-Zr filler powders. Through comprehensive characterization of porosity characteristics, microstructural evolution, and mechanical properties of the welded joints, this work provides fundamental insights into the process-microstructure-property relationships, establishing a scientific foundation for optimizing welding and repair techniques for additively manufactured aluminum components in industrial applications. 2. Materials and experimental method 2.1 Materials The BM specimens for welding were fabricated as 50 × 25 × 3.0 mm sheets using an EOS M280 system in a high-purity argon atmosphere. The process employed gas-atomized AlSi10Mg powder (15–53 µm size) deposited on a 150 ℃ preheated aluminum substrate. Throughout the L-PBF process, the build chamber maintained oxygen levels below 200 ppm under continuous argon protection. To optimize the manufacturing process, each successive laser scanning path was rotated by 67° relative to the previous track. Through orthogonal experimentation, the optimal processing parameters were determined to be a laser power of 370 W, a scanning speed of 1300 mm/s, a hatching space of 0.19 mm, and a track thickness of 20 µm, which yielded specimens with high relative density. Density measurements were performed using both Archimedes' principle and ImageJ-based image analysis. Only L-PBF AlSi10Mg sheets demonstrating relative densities exceeding 99.5% were selected as BM specimens for subsequent LMD welding. 2.2 Welding experimental PWHT was performed on the BM specimens in a PZTH-150 vacuum furnace prior to welding. Four distinct heat treatment temperatures (250 ℃, 280 ℃, 300 ℃, and 400 ℃) were applied with a 2-hour holding time at each temperature, employing a controlled heating rate of 10 ℃/min followed by air cooling to ambient temperature. The resultant surface oxide films were subsequently removed through sequential mechanical grinding and chemical cleaning procedures. The LMD welding was conducted using a 4-kW TruDisk laser system (TRUMPF) integrated with an ABB 6 -axis robotic manipulator and a DPSF-2 automated powder feeder. Two variants of filler powders were utilized: commercially available gas-atomized AlSi10Mg powder and specially formulated AlSi10Mg-Er-Zr powder, both sieved to a particle size range of 45–105 µm. The AlSi10Mg-Er-Zr powder was produced via in-situ alloying during gas atomization, incorporating master alloys of Al-Mg, Al-Er, and Al-Zr with the AlSi10Mg powder. High-purity argon (99.99%) served as the shielding gas during welding, with precisely controlled flow rates of 2.5 L/min for powder delivery and 25.0 L/min for molten pool protection. The detailed chemical compositions of both the BM and filler powders are systematically presented in Table 1 . Table 1 Chemical compositions of the BM and filler powder (wt.%) Fe Mg Mn Si Zn Ti Er Zr Al Base metal 0.15 0.39 ≤ 0.01 10.51 ≤ 0.01 0.14 - - Bal. AlSi10Mg 0.20 0.30 ≤ 0.01 10.53 ≤ 0.01 0.12 - - Bal. AlSi10Mg-Er-Zr 0.11 0.40 ≤ 0.01 10.05 ≤ 0.01 0.11 0.65 0.23 Bal. Full-penetration butt joints were obtained through five-track LMD processes using both AlSi10Mg and AlSi10Mg-Er-Zr filler powders. The BM specimens were prepared in a 40° V-groove configuration without root gap, as shown in Fig. 1 . During the LMD welding, the filler powder was continuously and uniformly fed into the groove, where it was melted and sequentially deposited track-by-track to form the weld joint. The welding parameters were optimized through a series of preliminary trials, maintaining constant values for laser power of 1500 W, focus spot diameter of 2.2 mm, defocus distance of 0 mm, and powder feeding rate of 3.15 g/min. The travel speed was set at 10 mm/s for the first deposition track and increased to 15 mm/s for the subsequent four tracks. This variation in travel speeds produced different linear heat inputs, with 150 J/mm for the initial track and 100 J/mm for the remaining tracks. 2.3 Microstructure, porosity characterization and mechanical testing Following LMD welding, cross-sectional specimens were extracted from the welded joints for metallographic characterization. The preparation sequence consisted of mounting, mechanical polishing, and chemical etching with Keller's reagent to delineate bead morphology and microstructural features. Microstructural analysis and pore characterization were performed using an Olympus LEXT OLS4100 laser scanning confocal microscope, with quantitative measurements of dendrite arm spacing, maximum pore diameter, and porosity percentage conducted using Image Pro Plus software. For higher-resolution examination, Si-rich eutectic morphology was investigated using a QUANTA FEG650 field-emission scanning electron microscope (FE-SEM) equipped with energy dispersive spectroscopy (EDS). Fractographic analysis of failed specimens was carried out employing both a Nikon SMZ800 stereomicroscope and a JEOL JSM-7400F high-resolution SEM. Electron backscatter diffraction (EBSD) analysis was conducted using a Hitachi S-3400N thermal field-emission SEM integrated with an Oxford Instruments Nordlys Nano detector system, operating at 20.0 kV. The EBSD specimens were sequentially prepared by mechanical polishing and subsequent electropolishing in a 30% nitric acid-methanol solution at -25 ℃, using an applied voltage of 20 V for 30 seconds. Orientation mapping was acquired in rectangular scan areas with a step size of 1.4 µm. All EBSD data were processed using HKL Technology's Channel 5 software. Microhardness measurements were performed using an HVS-1000 Vickers hardness tester with a 98 N load applied for 15 seconds. The reported hardness values represent the average of at least five indentations per test condition. Tensile specimens were machined according to ASTM E8-04 specifications and tested at room temperature using an Instron universal testing machine with a constant crosshead speed of 0.5 mm/min. To ensure statistical reliability, three specimens were tested for each joint condition, with the average value reported as the representative result. 3.Results 3.1 Weld shape and porosity characteristics Figure 2 presents cross-sectional views of the resulting welded joints produced using LMD welding with both AlSi10Mg and AlSi10Mg-Er-Zr filler powders for the BM subjected to 250 ℃, 280 ℃, 300 ℃, and 400 ℃ PWHT. The BM microstructure displayed typical columnar grain structures with widths varying between 20–200 µm and diverse melt pool dimensions, aligning with established literature findings [ 4 , 27 ]. However, microstructural examination at this magnification revealed no heat-affected zone (HAZ), as indicated by the identical melt pool morphology between the welded region and unaffected BM. All joints achieved full penetration with comparable geometries, indicating that neither PWHT temperature variations nor filler powder selection significantly affected weld shape. Additionally, microstructural analysis revealed distinct welding modes between different regions of the welds. The bottom region (first deposition track) displayed a characteristic keyhole mode morphology with a high depth-to-width ratio, while the upper region (tracks 2–5) exhibited shallower, hemispherical penetration profiles typical of conduction mode welding. Regarding the porosity defects, as shown in Fig. 3 , the welds exhibited characteristic spherical porosity in both upper and bottom regions. The porosity measurements demonstrated a strong temperature dependence for AlSi10Mg welds, decreasing progressively from 3.2% at 250 ℃ PWHT to 0.7% at 400 ℃ PWHT with increasing PWHT temperatures. Additionally, a substantial 52.7% reduction in maximum pore diameter was observed, decreasing from 276.7 µm to 131.0 µm over the investigated temperature range. However, the PWHT at elevated temperatures of 300 ℃ and 400 ℃ substantially decreased weld porosity compared to lower temperature treatments of 250 ℃ and 280 ℃. Moreover, quantitative analysis revealed the upper region possessed significantly lower porosity and smaller hydrogen pore diameters compared to the bottom region. This phenomenon results from the lower heat input of 100 J/mm in the upper layers relative to the 150 J/mm in the bottom layer, reducing the dilution rate and thereby restricting hydrogen absorption into the weld pool [ 12 , 15 ]. Regarding the influence of filler powders, the AlSi10Mg-Er-Zr welds exhibited reduced porosity values of 3.0% at 250 ℃ and 2.7% at 280℃, representing a slight decrease compared to the conventional AlSi10Mg welds. As a result, the PWHT temperature emerged as the dominant factor controlling hydrogen pore formation of the welds, whereas the addition of Er and Zr to the AlSi10Mg filler powder exhibited minimal impact on porosity when PWHT temperature remained constant. 3.2 Mechanical properties of welded joints 3.2.1 Microhardness distribution The microhardness distribution across various regions of the welded joints produced with four different PWHT temperatures, including the BM, HAZ, fusion boundary, and weld metal, is presented in Fig. 4 . The as-built BM exhibited a hardness of approximately 120 HV, resulting from the fine microstructure developed during the L-PBF process [ 1 , 2 ]. Following PWHT, the BM showed a progressive decrease in average hardness with increasing temperatures from 250 ℃ to 400 ℃. The measured hardness values were 89.2 HV at 250 ℃, decreasing to 82.1 HV at 280 ℃, 75.9 HV at 300 ℃, and reaching a minimum of 61.9 HV at 400 ℃. This temperature-dependent softening behavior can be attributed to microstructural changes in the Si-rich eutectic structure. As reported in previous studies, increasing the annealing temperature or duration leads to the breakdown of the cellular Si network structure and promotes Si precipitation in the α-Al matrix, thereby reducing the solid solution strengthening effect [ 28 , 29 ]. While the network structure remains largely intact at 260℃, complete decomposition occurs at 300 ℃ [ 30 ]. These microstructural transformations explain the significant hardness reduction observed in BM samples subjected to PWHT at 400 ℃. The pronounced BM softening led to different hardness distribution patterns in welded joints treated at different temperatures. In the bottom region, as shown in Fig. 4 (a), the welded joints subjected to 280 ℃, 300 ℃, and 400 ℃ PWHT developed a distinct inverted U-shaped hardness profile, while the welded joints treated at the lowest temperature of 250 ℃ maintained an U-shaped hardness distribution using whether AlSi10Mg powder or AlSi10Mg-Er-Zr one. This difference in the hardness profile of the welded joints subjected to PWHT between the lowest temperature at 250 ℃ and the higher temperatures was mainly attributable to the different degree of BM softening. In the upper regions, all the welded joints consistently showed inverted U-shaped profiles regardless of PWHT temperatures, since the hardness in the WM consistently surpassed BM values, as shown in Fig. 4 (b). In addition, detailed measurements revealed hardness ranges of 76.2–87.8 HV in the bottom region and 89.2-103.7 HV in the upper region, indicating systematically higher hardness values in the upper regions. Notably, peak hardness of 103.7 HV was achieved in the upper region of the welds treated at 280℃ with AlSi10Mg-Er-Zr filler powder. It also can be seen from Fig. 4 that a sharp hardness reduction in the HAZ along the transition from the BM to the WM. However, the location of minimum hardness differed significantly between lower-temperature at 250℃, 280 ℃, 300 ℃ PWHT and higher-temperature at 400 ℃ PWHT for both filler powders. The hardness minimum consistently occurred at the fusion boundary for joints subjected to 250 ℃, 280 ℃, and 300 ℃ PWHT, while the lowest hardness values occurred within the BM for the joints at 400 ℃ PWHT, irrespective of AlSi10Mg or AlSi10Mg-Er-Zr filler powders. 3.2.2 Tensile property of the welded joints The tensile testing was conducted on the welded joints subjected to 250 ℃, 280 ℃, 300 ℃, and 400 ℃ PWHT with AlSi10Mg and AlSi10Mg-Er-Zr filler powders. The stress-strain curves of these joint specimens are presented in Fig. 5 , with the corresponding tensile test results summarized in Table 2 . The UTS values of the welded joints showed clear variations with PWHT temperatures, measuring 230.3 MPa at 250 ℃, 241.5 MPa at 280 ℃, 231.1 MPa at 300 ℃, and 175.5 MPa at 400 ℃. This demonstrates that joints treated at 250–300 ℃ maintained relatively high strength levels between 230–240 MPa, but those subjected to 400 ℃ PWHT experienced a significant reduction in UTS. The highest UTS of 259.0 MPa was achieved at 280℃ PWHT combined with AlSi10Mg-Er-Zr filler powder. Specially, the elongation at fracture (EF) exhibited a progressive increase with rising PWHT temperatures, ranging from 5.3% at 250 ℃ to 19.1% at 400 ℃, representing a substantial 3.6-fold improvement in ductility. This highlights the important influence of PWHT temperatures on the ductility of the welded joints in L-PBF AlSi10Mg alloys. Table 2 Tensile properties of the welded joints at various PWHT temperatures with AlSi10Mg and AlSi10Mg-Er-Zr filler powder Welded joints UTS /MPa EF /% Welded joints UTS /MPa EF /% 250 ℃ 230.3 5.3 250°C + Er/Zr 240.1 5.5 280 ℃ 241.5 5.6 280°C + Er/Zr 259.0 10.9 300 ℃ 231.1 9.0 - - - 400 ℃ 175.5 19.1 - - - The welded joints produced with AlSi10Mg-Er-Zr filler powder demonstrated markedly enhanced tensile properties compared to those employing AlSi10Mg powder. These joints at 280 ℃ PWHT exhibited a 7.2% increase in UTS accompanied by a remarkable 94.6% improvement in EF. The optimal mechanical performance was attained in joints treated with both 280 ℃ PWHT and AlSi10Mg-Er-Zr filler powder, achieving peak values of 259.0 MPa UTS and favorable 10.9% elongation. These property enhancements stem from the synergistic interaction between PWHT and the AlSi10Mg-Er-Zr filler powder, which promoted favorable microstructural alterations in the welds of L-PBF AlSi10Mg alloys. 3.2.3 Fracture behavior of the welded joints The welded joints subjected to 250 ℃, 280 ℃, and 300 ℃ PWHT consistently fractured near the fusion boundary, while those treated at 400 ℃ PWHT fractured in the BM, in agreement with the corresponding microhardness profiles shown in Fig. 6 . Tensile testing results showed distinct fracture behaviors based on PWHT temperatures. The fracture surfaces of failed joints treated between 250 ℃ and 300 ℃ displayed similar morphological characteristics, showing neither visible necking nor pronounced plastic deformation. These surfaces revealed distinct hydrogen porosity features with maximum pore diameters under 150 µm, as illustrated in Fig. 6 (a)-(e). Fracture mechanisms in these joints were primarily controlled by microstructural factors and pre-existing porosity, depending on the pore size distribution and morphological characteristics. Of particular significance, welded joints treated at 280 ℃ PWHT with AlSi10Mg-Er-Zr filler powder exhibited superior fracture surface quality, characterized by reduced porosity density and diminished presence of large pores, evident in Fig. 6 (d). In contrast, the joints treated at 400 ℃ PWHT presented fundamentally different fracture characteristics, shown in Fig. 6 (f). These surfaces completely lacked hydrogen pores while demonstrating unambiguous macroscopic necking, providing definitive evidence of considerable plastic deformation preceding final fracture. Higher-magnification SEM images of fracture surfaces of the welded joints treated at 280 ℃ and 400 ℃ PWHT are shown in Fig. 6 (g) and Fig. 6 (h), respectively. The welded joints at 280℃ PWHT exhibited numerous spherical pores with 10–20 µm in diameter surrounded by elongated dimples on their fracture surfaces. These pores promoted local brittle fracture behavior through stress concentration during loading, initiating microcracks that preferentially formed at pore sites under maximum stress and propagated to adjacent pores [ 9 ]. The significant hydrogen porosity also reduced the effective load-bearing cross-section, creating localized stress concentrations that accelerated crack propagation [ 27 ] and consequently diminished both strength and ductility. While the influence of initial porosity on damage mechanisms in L-PBF AlSi10Mg alloys and their welded joints remains incompletely understood, current evidence suggests that it depends critically on pore size and morphology [ 4 ]. In contrast, the welded joints subjected to 400 ℃ PWHT exhibited completely different fracture morphology, characterized by uniformly distributed large equiaxed dimples covering the entire fracture surface, which clearly confirmed highly ductile fracture behavior. 3.3 PWHT influence on joint microstructures 3.3.1 Si-rich eutectic of the welds The critical influence of silicon (Si) element on both the printability and mechanical properties of L-PBF AlSi10Mg alloys has been well documented [ 2 , 4 ]. In as-built state, the BM displays a homogeneous distribution of continuous Si-rich eutectic networks in aluminum (Al) matrix, featuring fine α-Al cells measuring 0.5 ~ 1 µm in size. Following welding, the welds revealed distinct modifications to the Si-rich eutectic networks across the different PWHT temperatures, as shown in Fig. 7 . It can be observed that both the bottom and upper regions of the welds were composed of α-Al phases and Si-rich eutectic phases, maintaining the same phase composition as the BM. However, differences in the morphology and connectivity of the Si-rich eutectic between the bottom region and the upper region was significant. As shown in Fig. 7 (a)-(d), the bottom region subjected to all the PWHT temperatures of 250 ℃~400 ℃ contained plate-like or rod-like Si-rich eutectic morphology, demonstrating significant breakdown of the original cellular networks. In contrast, the upper regions shown in Fig. 7 (a')-(d') preserved well-interconnected Si-rich networks, displaying finer morphological features and enhanced continuity relative to the bottom regions. This microstructural variation stems from different cooling conditions during the LMD welding owing to an increased travel speed of 15 mm/s in the upper region compared to 10 mm/s in the bottom region. As documented in previous studies [ 2 , 31 , 32 ], such higher welding speeds accelerated cooling promotes refinement of Si-rich networks in L-PBF AlSi10Mg alloys. High-resolution characterization revealed that α-Al cells in the upper region retained an elliptical morphology with characteristic dimensions of 2–6 µm, irrespective of applied PWHT temperatures. Conversely, the bottom region exhibited fragmented Si-rich networks featuring plate-like or rod-like structures across all PWHT temperatures. Remarkably, the Si-rich eutectic morphology in both regions demonstrated minimal dependence on PWHT temperature variations. These observations collectively suggest that welding parameters, rather than PWHT, served as the dominant factor controlling microstructural evolution in both the bottom and upper regions. 3.3.2 Grain size and grain boundary feature of the welds Figure 8 presents the inverse pole figure (IPF) maps and grain size distribution of α-Al grains in the upper region of welds produced with AlSi10Mg filler powder subjected to 280 ℃ and 400 ℃ PWHT. As shown in Fig. 8 (a), the weld microstructure at 280 ℃ PWHT primarily consisted of randomly oriented equiaxed and ultrafine grains. In contrast, the weld treated at 400 ℃ exhibited predominantly columnar grains, as seen in Fig. 8 (b). The average equivalent grain sizes were 84.2 µm and 116.3 µm for the welds at 280 ℃ and 400 ℃ PWHT, respectively. Additionally, the weld treated at 280 ℃ contained a significantly higher fraction of fine grains below 50 µm, measuring 81.9%, compared to 74.2% in the weld at 400 ℃ PWHT. This grain size distribution clearly demonstrates the superior grain refinement efficacy of lower temperature PWHT at 280 ℃ relative to higher temperature PWHT at 400 ℃. Figure 9 displays the grain boundary distribution and misorientation angle histograms for welds subjected to PWHT at 280 ℃ and 400 ℃, where high-angle grain boundaries (HAGBs, > 15°) and low-angle grain boundaries (LAGBs, 2°–15°) are marked by blue and red lines, respectively. The weld at 280 ℃ PWHT, as shown in Fig. 9 (a), and Fig. 9 (b), revealed a predominance of HAGBs of 54.0%. In contrast, the 400 ℃ PWHT weld in Fig. 9 (c) and Fig. 9 (d) exhibited a higher concentration of LAGBs, resulting in a reduced HAGB fraction of only 40.8%. This microstructural evolution was further quantified through misorientation angle analysis, where the average angle decreased significantly from 29.6° for the 280 ℃ PWHT to 19.2° at 400 ℃ PWHT, demonstrating a pronounced temperature dependence of grain boundary characteristics. The increased HAGB fraction in the weld at 280 ℃ PWHT can be attributed to the prevalence of finely equiaxed grains, suggesting that lower-temperature PWHT promotes the formation of such boundaries. This is particularly noteworthy since HAGBs are known to effectively hinder crack propagation [ 21 , 33 ], which explains one of the reasons why the enhanced mechanical properties observed in the weld at 280 ℃ PWHT compared to those treated at 400 ℃. 3.4 Role of Er and Zr addition in the welded joints 3.4.1 Microstructure characteristics at fusion boundary As demonstrated in previous results, all welded joints subjected to 250 ℃, 280 ℃ and 300 ℃ PWHT consistently fractured near the fusion boundary within the weld, confirming this area as the joint's mechanically weakest region. Figure 10 provides a comprehensive microstructural analysis of the fusion boundary in welded joints produced with both AlSi10Mg and AlSi10Mg-Er-Zr filler powders at 280 ℃ PWHT. The microstructural examination revealed distinct features along the fusion boundary of the welds produced with both AlSi10Mg and AlSi10Mg-Er-Zr filler powders. Contrary to conventional expectations, epitaxial growth from the BM was not observed. Instead, a well-defined narrow band of non-dendritic equiaxed grains (EQZ) formed consistently along the fusion boundary, with measured widths ranging from 40.2 to 78.1 µm for both welds with the different filler powders, as shown in Fig. 10 (a) and Fig. 10 (b). This EQZ formation resulted primarily from heterogeneous nucleation on Al 3 Ti, Al 3 Er, and Al 3 (Er, Zr) particles originating from both the BM and filler powders [ 26 ]. The comparable EQZ widths between both filler powders was similar, indicating that Er and Zr additions primarily influenced microstructural refinement rather than EQZ formation itself. Beyond the EQZ region, both welds close to the EQZ exhibited comparable microstructural characteristics, featuring cellular dendritic and equiaxed structures. However, the welds produced with AlSi10Mg-Er-Zr filler powder displayed significantly refined dendritic structures near the EQZ boundary, clearly demonstrating the grain-refining capability of Er and Zr additions. More importantly, welds produced with the AlSi10Mg-Er-Zr filler powder showed a substantial reduction in both porosity density and pore size distribution along the fusion boundary. These synergistic microstructural enhancements, which include improved grain refinement at the fusion boundary and superior pore morphology control, collectively partially contributed to the observed increase in tensile strength of the welded joints produced using the AlSi10Mg-Er-Zr filler powder. 3.4.2 Optical microstructure of the welds The influence of Er and Zr additions on Si-rich eutectic of the welds was investigated comparing the eutectic morphologies in the bottom and upper regions obtained with AlSi10Mg and AlSi10Mg-Er-Zr filler powders at 280°C PWHT. As shown in Fig. 11 , a clear morphological difference in the Si-rich eutectic phases was observed between the upper and lower regions. In the bottom region, most of the eutectic Si-rich networks were fragmented in the Al matrix due to thermal accumulation experienced from the 2nd to 5th tracks in the upper region. The fragmented Si-rich phases predominantly exhibited coarse rod-like or bar-like structures, with a minor fraction displaying particulate morphology. In contrast, the upper region exhibited a well-developed cellular dendrite consisting of large α-Al cells and finer Si-rich eutectic. The Si-rich eutectic in this region appeared as significantly refined particulates distributed along the α-Al dendrites. Consequently, both the Si-rich eutectic and α-Al cells underwent substantial refinement in the upper region compared to the bottom region of the welds. Notably, the addition of Er and Zr in the AlSi10Mg powders had a limited influence on the cellular dendritic structures in the welds. Both in the bottom and upper regions, the microstructures produced using AlSi10Mg filler powder were remarkably similar to those obtained with AlSi10Mg-Er-Zr filler powders. This observation indicated that regional variations in welding heat input exerted a more significant influence on microstructure evolution than did the composition of the filler powder. However, the secondary dendrite arm spacing (SDAS) in the upper region produced with AlSi10Mg-Er-Zr powder was smaller than those made with AlSi10Mg powder, indicating that the addition of Er and Zr in the filler powder promoted dendritic refinement under identical welding heat input conditions. 3.4.3 Si-rich eutectic feature in the welds High-magnification SEM micrographs were analyzed to compare the Si-rich eutectic structures in welds produced at 280°C PWHT using AlSi10Mg and AlSi10Mg-Er-Zr filler powders, as shown in Fig. 12 . The bottom regions of both welds exhibited similar morphologies of coarse rod-like and plate-like Si-rich eutectic structures, confirming the optical observations. However, the Si-rich eutectic revealed markedly different spacing characteristics in both the bottom and upper regions. In the bottom regions, welds produced with AlSi10Mg filler powder showed significantly coarser spacing compared to those produced with AlSi10Mg-Er-Zr filler powder. While the welds produced with AlSi10Mg filler powder produced an average spacing of 10.1 µm, the welds with AlSi10Mg-Er-Zr filler powder resulted in substantially finer spacing of 5.9 µm. In the upper regions, both filler powders formed continuous Si-rich eutectic networks, presenting a distinct contrast to the bottom region's microstructure, proving the optical observational results. Corresponding α-Al cell sizes measured 6.1 µm for AlSi10Mg powder and 4.2 µm for AlSi10Mg-Er-Zr filler. Thus, the α-Al cell sizes followed a similar refinement trend, with the AlSi10Mg-Er-Zr filler powder producing finer cellular structures. Therefore, it is clear that the AlSi10Mg-Er-Zr filler powder effectively refined the α-Al and Si-rich eutectic phases in both the bottom region and the upper regions of the welds. Notably, the Si-rich eutectic particles were predominantly distributed along α-Al cell boundaries, forming three-dimensional networks characteristic of L-PBF AlSi10Mg alloys [ 20 , 28 ]. These results demonstrate that while Er and Zr additions effectively refined both the Si-rich eutectic and α-Al structure, they did not substantially alter the fundamental morphology or connectivity of the Si-rich eutectic within corresponding weld regions. 3.4.4 Grain boundary characteristics The grain boundary distribution of α-Al structures in the welds produced using AlSi10Mg and AlSi10Mg-Er-Zr filler powders was analyzed. The grain boundary distribution and misorientation angle histograms for the upper region of both welds subjected to 280 ℃ PWHT are presented in Fig. 13 , where the low-angle grain boundaries (LAGBs, 2°–15°) were marked in red, and high-angle grain boundaries (HAGBs, 15°–65°) were highlighted in blue. As shown in Fig. 13 (a) and Fig. 13 (b), both welds exhibit a predominance of blue-colored boundaries, indicating that HAGBs dominate the microstructure. However, the fraction of HAGBs differs between the two welds: 53.9% for the AlSi10Mg weld and 61.3% for the AlSi10Mg-Er-Zr weld, demonstrating an increase in HAGBs for the welds produced with AlSi10Mg-Er-Zr filler powders. Furthermore, the average grain misorientation angles were 29.6° and 33.8° for the AlSi10Mg and AlSi10Mg-Er-Zr welds, respectively. These results clearly indicate that the AlSi10Mg-Er-Zr filler powder promotes higher grain boundary misorientation angles. Recent studies have recognized grain boundary engineering (GBE) as a promising strategy for improving the properties of LPBF components [ 34 ]. Among various grain boundary characteristics, coincidence site lattice (CSL) boundaries are particularly noteworthy due to their substantial impact on the mechanical performance of polycrystalline materials [ 35 ]. As shown in Fig. 13 (c) and Fig. 13 (d), both welds exhibited a pronounced peak at approximately 60°, corresponding to the Σ3 (60° ) grain boundary, a common CSL boundary in face-centered cubic (FCC) aluminum alloys. This indicates an increased fraction of CSL boundaries in the welds using both AlSi10Mg and AlSi10Mg-Er-Zr filler powders. Quantitative analysis revealed that the Σ3 boundary fraction reached 17.2% in the AlSi10Mg weld and 24.9% in the AlSi10Mg-Er-Zr weld, demonstrating a significant enhancement of Σ3 grain boundaries in the welds as a result of the addition of Er and Zr in the AlSi10Mg filler powder. 3.4.5 Schmid factor The Schmid factor (SF) is a crucial parameter governing slip system activation during tensile loading, with lower SF values indicating greater resistance to dislocation motion [ 36 ]. Generally, grains with higher SF values possess softer orientations that promote crack propagation [ 37 ]. Figure 14 presents the SF distributions for welds produced using AlSi10Mg and AlSi10Mg-Er-Zr filler powders, showing a predominant concentration of SF values in the range of 0.3–0.5. Detailed microstructural analysis demonstrated a marked difference in grain orientation characteristics of the welds between the AlSi10Mg and AlSi10Mg-Er-Zr filler powders. In the AlSi10Mg welds, merely 14.1% of grains exhibited SFs within the low range of 0.3 to 0.4. This value underwent a remarkable increase to 34.8% when examining welds produced with the AlSi10Mg-Er-Zr filler powder. Such a pronounced enhancement in the population of grains with low SFs provided conclusive evidence that the addition of Er and Zr in the AlSi10Mg filler powder during LMD welding effectively promotes the development of crystallographically hard-oriented grains in the welds. The increased population of hard-oriented grains enhances both load-bearing capacity through improved stress distribution and resistance to slip system activation during deformation. 4. Conclusions Butt joints of 3.0 mm thick sheets of L-PBF AlSi10Mg alloys have been produced using LMD welding with both AlSi10Mg and AlSi10Mg-Er-Zr filler powders. The porosity characteristics, mechanical performance, hardness distribution and microstructural evolution of the welded joints were investigated. On the basis of the present results, the following conclusions can be reached. (1) All welded joints produced a fully penetrated weld with different welding modes between the upper and lower regions. As the PWHT temperature increased from 250 ℃ to 400 ℃, the hydrogen porosity in the welds gradually decreased from 3.2–0.7%, and the maximum pore diameter also decreased from 276.7 µm to 131.0 µm. The AlSi10Mg-Er-Zr filler powder slightly reduced the porosity in the welds. Moreover, the upper region showed lower porosity and smaller hydrogen pore sizes compared to the bottom region. (2) The welded joints subjected to PWHT at 280 ℃, 300 ℃, and 400 ℃ in the bottom region show an inverted U-shaped hardness profile, while a U-shaped hardness profile was observed for the joints treated at 250 ℃ PWHT due to varying degrees of hardness reduction in the BM. The upper region of the welds subjected to different PWHT temperatures had higher hardness compared to the lower region. When PWHT was conducted at 250 ℃, 280 ℃, and 300 ℃, the resulting welded joints produced with both AlSi10Mg and AlSi10Mg-Er-Zr filler powders exhibited different locations of minimum hardness from those treated at 400 ℃. (3) Welded joints subjected to 250 ℃, 280 ℃ and 300 ℃ PWHT showed high UTS ranging from 230 MPa to 240 MPa. However, the UTS significantly decreased for the joints treated at 400 ℃ PWHT. The EF increased significantly from 5.3–19.1% as the PWHT temperature rose from 250 ℃ to 400 ℃. Compared to AlSi10Mg powder, those joints at 250 ℃ and 280 ℃ PWHT with AlSi10Mg-Er-Zr powder demonstrated an increased UTS by 4.3% and 7.2%, and improvements in EF by 3.8% and 94.6%, respectively. The welded joints at 280°C PWHT in combination with AlSi10Mg-Er-Zr filler powder exhibited the highest UTS of 259.0 MPa along with an excellent EF of 10.9%. The tensile fracture of the welded joints at 250 ℃, 280 ℃, and 300 ℃ PWHT was governed by microstructure and initial porosity, while the welded joints at 400 ℃ PWHT was characterized by highly ductile fracture behavior. (4) The Si-rich eutectic in the upper region of the welds subjected to 250 ~ 400 ℃ PWHT was well interconnected, exhibiting finer and better connectivity of the Si-rich eutectic compared to the bottom region. Despite this, the morphology and size of the Si-rich eutectic in the bottom and upper regions of the welds did not significantly change with increasing PWHT temperatures. Both the lower and upper regions of the welds produced with AlSi10Mg-Er-Zr powder exhibited a more uniform and finer Si-rich eutectic structure along with refined α-Al cells. The welded joints subjected to 280 ℃ PWHT demonstrated superior grain refinement efficacy compared to 400 ℃ PWHT, producing a significantly higher fraction of HAGBs at 54.0% relative to 40.8% at the elevated temperature. This substantial difference establishes a clear inverse relationship between PWHT temperature and HAGB formation. (5) A distinct narrow band of the EQZ ranging from 40.2 to 78.1 µm was clearly observed along fusion boundary of the welded joints, while LMD welding using AlSi10Mg-Er-Zr filler powder produced minimal influence on the EQZ width. The welds produced with AlSi10Mg-Er-Zr filler powder displayed significantly refined dendritic structures near the EQZ boundary, clearly demonstrating the grain-refining capability of Er and Zr additions. Notably, a substantial reduction in both porosity density and pore size near the fusion boundary was observed when using AlSi10Mg-Er-Zr filler powder. (6) The bottom regions of both welds exhibited similar morphologies of coarse rod-like and plate-like Si-rich eutectic structures, while both filler powders formed continuous Si-rich eutectic networks in the upper regions. Corresponding α-Al cell sizes in the bottom region was 10.1 µm and 5.9 µm, and measured 6.1 µm and 4.2 µm in the upper region, for AlSi10Mg powder for AlSi10Mg-Er-Zr filler powder, respectively. Therefore, Er and Zr additions in the AlSi10Mg filler powder effectively refined both the Si-rich eutectic and α-Al structure, whereas they did not substantially alter the fundamental morphology or connectivity of the Si-rich eutectic within corresponding weld regions. (7) The fraction of HAGBs was 53.9% and 61.3%, and the average grain misorientation angles were 29.6° and 33.8°, for the AlSi10Mg and AlSi10Mg-Er-Zr weld, respectively. The Σ3 boundary fraction reached 17.2% and 24.9% in the AlSi10Mg and AlSi10Mg-Er-Zr weld, respectively. Moreover, merely 14.1% of grains exhibited SFs within the low range of 0.3 to 0.4 in AlSi10Mg welds, while this value underwent a remarkable increase to 34.8% in the welds produced with the AlSi10Mg-Er-Zr filler powder. Thus, the AlSi10Mg-Er-Zr filler powder effectively increased proportions of HAGBs, Σ3 grain boundaries, and promoted the development of hard-oriented grains. Declarations Competing Interests The authors have no relevant financial or nonfinancial interests to disclose. Funding This work was supported by the National Natural Science Foundation of China (Grant number 52271018). Author contributions Yingying Liu: Formal analysis, Writing-original draft. Jingchuan Li: Formal analysis, Validation. Zhaotong Li: Formal analysis, Validation, review & editing. Li Cui: Conceptualization, Writing-review & editing, Funding acquisition. Dingyong He: Conceptualization, Supervision. Jie Xu: Conceptualization, Validation. Data availability Statement The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. References Gu DD, Meiners W, Wissenbach K, Poprawe R (2012) Laser additive manufacturing of metallic components: materials, processes and mechanisms. Int Mater Rev 57(3):133–164. https://doi.org/10.1179/1743280411Y.0000000014 Martin JH, Yahata BD, Hundley JM, Martin JH, Yahata BD, Hundley JM, Mayer JA, Schaedler TA, Pollock TM (2017) 3D printing of high-strength aluminium alloys. Nature 549(7672):365–369. https://doi.org/10.1038/nature23894 Khairallah SA, Martin AA, Lee JRI, Martin AA, Guss G, King WE (2020) Controlling interdependent meso-nanosecond dynamics and defect generation in metal 3D printing. Science 368(6491):660–665. https://doi.org/10.1126/science.aay7830 Zhao L, Song LB, Macías JGS, Zhu YX, Huang MS, Simar A, Li ZH (2022) Review on the correlation between microstructure and mechanical performance for laser powder bed fusion AlSi10Mg. Addit Manuf 56:102914. https://doi.org/10.1016/j.addma.2022.102914 Wu DF, Cui L, Wu X, Guo XY, Shao W, Tan Z, He DY (2022) Cold metal transfer welding of AlSi7Mg alloy sheets prepared by selective laser melting. J Manuf Process 84:282–297. https://doi.org/10.1016/j.jmapro.2022.10.013 Michler M, Hollmann P, Zenker R, Buchwalder A (2021) Investigation of pore reduction in hybrid joints of conventionally and additively manufactured AlSi10Mg using electron beam welding. Adv Eng Mater 23(6):2001325. https://doi.org/10.1002/adem.202001325 Nahmany M, Rosenthal I, Benishti I, Frage N, Stern A (2015) Electron beam welding of AlSi10Mg workpieces produced by selected laser melting additive manufacturing technology. Addit Manuf 8:63–70. https://doi.org/10.1016/j.addma.2015.08.002 Mäkikangas J, Rautio T, Mustakangas A, Mäntyjärvi K (2019) Laser welding of AlSi10Mg aluminium-based alloy produced by selective laser melting (SLM). Procedia Manuf 36:88–94. https://doi.org/10.1016/j.promfg.2019.08.013 Biffi C, Fiocchi J, Tuissi A (2019) Laser weldability of AlSi10Mg alloy produced by selective laser melting: microstructure and mechanical behavior. J Mater Eng Perform 28:6714–6719. https://doi.org/10.1007/s11665-019-04402-7 Zhang CC, Bao YF, Zhu HH, Nie XJ, Zhang WQ, Zhang SS, Zeng XY (2019) A comparison between laser and TIG welding of selective laser melted AlSi10Mg. Opt Laser Technol 120:105696. https://doi.org/10.1016/j.optlastec.2019.105696 Emmelmann C, Beckmann F (2017) Optimization of laser welding process for laser additive manufactured aluminum parts by means of beam oscillation and process-oriented component design. In: Proceedings of Lasers in Manufacturing Conference. https://www.wlt.de/lim/Proceedings2017/Data/PDF/Contribution187_final.pdf Beckmann F, Emmelmann C (2019) Optimization of the weldability of laser additive manufactured aluminum by means of hydrogen minimization in the component and welding parameter optimization. In: Proceedings of Lasers in Manufacturing Conference. https://www.wlt.de/lim/Proceedings2019/data/PDF/Contribution_327_ final. pdf Schwarz A, Schleser M, Gerhards B, Popoola P, Gebhardt A (2021) Welding of additive manufactured AlSi10Mg: using laser welding in a vacuum for high quality weld seams-a new approach to welding LPBF manufactured AlSi10Mg. S Afr J Ind Eng 32:99–112. http://dx.doi.org/10.7166/32-4-2513 Chen NN, Wan ZX, Wang HP, Li JJ, Yang BX, Solomon J, Carlson B (2022) Effect of ambient pressure on laser welding of AlSi10Mg fabricated by selected laser melting. Mater Des 215:110427. https://doi.org/10.1016/j.matdes.2022.110427 Cui L, Peng ZB, Chang Y, He DY, Cao Q, Guo XY, Zeng Y (2022) Porosity, microstructure and mechanical property of welded joints produced by different laser welding processes in selective laser melting AlSi10Mg alloys. Opt Laser Technol 150:107952. https://doi.org/10.1016/j.optlastec.2022.107952 Wang C, He DY, Cui L, Shi ZF, Guo XY, Wu X, Zhai YY, Bu FH (2023) Improved tensile strength of welded AlSi10Mg alloys by using laser metal deposition. Sci Technol Weld Jo 28(9):946–954. https://doi.org/10.1080/13621718.2023.2253591 Li JC, Cui L, Wu DF, Wang C, Shi ZF, He DY, Cao Q (2024) Welding of additive manufacturing AlSi10Mg alloys using a laser metal deposition process with different heat input. Weld World 68(8):2083–2101. https://doi.org/10.1007/s40194-024-01764-y Chen WH, Liu FC, Niu PL, You QF, Liu FG, Wang ZT (2024) Effect of prewelding solution heat treatment on pore defect in laser welding of AlSi10Mg alloy prepared by selective laser melting. Rare Metal Mater Eng 53(04):1111–1118. https://doi.org/10.12442/j.issn.1002-185X.20230070 Wang C, He DY, He CX, Han FY, Li JC, Cui L (2024) Effect of heat treatments on microstructure and mechanical behavior of laser welded LPBF AlSi10Mg alloys. Mater Today Commun 108478. https://doi.org/10.1016/j.mtcomm.2024.108478 Zheng Y, Zhao ZH, Xiong RZ, Ren GC, Yao MH, Liu W, Zang LB (2024) Effect of post heat treatment on microstructure, mechanical property and corrosion behavior of AlSi10Mg alloy fabricated by selective laser melting. Prog Nat Sci-Mater 34(1):89–101. https://doi.org/10.1016/j.pnsc.2024.02.001 Long J, Zhang LZ, Zhang LL, Wang X, Zhang GF, Zhang XJ, Na SJ (2020) Effects of minor Zr addition on the microstructure and mechanical properties of laser welded joint of Al/SiCp metal-matrix composite. J Manuf Process 49:373–384. https://doi.org/10.1016/j.jmapro.2019.12.004 Spierings AB, Dawson K, Kern K, Palm F, Wegener K (2017) SLM-processed Sc- and Zr- modified Al-Mg alloy: Mechanical properties and microstructural effects of heat treatment. Mater Sci Eng A 701:264–273. https://doi.org/10.1016/j.msea.2017.06.089 Li RD, Wang MB, Yuan TC, Song B, Chen C, Zhou KC, Cao P (2017) Selective laser melting of a novel Sc and Zr modified Al-6.2Mg alloy: processing, microstructure, and properties. Powder Technol 319:117–128. https://doi.org/10.1016/j.powtec.2017.06.050 Chen B, Wang FZ, Ren SM, Xia P, Wang ZP, Hao JF, Chen ZP, Li RX, Jiang HX (2024) Influence of erbium addition on microstructure and performances of AlSi10Mg alloy prepared by selective laser melting. J Rare Earths 42(10):1988–1994. https://doi.org/10.1016/j.jre.2023.09.017 Yang TY, Cui L, He DY, Huang H (2022) Enhancement of microstructure and mechanical property of AlSi10Mg-Er-Zr alloys fabricated by selective laser melting. Acta Metall Sin 58(9):1108–1117. https://www.ams.org.cn/EN/ 10.11900/0412.1961.2021.00085 Peng ZB, Cui L, He DY, Guo XY, Zeng Y, Cao Q, Huang H (2022) Effect of Er and Zr addition on laser weldability of AlSi10Mg alloys fabricated by selective laser melting. Mater Charact 190:112070. https://doi.org/10.1016/j.matchar.2022.112070 Wen H, Yves N, Matthew F, Kan WH, Nadot Y, Foley M, Ridosz L, Proust G, Cairney JM (2019) Factors that affect the properties of additively-manufactured AlSi10Mg: Porosity versus microstructure. Addit Manuf 29:100805. https://doi.org/10.1016/j.addma.2019.100805 Nunes R, Faes K, Verlinde W, Probst F, Freese JD, Arnhold J, Waele WD, Lezaack M, Simar A (2025) Welding of additively manufactured aluminium alloy components: challenges and mitigation strategies. Weld World 69:1003–1021. https://doi.org/10.1007/s40194-025-01923-9 Wang CG, Zhu JX, Wang GW, Qin Y, Sun MY, Yang JL, Shen XF, Huang SK (2022) Effect of building orientation and heat treatment on the anisotropic tensile properties of AlSi10Mg fabricated by selective laser melting. J Alloy Compd 895:162665. https://doi.org/10.1016/j.jallcom.2021.162665 Albu M, Krisper R, Lammer J, Kothleitner G, Fiocchi J, Bassani P (2020) Microstructure evolution during in-situ heating of AlSi10Mg alloy powders and additive manufactured parts. Addit Manuf 36:101605. https://doi.org/10.1016/j.addma.2020.101605 Chen T, Wang L, Tan S (2016) Effects of vacuum annealing treatment on microstructures and residual stress of AlSi10Mg parts produced by selective laser melting process. Mod Phys Lett B 30(19):1650255. https://doi.org/10.1142/S0217984916502559 Wang M, Song B, Wei Q, Zhang Y, Shi Y (2019) Effects of annealing on the microstructure and mechanical properties of selective laser melted AlSi7Mg alloy. Mater Sci Eng A 739:463–472. https://doi.org/10.1016/j.msea.2018.10.047 Kotadia HR, Gibbons G, Das A, Howes PD (2021) A review of laser powder bed fusion additive manufacturing of aluminium alloys: microstructure and properties. Addit Manuf 46:102155. https://doi.org/10.1016/j.addma.2021.102155 Snopinski P, Liveric L, Beniak J (2024) Annealing effects on microstructure and texture in KOBO-Processed LPBF AlSi10Mg alloy: elucidating CSL boundary formation. Symmetry 16:1663. https://doi.org/10.3390/sym16121663 Cabrera-Correa L, Gonzalez-Rovira L, De DLJ, Castillo-Rodriguez M, Botana FJ (2023) Effect of the heat treatment on the mechanical properties and microstructure of scalmalloy manufactured by selective laser melting (SLM) under certified conditions. Mater Charact 196:112549. https://doi.org/10.1016/j.matchar.2022.112549 Ma M, Wang B, Liu H, Yi D, Shen F, Zhai T (2020) Investigation of fatigue crack propagation behavior of 5083 aluminum alloy under various stress ratios: role of grain boundary and Schmid factor. Mater Sci Eng A 773:138871. https://doi.org/10.1016/j.msea.2019.138871 Sun G, Wang C, Wei X, Shang D, Chen S (2019) Study on small fatigue crack initiation and growth for friction stir welded joints. Mater Sci Eng A 739:71–85. https://doi.org/10.1016/j.msea.2018.10.002 Cite Share Download PDF Status: Posted Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. Our growing team is made up of researchers and industry professionals working together to solve the most critical problems facing scientific publishing. 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20:45:11","extension":"html","order_by":34,"title":"","display":"","copyAsset":false,"role":"acdc-reference","size":145269,"visible":true,"origin":"","legend":"","description":"","filename":"earlyproof.html","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/fe05e18c8e48c0c48de69fc6.html"},{"id":92899360,"identity":"b01e202e-071c-43d4-b03f-f46f16cbea48","added_by":"auto","created_at":"2025-10-06 20:53:10","extension":"jpeg","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":170751,"visible":true,"origin":"","legend":"\u003cp\u003eSchematic of LMD welding for L-PBF AlSi10Mg alloys: (a) Welding setup;(b) Five-track LMD welding; (c) Tensile specimen dimensions of the welded joints\u003c/p\u003e","description":"","filename":"floatimage1.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/fac8e1cbcec6524ad0927ff3.jpeg"},{"id":92899242,"identity":"b1831094-2e73-48ec-8e75-5c1accb9c31b","added_by":"auto","created_at":"2025-10-06 20:45:10","extension":"jpeg","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":405282,"visible":true,"origin":"","legend":"\u003cp\u003eCross sections of the welded joints at varying PWHT temperatures: (a) 250 ℃, AlSi10Mg; (b) 250 ℃, AlSi10Mg-Er-Zr; (c) 280 ℃, AlSi10Mg; (d) 280 ℃, AlSi10Mg-Er-Zr; (e) 300 ℃, AlSi10Mg; (f) 400 ℃, AlSi10Mg\u003c/p\u003e","description":"","filename":"floatimage2.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/ddf9eae2806191ddb53185fb.jpeg"},{"id":92898891,"identity":"53dcdd42-1212-499f-a3c6-167af35c4753","added_by":"auto","created_at":"2025-10-06 20:37:10","extension":"jpeg","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":136096,"visible":true,"origin":"","legend":"\u003cp\u003ePorosity and maximum pore diameter generated in the welds as a function of PWHT temperatures withdifferent filler powders\u003c/p\u003e","description":"","filename":"floatimage3.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/97dac7e8002a39923967756f.jpeg"},{"id":92899799,"identity":"ba5e9769-19a2-4916-9cba-2b6458b00e1e","added_by":"auto","created_at":"2025-10-06 21:01:10","extension":"jpeg","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":227185,"visible":true,"origin":"","legend":"\u003cp\u003eMicrohardness distribution across LMD welded joints for AlSi10Mg and AlSi10Mg-Er-Zr fillers at varying PWHT temperatures: (a) The bottom region; (b) The upper region\u003c/p\u003e","description":"","filename":"floatimage4.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/7e7ca5384b9c747b3f13c111.jpeg"},{"id":92899244,"identity":"26e97b6e-c511-4c48-84ed-9f218df4be1b","added_by":"auto","created_at":"2025-10-06 20:45:10","extension":"jpeg","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":66002,"visible":true,"origin":"","legend":"\u003cp\u003eStress-strain curves of the welded joints at varying PWHT temperatures withdifferent filler powder\u003c/p\u003e","description":"","filename":"floatimage5.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/fd8aeb6129ba4423e90b29c1.jpeg"},{"id":92898906,"identity":"04e73f65-af52-448b-88f1-85bc6e7531d2","added_by":"auto","created_at":"2025-10-06 20:37:10","extension":"jpeg","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":1385108,"visible":true,"origin":"","legend":"\u003cp\u003eSurface morphology of the fractured joints at various PWHT temperatures with different filler powders: (a) 250 ℃, AlSi10Mg; (b) 250 ℃, AlSi10Mg-Er-Zr; (c) 280 ℃, AlSi10Mg; (d) 280 ℃, AlSi10Mg-Er-Zr; (e) 300 ℃, AlSi10Mg; (f) 400 ℃, AlSi10Mg;(g) High magnified micrograph at 280 ℃, AlSi10Mg-Er-Zr; (h) High magnified micrograph at 400 ℃, AlSi10Mg.\u003c/p\u003e","description":"","filename":"floatimage6.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/5df2e67539eb8667b253e1b9.jpeg"},{"id":92899361,"identity":"3a4a7d1d-f937-4ebc-84d8-ebb50dd9e326","added_by":"auto","created_at":"2025-10-06 20:53:10","extension":"jpeg","order_by":7,"title":"Figure 7","display":"","copyAsset":false,"role":"figure","size":1134154,"visible":true,"origin":"","legend":"\u003cp\u003eSi-rich eutectic morphology in the bottom and upper region of the welds at various PWHT temperatures: (a, a’)250 ℃; (b, b’) 280 ℃; (c, c’) 300 ℃; (d, d’) 400 ℃.\u003c/p\u003e","description":"","filename":"floatimage7.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/b1f2fa2d149e6066f0abdb18.jpeg"},{"id":92898905,"identity":"4468b5ae-b74a-4157-b27c-046c4b7735fe","added_by":"auto","created_at":"2025-10-06 20:37:10","extension":"jpeg","order_by":8,"title":"Figure 8","display":"","copyAsset":false,"role":"figure","size":578817,"visible":true,"origin":"","legend":"\u003cp\u003eIPF maps and grain size distributions in the upper region of the welds at different PWHT temperatures :(a, c) 280 ℃; (b, d) 400 ℃.\u003c/p\u003e","description":"","filename":"floatimage8.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/adefe87a9a21a64a295df1cc.jpeg"},{"id":92899246,"identity":"3b7090a0-9d76-487a-9097-c543ab2acdef","added_by":"auto","created_at":"2025-10-06 20:45:10","extension":"jpeg","order_by":9,"title":"Figure 9","display":"","copyAsset":false,"role":"figure","size":580834,"visible":true,"origin":"","legend":"\u003cp\u003eDistribution of the grain boundaries and misorientation angles in the upper region of the welds at different PWHT temperatures :(a, c) 280 ℃; (b, d) 400 ℃.\u003c/p\u003e","description":"","filename":"floatimage9.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/841ca49e68f0aa6e5120bef4.jpeg"},{"id":92898900,"identity":"3324551a-3547-4c70-af59-1231b7072815","added_by":"auto","created_at":"2025-10-06 20:37:10","extension":"jpeg","order_by":10,"title":"Figure 10","display":"","copyAsset":false,"role":"figure","size":420207,"visible":true,"origin":"","legend":"\u003cp\u003eFusion boundary microstructure of the welded joints treated at 280 ℃ PWHT with different filler powder: (a) AlSi10Mg; (b) AlSi10Mg-Er-Zr.\u003c/p\u003e","description":"","filename":"floatimage10.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/061e72d11c355cd67fe3c215.jpeg"},{"id":92899363,"identity":"d052b005-81a4-4bd5-b9ed-08a0fb4ef8cf","added_by":"auto","created_at":"2025-10-06 20:53:10","extension":"jpeg","order_by":11,"title":"Figure 11","display":"","copyAsset":false,"role":"figure","size":512870,"visible":true,"origin":"","legend":"\u003cp\u003eMicrostructure in the bottom and upper region of the welds at 280 ℃ PWHT with different filler powder: (a, b) AlSi10Mg; (c, d) AlSi10Mg-Er-Zr.\u003c/p\u003e","description":"","filename":"floatimage11.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/fbe9b75940babef77f328dd5.jpeg"},{"id":92899800,"identity":"844f633a-7fe1-4f63-b7e5-11e7138b6a9c","added_by":"auto","created_at":"2025-10-06 21:01:10","extension":"jpeg","order_by":12,"title":"Figure 12","display":"","copyAsset":false,"role":"figure","size":589085,"visible":true,"origin":"","legend":"\u003cp\u003eSi-rich eutectic morphology in the bottom and upper region of the welds at 280 ℃ PWHT with different filler powder: (a, b) AlSi10Mg; (c, d) AlSi10Mg-Er-Zr.\u003c/p\u003e","description":"","filename":"floatimage12.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/55ea70f8ce2347d55c1a43a8.jpeg"},{"id":92899253,"identity":"b277c2a6-8c26-4ee5-8707-f085a662e535","added_by":"auto","created_at":"2025-10-06 20:45:10","extension":"jpeg","order_by":13,"title":"Figure 13","display":"","copyAsset":false,"role":"figure","size":742345,"visible":true,"origin":"","legend":"\u003cp\u003eEffect of Er and Zr addition on grain boundary distribution and distribution of misorientation angles in the welds at 400 ℃ PWHT with different filler powder: (a, c) AlSi10Mg; (b, d) AlSi10Mg-Er-Zr.\u003c/p\u003e","description":"","filename":"floatimage13.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/230a8d450744ccdac7b42822.jpeg"},{"id":92899247,"identity":"a0eca216-1f20-4d83-a278-619a3c36b167","added_by":"auto","created_at":"2025-10-06 20:45:10","extension":"jpeg","order_by":14,"title":"Figure 14","display":"","copyAsset":false,"role":"figure","size":550697,"visible":true,"origin":"","legend":"\u003cp\u003eSF map and distribution of the grains in the upper region of the welds at 400 ℃ PWHT with different filler powder: (a, c) AlSi10Mg; (b, d) AlSi10Mg-Er-Zr.\u003c/p\u003e","description":"","filename":"floatimage14.jpeg","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/2a7446b42c5eee9dbe4fdbd9.jpeg"},{"id":95529417,"identity":"206cf3fd-b9d9-477b-a4e2-bf1a40cc5abc","added_by":"auto","created_at":"2025-11-10 10:17:05","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":8413483,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-7438037/v1/1a08f2be-607d-4bf9-9343-c8ebd792df34.pdf"}],"financialInterests":"","formattedTitle":"Influence of pre-weld heat treatment temperatures and AlSi10Mg-Er-Zr filler powder on microstructure and mechanical properties of welded joints produced using laser metal deposition for L-PBF AlSi10Mg alloys","fulltext":[{"header":"1. Introduction","content":"\u003cp\u003eAlSi10Mg alloys fabricated via the laser powder bed fusion (L-PBF) process have gained significant attention in the aerospace, aeronautics, and automotive industries owing to their excellent printability, lightweight properties, and relatively low production costs [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e]. Nevertheless, the production of complex and large-scale components using the L-PBF process remains impractical due to limitations in build chamber size [\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e, \u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e]. To overcome the size limitations of L-PBF components, welding techniques have been increasingly employed to join additively manufactured parts with other L-PBF components or with conventionally fabricated structures [\u003cspan additionalcitationids=\"CR6\" citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e]. This approach effectively circumvents the build volume constraints inherent to the L-PBF process while maintaining its advantages.\u003c/p\u003e\u003cp\u003eHowever, the weldability of L-PBF AlSi10Mg alloys using fusion welding techniques has been limitedly reported in literature. It was found that fusion welding of L-PBF AlSi10Mg alloys remains particularly challenging [\u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e, \u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e], primarily due to the significantly higher hydrogen pore susceptibility generated in the welds compared to casting materials of the same alloy. The extensive porosity was directly attributable to the melting of the base metal (BM) during the welding, releasing very high content of hydrogen gas and entering the welding pool [\u003cspan additionalcitationids=\"CR11\" citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e], which can be the hydrogen source for the nucleation of hydrogen pores. As a result, lowering the weld dilution reduces hydrogen pickup from the BM, thereby mitigating porosity. In this regard, some special welding processes, such as high-pressure or high-vacuum laser welding [\u003cspan citationid=\"CR13\" class=\"CitationRef\"\u003e13\u003c/span\u003e, \u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e], laser metal deposition (LMD) process [\u003cspan additionalcitationids=\"CR16\" citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e] were recently introduced to reduce the hydrogen porosity when welding of L-PBF AlSi10Mg alloys. However, the hydrogen porosity of the welds in L-PBF AlSi10Mg alloys is still higher than compared to conventional Al-based alloys, and it is unlikely that the porosity can be entirely avoided by optimization of the welding processes only.\u003c/p\u003e\u003cp\u003ePre-weld heat treatment (PWHT) is a widely adopted process during welding of metal and its alloys, serving to relieve residual stresses, mitigate porosity and crack formation, and consequently enhance microstructural and mechanical properties. However, research on the influence of PWHT on weldability of L-PBF AlSi10Mg alloys remains limited. M\u0026auml;kikangas et al. implemented stress relief annealing at 300℃ before welding [\u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e]. Although the study highlighted significant reduction of pores in the laser-welded AlSi10Mg alloys, the PWHT resulted in a strength reduction of approximately 19 MPa compared to non-heat-treated conditions. Chen et al. investigated the influence of solution heat treatment (SHT) at 520 ℃ on pore distribution, microstructural evolution, and mechanical properties of laser-welded joints [\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e]. Their findings revealed that vacuum SHT effectively reduced weld porosity to 0.14%. However, the heat treatment resulted in a substantial reduction in tensile strength, decreasing to 143 MPa, which was significantly lower than that observed in non-heat-treated joints. This strength deterioration primarily stemmed from microstructural changes, as elevated annealing temperatures and extended processing times promoted coarsening of the eutectic Si-rich phases while simultaneously decreasing their number density [\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e].\u003c/p\u003e\u003cp\u003eThe aforementioned studies demonstrate that pre-weld vacuum SHT effectively minimizes porosity in laser welded L-PBF AlSi10Mg alloys, whereas the PWHT at 520 ℃ substantially degrades joint strength. Notably, a desired microstructure can be achieved using lower temperatures and shorter durations than those prescribed by ASTM T6 standards [\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e], suggesting that microstructural optimization in welded L-PBF AlSi10Mg alloys is feasible through controlled PWHT at lower temperatures. To bridge this knowledge gap, the present study systematically explores PWHT at temperatures below 520 ℃, with the dual objectives of significantly suppressing hydrogen porosity while preserving the tensile strength of the welded joints.\u003c/p\u003e\u003cp\u003eThe weld composition in welded joints is typically designed to match the base metal (BM), but may be intentionally modified to achieve compatible properties. Recent advances have shown that, the principle of grain refinement through inoculate addition with scandium (Sc) and zirconium (Zr) element, well-established in aluminum alloy casting [\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e], has recently been extended to additive manufacturing research. While the Sc and Zr combination has demonstrated significant grain-refining effects [\u003cspan citationid=\"CR22\" class=\"CitationRef\"\u003e22\u003c/span\u003e, \u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e], erbium (Er) has emerged as a cost-effective rare earth alternative, exhibiting microalloying effects comparable to Sc. Numerous studies have confirmed Er's strengthening potential in aluminum alloys [\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e]. Our prior research revealed that minor additions of Er and Zr substantially enhance the mechanical performance of both L-PBF AlSi10Mg alloys and welded L-PBF AlSi10Mg alloys [\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e, \u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e]. This improvement stems from the formation of nanoscale coherent Al₃Er and Al₃(Zr,Er) precipitates, achieving an exceptional strength-ductility balance. However, the synergistic effects of Er and Zr microalloying and PWHT on microstructural evolution and mechanical properties in welded L-PBF AlSi10Mg remain insufficiently understood.\u003c/p\u003e\u003cp\u003eThis study systematically investigates the effects of low PWHT temperatures (250\u0026ndash;400 ℃) and AlSi10Mg-Er-Zr filler powder on the LMD welded L-PBF AlSi10Mg alloys. The research methodology involved conducting vacuum PWHT on L-PBF AlSi10Mg sheets at four specific temperatures (250 ℃, 280 ℃, 300 ℃, and 400 ℃) followed by LMD welding in a butt joint configuration using both conventional AlSi10Mg and the prepared AlSi10Mg-Er-Zr filler powders. Through comprehensive characterization of porosity characteristics, microstructural evolution, and mechanical properties of the welded joints, this work provides fundamental insights into the process-microstructure-property relationships, establishing a scientific foundation for optimizing welding and repair techniques for additively manufactured aluminum components in industrial applications.\u003c/p\u003e"},{"header":"2. Materials and experimental method","content":"\u003cdiv id=\"Sec3\" class=\"Section2\"\u003e\u003ch2\u003e2.1 Materials\u003c/h2\u003e\u003cp\u003eThe BM specimens for welding were fabricated as 50 \u0026times; 25 \u0026times; 3.0 mm sheets using an EOS M280 system in a high-purity argon atmosphere. The process employed gas-atomized AlSi10Mg powder (15\u0026ndash;53 \u0026micro;m size) deposited on a 150 ℃ preheated aluminum substrate. Throughout the L-PBF process, the build chamber maintained oxygen levels below 200 ppm under continuous argon protection. To optimize the manufacturing process, each successive laser scanning path was rotated by 67\u0026deg; relative to the previous track. Through orthogonal experimentation, the optimal processing parameters were determined to be a laser power of 370 W, a scanning speed of 1300 mm/s, a hatching space of 0.19 mm, and a track thickness of 20 \u0026micro;m, which yielded specimens with high relative density. Density measurements were performed using both Archimedes' principle and ImageJ-based image analysis. Only L-PBF AlSi10Mg sheets demonstrating relative densities exceeding 99.5% were selected as BM specimens for subsequent LMD welding.\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec4\" class=\"Section2\"\u003e\u003ch2\u003e2.2 Welding experimental\u003c/h2\u003e\u003cp\u003ePWHT was performed on the BM specimens in a PZTH-150 vacuum furnace prior to welding. Four distinct heat treatment temperatures (250 ℃, 280 ℃, 300 ℃, and 400 ℃) were applied with a 2-hour holding time at each temperature, employing a controlled heating rate of 10 ℃/min followed by air cooling to ambient temperature. The resultant surface oxide films were subsequently removed through sequential mechanical grinding and chemical cleaning procedures. The LMD welding was conducted using a 4-kW TruDisk laser system (TRUMPF) integrated with an ABB \u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e-axis robotic manipulator and a DPSF-2 automated powder feeder. Two variants of filler powders were utilized: commercially available gas-atomized AlSi10Mg powder and specially formulated AlSi10Mg-Er-Zr powder, both sieved to a particle size range of 45\u0026ndash;105 \u0026micro;m. The AlSi10Mg-Er-Zr powder was produced via in-situ alloying during gas atomization, incorporating master alloys of Al-Mg, Al-Er, and Al-Zr with the AlSi10Mg powder. High-purity argon (99.99%) served as the shielding gas during welding, with precisely controlled flow rates of 2.5 L/min for powder delivery and 25.0 L/min for molten pool protection. The detailed chemical compositions of both the BM and filler powders are systematically presented in Table\u0026nbsp;\u003cspan refid=\"Tab1\" class=\"InternalRef\"\u003e1\u003c/span\u003e.\u003c/p\u003e\u003cp\u003e\u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"Yes\" id=\"Tab1\" border=\"1\"\u003e\u003ccaption language=\"En\"\u003e\u003cdiv class=\"CaptionNumber\"\u003eTable 1\u003c/div\u003e\u003cdiv class=\"CaptionContent\"\u003e\u003cp\u003eChemical compositions of the BM and filler powder (wt.%)\u003c/p\u003e\u003c/div\u003e\u003c/caption\u003e\u003ccolgroup cols=\"10\"\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c3\" colnum=\"3\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c4\" colnum=\"4\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c5\" colnum=\"5\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c6\" colnum=\"6\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c7\" colnum=\"7\"\u003e\u003c/div\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c8\" colnum=\"8\"\u003e\u003c/div\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c9\" colnum=\"9\"\u003e\u003c/div\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c10\" colnum=\"10\"\u003e\u003c/div\u003e\u003cthead\u003e\u003ctr\u003e\u003cth align=\"left\" colname=\"c1\"\u003e\u0026nbsp;\u003c/th\u003e\u003cth align=\"left\" colname=\"c2\"\u003e\u003cp\u003eFe\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c3\"\u003e\u003cp\u003eMg\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c4\"\u003e\u003cp\u003eMn\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c5\"\u003e\u003cp\u003eSi\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c6\"\u003e\u003cp\u003eZn\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c7\"\u003e\u003cp\u003eTi\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c8\"\u003e\u003cp\u003eEr\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c9\"\u003e\u003cp\u003eZr\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c10\"\u003e\u003cp\u003eAl\u003c/p\u003e\u003c/th\u003e\u003c/tr\u003e\u003c/thead\u003e\u003ctbody\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003eBase metal\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e0.15\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e0.39\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e\u003cp\u003e\u0026le;\u0026thinsp;0.01\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e\u003cp\u003e10.51\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c6\"\u003e\u003cp\u003e\u0026le;\u0026thinsp;0.01\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c7\"\u003e\u003cp\u003e0.14\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c8\"\u003e\u003cp\u003e-\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c9\"\u003e\u003cp\u003e-\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c10\"\u003e\u003cp\u003eBal.\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003eAlSi10Mg\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e0.20\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e0.30\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e\u003cp\u003e\u0026le;\u0026thinsp;0.01\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e\u003cp\u003e10.53\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c6\"\u003e\u003cp\u003e\u0026le;\u0026thinsp;0.01\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c7\"\u003e\u003cp\u003e0.12\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c8\"\u003e\u003cp\u003e-\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c9\"\u003e\u003cp\u003e-\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c10\"\u003e\u003cp\u003eBal.\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003eAlSi10Mg-Er-Zr\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e0.11\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e0.40\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e\u003cp\u003e\u0026le;\u0026thinsp;0.01\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e\u003cp\u003e10.05\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c6\"\u003e\u003cp\u003e\u0026le;\u0026thinsp;0.01\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c7\"\u003e\u003cp\u003e0.11\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c8\"\u003e\u003cp\u003e0.65\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c9\"\u003e\u003cp\u003e0.23\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c10\"\u003e\u003cp\u003eBal.\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003c/tbody\u003e\u003c/colgroup\u003e\u003c/table\u003e\u003c/div\u003e\u003c/p\u003e\u003cp\u003eFull-penetration butt joints were obtained through five-track LMD processes using both AlSi10Mg and AlSi10Mg-Er-Zr filler powders. The BM specimens were prepared in a 40\u0026deg; V-groove configuration without root gap, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e. During the LMD welding, the filler powder was continuously and uniformly fed into the groove, where it was melted and sequentially deposited track-by-track to form the weld joint. The welding parameters were optimized through a series of preliminary trials, maintaining constant values for laser power of 1500 W, focus spot diameter of 2.2 mm, defocus distance of 0 mm, and powder feeding rate of 3.15 g/min. The travel speed was set at 10 mm/s for the first deposition track and increased to 15 mm/s for the subsequent four tracks. This variation in travel speeds produced different linear heat inputs, with 150 J/mm for the initial track and 100 J/mm for the remaining tracks.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec5\" class=\"Section2\"\u003e\u003ch2\u003e2.3 Microstructure, porosity characterization and mechanical testing\u003c/h2\u003e\u003cp\u003eFollowing LMD welding, cross-sectional specimens were extracted from the welded joints for metallographic characterization. The preparation sequence consisted of mounting, mechanical polishing, and chemical etching with Keller's reagent to delineate bead morphology and microstructural features. Microstructural analysis and pore characterization were performed using an Olympus LEXT OLS4100 laser scanning confocal microscope, with quantitative measurements of dendrite arm spacing, maximum pore diameter, and porosity percentage conducted using Image Pro Plus software.\u003c/p\u003e\u003cp\u003eFor higher-resolution examination, Si-rich eutectic morphology was investigated using a QUANTA FEG650 field-emission scanning electron microscope (FE-SEM) equipped with energy dispersive spectroscopy (EDS). Fractographic analysis of failed specimens was carried out employing both a Nikon SMZ800 stereomicroscope and a JEOL JSM-7400F high-resolution SEM. Electron backscatter diffraction (EBSD) analysis was conducted using a Hitachi S-3400N thermal field-emission SEM integrated with an Oxford Instruments Nordlys Nano detector system, operating at 20.0 kV. The EBSD specimens were sequentially prepared by mechanical polishing and subsequent electropolishing in a 30% nitric acid-methanol solution at -25 ℃, using an applied voltage of 20 V for 30 seconds. Orientation mapping was acquired in rectangular scan areas with a step size of 1.4 \u0026micro;m. All EBSD data were processed using HKL Technology's Channel 5 software.\u003c/p\u003e\u003cp\u003eMicrohardness measurements were performed using an HVS-1000 Vickers hardness tester with a 98 N load applied for 15 seconds. The reported hardness values represent the average of at least five indentations per test condition. Tensile specimens were machined according to ASTM E8-04 specifications and tested at room temperature using an Instron universal testing machine with a constant crosshead speed of 0.5 mm/min. To ensure statistical reliability, three specimens were tested for each joint condition, with the average value reported as the representative result.\u003c/p\u003e\u003c/div\u003e"},{"header":"3.Results","content":"\u003cdiv id=\"Sec7\" class=\"Section2\"\u003e\u003ch2\u003e3.1 Weld shape and porosity characteristics\u003c/h2\u003e\u003cp\u003eFigure \u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e presents cross-sectional views of the resulting welded joints produced using LMD welding with both AlSi10Mg and AlSi10Mg-Er-Zr filler powders for the BM subjected to 250 ℃, 280 ℃, 300 ℃, and 400 ℃ PWHT. The BM microstructure displayed typical columnar grain structures with widths varying between 20\u0026ndash;200 \u0026micro;m and diverse melt pool dimensions, aligning with established literature findings [\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e, \u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e]. However, microstructural examination at this magnification revealed no heat-affected zone (HAZ), as indicated by the identical melt pool morphology between the welded region and unaffected BM. All joints achieved full penetration with comparable geometries, indicating that neither PWHT temperature variations nor filler powder selection significantly affected weld shape. Additionally, microstructural analysis revealed distinct welding modes between different regions of the welds. The bottom region (first deposition track) displayed a characteristic keyhole mode morphology with a high depth-to-width ratio, while the upper region (tracks 2\u0026ndash;5) exhibited shallower, hemispherical penetration profiles typical of conduction mode welding.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003eRegarding the porosity defects, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e, the welds exhibited characteristic spherical porosity in both upper and bottom regions. The porosity measurements demonstrated a strong temperature dependence for AlSi10Mg welds, decreasing progressively from 3.2% at 250 ℃ PWHT to 0.7% at 400 ℃ PWHT with increasing PWHT temperatures. Additionally, a substantial 52.7% reduction in maximum pore diameter was observed, decreasing from 276.7 \u0026micro;m to 131.0 \u0026micro;m over the investigated temperature range. However, the PWHT at elevated temperatures of 300 ℃ and 400 ℃ substantially decreased weld porosity compared to lower temperature treatments of 250 ℃ and 280 ℃. Moreover, quantitative analysis revealed the upper region possessed significantly lower porosity and smaller hydrogen pore diameters compared to the bottom region. This phenomenon results from the lower heat input of 100 J/mm in the upper layers relative to the 150 J/mm in the bottom layer, reducing the dilution rate and thereby restricting hydrogen absorption into the weld pool [\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e, \u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e]. Regarding the influence of filler powders, the AlSi10Mg-Er-Zr welds exhibited reduced porosity values of 3.0% at 250 ℃ and 2.7% at 280℃, representing a slight decrease compared to the conventional AlSi10Mg welds. As a result, the PWHT temperature emerged as the dominant factor controlling hydrogen pore formation of the welds, whereas the addition of Er and Zr to the AlSi10Mg filler powder exhibited minimal impact on porosity when PWHT temperature remained constant.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec8\" class=\"Section2\"\u003e\u003ch2\u003e3.2 Mechanical properties of welded joints\u003c/h2\u003e\u003cdiv id=\"Sec9\" class=\"Section3\"\u003e\u003ch2\u003e3.2.1 Microhardness distribution\u003c/h2\u003e\u003cp\u003eThe microhardness distribution across various regions of the welded joints produced with four different PWHT temperatures, including the BM, HAZ, fusion boundary, and weld metal, is presented in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e. The as-built BM exhibited a hardness of approximately 120 HV, resulting from the fine microstructure developed during the L-PBF process [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e]. Following PWHT, the BM showed a progressive decrease in average hardness with increasing temperatures from 250 ℃ to 400 ℃. The measured hardness values were 89.2 HV at 250 ℃, decreasing to 82.1 HV at 280 ℃, 75.9 HV at 300 ℃, and reaching a minimum of 61.9 HV at 400 ℃. This temperature-dependent softening behavior can be attributed to microstructural changes in the Si-rich eutectic structure. As reported in previous studies, increasing the annealing temperature or duration leads to the breakdown of the cellular Si network structure and promotes Si precipitation in the α-Al matrix, thereby reducing the solid solution strengthening effect [\u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e, \u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e]. While the network structure remains largely intact at 260℃, complete decomposition occurs at 300 ℃ [\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e]. These microstructural transformations explain the significant hardness reduction observed in BM samples subjected to PWHT at 400 ℃. The pronounced BM softening led to different hardness distribution patterns in welded joints treated at different temperatures.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003eIn the bottom region, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e(a), the welded joints subjected to 280 ℃, 300 ℃, and 400 ℃ PWHT developed a distinct inverted U-shaped hardness profile, while the welded joints treated at the lowest temperature of 250 ℃ maintained an U-shaped hardness distribution using whether AlSi10Mg powder or AlSi10Mg-Er-Zr one. This difference in the hardness profile of the welded joints subjected to PWHT between the lowest temperature at 250 ℃ and the higher temperatures was mainly attributable to the different degree of BM softening. In the upper regions, all the welded joints consistently showed inverted U-shaped profiles regardless of PWHT temperatures, since the hardness in the WM consistently surpassed BM values, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e(b). In addition, detailed measurements revealed hardness ranges of 76.2\u0026ndash;87.8 HV in the bottom region and 89.2-103.7 HV in the upper region, indicating systematically higher hardness values in the upper regions. Notably, peak hardness of 103.7 HV was achieved in the upper region of the welds treated at 280℃ with AlSi10Mg-Er-Zr filler powder. It also can be seen from Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e that a sharp hardness reduction in the HAZ along the transition from the BM to the WM. However, the location of minimum hardness differed significantly between lower-temperature at 250℃, 280 ℃, 300 ℃ PWHT and higher-temperature at 400 ℃ PWHT for both filler powders. The hardness minimum consistently occurred at the fusion boundary for joints subjected to 250 ℃, 280 ℃, and 300 ℃ PWHT, while the lowest hardness values occurred within the BM for the joints at 400 ℃ PWHT, irrespective of AlSi10Mg or AlSi10Mg-Er-Zr filler powders.\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec10\" class=\"Section3\"\u003e\u003ch2\u003e3.2.2 Tensile property of the welded joints\u003c/h2\u003e\u003cp\u003eThe tensile testing was conducted on the welded joints subjected to 250 ℃, 280 ℃, 300 ℃, and 400 ℃ PWHT with AlSi10Mg and AlSi10Mg-Er-Zr filler powders. The stress-strain curves of these joint specimens are presented in Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003e, with the corresponding tensile test results summarized in Table\u0026nbsp;\u003cspan refid=\"Tab2\" class=\"InternalRef\"\u003e2\u003c/span\u003e. The UTS values of the welded joints showed clear variations with PWHT temperatures, measuring 230.3 MPa at 250 ℃, 241.5 MPa at 280 ℃, 231.1 MPa at 300 ℃, and 175.5 MPa at 400 ℃. This demonstrates that joints treated at 250\u0026ndash;300 ℃ maintained relatively high strength levels between 230\u0026ndash;240 MPa, but those subjected to 400 ℃ PWHT experienced a significant reduction in UTS. The highest UTS of 259.0 MPa was achieved at 280℃ PWHT combined with AlSi10Mg-Er-Zr filler powder. Specially, the elongation at fracture (EF) exhibited a progressive increase with rising PWHT temperatures, ranging from 5.3% at 250 ℃ to 19.1% at 400 ℃, representing a substantial 3.6-fold improvement in ductility. This highlights the important influence of PWHT temperatures on the ductility of the welded joints in L-PBF AlSi10Mg alloys.\u003c/p\u003e\u003cp\u003e\u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"Yes\" id=\"Tab2\" border=\"1\"\u003e\u003ccaption language=\"En\"\u003e\u003cdiv class=\"CaptionNumber\"\u003eTable 2\u003c/div\u003e\u003cdiv class=\"CaptionContent\"\u003e\u003cp\u003eTensile properties of the welded joints at various PWHT temperatures with AlSi10Mg and AlSi10Mg-Er-Zr filler powder\u003c/p\u003e\u003c/div\u003e\u003c/caption\u003e\u003ccolgroup cols=\"6\"\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c3\" colnum=\"3\"\u003e\u003c/div\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c4\" colnum=\"4\"\u003e\u003c/div\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c5\" colnum=\"5\"\u003e\u003c/div\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c6\" colnum=\"6\"\u003e\u003c/div\u003e\u003cthead\u003e\u003ctr\u003e\u003cth align=\"left\" colname=\"c1\"\u003e\u003cp\u003eWelded joints\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c2\"\u003e\u003cp\u003eUTS /MPa\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c3\"\u003e\u003cp\u003eEF /%\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c4\"\u003e\u003cp\u003eWelded joints\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c5\"\u003e\u003cp\u003eUTS /MPa\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c6\"\u003e\u003cp\u003eEF /%\u003c/p\u003e\u003c/th\u003e\u003c/tr\u003e\u003c/thead\u003e\u003ctbody\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e250 ℃\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e230.3\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e5.3\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c4\"\u003e\u003cp\u003e250\u0026deg;C\u0026thinsp;+\u0026thinsp;Er/Zr\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c5\"\u003e\u003cp\u003e240.1\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c6\"\u003e\u003cp\u003e5.5\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e280 ℃\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e241.5\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e5.6\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c4\"\u003e\u003cp\u003e280\u0026deg;C\u0026thinsp;+\u0026thinsp;Er/Zr\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c5\"\u003e\u003cp\u003e259.0\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c6\"\u003e\u003cp\u003e10.9\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e300 ℃\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e231.1\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e9.0\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c4\"\u003e\u003cp\u003e-\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c5\"\u003e\u003cp\u003e-\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c6\"\u003e\u003cp\u003e-\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e400 ℃\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e175.5\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e19.1\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c4\"\u003e\u003cp\u003e-\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c5\"\u003e\u003cp\u003e-\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c6\"\u003e\u003cp\u003e-\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003c/tbody\u003e\u003c/colgroup\u003e\u003c/table\u003e\u003c/div\u003e\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003eThe welded joints produced with AlSi10Mg-Er-Zr filler powder demonstrated markedly enhanced tensile properties compared to those employing AlSi10Mg powder. These joints at 280 ℃ PWHT exhibited a 7.2% increase in UTS accompanied by a remarkable 94.6% improvement in EF. The optimal mechanical performance was attained in joints treated with both 280 ℃ PWHT and AlSi10Mg-Er-Zr filler powder, achieving peak values of 259.0 MPa UTS and favorable 10.9% elongation. These property enhancements stem from the synergistic interaction between PWHT and the AlSi10Mg-Er-Zr filler powder, which promoted favorable microstructural alterations in the welds of L-PBF AlSi10Mg alloys.\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec11\" class=\"Section3\"\u003e\u003ch2\u003e3.2.3 Fracture behavior of the welded joints\u003c/h2\u003e\u003cp\u003eThe welded joints subjected to 250 ℃, 280 ℃, and 300 ℃ PWHT consistently fractured near the fusion boundary, while those treated at 400 ℃ PWHT fractured in the BM, in agreement with the corresponding microhardness profiles shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e. Tensile testing results showed distinct fracture behaviors based on PWHT temperatures. The fracture surfaces of failed joints treated between 250 ℃ and 300 ℃ displayed similar morphological characteristics, showing neither visible necking nor pronounced plastic deformation. These surfaces revealed distinct hydrogen porosity features with maximum pore diameters under 150 \u0026micro;m, as illustrated in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e(a)-(e). Fracture mechanisms in these joints were primarily controlled by microstructural factors and pre-existing porosity, depending on the pore size distribution and morphological characteristics. Of particular significance, welded joints treated at 280 ℃ PWHT with AlSi10Mg-Er-Zr filler powder exhibited superior fracture surface quality, characterized by reduced porosity density and diminished presence of large pores, evident in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e(d). In contrast, the joints treated at 400 ℃ PWHT presented fundamentally different fracture characteristics, shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e(f). These surfaces completely lacked hydrogen pores while demonstrating unambiguous macroscopic necking, providing definitive evidence of considerable plastic deformation preceding final fracture.\u003c/p\u003e\u003cp\u003eHigher-magnification SEM images of fracture surfaces of the welded joints treated at 280 ℃ and 400 ℃ PWHT are shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e (g) and Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e (h), respectively. The welded joints at 280℃ PWHT exhibited numerous spherical pores with 10\u0026ndash;20 \u0026micro;m in diameter surrounded by elongated dimples on their fracture surfaces. These pores promoted local brittle fracture behavior through stress concentration during loading, initiating microcracks that preferentially formed at pore sites under maximum stress and propagated to adjacent pores [\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e]. The significant hydrogen porosity also reduced the effective load-bearing cross-section, creating localized stress concentrations that accelerated crack propagation [\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e] and consequently diminished both strength and ductility. While the influence of initial porosity on damage mechanisms in L-PBF AlSi10Mg alloys and their welded joints remains incompletely understood, current evidence suggests that it depends critically on pore size and morphology [\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e]. In contrast, the welded joints subjected to 400 ℃ PWHT exhibited completely different fracture morphology, characterized by uniformly distributed large equiaxed dimples covering the entire fracture surface, which clearly confirmed highly ductile fracture behavior.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003c/div\u003e\u003cdiv id=\"Sec12\" class=\"Section2\"\u003e\u003ch2\u003e3.3 PWHT influence on joint microstructures\u003c/h2\u003e\u003cdiv id=\"Sec13\" class=\"Section3\"\u003e\u003ch2\u003e3.3.1 Si-rich eutectic of the welds\u003c/h2\u003e\u003cp\u003eThe critical influence of silicon (Si) element on both the printability and mechanical properties of L-PBF AlSi10Mg alloys has been well documented [\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e, \u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e]. In as-built state, the BM displays a homogeneous distribution of continuous Si-rich eutectic networks in aluminum (Al) matrix, featuring fine α-Al cells measuring 0.5\u0026thinsp;~\u0026thinsp;1 \u0026micro;m in size. Following welding, the welds revealed distinct modifications to the Si-rich eutectic networks across the different PWHT temperatures, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e. It can be observed that both the bottom and upper regions of the welds were composed of α-Al phases and Si-rich eutectic phases, maintaining the same phase composition as the BM. However, differences in the morphology and connectivity of the Si-rich eutectic between the bottom region and the upper region was significant. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e (a)-(d), the bottom region subjected to all the PWHT temperatures of 250 ℃~400 ℃ contained plate-like or rod-like Si-rich eutectic morphology, demonstrating significant breakdown of the original cellular networks. In contrast, the upper regions shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e (a')-(d') preserved well-interconnected Si-rich networks, displaying finer morphological features and enhanced continuity relative to the bottom regions. This microstructural variation stems from different cooling conditions during the LMD welding owing to an increased travel speed of 15 mm/s in the upper region compared to 10 mm/s in the bottom region. As documented in previous studies [\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e, \u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e, \u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e], such higher welding speeds accelerated cooling promotes refinement of Si-rich networks in L-PBF AlSi10Mg alloys. High-resolution characterization revealed that α-Al cells in the upper region retained an elliptical morphology with characteristic dimensions of 2\u0026ndash;6 \u0026micro;m, irrespective of applied PWHT temperatures. Conversely, the bottom region exhibited fragmented Si-rich networks featuring plate-like or rod-like structures across all PWHT temperatures. Remarkably, the Si-rich eutectic morphology in both regions demonstrated minimal dependence on PWHT temperature variations. These observations collectively suggest that welding parameters, rather than PWHT, served as the dominant factor controlling microstructural evolution in both the bottom and upper regions.\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec14\" class=\"Section3\"\u003e\u003ch2\u003e3.3.2 Grain size and grain boundary feature of the welds\u003c/h2\u003e\u003cp\u003eFigure \u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003e presents the inverse pole figure (IPF) maps and grain size distribution of α-Al grains in the upper region of welds produced with AlSi10Mg filler powder subjected to 280 ℃ and 400 ℃ PWHT. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003e(a), the weld microstructure at 280 ℃ PWHT primarily consisted of randomly oriented equiaxed and ultrafine grains. In contrast, the weld treated at 400 ℃ exhibited predominantly columnar grains, as seen in Fig.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003e(b). The average equivalent grain sizes were 84.2 \u0026micro;m and 116.3 \u0026micro;m for the welds at 280 ℃ and 400 ℃ PWHT, respectively. Additionally, the weld treated at 280 ℃ contained a significantly higher fraction of fine grains below 50 \u0026micro;m, measuring 81.9%, compared to 74.2% in the weld at 400 ℃ PWHT. This grain size distribution clearly demonstrates the superior grain refinement efficacy of lower temperature PWHT at 280 ℃ relative to higher temperature PWHT at 400 ℃.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003eFigure \u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003e displays the grain boundary distribution and misorientation angle histograms for welds subjected to PWHT at 280 ℃ and 400 ℃, where high-angle grain boundaries (HAGBs, \u0026gt;\u0026thinsp;15\u0026deg;) and low-angle grain boundaries (LAGBs, 2\u0026deg;\u0026ndash;15\u0026deg;) are marked by blue and red lines, respectively. The weld at 280 ℃ PWHT, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003e (a), and Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003e (b), revealed a predominance of HAGBs of 54.0%. In contrast, the 400 ℃ PWHT weld in Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003e(c) and Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003e(d) exhibited a higher concentration of LAGBs, resulting in a reduced HAGB fraction of only 40.8%. This microstructural evolution was further quantified through misorientation angle analysis, where the average angle decreased significantly from 29.6\u0026deg; for the 280 ℃ PWHT to 19.2\u0026deg; at 400 ℃ PWHT, demonstrating a pronounced temperature dependence of grain boundary characteristics. The increased HAGB fraction in the weld at 280 ℃ PWHT can be attributed to the prevalence of finely equiaxed grains, suggesting that lower-temperature PWHT promotes the formation of such boundaries. This is particularly noteworthy since HAGBs are known to effectively hinder crack propagation [\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e, \u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e], which explains one of the reasons why the enhanced mechanical properties observed in the weld at 280 ℃ PWHT compared to those treated at 400 ℃.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003c/div\u003e\u003cdiv id=\"Sec15\" class=\"Section2\"\u003e\u003ch2\u003e3.4 Role of Er and Zr addition in the welded joints\u003c/h2\u003e\u003cdiv id=\"Sec16\" class=\"Section3\"\u003e\u003ch2\u003e3.4.1 Microstructure characteristics at fusion boundary\u003c/h2\u003e\u003cp\u003eAs demonstrated in previous results, all welded joints subjected to 250 ℃, 280 ℃ and 300 ℃ PWHT consistently fractured near the fusion boundary within the weld, confirming this area as the joint's mechanically weakest region. Figure\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003e provides a comprehensive microstructural analysis of the fusion boundary in welded joints produced with both AlSi10Mg and AlSi10Mg-Er-Zr filler powders at 280 ℃ PWHT.\u003c/p\u003e\u003cp\u003eThe microstructural examination revealed distinct features along the fusion boundary of the welds produced with both AlSi10Mg and AlSi10Mg-Er-Zr filler powders. Contrary to conventional expectations, epitaxial growth from the BM was not observed. Instead, a well-defined narrow band of non-dendritic equiaxed grains (EQZ) formed consistently along the fusion boundary, with measured widths ranging from 40.2 to 78.1 \u0026micro;m for both welds with the different filler powders, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003e (a) and Fig.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003e (b). This EQZ formation resulted primarily from heterogeneous nucleation on Al\u003csub\u003e3\u003c/sub\u003eTi, Al\u003csub\u003e3\u003c/sub\u003eEr, and Al\u003csub\u003e3\u003c/sub\u003e(Er, Zr) particles originating from both the BM and filler powders [\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e]. The comparable EQZ widths between both filler powders was similar, indicating that Er and Zr additions primarily influenced microstructural refinement rather than EQZ formation itself. Beyond the EQZ region, both welds close to the EQZ exhibited comparable microstructural characteristics, featuring cellular dendritic and equiaxed structures. However, the welds produced with AlSi10Mg-Er-Zr filler powder displayed significantly refined dendritic structures near the EQZ boundary, clearly demonstrating the grain-refining capability of Er and Zr additions. More importantly, welds produced with the AlSi10Mg-Er-Zr filler powder showed a substantial reduction in both porosity density and pore size distribution along the fusion boundary. These synergistic microstructural enhancements, which include improved grain refinement at the fusion boundary and superior pore morphology control, collectively partially contributed to the observed increase in tensile strength of the welded joints produced using the AlSi10Mg-Er-Zr filler powder.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec17\" class=\"Section3\"\u003e\u003ch2\u003e3.4.2 Optical microstructure of the welds\u003c/h2\u003e\u003cp\u003eThe influence of Er and Zr additions on Si-rich eutectic of the welds was investigated comparing the eutectic morphologies in the bottom and upper regions obtained with AlSi10Mg and AlSi10Mg-Er-Zr filler powders at 280\u0026deg;C PWHT. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003e, a clear morphological difference in the Si-rich eutectic phases was observed between the upper and lower regions. In the bottom region, most of the eutectic Si-rich networks were fragmented in the Al matrix due to thermal accumulation experienced from the 2nd to 5th tracks in the upper region. The fragmented Si-rich phases predominantly exhibited coarse rod-like or bar-like structures, with a minor fraction displaying particulate morphology. In contrast, the upper region exhibited a well-developed cellular dendrite consisting of large α-Al cells and finer Si-rich eutectic. The Si-rich eutectic in this region appeared as significantly refined particulates distributed along the α-Al dendrites. Consequently, both the Si-rich eutectic and α-Al cells underwent substantial refinement in the upper region compared to the bottom region of the welds. Notably, the addition of Er and Zr in the AlSi10Mg powders had a limited influence on the cellular dendritic structures in the welds. Both in the bottom and upper regions, the microstructures produced using AlSi10Mg filler powder were remarkably similar to those obtained with AlSi10Mg-Er-Zr filler powders. This observation indicated that regional variations in welding heat input exerted a more significant influence on microstructure evolution than did the composition of the filler powder. However, the secondary dendrite arm spacing (SDAS) in the upper region produced with AlSi10Mg-Er-Zr powder was smaller than those made with AlSi10Mg powder, indicating that the addition of Er and Zr in the filler powder promoted dendritic refinement under identical welding heat input conditions.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec18\" class=\"Section3\"\u003e\u003ch2\u003e3.4.3 Si-rich eutectic feature in the welds\u003c/h2\u003e\u003cp\u003eHigh-magnification SEM micrographs were analyzed to compare the Si-rich eutectic structures in welds produced at 280\u0026deg;C PWHT using AlSi10Mg and AlSi10Mg-Er-Zr filler powders, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig12\" class=\"InternalRef\"\u003e12\u003c/span\u003e. The bottom regions of both welds exhibited similar morphologies of coarse rod-like and plate-like Si-rich eutectic structures, confirming the optical observations. However, the Si-rich eutectic revealed markedly different spacing characteristics in both the bottom and upper regions. In the bottom regions, welds produced with AlSi10Mg filler powder showed significantly coarser spacing compared to those produced with AlSi10Mg-Er-Zr filler powder. While the welds produced with AlSi10Mg filler powder produced an average spacing of 10.1 \u0026micro;m, the welds with AlSi10Mg-Er-Zr filler powder resulted in substantially finer spacing of 5.9 \u0026micro;m. In the upper regions, both filler powders formed continuous Si-rich eutectic networks, presenting a distinct contrast to the bottom region's microstructure, proving the optical observational results. Corresponding α-Al cell sizes measured 6.1 \u0026micro;m for AlSi10Mg powder and 4.2 \u0026micro;m for AlSi10Mg-Er-Zr filler. Thus, the α-Al cell sizes followed a similar refinement trend, with the AlSi10Mg-Er-Zr filler powder producing finer cellular structures. Therefore, it is clear that the AlSi10Mg-Er-Zr filler powder effectively refined the α-Al and Si-rich eutectic phases in both the bottom region and the upper regions of the welds. Notably, the Si-rich eutectic particles were predominantly distributed along α-Al cell boundaries, forming three-dimensional networks characteristic of L-PBF AlSi10Mg alloys [\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e, \u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e]. These results demonstrate that while Er and Zr additions effectively refined both the Si-rich eutectic and α-Al structure, they did not substantially alter the fundamental morphology or connectivity of the Si-rich eutectic within corresponding weld regions.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec19\" class=\"Section3\"\u003e\u003ch2\u003e3.4.4 Grain boundary characteristics\u003c/h2\u003e\u003cp\u003eThe grain boundary distribution of α-Al structures in the welds produced using AlSi10Mg and AlSi10Mg-Er-Zr filler powders was analyzed. The grain boundary distribution and misorientation angle histograms for the upper region of both welds subjected to 280 ℃ PWHT are presented in Fig.\u0026nbsp;\u003cspan refid=\"Fig13\" class=\"InternalRef\"\u003e13\u003c/span\u003e, where the low-angle grain boundaries (LAGBs, 2\u0026deg;\u0026ndash;15\u0026deg;) were marked in red, and high-angle grain boundaries (HAGBs, 15\u0026deg;\u0026ndash;65\u0026deg;) were highlighted in blue. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig13\" class=\"InternalRef\"\u003e13\u003c/span\u003e(a) and Fig.\u0026nbsp;\u003cspan refid=\"Fig13\" class=\"InternalRef\"\u003e13\u003c/span\u003e (b), both welds exhibit a predominance of blue-colored boundaries, indicating that HAGBs dominate the microstructure. However, the fraction of HAGBs differs between the two welds: 53.9% for the AlSi10Mg weld and 61.3% for the AlSi10Mg-Er-Zr weld, demonstrating an increase in HAGBs for the welds produced with AlSi10Mg-Er-Zr filler powders. Furthermore, the average grain misorientation angles were 29.6\u0026deg; and 33.8\u0026deg; for the AlSi10Mg and AlSi10Mg-Er-Zr welds, respectively. These results clearly indicate that the AlSi10Mg-Er-Zr filler powder promotes higher grain boundary misorientation angles.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003eRecent studies have recognized grain boundary engineering (GBE) as a promising strategy for improving the properties of LPBF components [\u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e]. Among various grain boundary characteristics, coincidence site lattice (CSL) boundaries are particularly noteworthy due to their substantial impact on the mechanical performance of polycrystalline materials [\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e]. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig13\" class=\"InternalRef\"\u003e13\u003c/span\u003e(c) and Fig.\u0026nbsp;\u003cspan refid=\"Fig13\" class=\"InternalRef\"\u003e13\u003c/span\u003e (d), both welds exhibited a pronounced peak at approximately 60\u0026deg;, corresponding to the Σ3 (60\u0026deg; \u0026lt;111\u0026gt;) grain boundary, a common CSL boundary in face-centered cubic (FCC) aluminum alloys. This indicates an increased fraction of CSL boundaries in the welds using both AlSi10Mg and AlSi10Mg-Er-Zr filler powders. Quantitative analysis revealed that the Σ3 boundary fraction reached 17.2% in the AlSi10Mg weld and 24.9% in the AlSi10Mg-Er-Zr weld, demonstrating a significant enhancement of Σ3 grain boundaries in the welds as a result of the addition of Er and Zr in the AlSi10Mg filler powder.\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec20\" class=\"Section3\"\u003e\u003ch2\u003e3.4.5 Schmid factor\u003c/h2\u003e\u003cp\u003eThe Schmid factor (SF) is a crucial parameter governing slip system activation during tensile loading, with lower SF values indicating greater resistance to dislocation motion [\u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e]. Generally, grains with higher SF values possess softer orientations that promote crack propagation [\u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e]. Figure\u0026nbsp;\u003cspan refid=\"Fig14\" class=\"InternalRef\"\u003e14\u003c/span\u003e presents the SF distributions for welds produced using AlSi10Mg and AlSi10Mg-Er-Zr filler powders, showing a predominant concentration of SF values in the range of 0.3\u0026ndash;0.5.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003eDetailed microstructural analysis demonstrated a marked difference in grain orientation characteristics of the welds between the AlSi10Mg and AlSi10Mg-Er-Zr filler powders. In the AlSi10Mg welds, merely 14.1% of grains exhibited SFs within the low range of 0.3 to 0.4. This value underwent a remarkable increase to 34.8% when examining welds produced with the AlSi10Mg-Er-Zr filler powder. Such a pronounced enhancement in the population of grains with low SFs provided conclusive evidence that the addition of Er and Zr in the AlSi10Mg filler powder during LMD welding effectively promotes the development of crystallographically hard-oriented grains in the welds. The increased population of hard-oriented grains enhances both load-bearing capacity through improved stress distribution and resistance to slip system activation during deformation.\u003c/p\u003e\u003c/div\u003e\u003c/div\u003e"},{"header":"4. Conclusions","content":"\u003cp\u003eButt joints of 3.0 mm thick sheets of L-PBF AlSi10Mg alloys have been produced using LMD welding with both AlSi10Mg and AlSi10Mg-Er-Zr filler powders. The porosity characteristics, mechanical performance, hardness distribution and microstructural evolution of the welded joints were investigated. On the basis of the present results, the following conclusions can be reached.\u003c/p\u003e\u003cp\u003e(1) All welded joints produced a fully penetrated weld with different welding modes between the upper and lower regions. As the PWHT temperature increased from 250 ℃ to 400 ℃, the hydrogen porosity in the welds gradually decreased from 3.2\u0026ndash;0.7%, and the maximum pore diameter also decreased from 276.7 \u0026micro;m to 131.0 \u0026micro;m. The AlSi10Mg-Er-Zr filler powder slightly reduced the porosity in the welds. Moreover, the upper region showed lower porosity and smaller hydrogen pore sizes compared to the bottom region.\u003c/p\u003e\u003cp\u003e(2) The welded joints subjected to PWHT at 280 ℃, 300 ℃, and 400 ℃ in the bottom region show an inverted U-shaped hardness profile, while a U-shaped hardness profile was observed for the joints treated at 250 ℃ PWHT due to varying degrees of hardness reduction in the BM. The upper region of the welds subjected to different PWHT temperatures had higher hardness compared to the lower region. When PWHT was conducted at 250 ℃, 280 ℃, and 300 ℃, the resulting welded joints produced with both AlSi10Mg and AlSi10Mg-Er-Zr filler powders exhibited different locations of minimum hardness from those treated at 400 ℃.\u003c/p\u003e\u003cp\u003e(3) Welded joints subjected to 250 ℃, 280 ℃ and 300 ℃ PWHT showed high UTS ranging from 230 MPa to 240 MPa. However, the UTS significantly decreased for the joints treated at 400 ℃ PWHT. The EF increased significantly from 5.3\u0026ndash;19.1% as the PWHT temperature rose from 250 ℃ to 400 ℃. Compared to AlSi10Mg powder, those joints at 250 ℃ and 280 ℃ PWHT with AlSi10Mg-Er-Zr powder demonstrated an increased UTS by 4.3% and 7.2%, and improvements in EF by 3.8% and 94.6%, respectively. The welded joints at 280\u0026deg;C PWHT in combination with AlSi10Mg-Er-Zr filler powder exhibited the highest UTS of 259.0 MPa along with an excellent EF of 10.9%. The tensile fracture of the welded joints at 250 ℃, 280 ℃, and 300 ℃ PWHT was governed by microstructure and initial porosity, while the welded joints at 400 ℃ PWHT was characterized by highly ductile fracture behavior.\u003c/p\u003e\u003cp\u003e(4) The Si-rich eutectic in the upper region of the welds subjected to 250\u0026thinsp;~\u0026thinsp;400 ℃ PWHT was well interconnected, exhibiting finer and better connectivity of the Si-rich eutectic compared to the bottom region. Despite this, the morphology and size of the Si-rich eutectic in the bottom and upper regions of the welds did not significantly change with increasing PWHT temperatures. Both the lower and upper regions of the welds produced with AlSi10Mg-Er-Zr powder exhibited a more uniform and finer Si-rich eutectic structure along with refined α-Al cells. The welded joints subjected to 280 ℃ PWHT demonstrated superior grain refinement efficacy compared to 400 ℃ PWHT, producing a significantly higher fraction of HAGBs at 54.0% relative to 40.8% at the elevated temperature. This substantial difference establishes a clear inverse relationship between PWHT temperature and HAGB formation.\u003c/p\u003e\u003cp\u003e(5) A distinct narrow band of the EQZ ranging from 40.2 to 78.1 \u0026micro;m was clearly observed along fusion boundary of the welded joints, while LMD welding using AlSi10Mg-Er-Zr filler powder produced minimal influence on the EQZ width. The welds produced with AlSi10Mg-Er-Zr filler powder displayed significantly refined dendritic structures near the EQZ boundary, clearly demonstrating the grain-refining capability of Er and Zr additions. Notably, a substantial reduction in both porosity density and pore size near the fusion boundary was observed when using AlSi10Mg-Er-Zr filler powder.\u003c/p\u003e\u003cp\u003e(6) The bottom regions of both welds exhibited similar morphologies of coarse rod-like and plate-like Si-rich eutectic structures, while both filler powders formed continuous Si-rich eutectic networks in the upper regions. Corresponding α-Al cell sizes in the bottom region was 10.1 \u0026micro;m and 5.9 \u0026micro;m, and measured 6.1 \u0026micro;m and 4.2 \u0026micro;m in the upper region, for AlSi10Mg powder for AlSi10Mg-Er-Zr filler powder, respectively. Therefore, Er and Zr additions in the AlSi10Mg filler powder effectively refined both the Si-rich eutectic and α-Al structure, whereas they did not substantially alter the fundamental morphology or connectivity of the Si-rich eutectic within corresponding weld regions.\u003c/p\u003e\u003cp\u003e(7) The fraction of HAGBs was 53.9% and 61.3%, and the average grain misorientation angles were 29.6\u0026deg; and 33.8\u0026deg;, for the AlSi10Mg and AlSi10Mg-Er-Zr weld, respectively. The Σ3 boundary fraction reached 17.2% and 24.9% in the AlSi10Mg and AlSi10Mg-Er-Zr weld, respectively. Moreover, merely 14.1% of grains exhibited SFs within the low range of 0.3 to 0.4 in AlSi10Mg welds, while this value underwent a remarkable increase to 34.8% in the welds produced with the AlSi10Mg-Er-Zr filler powder. Thus, the AlSi10Mg-Er-Zr filler powder effectively increased proportions of HAGBs, Σ3 grain boundaries, and promoted the development of hard-oriented grains.\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003ch2\u003eCompeting Interests\u003c/h2\u003e\u003cp\u003eThe authors have no relevant financial or nonfinancial interests to disclose.\u003c/p\u003e\u003c/p\u003e\u003ch2\u003eFunding\u003c/h2\u003e\u003cp\u003eThis work was supported by the National Natural Science Foundation of China (Grant number 52271018).\u003c/p\u003e\u003ch2\u003eAuthor contributions\u003c/h2\u003e\u003cp\u003eYingying Liu: Formal analysis, Writing-original draft. Jingchuan Li: Formal analysis, Validation. Zhaotong Li: Formal analysis, Validation, review \u0026amp; editing. Li Cui: Conceptualization, Writing-review \u0026amp; editing, Funding acquisition. Dingyong He: Conceptualization, Supervision. Jie Xu: Conceptualization, Validation.\u003c/p\u003e\u003ch2\u003eData availability Statement\u003c/h2\u003e\u003cp\u003eThe raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\u003cli\u003e\u003cspan\u003eGu DD, Meiners W, Wissenbach K, Poprawe R (2012) Laser additive manufacturing of metallic components: materials, processes and mechanisms. Int Mater Rev 57(3):133\u0026ndash;164. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1179/1743280411Y.0000000014\u003c/span\u003e\u003cspan address=\"10.1179/1743280411Y.0000000014\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eMartin JH, Yahata BD, Hundley JM, Martin JH, Yahata BD, Hundley JM, Mayer JA, Schaedler TA, Pollock TM (2017) 3D printing of high-strength aluminium alloys. Nature 549(7672):365\u0026ndash;369. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1038/nature23894\u003c/span\u003e\u003cspan address=\"10.1038/nature23894\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eKhairallah SA, Martin AA, Lee JRI, Martin AA, Guss G, King WE (2020) Controlling interdependent meso-nanosecond dynamics and defect generation in metal 3D printing. Science 368(6491):660\u0026ndash;665. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1126/science.aay7830\u003c/span\u003e\u003cspan address=\"10.1126/science.aay7830\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eZhao L, Song LB, Mac\u0026iacute;as JGS, Zhu YX, Huang MS, Simar A, Li ZH (2022) Review on the correlation between microstructure and mechanical performance for laser powder bed fusion AlSi10Mg. Addit Manuf 56:102914. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.addma.2022.102914\u003c/span\u003e\u003cspan address=\"10.1016/j.addma.2022.102914\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eWu DF, Cui L, Wu X, Guo XY, Shao W, Tan Z, He DY (2022) Cold metal transfer welding of AlSi7Mg alloy sheets prepared by selective laser melting. J Manuf Process 84:282\u0026ndash;297. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.jmapro.2022.10.013\u003c/span\u003e\u003cspan address=\"10.1016/j.jmapro.2022.10.013\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eMichler M, Hollmann P, Zenker R, Buchwalder A (2021) Investigation of pore reduction in hybrid joints of conventionally and additively manufactured AlSi10Mg using electron beam welding. Adv Eng Mater 23(6):2001325. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1002/adem.202001325\u003c/span\u003e\u003cspan address=\"10.1002/adem.202001325\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eNahmany M, Rosenthal I, Benishti I, Frage N, Stern A (2015) Electron beam welding of AlSi10Mg workpieces produced by selected laser melting additive manufacturing technology. Addit Manuf 8:63\u0026ndash;70. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.addma.2015.08.002\u003c/span\u003e\u003cspan address=\"10.1016/j.addma.2015.08.002\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eM\u0026auml;kikangas J, Rautio T, Mustakangas A, M\u0026auml;ntyj\u0026auml;rvi K (2019) Laser welding of AlSi10Mg aluminium-based alloy produced by selective laser melting (SLM). Procedia Manuf 36:88\u0026ndash;94. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.promfg.2019.08.013\u003c/span\u003e\u003cspan address=\"10.1016/j.promfg.2019.08.013\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eBiffi C, Fiocchi J, Tuissi A (2019) Laser weldability of AlSi10Mg alloy produced by selective laser melting: microstructure and mechanical behavior. J Mater Eng Perform 28:6714\u0026ndash;6719. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1007/s11665-019-04402-7\u003c/span\u003e\u003cspan address=\"10.1007/s11665-019-04402-7\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eZhang CC, Bao YF, Zhu HH, Nie XJ, Zhang WQ, Zhang SS, Zeng XY (2019) A comparison between laser and TIG welding of selective laser melted AlSi10Mg. Opt Laser Technol 120:105696. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.optlastec.2019.105696\u003c/span\u003e\u003cspan address=\"10.1016/j.optlastec.2019.105696\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eEmmelmann C, Beckmann F (2017) Optimization of laser welding process for laser additive manufactured aluminum parts by means of beam oscillation and process-oriented component design. In: Proceedings of Lasers in Manufacturing Conference. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://www.wlt.de/lim/Proceedings2017/Data/PDF/Contribution187_final.pdf\u003c/span\u003e\u003cspan address=\"https://www.wlt.de/lim/Proceedings2017/Data/PDF/Contribution187_final.pdf\" targettype=\"URL\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eBeckmann F, Emmelmann C (2019) Optimization of the weldability of laser additive manufactured aluminum by means of hydrogen minimization in the component and welding parameter optimization. In: Proceedings of Lasers in Manufacturing Conference. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://www.wlt.de/lim/Proceedings2019/data/PDF/Contribution_327_\u003c/span\u003e\u003cspan address=\"https://www.wlt.de/lim/Proceedings2019/data/PDF/Contribution_327_\" targettype=\"URL\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e final. pdf\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eSchwarz A, Schleser M, Gerhards B, Popoola P, Gebhardt A (2021) Welding of additive manufactured AlSi10Mg: using laser welding in a vacuum for high quality weld seams-a new approach to welding LPBF manufactured AlSi10Mg. S Afr J Ind Eng 32:99\u0026ndash;112. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttp://dx.doi.org/10.7166/32-4-2513\u003c/span\u003e\u003cspan address=\"10.7166/32-4-2513\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eChen NN, Wan ZX, Wang HP, Li JJ, Yang BX, Solomon J, Carlson B (2022) Effect of ambient pressure on laser welding of AlSi10Mg fabricated by selected laser melting. Mater Des 215:110427. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.matdes.2022.110427\u003c/span\u003e\u003cspan address=\"10.1016/j.matdes.2022.110427\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eCui L, Peng ZB, Chang Y, He DY, Cao Q, Guo XY, Zeng Y (2022) Porosity, microstructure and mechanical property of welded joints produced by different laser welding processes in selective laser melting AlSi10Mg alloys. Opt Laser Technol 150:107952. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.optlastec.2022.107952\u003c/span\u003e\u003cspan address=\"10.1016/j.optlastec.2022.107952\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eWang C, He DY, Cui L, Shi ZF, Guo XY, Wu X, Zhai YY, Bu FH (2023) Improved tensile strength of welded AlSi10Mg alloys by using laser metal deposition. Sci Technol Weld Jo 28(9):946\u0026ndash;954. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1080/13621718.2023.2253591\u003c/span\u003e\u003cspan address=\"10.1080/13621718.2023.2253591\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eLi JC, Cui L, Wu DF, Wang C, Shi ZF, He DY, Cao Q (2024) Welding of additive manufacturing AlSi10Mg alloys using a laser metal deposition process with different heat input. Weld World 68(8):2083\u0026ndash;2101. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1007/s40194-024-01764-y\u003c/span\u003e\u003cspan address=\"10.1007/s40194-024-01764-y\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eChen WH, Liu FC, Niu PL, You QF, Liu FG, Wang ZT (2024) Effect of prewelding solution heat treatment on pore defect in laser welding of AlSi10Mg alloy prepared by selective laser melting. Rare Metal Mater Eng 53(04):1111\u0026ndash;1118. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.12442/j.issn.1002-185X.20230070\u003c/span\u003e\u003cspan address=\"10.12442/j.issn.1002-185X.20230070\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eWang C, He DY, He CX, Han FY, Li JC, Cui L (2024) Effect of heat treatments on microstructure and mechanical behavior of laser welded LPBF AlSi10Mg alloys. Mater Today Commun 108478. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.mtcomm.2024.108478\u003c/span\u003e\u003cspan address=\"10.1016/j.mtcomm.2024.108478\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eZheng Y, Zhao ZH, Xiong RZ, Ren GC, Yao MH, Liu W, Zang LB (2024) Effect of post heat treatment on microstructure, mechanical property and corrosion behavior of AlSi10Mg alloy fabricated by selective laser melting. Prog Nat Sci-Mater 34(1):89\u0026ndash;101. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.pnsc.2024.02.001\u003c/span\u003e\u003cspan address=\"10.1016/j.pnsc.2024.02.001\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eLong J, Zhang LZ, Zhang LL, Wang X, Zhang GF, Zhang XJ, Na SJ (2020) Effects of minor Zr addition on the microstructure and mechanical properties of laser welded joint of Al/SiCp metal-matrix composite. J Manuf Process 49:373\u0026ndash;384. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.jmapro.2019.12.004\u003c/span\u003e\u003cspan address=\"10.1016/j.jmapro.2019.12.004\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eSpierings AB, Dawson K, Kern K, Palm F, Wegener K (2017) SLM-processed Sc- and Zr- modified Al-Mg alloy: Mechanical properties and microstructural effects of heat treatment. Mater Sci Eng A 701:264\u0026ndash;273. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.msea.2017.06.089\u003c/span\u003e\u003cspan address=\"10.1016/j.msea.2017.06.089\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eLi RD, Wang MB, Yuan TC, Song B, Chen C, Zhou KC, Cao P (2017) Selective laser melting of a novel Sc and Zr modified Al-6.2Mg alloy: processing, microstructure, and properties. Powder Technol 319:117\u0026ndash;128. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.powtec.2017.06.050\u003c/span\u003e\u003cspan address=\"10.1016/j.powtec.2017.06.050\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eChen B, Wang FZ, Ren SM, Xia P, Wang ZP, Hao JF, Chen ZP, Li RX, Jiang HX (2024) Influence of erbium addition on microstructure and performances of AlSi10Mg alloy prepared by selective laser melting. J Rare Earths 42(10):1988\u0026ndash;1994. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.jre.2023.09.017\u003c/span\u003e\u003cspan address=\"10.1016/j.jre.2023.09.017\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eYang TY, Cui L, He DY, Huang H (2022) Enhancement of microstructure and mechanical property of AlSi10Mg-Er-Zr alloys fabricated by selective laser melting. Acta Metall Sin 58(9):1108\u0026ndash;1117. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://www.ams.org.cn/EN/\u003c/span\u003e\u003cspan address=\"https://www.ams.org.cn/EN/\" targettype=\"URL\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003e10.11900/0412.1961.2021.00085\u003c/span\u003e\u003cspan address=\"10.11900/0412.1961.2021.00085\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003ePeng ZB, Cui L, He DY, Guo XY, Zeng Y, Cao Q, Huang H (2022) Effect of Er and Zr addition on laser weldability of AlSi10Mg alloys fabricated by selective laser melting. Mater Charact 190:112070. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.matchar.2022.112070\u003c/span\u003e\u003cspan address=\"10.1016/j.matchar.2022.112070\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eWen H, Yves N, Matthew F, Kan WH, Nadot Y, Foley M, Ridosz L, Proust G, Cairney JM (2019) Factors that affect the properties of additively-manufactured AlSi10Mg: Porosity versus microstructure. Addit Manuf 29:100805. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.addma.2019.100805\u003c/span\u003e\u003cspan address=\"10.1016/j.addma.2019.100805\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eNunes R, Faes K, Verlinde W, Probst F, Freese JD, Arnhold J, Waele WD, Lezaack M, Simar A (2025) Welding of additively manufactured aluminium alloy components: challenges and mitigation strategies. Weld World 69:1003\u0026ndash;1021. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1007/s40194-025-01923-9\u003c/span\u003e\u003cspan address=\"10.1007/s40194-025-01923-9\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eWang CG, Zhu JX, Wang GW, Qin Y, Sun MY, Yang JL, Shen XF, Huang SK (2022) Effect of building orientation and heat treatment on the anisotropic tensile properties of AlSi10Mg fabricated by selective laser melting. J Alloy Compd 895:162665. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.jallcom.2021.162665\u003c/span\u003e\u003cspan address=\"10.1016/j.jallcom.2021.162665\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eAlbu M, Krisper R, Lammer J, Kothleitner G, Fiocchi J, Bassani P (2020) Microstructure evolution during in-situ heating of AlSi10Mg alloy powders and additive manufactured parts. Addit Manuf 36:101605. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.addma.2020.101605\u003c/span\u003e\u003cspan address=\"10.1016/j.addma.2020.101605\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eChen T, Wang L, Tan S (2016) Effects of vacuum annealing treatment on microstructures and residual stress of AlSi10Mg parts produced by selective laser melting process. Mod Phys Lett B 30(19):1650255. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1142/S0217984916502559\u003c/span\u003e\u003cspan address=\"10.1142/S0217984916502559\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eWang M, Song B, Wei Q, Zhang Y, Shi Y (2019) Effects of annealing on the microstructure and mechanical properties of selective laser melted AlSi7Mg alloy. Mater Sci Eng A 739:463\u0026ndash;472. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.msea.2018.10.047\u003c/span\u003e\u003cspan address=\"10.1016/j.msea.2018.10.047\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eKotadia HR, Gibbons G, Das A, Howes PD (2021) A review of laser powder bed fusion additive manufacturing of aluminium alloys: microstructure and properties. Addit Manuf 46:102155. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.addma.2021.102155\u003c/span\u003e\u003cspan address=\"10.1016/j.addma.2021.102155\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eSnopinski P, Liveric L, Beniak J (2024) Annealing effects on microstructure and texture in KOBO-Processed LPBF AlSi10Mg alloy: elucidating CSL boundary formation. Symmetry 16:1663. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.3390/sym16121663\u003c/span\u003e\u003cspan address=\"10.3390/sym16121663\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eCabrera-Correa L, Gonzalez-Rovira L, De DLJ, Castillo-Rodriguez M, Botana FJ (2023) Effect of the heat treatment on the mechanical properties and microstructure of scalmalloy manufactured by selective laser melting (SLM) under certified conditions. Mater Charact 196:112549. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.matchar.2022.112549\u003c/span\u003e\u003cspan address=\"10.1016/j.matchar.2022.112549\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eMa M, Wang B, Liu H, Yi D, Shen F, Zhai T (2020) Investigation of fatigue crack propagation behavior of 5083 aluminum alloy under various stress ratios: role of grain boundary and Schmid factor. Mater Sci Eng A 773:138871. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.msea.2019.138871\u003c/span\u003e\u003cspan address=\"10.1016/j.msea.2019.138871\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003cli\u003e\u003cspan\u003eSun G, Wang C, Wei X, Shang D, Chen S (2019) Study on small fatigue crack initiation and growth for friction stir welded joints. Mater Sci Eng A 739:71\u0026ndash;85. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.msea.2018.10.002\u003c/span\u003e\u003cspan address=\"10.1016/j.msea.2018.10.002\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e\u003c/span\u003e\u003c/li\u003e\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":false,"hideJournal":true,"highlight":"","institution":"","isAcceptedByJournal":false,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true},"keywords":"L-PBF AlSi10Mg alloys, Laser metal deposition (LMD) welding, Pre-weld heat treatment (PWHT), AlSi10Mg-Er-Zr filler powder, hydrogen porosity, microstructure improvement","lastPublishedDoi":"10.21203/rs.3.rs-7438037/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-7438037/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eThis study investigates the influence of pre-weld heat treatment (PWHT) temperatures (250\u0026thinsp;~\u0026thinsp;400 ℃) on laser metal deposition (LMD) welding of laser powder bed fusion (L-PBF) AlSi10Mg alloys using novel AlSi10Mg-Er-Zr filler powders, with particular emphasis on porosity characteristics, mechanical performance, and microstructural evolution of the welded joints. Results show that elevated PWHT temperatures effectively mitigated hydrogen porosity, reducing both pore density and maximum diameter. The synergistic combination of low PWHT temperatures at 250\u0026thinsp;~\u0026thinsp;280 ℃ using AlSi10Mg-Er-Zr filler powder demonstrated superior mechanical enhancement, achieving 4.3\u0026thinsp;~\u0026thinsp;7.2% improvement in ultimate tensile strength (UTS) and 3.8\u0026thinsp;~\u0026thinsp;94.6% increase in elongation at fracture (EF) compared to the conventional AlSi10Mg filler powder. Welded joints produced using AlSi10Mg-Er-Zr filler powder at 280 ℃ PWHT yielded the optimal mechanical performance, achieving a remarkable balance between strength and ductility with an UTS of 259.0 MPa and EF reaching 10.9%. The synergistic combination of AlSi10Mg-Er-Zr filler powder and PWHT not only enhanced microstructure improvement through substantial grain refinement, increased proportions of high-angle grain boundaries (HAGBs), Σ3 grain boundaries, and hard-oriented grains, but also reduced hydrogen porosity in the welds, collectively contributing to the superior mechanical performance of L-PBF AlSi10Mg welded joints.\u003c/p\u003e","manuscriptTitle":"Influence of pre-weld heat treatment temperatures and AlSi10Mg-Er-Zr filler powder on microstructure and mechanical properties of welded joints produced using laser metal deposition for L-PBF AlSi10Mg alloys","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-10-06 20:37:05","doi":"10.21203/rs.3.rs-7438037/v1","editorialEvents":[{"type":"communityComments","content":0}],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true}}],"origin":"","ownerIdentity":"54b79674-38b1-4829-89ed-9b41e1f0656e","owner":[],"postedDate":"October 6th, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"posted","subjectAreas":[],"tags":[],"updatedAt":"2025-11-10T05:40:06+00:00","versionOfRecord":[],"versionCreatedAt":"2025-10-06 20:37:05","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-7438037","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-7438037","identity":"rs-7438037","version":["v1"]},"buildId":"8U1c8b4HqxoKbykW_rLl7","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}