Strengthening and toughening mechanism of Mg–RE–Al alloy repair welded joints by suppressing abnormal grain growth during heat treatment

Strengthening and toughening mechanism of Mg–RE–Al alloy repair welded joints by suppressing abnormal grain growth during heat treatment | 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 Strengthening and toughening mechanism of Mg–RE–Al alloy repair welded joints by suppressing abnormal grain growth during heat treatment Zhengtao Liu, Lei Wang, Hongyang Duan, Sicong Zhao, Erjun Guo, and 1 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-7018798/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 Mg–RE alloys refined with Al demonstrate excellent grain thermal stability, enhancing the mechanical properties of repair welded joints after heat treatment. This research examines the impact of heat treatment on the microstructure and mechanical properties of repair welded joints composed of Mg–4.02Y–3.06Nd–1.56Al alloy. The findings indicate that the microstructure of the repaired joint comprises primarily of an equiaxed α–Mg matrix, an Mg–RE eutectic structure, particle and needle–like Al–RE phases. The particle and needle–like Al–RE phases have a high–melting–point and do not readily dissolve during the heat treatment process. Eliminating the Mg–RE eutectic structure is crucial for improving the mechanical properties of the repair welded joint. If the solution temperature is too low (500°C), an excessive amount of the residual Mg–RE eutectic structure will hinder the improvement of mechanical properties. At a solution temperature of 550°C, although the pinning effect of Al 2 RE on the grain boundary inhibits abnormal grain growth (AGG), localized AGG still occurs at the FZ edge. At 525°C, the eutectic structure fully dissolved, and no AGG occurred, achieving an optimal solid solution influence. Subsequently, after aging at 200°C for 16 h, a high density β' strengthening phase uniformly precipitated in the matrix, markedly enhancing the alloy's strength. As a result, the joint's UTS (292 MPa), YS (215 MPa), and EL (5.8%) were comparable to those of the original, unrepaired alloy. Magnesium alloy Repair welding Heat treatment Mechanical properties Grain thermal stability 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 1. Introduction Magnesium alloy as the lightest metal structure material, is known for its low density and high specific strength [ 1 , 2 ]. In particular, the incorporation of rare earth (RE) elements further enhances the mechanical properties and high–temperature heat resistance of magnesium alloys, making them widely used in aerospace [ 3 , 4 ]. In industrial production, large and complex Mg–RE alloy components are commonly produced via sand casting [ 5 , 6 ]. Due to the shortcomings of sand casting, such as slow heat dissipation and long solidification time, as well as the high density and easy segregation of RE elements, complex Mg–RE alloy castings are susceptible to defects [ 7 , 8 ]. Therefore, to improve casting yield, it is crucial to repair defects through welding. Tungsten inert gas (TIG) welding is commonly employed for defect repair in large alloy castings due to its simplicity, high flexibility, and low cost [ 9 , 10 ]. Meng et al. [ 11 ] applied TIG welding to repair casting defects in Mg–Gd–Y–Zr alloy, achieving a maximum joint efficiency (the strength ratio of the repair weld to the base metal) of 99.7%. To further enhance the properties of the repair welded joint, Tong et al. [ 9 ] optimized the welding current, resulting in the Mg–4Y–2Nd–1Gd–0.5Zr alloy repair welded joint with no apparent repair defects at a current of 170 A. The ultimate tensile strength (UTS) was 195 MPa, and the elongation (EL) was 5.1%. However, continuous heat input promotes the formation of more continuous eutectic structure at the grain boundary of the heat–affected zone (HAZ), which increases the likelihood of crack initiation and causes a significant decline of mechanical properties in the repair welded joint. Subsequently, Wang et al. [ 12 ] applied pulse current during the repair welding of Mg–3Nd–3Gd–0.2Zn–0.5Zr alloy. The results indicate that applying pulse current further refines the grain size within the fusion zone (FZ). When the pulse current frequency is 10 Hz, the repair welded joint achieves excellent comprehensive mechanical properties are obtained, UTS, yield strength (YS), and EL are 220 ± 14 MPa, 129 ± 9 MPa, and 11.9 ± 1.2%. The maximum joint efficiency reaches 88.4%. However, the eutectic structure at the HAZ grain boundary remains coarse and continuous, negatively affecting the mechanical properties of the repair welded joints. Consequently, optimizing the welding process can substantially enhance the mechanical properties of the repair welded joint. However, for large sand–cast Mg–RE alloys, a slow cooling rate during casting promotes the formation of coarse and sharp eutectic structures at the alloy grain boundaries. These eutectic structures are prone to stress concentration during tensile testing, leading to brittle fracture and limiting the improvement of alloy properties [ 13 , 14 ]. Particularly after TIG welding, the eutectic structure in the HAZ becomes more continuous, increasing the probability of crack initiation and further degrading the properties of the repair welded joint. Therefore, eliminating eutectic structures at the grain boundaries is essential for amplifying the overall mechanical properties of repair welded joints. Solid solution and aging treatment are conventional methods for eliminating brittle eutectic structures and enhancing the mechanical properties of as–cast Mg–RE alloys [ 15 , 16 ]. Therefore, studying the heat treatment parameters of Mg–RE alloy repair welded joints is crucial for enhancing joint properties. Zhang et al. [ 17 ] investigated the heat treatment parameters of Mg–4Y–3RE–0.5Zr alloy repair welded joints. The results indicate that the alloy exhibits poor grain thermal stability. Although the eutectic structure in the joint can be thoroughly dissolved by a 520°C × 8 h T4 treatment, some grains in the molten pool region undergo abnormal grain growth (AGG). These coarse grains are prone to cracking during the tensile process, causing a sharp decrease in the joint's EL to 1.8%. To address the conflict between eutectic dissolution and AGG, Tong et al. [ 18 ] developed a combined solid solution treatment process, which involves 520°C × 8 h treatment before welding and 520°C × 2 h treatment after welding, successfully preparing the Mg–4Y–2Nd–1Gd–0.5Zr alloy repair welded joint without residual eutectic or AGG. After aging treatment, the UTS of the joint reached 314 MPa, and the EL increased to 4.4%. However, for mass–produced Mg–RE alloy castings, this combined solution treatment process is both cumbersome and costly, hindering its practical industrial application. Therefore, demonstrating the feasibility of a single–stage solid solution process is crucial, while addressing the conflict between AGG formation in the molten pool and the dissolution of the eutectic structure. This is beneficial for enhancing the properties of repaired joints and reducing industrial production costs. Previous research has indicated that the addition of Al refines cast Mg–RE alloys, under elevated temperatures (preceding α–Mg solidification), the in–situ formation of Al 2 RE particles serves as heterogeneous nucleation sites, successfully refining the grain size and achieving a refinement effect similar to that of Zr [ 19 – 23 ]. Additionally, numerous high–melting–point Al–RE phases (such as Al 2 RE, Al 11 RE 3 ) form near the grain boundaries of Mg–RE alloys during the eutectic stage [ 24 , 25 ]. According to Zener pinning theory [ 26 ], these Al–RE phases can pin the grain boundaries, hindering their movement and thus improving the alloy's thermal stability. Therefore, this research aims to utilize the excellent thermal stability of Mg–RE–Al alloys to address the challenge of balancing AGG and eutectic structure dissolution imbalance caused by the inadequate thermal stability of traditional Mg–RE alloy repair welded joints in single-stage solution treatment. In our previous research [ 27 ], we successfully utilized the excellent thermal stability of the Mg–4Y–3Nd–1.5Al alloy (developed from the WE43 alloy by substituting Zr with Al as the alloy refiner) to mitigate grain growth in the HAZ during repair welding, thus preventing significant property loss. Therefore, this study systematically investigates the influences of various heat treatment parameters on the microstructure and mechanical properties of Mg–4Y–3Nd–1.5Al alloy repair welded joints. The aim is to resolve the conflict between AGG formation and eutectic structure dissolution during heat treatment of traditional Mg–RE alloys by leveraging the excellent thermal stability of repair welded joints. This finding is expected to have significant implications for the design of high thermal stability alloys suitable for post-weld heat treatment. 2. Experimental process 2.1 Sample preparation This study selected metal mold cast Mg–4.02Y–3.06Nd–1.56Al alloy sheet with dimensions of 100 mm × 50 mm × 15 mm as the raw material. The alloy composition was determined by Inductively Coupled Plasma luminescence spectrometer (ICP–6300). A φ 5 mm × 100 mm filler rod and a 70 mm × 40 mm × 10 mm substrate were machined from the cast plate employing wire cutting. Three sets of 30 mm × 10 mm × 5 mm grooves were processed on the substrate using a pneumatic milling cutter for the repair welding experiments. The schematic image of repair welding process is depicted in Fig. 1 a. Before the experiment, the surface of the substrate and filler rod were polished with 800 # grit sandpaper. TIG repair welding was used, and the welding process parameters are provided in Table 1 . To further improve the mechanical properties of the repaired joints, they were subjected to solution treatment (T4) and aging treatment (T6). T4 treatment was performed at 500 ℃, 525 ℃, and 550 ℃ for 24 hours, subsequently water quenching. The repair welded joints after T4 treatment were assigned as T4–1, T4–2, and T4–3, respectively. The T4 treated repaired joints were then exposed to aging treatment at 200 ℃ for 0–100 hours. After BM peak aging treatment, the samples are labeled as T6–1 (500 ℃ × 24 h + 200 ℃ × 18 h), T6–2 (525 ℃ × 24 h + 200 ℃ × 16 h), and T6–3 (550 ℃ × 24 h + 200 ℃ × 14 h). 2.2 Characterization X–ray computed tomography (CT, AX7900 X–ray) was applied to detect defects in the repaired joints. Figure 1 b depicts the processing location of the microstructure sample. After polishing, the microstructure sample was etched with a solution of 10 g picric acid, 8 ml acetic acid, 20 ml deionized water, and 70 ml ethanol for 15 seconds. Optical microscopy (OM, DSX1000) was employed to examine the metallographic structures of various repair welded joints. The second phase dissolution situation and element distribution were examined by a scanning electron microscope (SEM, Apreo C) fitted with an energy–dispersive spectrometer (EDS). The angular distribution of texture strength and orientation differences was examined utilizing electron backscatter diffraction (EBSD, C–Nano). The EBSD sample was electropolished applying a polishing solution consisting of 24 ml nitric acid, 36 ml hydrochloric acid, and 250 ml alcohol. The polishing voltage was adjusted to 22 V, the current to 1.5 A, and the electrolysis time to 22 seconds. Transmission electron microscopy (TEM, Talos F200X) was utilized to examine the second phase composition of solid solution samples and identify precipitated phases in aged samples. After mechanically grinding the transmission sample to 50 µm, the Gatan 695 ion thinning instrument was used for thinning, with an ion thinning voltage range of 0.1 keV to 4.5 keV and an angle range of − 8 ° to 8 °. Figure 1 b illustrates the processing area (yellow region) and size of the tensile samples. To reduce experimental error, two sets of four tensile samples were cut from top to bottom, 1 mm from the upper surface of the sample for each heat treatment condition. Prior to the tensile test, the molten pool position of each sample was calibrated to determine the fracture location during testing. The tensile properties of the repair welded joints were evaluated at room temperature by means of a tensile machine (MTS E44.304) with a constant tensile rate of 1 mm/min. Microhardness measurements were conducted utilizing a Vickers hardness tester (DHV–1000). The load applied was 49 N, with a holding time of 15 seconds. Table 1 Repair welding process parameters. Welding current (A) Tungsten electrode diameter (mm) Filler rod diameter (mm) Torch length (mm) The flow rate of the Ar gas (L/min) 200 3.2 5 2 10 3. Conclusion 3.1 Microstructure of repair welded joints Figure 2 a presents the X–ray computed tomography analysis of the repair welded joint in the Mg–4.02Y–3.06Nd–1.56Al alloy. The X–ray penetrates the entire sample from the upper surface, revealing no apparent repair defects, indicating a well formed joint. As depicted in Fig. 2 b, the repair welded joint comprises three distinct regions: the FZ, HAZ, and base material (BM). Due to the continuous heat transfer from the FZ to the BM, the FZ experiences a higher cooling rate, resulting in a finer grain structure with an average grain size of 14 µm. Conversely, the HAZ, which is adjacent to the FZ, undergoes the highest thermal input for an extended duration, leading to a slight rise in grain size (51 µm) in comparison with the BM (47 µm) (Fig. 2 c). Further measurements indicate that the depth and width of the FZ are 5.5 mm and 14 mm, correspondingly, while the average width of the HAZ is 2.5 mm (Fig. 2 b). As shown in Figs. 2 d and g, the FZ displays a finer microstructure in comparison with the HAZ. The secondary phase in both the FZ and HAZ consists of a particle phase enriched in Al and RE elements (Y, Nd) within the grains, a reticular phase rich in RE (Y, Nd), and a needle–like phase rich in Al and RE (Y, Nd) at the grain boundaries ( Figs. 2 e, f, h, and i). Additionally, under continuous thermal input, a small quantity of linear phase precipitates within the grains of the HAZ (Fig. 2 h). Previous research has identified the reticular phase as an Mg–RE eutectic structure, whereas the needle–like, particulate, and linear phases correspond to Al 2 RE phases [ 28 ]. 3.2 Microstructure of T4 treated joints after welding Figure 3 presents OM images of the repair welded joint following T4 treatment. After T4 treatment, the FZ area in the T4–1 and T4–2 samples are still composed of fine equiaxed grains, whereas a small quantity of grains in the FZ edge area of the T4–3 sample exhibit the AGG phenomenon (Figs. 3 a, d and g). As depicted in Figs. 3 b, e, and h, the FZ’s grain size elevates progressively with rising T4 temperature, while the HAZ’s grain size doesn't exhibit significant variation. As illustrated in Figs. 3 c, f, and i, the number of second phase in the repaired joint gradually reduces as the T4 temperature increases. Figure 4 depicts the grain size distribution in the FZ and HAZ of the repair welded joints subjected to various T4 treatments. Compared to the repaired joints without heat treatment (WH), the FZ’s grain size of the T4–1 and T4–2 samples increased to 17 µm and 20 µm, correspondingly, whereas the HAZ’s grain size showed little variation, averaging 52 µm (Figs. 4 a–d). When the repair welded joint underwent T4 treatment at 550°C × 24 h, considering the occurrence of AGG in a small portion of the grains at the FZ edge in the T4–3 sample, the average grain size ( \(\:GS\) ) is evaluated as follows [ 29 ]: $$\:GS=xG{S}_{1}+\left(1-x\right)G{S}_{2}$$ 1 where \(\:G{S}_{1}\) and \(\:G{S}_{2}\) are the average sizes of abnormally grown grains and normal grains in FZ, respectively, and \(\:x\) is the region fraction of abnormally grown grains. As shown in Fig. 4 e, statistical analysis reveals that the values of \(\:G{S}_{1}\) , \(\:G{S}_{2}\) , and \(\:x\) were 115 µm, 24 µm, and 5.4%, respectively (Fig. 4 e). Based on calculations, the FZ’s average grain size in the T4–3 sample exhibited a substantial increase to 29 µm, while the HAZ’s grain size marginally rose to 54 µm (Fig. 4 f). Zhang et al. [ 17 ] depicted that after 8 hours of T4 treatment at 520°C for the traditional Mg–4Nd–2Y–1Gd–0.5Zr alloy repair welded joint, a pronounced AGG phenomenon occurred in the FZ region, causing a significant increase in the FZ’s grain size from 13 µm to 656.4 µm. Therefore, compared to the traditional Mg–4Nd–2Y–1Gd–0.5Zr alloy repair welded joint, the Mg–4.02Y–3.06Nd–1.56Al alloy repair welded joint displays superior grain thermal stability during post weld heat treatment. Figure 5 displays the SEM and EDS analyses of repair welded joints subjected to different T4 treatment processes. Owing to the high–melting–point of the Al 2 RE phase, it remains undissolved in all samples across the T4 treatment conditions. After the joint underwent T4 treatment at 500°C for 24 h, some eutectic structures remain undissolved in the T4–1 sample, and a limited amount of linear phase precipitates within the grains of the FZ. At a T4 temperature of 525°C, the eutectic structure in the T4–2 sample is completely dissolved, the number of linear phases increases, and their size enlarges. As the T4 treatment temperature increases to 550°C, the linear phase in the T4–3 sample nearly disappears. As illustrated in Figs. 5 j and k, line scan (line 1) and energy spectrum (point 1) analyses were conducted to examine the newly precipitated linear phase in the FZ region. The findings demonstrate that the linear phase mainly comprises of Y, Nd, and Al elements, with an atomic ratio of Al to RE (Y and Nd) of about 2:1. Based on this, it can be inferred that this phase is identical to the linear phase in the HAZ region and is classified as the Al 2 RE phase. To further verify the composition of the second phase in the FZ after T4 treatment, transmission electron microscopy (TEM) was utilized to investigate the second phase in the FZ of the T4–2 sample. As depicted in Fig. 6 , three distinct morphologies of secondary phases were detected in the FZ: particulate, needle–like, and linear phases (Figs. 6 a, d, and g). EDS analysis confirms that these three phases predominantly include Al, Y, and Nd elements (Figs. 6 b, e, and h). Further selected area electron diffraction (SAED) analysis (Figs. 6 c, f, and i) establishes that all three phases exhibit face–centered cubic structures with an identical lattice constant (a ≈ 0.790 nm), confirming their classification as Al 2 RE phases. 3.3 Age hardening behavior of repair welded joints Figures 7 a–c illustrates the aging hardening behavior of repair welded joints following various T4 treatments. The repair welded joints subjected to various T4 treatments exhibit similar aging hardening trends. The FZ consistently reaches the peak aging and over aging states ahead of the HAZ. Similarly, the HAZ reaches these two states before the BM. This results from the high cooling rate in the FZ, resulting in significant grain refinement, which significantly increases the grain boundary area. The increased grain boundary area provides additional nucleation sites for precipitates, thereby promoting their nucleation and growth [ 30 ]. Consequently, the FZ exhibits a greater tendency for precipitation compared to other regions. Additionally, because the HAZ is situated between the FZ and BM, it undergoes rapid heating and cooling during repair welding, expanding during heating and contracting during cooling. Uneven temperature changes during heating and cooling cause inconsistent expansion and contraction at different positions within the HAZ, leading to large residual stresses [ 31 ]. These residual stresses increase the diffusion rate of solute atoms and accelerate the emergence of precipitates, which leads to the precipitation tendency in HAZ is greater than that in BM [ 18 , 32 , 33 ]. As the T4 temperature rises, the comprehensive aging of the repair welded joint increases, and the time necessary to achieve the peak aging state decreases. At a T4 temperature of 500 ℃, the peak aging state is reached in the FZ, HAZ, and BM regions at 12 h, 14 h, and 18 h, with corresponding hardness values of 84 HV, 83 HV, and 82 HV. As the T4 temperature increases to 525 ℃, the peak aging times in the FZ, HAZ, and BM regions shorten to 10 h, 12 h, and 16 h, while the hardness increases to 89 HV, 88 HV, and 87 HV. Further increasing the T4 temperature to 550 ℃ reduces the peak aging times to 8 h, 10 h, and 14 h, while the hardness increases to 90 HV, 89 HV, and 88 HV, reaching its maximum. As illustrated in Fig. 7 d, the microhardness of the repaired joint at peak aging in the BM is selected to comprehensively evaluate its mechanical properties. At a T4 temperature of 500 ℃, the T6–1 sample exhibits low aging hardness, with microhardness values of 78 HV, 80 HV, and 82 HV for the FZ, HAZ, and BM areas, respectively. As the T4 temperature increases to 525 ℃, the average microhardness of the FZ, HAZ, and BM regions of the T6–2 sample rises significantly to 85 HV, 86 HV, and 87 HV, respectively. When the T4 temperature elevates further to 550 ℃, the microhardness of the FZ, HAZ, and BM areas of the T6–3 sample increases slightly to 87 HV, 87 HV, and 88 HV, respectively. Aging hardening is closely related to the content of aging precipitates. With elevating T4 temperature, the dissolution of the second phase is significantly enhanced, leading to a higher concentration of solute atoms within the α–Mg matrix. This generates a stronger driving force for the nucleation and growth of precipitates, thereby promoting an increased precipitation tendency and improving aging hardness [ 34 , 35 ]. To further investigate the precipitation behavior of strengthening phases during aging, the FZ under different aging conditions was analyzed using TEM. The T4–1 sample was aged at 200°C for 12 h, revealing a small quantity of granular precipitates having a mean length of close to 6 nm, detected along the [11 \(\:\stackrel{-}{2}\) 0] α–Mg direction of the α–Mg matrix (Figs. 8 a–c). High–resolution transmission electron microscopy (HRTEM) combined with fast fourier transform (FFT) examination demonstrates that, alongside the α–Mg matrix, weak diffraction spots are detected at the 1/4, 2/4, and 3/4 positions of the [1 \(\:\stackrel{\text{-}}{\text{1}}\) 00] α–Mg , which are in agreement with the diffraction patterns characteristic of the β′ precipitate phase. The phase relationship between β′ precipitates and α–Mg matrix is as follows: [010] β′ // [01 \(\:\stackrel{\text{-}}{\text{1}}\) 0] α–Mg , (200) β′ // (2 \(\:\stackrel{\text{-}}{\text{1}}\stackrel{\text{-}}{\text{1}}\) 0) α–Mg [ 36 ]. Figures 8 d–l shown the TEM analyzed of the T4–2 sample under three aging conditions: under aging (200°C × 2 h), peak aging (200°C × 10 h), and over aging (200°C × 100 h). Following aging at 200°C for 2 h, a small amount of granular precipitates, averaging 5 nm in length, formed in the FZ. SAED and HRTEM revealed distinct weak diffraction spots at the 1/2 position of [1 \(\:\stackrel{\text{-}}{\text{1}}\) 00] α–Mg , confirming that the precipitate phase was β′′ [ 30 ] (Figs. 8 d–f). As the aging time elevates to 10 h, the number and size of precipitates increase, the morphology changes from granular to flake–like, and the average length increases to 10 nm (Figs. 8 g and h). Figure 8 i shows that the precipitates at this stage are β′ phase. The T4–2 sample, subjected to an elevated solution temperature, demonstrates a markedly higher precipitate density and larger precipitate size compared to the peak–aged T4–1 sample. As aging time is extended to 100 h, the quantity of precipitates reduces, and their average size elevates to 56 nm. The precipitates remain of the flaky β′ phase (Figs. 8 j–l). In Mg–RE alloys, both nano β′′ and β′ precipitates impede dislocation slip, thus improving the alloy's strength [ 37 ]. In conclusion, increasing the solution temperature promotes the formation of more precipitates in the FZ. Simultaneously, as aging time is extended, the number of precipitates in the FZ region initially increases and then decreases, while their size continues to grow. 3.4 Mechanical properties of heat treatment repair welded joints The peak aging treatment process of the BM was selected to conduct the aging treatment on the repair welded joint. As shown in Figs. 9 a and b, the UTS, YS, and EL of the repair welded joint WH are 242 MPa, 165 MPa, and 8.6%, correspondingly. Relative to the WH sample, the UTS and YS of theT4 sample reduced, while the EL elevated markedly. After the joint underwent T4 treatment at 500°C for 24 h, the UTS, YS, and EL of the T4–1 sample were 233 MPa, 151 MPa, and 10.3%, respectively. Upon elevating the T4 temperature to 525°C, the UTS and YS of the T4–2 sample declined to 230 MPa and 141 MPa, while the EL improved to 12.4%. When the T4 temperature was further elevated to 550°C, the UTS and YS of the T4–3 sample dropped to 226 MPa and 134 MPa, whereas the EL increased to 13.1%. In contrast, the UTS and YS of the aged samples elevated substantially, while EL reduced. At a T4 temperature of 500°C, the UTS, YS, and EL of the T6–1 sample were 271 MPa, 200 MPa, and 4.9%. Upon increasing the T4 temperature to 525°C, the UTS, YS, and EL of the T6–2 sample enhanced to 293 MPa, 215 MPa, and 5.8%, correspondingly. With an additional rise in T4 temperature to 550°C, the UTS and YS of the T6–3 sample slightly increased to 304 MPa and 225 MPa, whereas the EL decreased to 5.2%. It is evident that at a T4 temperature of 550°C, the existence of a small amount of AGG in the FZ edge does not cause a decline in the mechanical properties of the Mg–4.02Y–3.06Nd–1.56Al alloy repaired joints. Additionally, the mechanical properties of the repaired joints following T6–2 heat treatment (UTS: 292 MPa, YS: 215 MPa, EL: 5.8%) are comparable to those of the T6 (525°C × 24 h + 200°C × 16 h) unrepaired alloy (UTS: 293 MPa, YS: 218 MPa, EL: 6%), indicating that the repair process does not substantially compromise the mechanical properties of the Mg–4.02Y–3.06Nd–1.56Al alloy under conventional heat treatment. Figure 9 c and d presents a macro image and a low–magnification OM of the fracture location in the tensile sample. The fracture of the WH sample occurred in the HAZ, while the fracture locations of the heat–treated samples predominantly appeared in the BM, even with minor AGG formation in the FZ edge of the T4–3 or T6–3 treated samples. This indicates that AGG did not significantly impact the mechanical properties of the repair welded joint. Figures 10 a–h illustrates the fracture morphologies of the repair welded joints with WH and heat treatment. The sample fractures exhibit a combined characteristic of transgranular and intergranular fractures, combined with fracture dimples, tearing edges, and cleavage planes. Fractured Al 2 RE particles (demonstrated by the orange arrow) are visible on the fracture surface of the WH sample, indicating that Al 2 RE particles are one of the primary crack sources (Fig. 10 b). Following solution treatment, the amount of tearing edges and dimples in the repair welded joint increases, while the number of cleavage planes reduces, compared to the WH sample. Consequently, the solution–treated sample demonstrates higher ductility and lower strength. With increasing T4 temperature, the quantity of tearing edges and dimples progressively rises, while both the number and size of cleavage planes exhibit a gradual reduction (Figs. 10 c, e, and g). After aging treatment, the quantity and size of cleavage planes increase, whereas the number and size of tearing edges and dimples significantly decrease, compared to the solution treated repaired joint. Thus, the aged sample exhibits lower ductility and higher strength (Figs. 10 d, f, and h). Additionally, the fractured Al 2 RE particles remain visible after various heat treatments, indicating that these particles continue to be one of the primary crack sources in the heat–treated samples. Figure 11 illustrates OM images of the longitudinal cross–sections of tensile samples extracted from both the WH and the heat–treated repair welded joint. The fracture mode of the WH repair welded joint exhibits a mixture of transgranular and intergranular fractures, with intergranular fracture being predominant. Cracks originate within the eutectic structure and propagate either along the brittle eutectic or into the grains (Figs. 11 a and b). This occurs because of the reticular morphology of the eutectic structure at grain boundaries, which has large dimensions and sharp edges that readily induce stress concentration under loading [ 13 , 14 ]. After the joint underwent T4 treatment at 500°C for 24 h, the T4–1 sample still contains undissolved eutectic structures, making crack initiation more likely at these sites, predominantly resulting in intergranular fracture (Fig. 11 c). With the rise in T4 temperature to 525°C and 550°C, the eutectic structure located at the grain boundaries dissolves completely. Fractured Al 2 RE particles (indicated by the orange arrow) appear on the fracture surface, confirming that Al 2 RE particles become the primary crack initiation sites following the complete dissolution of the eutectic structure (Figs. 11 e and g). This transition is attributed to the increased grain boundary bonding strength in T4–2 and T4–3 samples resulting from the complete dissolution of the eutectic structure, reducing the likelihood of crack formation along grain boundaries. Therefore, cracks tend to form on the Al₂RE particles within the grains and propagate through them, leading to a shift in the primary fracture mode from intergranular to transgranular. Following aging treatment, the fracture morphology of the repaired joint closely resembles that observed in the T4 treated condition. Nevertheless, the precipitation phases within the grains strengthens the matrix and restricts plastic deformation, leading to a more uniform fracture surface [ 29 ] (Figs. 11 d f and h). 4 Discussion 4.1 Causes of AGG phenomenon and grain thermal stability Traditional Mg–RE alloy repair welded joints are susceptible to AGG in the molten pool during post–weld heat treatment, adversely affecting their mechanical properties [ 17 , 18 ]. In this study, the Mg–4.02Y–3.06Nd–1.56Al alloy repaired joint underwent T4 treatment at 550°C for 24 h, and only a few grains at the molten pool edge exhibited AGG. Compared to conventional Mg–RE–Zr alloy repaired joints, this alloy demonstrated superior thermal stability during post–weld heat treatment, significantly mitigating the detrimental effects of AGG on joint properties. Therefore, a comprehensive analysis of the mechanisms underlying AGG formation and grain thermal stability in Mg–4.02Y–3.06Nd–1.56Al alloy repair welded joints is essential for further improving their mechanical properties. During the T4 process, grain growth occurs to decrease interface energy, thus decreasing the system’s total energy and enhancing thermodynamic stability. This growth is typically spontaneous, influenced by atomic diffusion rates and interfacial energy [ 38 ]. As the T4 temperature increases, atomic diffusion accelerates, promoting grain coarsening. At 500°C and 525°C, atomic diffusion remains relatively slow due to lower temperatures, resulting in a low grain growth trend. However, at 550°C, increased atomic diffusion enhances the tendency for grain coarsening. Additionally, grain boundaries are typically irregular regions within the lattice structure, which can induce lattice distortion and result in high energy, leading to significant interfacial energy, which is more pronounced in smaller grains due to their larger boundary area [ 39 , 40 ]. Compared to other regions, the FZ exhibits smaller grain sizes, with the edge region (Zone 2) showing even finer grains than the center (Zone 1) (Figs. 12 a and d). This results in higher interface energy at the Zone 2, intensifying grain growth trend. Rapid cooling of the molten pool during repair welding generates substantial residual shrinkage stress, leading to an increased density of structural defects, for example vacancies and dislocations, therefore, during the solid solution treatment, grains in the molten pool are prone to thermodynamic instability, making them susceptible to coarsening [ 32 , 33 ]. This effect is particularly pronounced at the Zone 2, where the cooling rate is highest due to direct contact with the substrate. As a result, compared to the Zone 1, the Zone 2 experiences greater residual shrinkage stress, leading to a more prominent grain coarsening phenomenon. Based on the distribution and frequency of grain boundaries with varying misorientation angles shown in Figs. 13 c, f, and g, the fraction of grain boundaries with high misorientation angles (> 45°) is significantly greater at the Zone 2 than at the Zone 1. These high–angle grain boundaries possess elevated interfacial energy and enhanced atomic diffusion rates, accelerating grain migration and coarsening [ 41 , 42 ]. Consequently, under the combined effects of atomic diffusion, interface energy, and residual shrinkage stress, secondary recrystallization occurs at the molten pool edge in the T4–3 sample after treatment at 550°C for 24 h, causing the rapid disappearance of surrounding grain boundaries and the consumption of adjacent smaller grains, leading to AGG. However, only a limited number of grains exhibit AGG at the molten pool edge in the T4–3 sample, with the FZ’s average grain size increasing to 29 µm, still significantly smaller than in other regions. By contrast, in traditional Mg–4Nd–2Y–1Gd–0.5Zr alloy repair–welded joints subjected to T4 treatment at 520°C for 8 h, the FZ’s average grain size increases substantially to 656.4 µm [ 17 ]. This comparison demonstrates that Mg–4.02Y–3.06Nd–1.56Al alloy repair welded joints exhibit superior grain thermal stability during post–weld heat treatment. Therefore, further investigation into the mechanisms governing grain thermal stability in Mg–4.02Y–3.06Nd–1.56Al alloy repair welded joints is essential for optimizing their properties. It is widely proven that the presence of secondary phases markedly influences grain growth behavior. Based on the Zener pinning mechanism [ 26 ], stable secondary phases at grain boundaries can effectively hinder their movement at elevated temperatures, thereby restricting grain boundary migration and suppressing grain coarsening. In the Mg–4.02Y–3.06Nd–1.56Al alloy, a substantial amount of high–melting–point Al 2 RE phases is present, particularly in the form of needle–like Al 2 RE structures near grain boundaries. These phases effectively limit grain boundary migration and control grain coarsening during heat treatment in repair–welded joints. The Zener pinning influence is solely operative when the grain size does not exceed a critical value ( \(\:\text{D}\) ), which is defined as follows [ 26 , 43 ]: $$\:\text{D=}\frac{\text{4r}}{\text{3f}}$$ 2 where \(\:\text{r}\) and \(\:\text{f}\) are the diameter and volume fraction of needle–like Al 2 RE. According to the SEM observations, the data of \(\:\text{r}\) and \(\:\text{f}\) are calculated as be 1.32 µm and 1.8%. The determined " \(\:\text{D}\) " value is 97.8 µm, which is substantially exceeding the average grain size of 14 µm in the FZ. Therefore, at high temperatures, the presence of abundant high–melting–point needle–like Al 2 RE phases near the grain boundaries in the Mg–4.02Y–3.06Nd–1.56Al molten pool effectively restricts grain boundary migration during the repair welding process. This suppression of grain coarsening enhances the exceptional thermal stability of the Mg–4.02Y–3.06Nd–1.56Al alloy, ensuring superior structural integrity during post–weld heat treatment. 4.2 Evolution of microstructure and mechanical properties Figure 13 illustrates the microstructural evolution and fracture mechanisms of repaired joints subjected to various post–weld heat treatment conditions. As illustrated in Fig. 13 a, the WH sample contains a substantial amount of Mg–RE eutectic structures, along with needle–like and particulate Al 2 RE phases. During the tensile process, Mg–RE eutectic structures at grain boundaries and particle Al₂RE phases within the grains, which are relatively large and exhibit complex morphologies, contribute to stress concentration, ultimately leading to crack initiation and brittle fracture. These two phases serve as the primary sources of cracking in the repair welded joint of the Mg–4.02Y–3.06Nd–1.56Al alloy, thereby limiting further improvement in mechanical properties. Compared to particle Al₂RE, the Mg–RE eutectic structures at grain boundaries are larger and possess sharper morphologies, making them more prone to stress concentration. When eutectic structures are present, cracks tend to initiate and propagate preferentially within them, predominantly resulting in intergranular fracture. Given that the Al₂RE phase has a high–melting–point (approximately 1600°C), it is difficult to dissolve in the matrix. Therefore, eliminating the Mg–RE eutectic structure through heat treatment is essential for strengthening the overall properties of the repaired joint. As illustrated in Figs. 13 b and c, after T4 treatment at 500°C for 24 h, a small quantity of linear Al 2 RE phases formed inside the grains of the T4–1 sample, while some undissolved Mg–RE eutectic structures remained at the grain boundaries. This led to a low solid solubility of RE elements in the α–Mg matrix, limiting the quantity density of precipitates in the T6–1 sample. Consequently, only a small quantity of fine granular β′ precipitates formed, which had minimal impact on improving the sample's tensile strength. Additionally, the presence of undissolved Mg–RE eutectic structures resulted in a predominantly intergranular fracture mode. As the T4 temperature reached 525°C, the quantity and size of linear Al 2 RE phases within the grains of the T4–2 sample increased, and the Mg–RE eutectic structures at the grain boundaries were fully dissolved. This enhanced the solid solubility of RE elements in the α–Mg matrix, promoting the formation of numerous flake–like β′ precipitates in the T6–2 sample, significantly improving its tensile strength. With the complete dissolution of the Mg–RE eutectic structures, particulate Al 2 RE phases inside the grains became the primary crack initiation sites, leading to a transition from predominantly intergranular to transgranular fracture. At 550°C, the linear Al 2 RE phases inside the grains of the T4–3 sample dissolved back into the matrix, further increasing the solid solubility of RE elements in the α–Mg matrix. This resulted in an even higher density of flake–like β′ precipitates in the T6–3 sample, which had the most pronounced effect on enhancing tensile strength. However, owing to the existence of numerous undissolved particulate Al 2 RE phases, these particles remained the primary crack sources, maintaining a predominantly transgranular fracture mode. Notably, at 550°C, although some grains at the edge of the FZ exhibited AGG, fractures did not occur in the FZ region, indicating that AGG did not markedly impact the mechanical properties of the repaired joint. Thus, it is essential to use the Hall–Petch relationship to evaluate the grain boundary strengthening ( \(\:{\sigma\:}_{y}\) ) effect of repair welded joints with solid solution temperature of 550 ℃ [ 44 , 45 ]: $$\:{\sigma\:}_{y}={\sigma\:}_{0}+{K}_{y}{D}^{-\frac{1}{2}}$$ 3 where \(\:{\sigma\:}_{0}\) is the friction stress, \(\:{K}_{y}\) is a Hall-Petch coefficient (250 MPa µm1/2) [ 46 ], and \(\:D\) is the average grain size. Following T4 treatment at 550°C, the average grain sizes in the FZ and HAZ are 29 µm and 54 µm, correspondingly. Calculations show that the \(\:{\sigma\:}_{y}\) values in the FZ and HAZ are 58.4 MPa and 46 MPa, respectively. It can be inferred that the high strength retained in the FZ enhances its resistance to fracture during tensile deformation. However, considering the AGG phenomenon, the heat treatment process for the T6–2 sample (525°C × 24 h + 200°C × 16 h) is determined to be optimal. The mechanical properties of the T6–2 sample (UTS: 292 MPa, YS: 215 MPa, EL: 5.8%) are comparable to those of the T6 alloy without repair welding (UTS: 293 MPa, YS: 218 MPa, EL: 6%). In summary, the excellent thermal stability of the Mg–4.02Y–3.06Nd–1.56Al alloy successfully addresses the conflict between the occurrence of AGG and the dissolution of the eutectic structure in the repair weld, enabling the feasibility of single–stage solution treatment, thereby effectively reducing actual industrial production costs. Furthermore, a study of the heat treatment process for the alloy repair weld revealed that, after optimal treatment, the properties of the repaired joints was significantly improved to a level comparable to the without repair welding alloy, effectively addressing the properties degradation typically observed in traditional Mg–RE alloys after repair welding. 5. conclusion During heat treatment, the Mg–4.02Y–3.06Nd–1.56Al alloy repaired joint demonstrated notable thermal stability. With an elevate in solution temperature from 500°C to 525°C, the grain size in the FZ region exhibited only, a slight increase from 17 µm to 20 µm. An additional rise in solution temperature to 550°C, caused the occurrence of AGG in certain grains at the FZ edge, resulting in a grain size rise to 29 µm, which remained considerably smaller than in other regions. Additionally, despite the rise in solution temperature, the HAZ’s grain size region remained relatively stable (within the range of 52 µm to 54 µm). The particle and needle–like Al 2 RE phases possess a high–melting–point, making them challenging to dissolve within the matrix. Therefore, the elimination of the Mg–RE eutectic structure is essential for improving mechanical properties. After the joint underwent T4 treatment at 500°C for 24 h, portion of the Mg–RE eutectic structure remains undissolved, with a small limited amount of linear Al 2 RE phase precipitating within the grain. As the T4 temperature increases to 525°C, the Mg–RE eutectic structure undergoes complete dissolution, accompanied by a growth in in both the quantity and size of the linear Al 2 RE phase. With a further rise in T4 temperature to 550°C, the linear Al 2 RE phase is entirely eliminated, however, the AGG phenomenon appears at the FZ edge. Based on these observations, a T4 treatment at 525°C for 24 h is determined to be optimal. The duration of the aging treatment has a markedly influence on the type, morphology, and size of precipitates. When the T4–2 sample undergoes aging at 200°C for 2 h, a limited quantity of granular β′′ phase precipitates form in the FZ region, with an averaging about 5 nm in length. With an extension of aging time to 10 h, both the quantity and size of the precipitates increase, and the phase type transitions to a flake–like β′ phase with an average length of 10 nm. Further extending the aging time to 100 h, the precipitate type remained as flake–like β' phase, however, the number of precipitates decreases while their size increases to 56 nm. After T4 treatment at 525°C for 24 h and subsequent aging at 200°C for 16 h, the repair welded joint displayed excellent overall mechanical properties. The UTS, YS, and EL reached 292 MPa, 215 MPa, and 5.8%, correspondingly. These values are comparable to the mechanical properties of the T6 alloy in its original, unrepaired state, demonstrating that the optimized heat treatment effectively restores the strength and ductility of the repair welded joint. Declarations Author contribution Zhengtao Liu : Data curation, Investigation, Methodology, Writing – original draft. Lei Wang : Conceptualization, Funding acquisition, Writing – review & editing. Hongyang Duan : Investigation, Methodology, Validation. Sicong Zhao : Data curation, Formal analysis, Funding acquisition. Erjun Guo : Supervision, Validation. Yicheng Feng : Data curation. Funding This work was supported by the Heilongjiang Province Postdoctoral Science Foundation (LBH-Z24308). Data availability Data in the present work can be obtained from the corresponding author on request. Ethics approval The manuscript was approved by all authors for publication. Consent to participate: Not applicable. Consent for publication: Not applicable. Competing interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. 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Mater Sci Eng A 677(20):411–420. https://doi.org/10.1016/j.msea.2016.09.044 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. Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-7018798","acceptedTermsAndConditions":true,"allowDirectSubmit":true,"archivedVersions":[],"articleType":"Research Article","associatedPublications":[],"authors":[{"id":490328766,"identity":"c281308d-823b-4f0d-9b7f-bfb98dfbcdc3","order_by":0,"name":"Zhengtao Liu","email":"","orcid":"","institution":"Harbin University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Zhengtao","middleName":"","lastName":"Liu","suffix":""},{"id":490328767,"identity":"d8aec09a-3c57-4693-b7e6-dcdd48adf630","order_by":1,"name":"Lei Wang","email":"data:image/png;base64,iVBORw0KGgoAAAANSUhEUgAAAZAAAAAyAQMAAABI0h/eAAAABlBMVEX///8AAABVwtN+AAAACXBIWXMAAA7EAAAOxAGVKw4bAAAA6klEQVRIiWNgGAWjYBACPgbGBhAtZ8DAAxU6QEALG1SLMURLAlFaICBxA/Fa2JvbJH7uqE3fLn32mOTPHwxyfDcSGD8X4NPCc7BNsvfM8dydfXlp0jwJDMaSNxKYpWfg0yKR2HaDt+1Y7oYzPGbSQIclbriRwMbMg0+L/MO2m3/bjqUbALVI/khgqCesRYKx7TZvW00CSIsE0GEJBgS18CS2/5ZtO2C44QxfsjVPmoThzDMPm6XxaeFnP/7Y8G1bnbzBGd6DN3/Y2MjzHU8++BmfFig4DGNIADEkcgmBOmIUjYJRMApGwUgFAA4jSUHlKgoQAAAAAElFTkSuQmCC","orcid":"https://orcid.org/0000-0003-1310-8603","institution":"Harbin University of Science and Technology","correspondingAuthor":true,"prefix":"","firstName":"Lei","middleName":"","lastName":"Wang","suffix":""},{"id":490328768,"identity":"b8628147-3ace-4f1d-8a90-ffebc459d918","order_by":2,"name":"Hongyang Duan","email":"","orcid":"","institution":"Harbin University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Hongyang","middleName":"","lastName":"Duan","suffix":""},{"id":490328769,"identity":"e769ae53-6668-4c79-a95e-b2ad1e89e026","order_by":3,"name":"Sicong Zhao","email":"","orcid":"","institution":"Harbin University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Sicong","middleName":"","lastName":"Zhao","suffix":""},{"id":490328770,"identity":"dbc6ffe0-7c00-4927-beb5-fad4f9386eae","order_by":4,"name":"Erjun Guo","email":"","orcid":"","institution":"Harbin University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Erjun","middleName":"","lastName":"Guo","suffix":""},{"id":490328771,"identity":"b82a90dd-c3a5-4c42-bbe8-97fd1b95b934","order_by":5,"name":"Yicheng Feng","email":"","orcid":"","institution":"Harbin University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Yicheng","middleName":"","lastName":"Feng","suffix":""}],"badges":[],"createdAt":"2025-07-01 09:44:03","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-7018798/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-7018798/v1","draftVersion":[],"editorialEvents":[],"editorialNote":"","failedWorkflow":false,"files":[{"id":87766765,"identity":"b117fe86-c28f-47ce-bd00-2def55a9ee01","added_by":"auto","created_at":"2025-07-28 18:17:05","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":450532,"visible":true,"origin":"","legend":"\u003cp\u003e(a) Schematic image of repair welding process. (b) Processing location of microstructure sample, processing area and size of tensile sample.\u003c/p\u003e","description":"","filename":"floatimage1.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/e7ce88d8a9403685badb6c45.png"},{"id":87766764,"identity":"3b24376b-f6ff-47af-84d2-63955906afa6","added_by":"auto","created_at":"2025-07-28 18:17:05","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":2369450,"visible":true,"origin":"","legend":"\u003cp\u003eRepair welded joint ofMg–4.02Y–3.06Nd–1.56Al alloy: (a) X–ray computed tomography analysis, (b) OM image, (c) grain size distribution, (d) OM image of transition region, (e, f) SEM and EDS analysis of FZ region, and (g) SEM of transition region, (h, i) SEM and EDS analysis of HAZ.\u003c/p\u003e","description":"","filename":"floatimage2.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/ae190a7a7b2fe1929b046457.png"},{"id":87766808,"identity":"7895b456-074d-47d6-a2d5-972359d3812f","added_by":"auto","created_at":"2025-07-28 18:17:07","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":3343777,"visible":true,"origin":"","legend":"\u003cp\u003eOM images of repair welded joint after various T4 treatments: (a–c) T4–1, (d–f) T4–2, (g–i) T4–3.\u003c/p\u003e","description":"","filename":"floatimage3.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/67cd85c68435b655c427aceb.png"},{"id":87766766,"identity":"f2c6cce4-c84f-41af-b5ff-d702f3f2c0e1","added_by":"auto","created_at":"2025-07-28 18:17:05","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":776887,"visible":true,"origin":"","legend":"\u003cp\u003eGrain size distribution in FZ and HAZ regions of repair welded joints after various T4 treatments: (a, b) T4–1, (c, d) T4–2, (e, f) T4–3.\u003c/p\u003e","description":"","filename":"floatimage4.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/f760c5dd2b8af56fa862a5fc.png"},{"id":87767299,"identity":"d9ebb7e6-d33c-4053-a404-785dd3a2ba7e","added_by":"auto","created_at":"2025-07-28 18:25:06","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":2623515,"visible":true,"origin":"","legend":"\u003cp\u003eSEM and EDS analysis of repair welded joint of FZ and HAZ after different T4 treatments: (a–c) T4–1, (d–f) T4–2, (g–i) T4–3, (j) line 1 and (k) point 1.\u003c/p\u003e","description":"","filename":"floatimage5.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/0e2b52e47b97f91c0c4d5df0.png"},{"id":87766772,"identity":"640d8f7c-2e96-421c-a60e-84e2f7d81ce9","added_by":"auto","created_at":"2025-07-28 18:17:05","extension":"png","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":1838608,"visible":true,"origin":"","legend":"\u003cp\u003eTEM analysis of typical particle, needle–like and linear precipitated phase in FZ of T4–2: (a, d, g) bright field images, (b, e, h) EDS analysis, (c, f, l) SAED images.\u003c/p\u003e","description":"","filename":"floatimage6.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/e87107b6f5d534d57f825b96.png"},{"id":87766774,"identity":"acec5a40-5bea-4c97-b194-c7c725324728","added_by":"auto","created_at":"2025-07-28 18:17:06","extension":"png","order_by":7,"title":"Figure 7","display":"","copyAsset":false,"role":"figure","size":937526,"visible":true,"origin":"","legend":"\u003cp\u003eAge hardening image of repair welded joints at various T4 temperatures: (a) T4–1, (b) T4–2, (c) T4–3, and (d) comparison image of peak aging microhardness of T6–1, T6–2 and T6–3.\u003c/p\u003e","description":"","filename":"floatimage7.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/3e990756fcf3a86349e52c47.png"},{"id":87767305,"identity":"41a11ecf-2bf0-438c-ba38-9d8674ccccab","added_by":"auto","created_at":"2025-07-28 18:25:06","extension":"png","order_by":8,"title":"Figure 8","display":"","copyAsset":false,"role":"figure","size":2403539,"visible":true,"origin":"","legend":"\u003cp\u003eTEM analysis of peak aging (200 ℃ × 12 h) in the FZ of T4–1 sample and under aging (200 ℃ × 2 h), peak aging (200 ℃ × 10 h), over aging (200 ℃ × 100 h) in the FZ of T4–2 sample: (a, d, g, e) bright field images, (b, e, h, k) SAED images, (c, f, i, l) HRTEM and FFT pattern.\u003c/p\u003e","description":"","filename":"floatimage8.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/7e64b529d73e561916ace454.png"},{"id":87767315,"identity":"d2f49c6e-1963-43bf-a9b5-a3f8365d26db","added_by":"auto","created_at":"2025-07-28 18:25:07","extension":"png","order_by":9,"title":"Figure 9","display":"","copyAsset":false,"role":"figure","size":951962,"visible":true,"origin":"","legend":"\u003cp\u003eThe WH and heat-treated repair welded joints: (a) engineering stress–strain curve, (b) tensile property comparison image, and the fracture location of tensile samples of (c) macro image and (d) low–magnification OM.\u003c/p\u003e","description":"","filename":"floatimage9.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/625faed7ca96f6b0ded559f9.png"},{"id":87767612,"identity":"9e44a322-d708-4f60-9d10-ffec2da371d4","added_by":"auto","created_at":"2025-07-28 18:33:06","extension":"png","order_by":10,"title":"Figure 10","display":"","copyAsset":false,"role":"figure","size":1898347,"visible":true,"origin":"","legend":"\u003cp\u003eFracture morphology of the WH and heat–treated repair welded joints: (a, b) WH, (c) T4–1, (d) T6–1, (e) T4–2, (f) T6–2, (g) T4–3 and (h) T6–3.\u003c/p\u003e","description":"","filename":"floatimage10.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/7fb5770ac111e32b43bf4db6.png"},{"id":87767293,"identity":"2cd6f30b-eecd-4d84-b209-20ca6df7bade","added_by":"auto","created_at":"2025-07-28 18:25:05","extension":"png","order_by":11,"title":"Figure 11","display":"","copyAsset":false,"role":"figure","size":1648817,"visible":true,"origin":"","legend":"\u003cp\u003eOM images of the longitudinal cross–sections of WH and heat–treated repair welded joint: (a, b) WH, (c) T4–1, (d) T6–1, (e) T4–2, (f) T6–2, (g) T4–3 and (h) T6–3.\u003c/p\u003e","description":"","filename":"floatimage11.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/9c1b91a2daf98591887b5a1d.png"},{"id":87767606,"identity":"ad8eba4f-75ca-409a-a1b0-fd72c89ca240","added_by":"auto","created_at":"2025-07-28 18:33:06","extension":"png","order_by":12,"title":"Figure 12","display":"","copyAsset":false,"role":"figure","size":2219963,"visible":true,"origin":"","legend":"\u003cp\u003eThe FZ center and FZ edge: (a, d) IPF images, (b, e) KAM images, (c, f) distribution of grain boundaries with various misorientation angles, and (g) frequency of grain boundaries with various misorientation angles.\u003c/p\u003e","description":"","filename":"floatimage12.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/be0069029b1a055c8d8251ef.png"},{"id":87768140,"identity":"e97070b7-4cfd-4536-84a5-b97c418981ee","added_by":"auto","created_at":"2025-07-28 18:41:06","extension":"png","order_by":13,"title":"Figure 13","display":"","copyAsset":false,"role":"figure","size":1689674,"visible":true,"origin":"","legend":"\u003cp\u003eMicrostructure evolution and fracture mechanism images of the WH and heat–treated repair welded joints: (a) WH, (b) T4 state and (c) T6 state.\u003c/p\u003e","description":"","filename":"floatimage13.png","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/db94a2965af0d1e38a306fd2.png"},{"id":90498775,"identity":"5851178b-1a82-44f5-a157-ee44470d9c26","added_by":"auto","created_at":"2025-09-03 11:16:17","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":23967667,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-7018798/v1/55d28e92-9005-4b03-b7d4-16ed54084f25.pdf"}],"financialInterests":"","formattedTitle":"Strengthening and toughening mechanism of Mg–RE–Al alloy repair welded joints by suppressing abnormal grain growth during heat treatment","fulltext":[{"header":"1. Introduction","content":"\u003cp\u003eMagnesium alloy as the lightest metal structure material, is known for its low density and high specific strength [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e]. In particular, the incorporation of rare earth (RE) elements further enhances the mechanical properties and high\u0026ndash;temperature heat resistance of magnesium alloys, making them widely used in aerospace [\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e, \u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e]. In industrial production, large and complex Mg\u0026ndash;RE alloy components are commonly produced via sand casting [\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e, \u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e]. Due to the shortcomings of sand casting, such as slow heat dissipation and long solidification time, as well as the high density and easy segregation of RE elements, complex Mg\u0026ndash;RE alloy castings are susceptible to defects [\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e, \u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e]. Therefore, to improve casting yield, it is crucial to repair defects through welding.\u003c/p\u003e\u003cp\u003eTungsten inert gas (TIG) welding is commonly employed for defect repair in large alloy castings due to its simplicity, high flexibility, and low cost [\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e, \u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e]. Meng et al. [\u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e] applied TIG welding to repair casting defects in Mg\u0026ndash;Gd\u0026ndash;Y\u0026ndash;Zr alloy, achieving a maximum joint efficiency (the strength ratio of the repair weld to the base metal) of 99.7%. To further enhance the properties of the repair welded joint, Tong et al. [\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e] optimized the welding current, resulting in the Mg\u0026ndash;4Y\u0026ndash;2Nd\u0026ndash;1Gd\u0026ndash;0.5Zr alloy repair welded joint with no apparent repair defects at a current of 170 A. The ultimate tensile strength (UTS) was 195 MPa, and the elongation (EL) was 5.1%. However, continuous heat input promotes the formation of more continuous eutectic structure at the grain boundary of the heat\u0026ndash;affected zone (HAZ), which increases the likelihood of crack initiation and causes a significant decline of mechanical properties in the repair welded joint. Subsequently, Wang et al. [\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e] applied pulse current during the repair welding of Mg\u0026ndash;3Nd\u0026ndash;3Gd\u0026ndash;0.2Zn\u0026ndash;0.5Zr alloy. The results indicate that applying pulse current further refines the grain size within the fusion zone (FZ). When the pulse current frequency is 10 Hz, the repair welded joint achieves excellent comprehensive mechanical properties are obtained, UTS, yield strength (YS), and EL are 220\u0026thinsp;\u0026plusmn;\u0026thinsp;14 MPa, 129\u0026thinsp;\u0026plusmn;\u0026thinsp;9 MPa, and 11.9\u0026thinsp;\u0026plusmn;\u0026thinsp;1.2%. The maximum joint efficiency reaches 88.4%. However, the eutectic structure at the HAZ grain boundary remains coarse and continuous, negatively affecting the mechanical properties of the repair welded joints. Consequently, optimizing the welding process can substantially enhance the mechanical properties of the repair welded joint. However, for large sand\u0026ndash;cast Mg\u0026ndash;RE alloys, a slow cooling rate during casting promotes the formation of coarse and sharp eutectic structures at the alloy grain boundaries. These eutectic structures are prone to stress concentration during tensile testing, leading to brittle fracture and limiting the improvement of alloy properties [\u003cspan citationid=\"CR13\" class=\"CitationRef\"\u003e13\u003c/span\u003e, \u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e]. Particularly after TIG welding, the eutectic structure in the HAZ becomes more continuous, increasing the probability of crack initiation and further degrading the properties of the repair welded joint. Therefore, eliminating eutectic structures at the grain boundaries is essential for amplifying the overall mechanical properties of repair welded joints.\u003c/p\u003e\u003cp\u003eSolid solution and aging treatment are conventional methods for eliminating brittle eutectic structures and enhancing the mechanical properties of as\u0026ndash;cast Mg\u0026ndash;RE alloys [\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e, \u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e]. Therefore, studying the heat treatment parameters of Mg\u0026ndash;RE alloy repair welded joints is crucial for enhancing joint properties. Zhang et al. [\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e] investigated the heat treatment parameters of Mg\u0026ndash;4Y\u0026ndash;3RE\u0026ndash;0.5Zr alloy repair welded joints. The results indicate that the alloy exhibits poor grain thermal stability. Although the eutectic structure in the joint can be thoroughly dissolved by a 520\u0026deg;C \u0026times; 8 h T4 treatment, some grains in the molten pool region undergo abnormal grain growth (AGG). These coarse grains are prone to cracking during the tensile process, causing a sharp decrease in the joint's EL to 1.8%. To address the conflict between eutectic dissolution and AGG, Tong et al. [\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e] developed a combined solid solution treatment process, which involves 520\u0026deg;C \u0026times; 8 h treatment before welding and 520\u0026deg;C \u0026times; 2 h treatment after welding, successfully preparing the Mg\u0026ndash;4Y\u0026ndash;2Nd\u0026ndash;1Gd\u0026ndash;0.5Zr alloy repair welded joint without residual eutectic or AGG. After aging treatment, the UTS of the joint reached 314 MPa, and the EL increased to 4.4%. However, for mass\u0026ndash;produced Mg\u0026ndash;RE alloy castings, this combined solution treatment process is both cumbersome and costly, hindering its practical industrial application. Therefore, demonstrating the feasibility of a single\u0026ndash;stage solid solution process is crucial, while addressing the conflict between AGG formation in the molten pool and the dissolution of the eutectic structure. This is beneficial for enhancing the properties of repaired joints and reducing industrial production costs.\u003c/p\u003e\u003cp\u003ePrevious research has indicated that the addition of Al refines cast Mg\u0026ndash;RE alloys, under elevated temperatures (preceding α\u0026ndash;Mg solidification), the in\u0026ndash;situ formation of Al\u003csub\u003e2\u003c/sub\u003eRE particles serves as heterogeneous nucleation sites, successfully refining the grain size and achieving a refinement effect similar to that of Zr [\u003cspan additionalcitationids=\"CR20 CR21 CR22\" citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e]. Additionally, numerous high\u0026ndash;melting\u0026ndash;point Al\u0026ndash;RE phases (such as Al\u003csub\u003e2\u003c/sub\u003eRE, Al\u003csub\u003e11\u003c/sub\u003eRE\u003csub\u003e3\u003c/sub\u003e) form near the grain boundaries of Mg\u0026ndash;RE alloys during the eutectic stage [\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e, \u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e]. According to Zener pinning theory [\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e], these Al\u0026ndash;RE phases can pin the grain boundaries, hindering their movement and thus improving the alloy's thermal stability. Therefore, this research aims to utilize the excellent thermal stability of Mg\u0026ndash;RE\u0026ndash;Al alloys to address the challenge of balancing AGG and eutectic structure dissolution imbalance caused by the inadequate thermal stability of traditional Mg\u0026ndash;RE alloy repair welded joints in single-stage solution treatment.\u003c/p\u003e\u003cp\u003eIn our previous research [\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e], we successfully utilized the excellent thermal stability of the Mg\u0026ndash;4Y\u0026ndash;3Nd\u0026ndash;1.5Al alloy (developed from the WE43 alloy by substituting Zr with Al as the alloy refiner) to mitigate grain growth in the HAZ during repair welding, thus preventing significant property loss. Therefore, this study systematically investigates the influences of various heat treatment parameters on the microstructure and mechanical properties of Mg\u0026ndash;4Y\u0026ndash;3Nd\u0026ndash;1.5Al alloy repair welded joints. The aim is to resolve the conflict between AGG formation and eutectic structure dissolution during heat treatment of traditional Mg\u0026ndash;RE alloys by leveraging the excellent thermal stability of repair welded joints. This finding is expected to have significant implications for the design of high thermal stability alloys suitable for post-weld heat treatment.\u003c/p\u003e"},{"header":"2. Experimental process","content":"\u003cdiv id=\"Sec3\" class=\"Section2\"\u003e\u003ch2\u003e2.1 Sample preparation\u003c/h2\u003e\u003cp\u003eThis study selected metal mold cast Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy sheet with dimensions of 100 mm \u0026times; 50 mm \u0026times; 15 mm as the raw material. The alloy composition was determined by Inductively Coupled Plasma luminescence spectrometer (ICP\u0026ndash;6300). A φ 5 mm \u0026times; 100 mm filler rod and a 70 mm \u0026times; 40 mm \u0026times; 10 mm substrate were machined from the cast plate employing wire cutting. Three sets of 30 mm \u0026times; 10 mm \u0026times; 5 mm grooves were processed on the substrate using a pneumatic milling cutter for the repair welding experiments. The schematic image of repair welding process is depicted in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea. Before the experiment, the surface of the substrate and filler rod were polished with 800 # grit sandpaper. TIG repair welding was used, and the welding process parameters are provided in Table\u0026nbsp;\u003cspan refid=\"Tab1\" class=\"InternalRef\"\u003e1\u003c/span\u003e.\u003c/p\u003e\u003cp\u003eTo further improve the mechanical properties of the repaired joints, they were subjected to solution treatment (T4) and aging treatment (T6). T4 treatment was performed at 500 ℃, 525 ℃, and 550 ℃ for 24 hours, subsequently water quenching. The repair welded joints after T4 treatment were assigned as T4\u0026ndash;1, T4\u0026ndash;2, and T4\u0026ndash;3, respectively. The T4 treated repaired joints were then exposed to aging treatment at 200 ℃ for 0\u0026ndash;100 hours. After BM peak aging treatment, the samples are labeled as T6\u0026ndash;1 (500 ℃ \u0026times; 24 h\u0026thinsp;+\u0026thinsp;200 ℃ \u0026times; 18 h), T6\u0026ndash;2 (525 ℃ \u0026times; 24 h\u0026thinsp;+\u0026thinsp;200 ℃ \u0026times; 16 h), and T6\u0026ndash;3 (550 ℃ \u0026times; 24 h\u0026thinsp;+\u0026thinsp;200 ℃ \u0026times; 14 h).\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec4\" class=\"Section2\"\u003e\u003ch2\u003e2.2 Characterization\u003c/h2\u003e\u003cp\u003eX\u0026ndash;ray computed tomography (CT, AX7900 X\u0026ndash;ray) was applied to detect defects in the repaired joints. Figure\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb depicts the processing location of the microstructure sample. After polishing, the microstructure sample was etched with a solution of 10 g picric acid, 8 ml acetic acid, 20 ml deionized water, and 70 ml ethanol for 15 seconds. Optical microscopy (OM, DSX1000) was employed to examine the metallographic structures of various repair welded joints. The second phase dissolution situation and element distribution were examined by a scanning electron microscope (SEM, Apreo C) fitted with an energy\u0026ndash;dispersive spectrometer (EDS). The angular distribution of texture strength and orientation differences was examined utilizing electron backscatter diffraction (EBSD, C\u0026ndash;Nano). The EBSD sample was electropolished applying a polishing solution consisting of 24 ml nitric acid, 36 ml hydrochloric acid, and 250 ml alcohol. The polishing voltage was adjusted to 22 V, the current to 1.5 A, and the electrolysis time to 22 seconds. Transmission electron microscopy (TEM, Talos F200X) was utilized to examine the second phase composition of solid solution samples and identify precipitated phases in aged samples. After mechanically grinding the transmission sample to 50 \u0026micro;m, the Gatan 695 ion thinning instrument was used for thinning, with an ion thinning voltage range of 0.1 keV to 4.5 keV and an angle range of \u0026minus;\u0026thinsp;8 \u0026deg; to 8 \u0026deg;.\u003c/p\u003e\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb illustrates the processing area (yellow region) and size of the tensile samples. To reduce experimental error, two sets of four tensile samples were cut from top to bottom, 1 mm from the upper surface of the sample for each heat treatment condition. Prior to the tensile test, the molten pool position of each sample was calibrated to determine the fracture location during testing. The tensile properties of the repair welded joints were evaluated at room temperature by means of a tensile machine (MTS E44.304) with a constant tensile rate of 1 mm/min. Microhardness measurements were conducted utilizing a Vickers hardness tester (DHV\u0026ndash;1000). The load applied was 49 N, with a holding time of 15 seconds.\u003c/p\u003e\u003cp\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\u003eRepair welding process parameters.\u003c/p\u003e\u003c/div\u003e\u003c/caption\u003e\u003ccolgroup cols=\"5\"\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e\u003cdiv align=\"left\" 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\u003cthead\u003e\u003ctr\u003e\u003cth align=\"left\" colname=\"c1\"\u003e\u003cp\u003eWelding current (A)\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c2\"\u003e\u003cp\u003eTungsten electrode diameter (mm)\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c3\"\u003e\u003cp\u003eFiller rod diameter (mm)\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c4\"\u003e\u003cp\u003eTorch length (mm)\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c5\"\u003e\u003cp\u003eThe flow rate of the Ar gas (L/min)\u003c/p\u003e\u003c/th\u003e\u003c/tr\u003e\u003c/thead\u003e\u003ctbody\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e200\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c2\"\u003e\u003cp\u003e3.2\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c3\"\u003e\u003cp\u003e5\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c4\"\u003e\u003cp\u003e2\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"left\" colname=\"c5\"\u003e\u003cp\u003e10\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003c/tbody\u003e\u003c/colgroup\u003e\u003c/table\u003e\u003c/div\u003e\u003c/p\u003e\u003c/div\u003e"},{"header":"3. Conclusion","content":"\u003cdiv id=\"Sec6\" class=\"Section2\"\u003e\u003ch2\u003e3.1 Microstructure of repair welded joints\u003c/h2\u003e\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ea presents the X\u0026ndash;ray computed tomography analysis of the repair welded joint in the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy. The X\u0026ndash;ray penetrates the entire sample from the upper surface, revealing no apparent repair defects, indicating a well formed joint. As depicted in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003eb, the repair welded joint comprises three distinct regions: the FZ, HAZ, and base material (BM). Due to the continuous heat transfer from the FZ to the BM, the FZ experiences a higher cooling rate, resulting in a finer grain structure with an average grain size of 14 \u0026micro;m. Conversely, the HAZ, which is adjacent to the FZ, undergoes the highest thermal input for an extended duration, leading to a slight rise in grain size (51 \u0026micro;m) in comparison with the BM (47 \u0026micro;m) (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ec). Further measurements indicate that the depth and width of the FZ are 5.5 mm and 14 mm, correspondingly, while the average width of the HAZ is 2.5 mm (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003eb).\u003c/p\u003e\u003cp\u003eAs shown in Figs.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ed and g, the FZ displays a finer microstructure in comparison with the HAZ. The secondary phase in both the FZ and HAZ consists of a particle phase enriched in Al and RE elements (Y, Nd) within the grains, a reticular phase rich in RE (Y, Nd), and a needle\u0026ndash;like phase rich in Al and RE (Y, Nd) at the grain boundaries \u003cem\u003e(\u003c/em\u003eFigs.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ee, f, h, and i). Additionally, under continuous thermal input, a small quantity of linear phase precipitates within the grains of the HAZ (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003eh). Previous research has identified the reticular phase as an Mg\u0026ndash;RE eutectic structure, whereas the needle\u0026ndash;like, particulate, and linear phases correspond to Al\u003csub\u003e2\u003c/sub\u003eRE phases [\u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e].\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec7\" class=\"Section2\"\u003e\u003ch2\u003e3.2 Microstructure of T4 treated joints after welding\u003c/h2\u003e\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e presents OM images of the repair welded joint following T4 treatment. After T4 treatment, the FZ area in the T4\u0026ndash;1 and T4\u0026ndash;2 samples are still composed of fine equiaxed grains, whereas a small quantity of grains in the FZ edge area of the T4\u0026ndash;3 sample exhibit the AGG phenomenon (Figs.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea, d and g). As depicted in Figs.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb, e, and h, the FZ\u0026rsquo;s grain size elevates progressively with rising T4 temperature, while the HAZ\u0026rsquo;s grain size doesn't exhibit significant variation. As illustrated in Figs.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ec, f, and i, the number of second phase in the repaired joint gradually reduces as the T4 temperature increases.\u003c/p\u003e\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e depicts the grain size distribution in the FZ and HAZ of the repair welded joints subjected to various T4 treatments. Compared to the repaired joints without heat treatment (WH), the FZ\u0026rsquo;s grain size of the T4\u0026ndash;1 and T4\u0026ndash;2 samples increased to 17 \u0026micro;m and 20 \u0026micro;m, correspondingly, whereas the HAZ\u0026rsquo;s grain size showed little variation, averaging 52 \u0026micro;m (Figs.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea\u0026ndash;d). When the repair welded joint underwent T4 treatment at 550\u0026deg;C \u0026times; 24 h, considering the occurrence of AGG in a small portion of the grains at the FZ edge in the T4\u0026ndash;3 sample, the average grain size (\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:GS\\)\u003c/span\u003e\u003c/span\u003e) is evaluated as follows [\u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e]:\u003cdiv id=\"Equ1\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equ1\" name=\"EquationSource\"\u003e\n$$\\:GS=xG{S}_{1}+\\left(1-x\\right)G{S}_{2}$$\u003c/div\u003e\u003cdiv class=\"EquationNumber\"\u003e1\u003c/div\u003e\u003c/div\u003e\u003c/p\u003e\u003cp\u003ewhere \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:G{S}_{1}\\)\u003c/span\u003e\u003c/span\u003e and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:G{S}_{2}\\)\u003c/span\u003e\u003c/span\u003e are the average sizes of abnormally grown grains and normal grains in FZ, respectively, and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:x\\)\u003c/span\u003e\u003c/span\u003e is the region fraction of abnormally grown grains. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ee, statistical analysis reveals that the values of \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:G{S}_{1}\\)\u003c/span\u003e\u003c/span\u003e, \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:G{S}_{2}\\)\u003c/span\u003e\u003c/span\u003e, and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:x\\)\u003c/span\u003e\u003c/span\u003e were 115 \u0026micro;m, 24 \u0026micro;m, and 5.4%, respectively (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ee). Based on calculations, the FZ\u0026rsquo;s average grain size in the T4\u0026ndash;3 sample exhibited a substantial increase to 29 \u0026micro;m, while the HAZ\u0026rsquo;s grain size marginally rose to 54 \u0026micro;m (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ef). Zhang et al. [\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e] depicted that after 8 hours of T4 treatment at 520\u0026deg;C for the traditional Mg\u0026ndash;4Nd\u0026ndash;2Y\u0026ndash;1Gd\u0026ndash;0.5Zr alloy repair welded joint, a pronounced AGG phenomenon occurred in the FZ region, causing a significant increase in the FZ\u0026rsquo;s grain size from 13 \u0026micro;m to 656.4 \u0026micro;m. Therefore, compared to the traditional Mg\u0026ndash;4Nd\u0026ndash;2Y\u0026ndash;1Gd\u0026ndash;0.5Zr alloy repair welded joint, the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy repair welded joint displays superior grain thermal stability during post weld heat treatment.\u003c/p\u003e\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003e displays the SEM and EDS analyses of repair welded joints subjected to different T4 treatment processes. Owing to the high\u0026ndash;melting\u0026ndash;point of the Al\u003csub\u003e2\u003c/sub\u003eRE phase, it remains undissolved in all samples across the T4 treatment conditions. After the joint underwent T4 treatment at 500\u0026deg;C for 24 h, some eutectic structures remain undissolved in the T4\u0026ndash;1 sample, and a limited amount of linear phase precipitates within the grains of the FZ. At a T4 temperature of 525\u0026deg;C, the eutectic structure in the T4\u0026ndash;2 sample is completely dissolved, the number of linear phases increases, and their size enlarges. As the T4 treatment temperature increases to 550\u0026deg;C, the linear phase in the T4\u0026ndash;3 sample nearly disappears. As illustrated in Figs.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ej and k, line scan (line 1) and energy spectrum (point 1) analyses were conducted to examine the newly precipitated linear phase in the FZ region. The findings demonstrate that the linear phase mainly comprises of Y, Nd, and Al elements, with an atomic ratio of Al to RE (Y and Nd) of about 2:1. Based on this, it can be inferred that this phase is identical to the linear phase in the HAZ region and is classified as the Al\u003csub\u003e2\u003c/sub\u003eRE phase. To further verify the composition of the second phase in the FZ after T4 treatment, transmission electron microscopy (TEM) was utilized to investigate the second phase in the FZ of the T4\u0026ndash;2 sample. As depicted in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e, three distinct morphologies of secondary phases were detected in the FZ: particulate, needle\u0026ndash;like, and linear phases (Figs.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ea, d, and g). EDS analysis confirms that these three phases predominantly include Al, Y, and Nd elements (Figs.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003eb, e, and h). Further selected area electron diffraction (SAED) analysis (Figs.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ec, f, and i) establishes that all three phases exhibit face\u0026ndash;centered cubic structures with an identical lattice constant (a\u0026thinsp;\u0026asymp;\u0026thinsp;0.790 nm), confirming their classification as Al\u003csub\u003e2\u003c/sub\u003eRE phases.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec8\" class=\"Section2\"\u003e\u003ch2\u003e3.3 Age hardening behavior of repair welded joints\u003c/h2\u003e\u003cp\u003eFigures\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ea\u0026ndash;c illustrates the aging hardening behavior of repair welded joints following various T4 treatments. The repair welded joints subjected to various T4 treatments exhibit similar aging hardening trends. The FZ consistently reaches the peak aging and over aging states ahead of the HAZ. Similarly, the HAZ reaches these two states before the BM. This results from the high cooling rate in the FZ, resulting in significant grain refinement, which significantly increases the grain boundary area. The increased grain boundary area provides additional nucleation sites for precipitates, thereby promoting their nucleation and growth [\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e]. Consequently, the FZ exhibits a greater tendency for precipitation compared to other regions. Additionally, because the HAZ is situated between the FZ and BM, it undergoes rapid heating and cooling during repair welding, expanding during heating and contracting during cooling. Uneven temperature changes during heating and cooling cause inconsistent expansion and contraction at different positions within the HAZ, leading to large residual stresses [\u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e]. These residual stresses increase the diffusion rate of solute atoms and accelerate the emergence of precipitates, which leads to the precipitation tendency in HAZ is greater than that in BM [\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e, \u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e, \u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e]. As the T4 temperature rises, the comprehensive aging of the repair welded joint increases, and the time necessary to achieve the peak aging state decreases. At a T4 temperature of 500 ℃, the peak aging state is reached in the FZ, HAZ, and BM regions at 12 h, 14 h, and 18 h, with corresponding hardness values of 84 HV, 83 HV, and 82 HV. As the T4 temperature increases to 525 ℃, the peak aging times in the FZ, HAZ, and BM regions shorten to 10 h, 12 h, and 16 h, while the hardness increases to 89 HV, 88 HV, and 87 HV. Further increasing the T4 temperature to 550 ℃ reduces the peak aging times to 8 h, 10 h, and 14 h, while the hardness increases to 90 HV, 89 HV, and 88 HV, reaching its maximum. As illustrated in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ed, the microhardness of the repaired joint at peak aging in the BM is selected to comprehensively evaluate its mechanical properties. At a T4 temperature of 500 ℃, the T6\u0026ndash;1 sample exhibits low aging hardness, with microhardness values of 78 HV, 80 HV, and 82 HV for the FZ, HAZ, and BM areas, respectively. As the T4 temperature increases to 525 ℃, the average microhardness of the FZ, HAZ, and BM regions of the T6\u0026ndash;2 sample rises significantly to 85 HV, 86 HV, and 87 HV, respectively. When the T4 temperature elevates further to 550 ℃, the microhardness of the FZ, HAZ, and BM areas of the T6\u0026ndash;3 sample increases slightly to 87 HV, 87 HV, and 88 HV, respectively. Aging hardening is closely related to the content of aging precipitates. With elevating T4 temperature, the dissolution of the second phase is significantly enhanced, leading to a higher concentration of solute atoms within the α\u0026ndash;Mg matrix. This generates a stronger driving force for the nucleation and growth of precipitates, thereby promoting an increased precipitation tendency and improving aging hardness [\u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e, \u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e].\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003eTo further investigate the precipitation behavior of strengthening phases during aging, the FZ under different aging conditions was analyzed using TEM. The T4\u0026ndash;1 sample was aged at 200\u0026deg;C for 12 h, revealing a small quantity of granular precipitates having a mean length of close to 6 nm, detected along the [11\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\stackrel{-}{2}\\)\u003c/span\u003e\u003c/span\u003e0] \u003csub\u003eα\u0026ndash;Mg\u003c/sub\u003e direction of the α\u0026ndash;Mg matrix (Figs.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003ea\u0026ndash;c). High\u0026ndash;resolution transmission electron microscopy (HRTEM) combined with fast fourier transform (FFT) examination demonstrates that, alongside the α\u0026ndash;Mg matrix, weak diffraction spots are detected at the 1/4, 2/4, and 3/4 positions of the [1\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\stackrel{\\text{-}}{\\text{1}}\\)\u003c/span\u003e\u003c/span\u003e00] \u003csub\u003eα\u0026ndash;Mg\u003c/sub\u003e, which are in agreement with the diffraction patterns characteristic of the β\u0026prime; precipitate phase. The phase relationship between β\u0026prime; precipitates and α\u0026ndash;Mg matrix is as follows: [010] \u003csub\u003eβ\u0026prime;\u003c/sub\u003e // [01\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\stackrel{\\text{-}}{\\text{1}}\\)\u003c/span\u003e\u003c/span\u003e0] \u003csub\u003eα\u0026ndash;Mg\u003c/sub\u003e, (200) \u003csub\u003eβ\u0026prime;\u003c/sub\u003e // (2\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\stackrel{\\text{-}}{\\text{1}}\\stackrel{\\text{-}}{\\text{1}}\\)\u003c/span\u003e\u003c/span\u003e0) \u003csub\u003eα\u0026ndash;Mg\u003c/sub\u003e [\u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e]. Figures\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003ed\u0026ndash;l shown the TEM analyzed of the T4\u0026ndash;2 sample under three aging conditions: under aging (200\u0026deg;C \u0026times; 2 h), peak aging (200\u0026deg;C \u0026times; 10 h), and over aging (200\u0026deg;C \u0026times; 100 h). Following aging at 200\u0026deg;C for 2 h, a small amount of granular precipitates, averaging 5 nm in length, formed in the FZ. SAED and HRTEM revealed distinct weak diffraction spots at the 1/2 position of [1\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\stackrel{\\text{-}}{\\text{1}}\\)\u003c/span\u003e\u003c/span\u003e00] \u003csub\u003eα\u0026ndash;Mg\u003c/sub\u003e, confirming that the precipitate phase was β\u0026prime;\u0026prime; [\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e] (Figs.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003ed\u0026ndash;f). As the aging time elevates to 10 h, the number and size of precipitates increase, the morphology changes from granular to flake\u0026ndash;like, and the average length increases to 10 nm (Figs.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003eg and h). Figure\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003ei shows that the precipitates at this stage are β\u0026prime; phase. The T4\u0026ndash;2 sample, subjected to an elevated solution temperature, demonstrates a markedly higher precipitate density and larger precipitate size compared to the peak\u0026ndash;aged T4\u0026ndash;1 sample. As aging time is extended to 100 h, the quantity of precipitates reduces, and their average size elevates to 56 nm. The precipitates remain of the flaky β\u0026prime; phase (Figs.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003ej\u0026ndash;l). In Mg\u0026ndash;RE alloys, both nano β\u0026prime;\u0026prime; and β\u0026prime; precipitates impede dislocation slip, thus improving the alloy's strength [\u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e]. In conclusion, increasing the solution temperature promotes the formation of more precipitates in the FZ. Simultaneously, as aging time is extended, the number of precipitates in the FZ region initially increases and then decreases, while their size continues to grow.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec9\" class=\"Section2\"\u003e\u003ch2\u003e3.4 Mechanical properties of heat treatment repair welded joints\u003c/h2\u003e\u003cp\u003eThe peak aging treatment process of the BM was selected to conduct the aging treatment on the repair welded joint. As shown in Figs.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003ea and b, the UTS, YS, and EL of the repair welded joint WH are 242 MPa, 165 MPa, and 8.6%, correspondingly. Relative to the WH sample, the UTS and YS of theT4 sample reduced, while the EL elevated markedly. After the joint underwent T4 treatment at 500\u0026deg;C for 24 h, the UTS, YS, and EL of the T4\u0026ndash;1 sample were 233 MPa, 151 MPa, and 10.3%, respectively. Upon elevating the T4 temperature to 525\u0026deg;C, the UTS and YS of the T4\u0026ndash;2 sample declined to 230 MPa and 141 MPa, while the EL improved to 12.4%. When the T4 temperature was further elevated to 550\u0026deg;C, the UTS and YS of the T4\u0026ndash;3 sample dropped to 226 MPa and 134 MPa, whereas the EL increased to 13.1%. In contrast, the UTS and YS of the aged samples elevated substantially, while EL reduced. At a T4 temperature of 500\u0026deg;C, the UTS, YS, and EL of the T6\u0026ndash;1 sample were 271 MPa, 200 MPa, and 4.9%. Upon increasing the T4 temperature to 525\u0026deg;C, the UTS, YS, and EL of the T6\u0026ndash;2 sample enhanced to 293 MPa, 215 MPa, and 5.8%, correspondingly. With an additional rise in T4 temperature to 550\u0026deg;C, the UTS and YS of the T6\u0026ndash;3 sample slightly increased to 304 MPa and 225 MPa, whereas the EL decreased to 5.2%. It is evident that at a T4 temperature of 550\u0026deg;C, the existence of a small amount of AGG in the FZ edge does not cause a decline in the mechanical properties of the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy repaired joints. Additionally, the mechanical properties of the repaired joints following T6\u0026ndash;2 heat treatment (UTS: 292 MPa, YS: 215 MPa, EL: 5.8%) are comparable to those of the T6 (525\u0026deg;C \u0026times; 24 h\u0026thinsp;+\u0026thinsp;200\u0026deg;C \u0026times; 16 h) unrepaired alloy (UTS: 293 MPa, YS: 218 MPa, EL: 6%), indicating that the repair process does not substantially compromise the mechanical properties of the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy under conventional heat treatment.\u003c/p\u003e\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003ec and d presents a macro image and a low\u0026ndash;magnification OM of the fracture location in the tensile sample. The fracture of the WH sample occurred in the HAZ, while the fracture locations of the heat\u0026ndash;treated samples predominantly appeared in the BM, even with minor AGG formation in the FZ edge of the T4\u0026ndash;3 or T6\u0026ndash;3 treated samples. This indicates that AGG did not significantly impact the mechanical properties of the repair welded joint. Figures\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003ea\u0026ndash;h illustrates the fracture morphologies of the repair welded joints with WH and heat treatment. The sample fractures exhibit a combined characteristic of transgranular and intergranular fractures, combined with fracture dimples, tearing edges, and cleavage planes. Fractured Al\u003csub\u003e2\u003c/sub\u003eRE particles (demonstrated by the orange arrow) are visible on the fracture surface of the WH sample, indicating that Al\u003csub\u003e2\u003c/sub\u003eRE particles are one of the primary crack sources (Fig.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003eb). Following solution treatment, the amount of tearing edges and dimples in the repair welded joint increases, while the number of cleavage planes reduces, compared to the WH sample. Consequently, the solution\u0026ndash;treated sample demonstrates higher ductility and lower strength. With increasing T4 temperature, the quantity of tearing edges and dimples progressively rises, while both the number and size of cleavage planes exhibit a gradual reduction (Figs.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003ec, e, and g). After aging treatment, the quantity and size of cleavage planes increase, whereas the number and size of tearing edges and dimples significantly decrease, compared to the solution treated repaired joint. Thus, the aged sample exhibits lower ductility and higher strength (Figs.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003ed, f, and h). Additionally, the fractured Al\u003csub\u003e2\u003c/sub\u003eRE particles remain visible after various heat treatments, indicating that these particles continue to be one of the primary crack sources in the heat\u0026ndash;treated samples.\u003c/p\u003e\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003e illustrates OM images of the longitudinal cross\u0026ndash;sections of tensile samples extracted from both the WH and the heat\u0026ndash;treated repair welded joint. The fracture mode of the WH repair welded joint exhibits a mixture of transgranular and intergranular fractures, with intergranular fracture being predominant. Cracks originate within the eutectic structure and propagate either along the brittle eutectic or into the grains (Figs.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003ea and b). This occurs because of the reticular morphology of the eutectic structure at grain boundaries, which has large dimensions and sharp edges that readily induce stress concentration under loading [\u003cspan citationid=\"CR13\" class=\"CitationRef\"\u003e13\u003c/span\u003e, \u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e]. After the joint underwent T4 treatment at 500\u0026deg;C for 24 h, the T4\u0026ndash;1 sample still contains undissolved eutectic structures, making crack initiation more likely at these sites, predominantly resulting in intergranular fracture (Fig.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003ec). With the rise in T4 temperature to 525\u0026deg;C and 550\u0026deg;C, the eutectic structure located at the grain boundaries dissolves completely. Fractured Al\u003csub\u003e2\u003c/sub\u003eRE particles (indicated by the orange arrow) appear on the fracture surface, confirming that Al\u003csub\u003e2\u003c/sub\u003eRE particles become the primary crack initiation sites following the complete dissolution of the eutectic structure (Figs.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003ee and g). This transition is attributed to the increased grain boundary bonding strength in T4\u0026ndash;2 and T4\u0026ndash;3 samples resulting from the complete dissolution of the eutectic structure, reducing the likelihood of crack formation along grain boundaries. Therefore, cracks tend to form on the Al₂RE particles within the grains and propagate through them, leading to a shift in the primary fracture mode from intergranular to transgranular. Following aging treatment, the fracture morphology of the repaired joint closely resembles that observed in the T4 treated condition. Nevertheless, the precipitation phases within the grains strengthens the matrix and restricts plastic deformation, leading to a more uniform fracture surface [\u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e] (Figs.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003ed f and h).\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e"},{"header":"4 Discussion","content":"\u003cdiv id=\"Sec11\" class=\"Section2\"\u003e\u003ch2\u003e4.1 Causes of AGG phenomenon and grain thermal stability\u003c/h2\u003e\u003cp\u003eTraditional Mg\u0026ndash;RE alloy repair welded joints are susceptible to AGG in the molten pool during post\u0026ndash;weld heat treatment, adversely affecting their mechanical properties [\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e, \u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e]. In this study, the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy repaired joint underwent T4 treatment at 550\u0026deg;C for 24 h, and only a few grains at the molten pool edge exhibited AGG. Compared to conventional Mg\u0026ndash;RE\u0026ndash;Zr alloy repaired joints, this alloy demonstrated superior thermal stability during post\u0026ndash;weld heat treatment, significantly mitigating the detrimental effects of AGG on joint properties. Therefore, a comprehensive analysis of the mechanisms underlying AGG formation and grain thermal stability in Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy repair welded joints is essential for further improving their mechanical properties.\u003c/p\u003e\u003cp\u003eDuring the T4 process, grain growth occurs to decrease interface energy, thus decreasing the system\u0026rsquo;s total energy and enhancing thermodynamic stability. This growth is typically spontaneous, influenced by atomic diffusion rates and interfacial energy [\u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e]. As the T4 temperature increases, atomic diffusion accelerates, promoting grain coarsening. At 500\u0026deg;C and 525\u0026deg;C, atomic diffusion remains relatively slow due to lower temperatures, resulting in a low grain growth trend. However, at 550\u0026deg;C, increased atomic diffusion enhances the tendency for grain coarsening. Additionally, grain boundaries are typically irregular regions within the lattice structure, which can induce lattice distortion and result in high energy, leading to significant interfacial energy, which is more pronounced in smaller grains due to their larger boundary area [\u003cspan citationid=\"CR39\" class=\"CitationRef\"\u003e39\u003c/span\u003e, \u003cspan citationid=\"CR40\" class=\"CitationRef\"\u003e40\u003c/span\u003e]. Compared to other regions, the FZ exhibits smaller grain sizes, with the edge region (Zone 2) showing even finer grains than the center (Zone 1) (Figs.\u0026nbsp;\u003cspan refid=\"Fig12\" class=\"InternalRef\"\u003e12\u003c/span\u003ea and d). This results in higher interface energy at the Zone 2, intensifying grain growth trend. Rapid cooling of the molten pool during repair welding generates substantial residual shrinkage stress, leading to an increased density of structural defects, for example vacancies and dislocations, therefore, during the solid solution treatment, grains in the molten pool are prone to thermodynamic instability, making them susceptible to coarsening [\u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e, \u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e]. This effect is particularly pronounced at the Zone 2, where the cooling rate is highest due to direct contact with the substrate. As a result, compared to the Zone 1, the Zone 2 experiences greater residual shrinkage stress, leading to a more prominent grain coarsening phenomenon. Based on the distribution and frequency of grain boundaries with varying misorientation angles shown in Figs.\u0026nbsp;\u003cspan refid=\"Fig13\" class=\"InternalRef\"\u003e13\u003c/span\u003ec, f, and g, the fraction of grain boundaries with high misorientation angles (\u0026gt;\u0026thinsp;45\u0026deg;) is significantly greater at the Zone 2 than at the Zone 1. These high\u0026ndash;angle grain boundaries possess elevated interfacial energy and enhanced atomic diffusion rates, accelerating grain migration and coarsening [\u003cspan citationid=\"CR41\" class=\"CitationRef\"\u003e41\u003c/span\u003e, \u003cspan citationid=\"CR42\" class=\"CitationRef\"\u003e42\u003c/span\u003e]. Consequently, under the combined effects of atomic diffusion, interface energy, and residual shrinkage stress, secondary recrystallization occurs at the molten pool edge in the T4\u0026ndash;3 sample after treatment at 550\u0026deg;C for 24 h, causing the rapid disappearance of surrounding grain boundaries and the consumption of adjacent smaller grains, leading to AGG. However, only a limited number of grains exhibit AGG at the molten pool edge in the T4\u0026ndash;3 sample, with the FZ\u0026rsquo;s average grain size increasing to 29 \u0026micro;m, still significantly smaller than in other regions. By contrast, in traditional Mg\u0026ndash;4Nd\u0026ndash;2Y\u0026ndash;1Gd\u0026ndash;0.5Zr alloy repair\u0026ndash;welded joints subjected to T4 treatment at 520\u0026deg;C for 8 h, the FZ\u0026rsquo;s average grain size increases substantially to 656.4 \u0026micro;m [\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e]. This comparison demonstrates that Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy repair welded joints exhibit superior grain thermal stability during post\u0026ndash;weld heat treatment. Therefore, further investigation into the mechanisms governing grain thermal stability in Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy repair welded joints is essential for optimizing their properties.\u003c/p\u003e\u003cp\u003eIt is widely proven that the presence of secondary phases markedly influences grain growth behavior. Based on the Zener pinning mechanism [\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e], stable secondary phases at grain boundaries can effectively hinder their movement at elevated temperatures, thereby restricting grain boundary migration and suppressing grain coarsening. In the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy, a substantial amount of high\u0026ndash;melting\u0026ndash;point Al\u003csub\u003e2\u003c/sub\u003eRE phases is present, particularly in the form of needle\u0026ndash;like Al\u003csub\u003e2\u003c/sub\u003eRE structures near grain boundaries. These phases effectively limit grain boundary migration and control grain coarsening during heat treatment in repair\u0026ndash;welded joints. The Zener pinning influence is solely operative when the grain size does not exceed a critical value (\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\text{D}\\)\u003c/span\u003e\u003c/span\u003e), which is defined as follows [\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e, \u003cspan citationid=\"CR43\" class=\"CitationRef\"\u003e43\u003c/span\u003e]:\u003cdiv id=\"Equ2\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equ2\" name=\"EquationSource\"\u003e\n$$\\:\\text{D=}\\frac{\\text{4r}}{\\text{3f}}$$\u003c/div\u003e\u003cdiv class=\"EquationNumber\"\u003e2\u003c/div\u003e\u003c/div\u003e\u003c/p\u003e\u003cp\u003ewhere \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\text{r}\\)\u003c/span\u003e\u003c/span\u003e and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\text{f}\\)\u003c/span\u003e\u003c/span\u003e are the diameter and volume fraction of needle\u0026ndash;like Al\u003csub\u003e2\u003c/sub\u003eRE. According to the SEM observations, the data of \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\text{r}\\)\u003c/span\u003e\u003c/span\u003e and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\text{f}\\)\u003c/span\u003e\u003c/span\u003e are calculated as be 1.32 \u0026micro;m and 1.8%. The determined \"\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\text{D}\\)\u003c/span\u003e\u003c/span\u003e\" value is 97.8 \u0026micro;m, which is substantially exceeding the average grain size of 14 \u0026micro;m in the FZ. Therefore, at high temperatures, the presence of abundant high\u0026ndash;melting\u0026ndash;point needle\u0026ndash;like Al\u003csub\u003e2\u003c/sub\u003eRE phases near the grain boundaries in the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al molten pool effectively restricts grain boundary migration during the repair welding process. This suppression of grain coarsening enhances the exceptional thermal stability of the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy, ensuring superior structural integrity during post\u0026ndash;weld heat treatment.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e\u003cdiv id=\"Sec12\" class=\"Section2\"\u003e\u003ch2\u003e4.2 Evolution of microstructure and mechanical properties\u003c/h2\u003e\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig13\" class=\"InternalRef\"\u003e13\u003c/span\u003e illustrates the microstructural evolution and fracture mechanisms of repaired joints subjected to various post\u0026ndash;weld heat treatment conditions. As illustrated in Fig.\u0026nbsp;\u003cspan refid=\"Fig13\" class=\"InternalRef\"\u003e13\u003c/span\u003ea, the WH sample contains a substantial amount of Mg\u0026ndash;RE eutectic structures, along with needle\u0026ndash;like and particulate Al\u003csub\u003e2\u003c/sub\u003eRE phases. During the tensile process, Mg\u0026ndash;RE eutectic structures at grain boundaries and particle Al₂RE phases within the grains, which are relatively large and exhibit complex morphologies, contribute to stress concentration, ultimately leading to crack initiation and brittle fracture. These two phases serve as the primary sources of cracking in the repair welded joint of the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy, thereby limiting further improvement in mechanical properties. Compared to particle Al₂RE, the Mg\u0026ndash;RE eutectic structures at grain boundaries are larger and possess sharper morphologies, making them more prone to stress concentration. When eutectic structures are present, cracks tend to initiate and propagate preferentially within them, predominantly resulting in intergranular fracture. Given that the Al₂RE phase has a high\u0026ndash;melting\u0026ndash;point (approximately 1600\u0026deg;C), it is difficult to dissolve in the matrix. Therefore, eliminating the Mg\u0026ndash;RE eutectic structure through heat treatment is essential for strengthening the overall properties of the repaired joint.\u003c/p\u003e\u003cp\u003eAs illustrated in Figs.\u0026nbsp;\u003cspan refid=\"Fig13\" class=\"InternalRef\"\u003e13\u003c/span\u003eb and c, after T4 treatment at 500\u0026deg;C for 24 h, a small quantity of linear Al\u003csub\u003e2\u003c/sub\u003eRE phases formed inside the grains of the T4\u0026ndash;1 sample, while some undissolved Mg\u0026ndash;RE eutectic structures remained at the grain boundaries. This led to a low solid solubility of RE elements in the α\u0026ndash;Mg matrix, limiting the quantity density of precipitates in the T6\u0026ndash;1 sample. Consequently, only a small quantity of fine granular β\u0026prime; precipitates formed, which had minimal impact on improving the sample's tensile strength. Additionally, the presence of undissolved Mg\u0026ndash;RE eutectic structures resulted in a predominantly intergranular fracture mode. As the T4 temperature reached 525\u0026deg;C, the quantity and size of linear Al\u003csub\u003e2\u003c/sub\u003eRE phases within the grains of the T4\u0026ndash;2 sample increased, and the Mg\u0026ndash;RE eutectic structures at the grain boundaries were fully dissolved. This enhanced the solid solubility of RE elements in the α\u0026ndash;Mg matrix, promoting the formation of numerous flake\u0026ndash;like β\u0026prime; precipitates in the T6\u0026ndash;2 sample, significantly improving its tensile strength. With the complete dissolution of the Mg\u0026ndash;RE eutectic structures, particulate Al\u003csub\u003e2\u003c/sub\u003eRE phases inside the grains became the primary crack initiation sites, leading to a transition from predominantly intergranular to transgranular fracture. At 550\u0026deg;C, the linear Al\u003csub\u003e2\u003c/sub\u003eRE phases inside the grains of the T4\u0026ndash;3 sample dissolved back into the matrix, further increasing the solid solubility of RE elements in the α\u0026ndash;Mg matrix. This resulted in an even higher density of flake\u0026ndash;like β\u0026prime; precipitates in the T6\u0026ndash;3 sample, which had the most pronounced effect on enhancing tensile strength. However, owing to the existence of numerous undissolved particulate Al\u003csub\u003e2\u003c/sub\u003eRE phases, these particles remained the primary crack sources, maintaining a predominantly transgranular fracture mode. Notably, at 550\u0026deg;C, although some grains at the edge of the FZ exhibited AGG, fractures did not occur in the FZ region, indicating that AGG did not markedly impact the mechanical properties of the repaired joint. Thus, it is essential to use the Hall\u0026ndash;Petch relationship to evaluate the grain boundary strengthening (\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:{\\sigma\\:}_{y}\\)\u003c/span\u003e\u003c/span\u003e) effect of repair welded joints with solid solution temperature of 550 ℃ [\u003cspan citationid=\"CR44\" class=\"CitationRef\"\u003e44\u003c/span\u003e, \u003cspan citationid=\"CR45\" class=\"CitationRef\"\u003e45\u003c/span\u003e]:\u003cdiv id=\"Equ3\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equ3\" name=\"EquationSource\"\u003e\n$$\\:{\\sigma\\:}_{y}={\\sigma\\:}_{0}+{K}_{y}{D}^{-\\frac{1}{2}}$$\u003c/div\u003e\u003cdiv class=\"EquationNumber\"\u003e3\u003c/div\u003e\u003c/div\u003e\u003c/p\u003e\u003cp\u003ewhere \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:{\\sigma\\:}_{0}\\)\u003c/span\u003e\u003c/span\u003e is the friction stress, \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:{K}_{y}\\)\u003c/span\u003e\u003c/span\u003e is a Hall-Petch coefficient (250 MPa \u0026micro;m1/2) [\u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e], and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:D\\)\u003c/span\u003e\u003c/span\u003e is the average grain size. Following T4 treatment at 550\u0026deg;C, the average grain sizes in the FZ and HAZ are 29 \u0026micro;m and 54 \u0026micro;m, correspondingly. Calculations show that the \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:{\\sigma\\:}_{y}\\)\u003c/span\u003e\u003c/span\u003e values in the FZ and HAZ are 58.4 MPa and 46 MPa, respectively. It can be inferred that the high strength retained in the FZ enhances its resistance to fracture during tensile deformation. However, considering the AGG phenomenon, the heat treatment process for the T6\u0026ndash;2 sample (525\u0026deg;C \u0026times; 24 h\u0026thinsp;+\u0026thinsp;200\u0026deg;C \u0026times; 16 h) is determined to be optimal. The mechanical properties of the T6\u0026ndash;2 sample (UTS: 292 MPa, YS: 215 MPa, EL: 5.8%) are comparable to those of the T6 alloy without repair welding (UTS: 293 MPa, YS: 218 MPa, EL: 6%).\u003c/p\u003e\u003cp\u003eIn summary, the excellent thermal stability of the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy successfully addresses the conflict between the occurrence of AGG and the dissolution of the eutectic structure in the repair weld, enabling the feasibility of single\u0026ndash;stage solution treatment, thereby effectively reducing actual industrial production costs. Furthermore, a study of the heat treatment process for the alloy repair weld revealed that, after optimal treatment, the properties of the repaired joints was significantly improved to a level comparable to the without repair welding alloy, effectively addressing the properties degradation typically observed in traditional Mg\u0026ndash;RE alloys after repair welding.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003c/div\u003e"},{"header":"5. conclusion","content":"\u003cp\u003e\u003col\u003e\u003cspan\u003e\u003cli\u003e\u003cp\u003eDuring heat treatment, the Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy repaired joint demonstrated notable thermal stability. With an elevate in solution temperature from 500\u0026deg;C to 525\u0026deg;C, the grain size in the FZ region exhibited only, a slight increase from 17 \u0026micro;m to 20 \u0026micro;m. An additional rise in solution temperature to 550\u0026deg;C, caused the occurrence of AGG in certain grains at the FZ edge, resulting in a grain size rise to 29 \u0026micro;m, which remained considerably smaller than in other regions. Additionally, despite the rise in solution temperature, the HAZ\u0026rsquo;s grain size region remained relatively stable (within the range of 52 \u0026micro;m to 54 \u0026micro;m).\u003c/p\u003e\u003c/li\u003e\u003c/span\u003e\u003cspan\u003e\u003cli\u003e\u003cp\u003eThe particle and needle\u0026ndash;like Al\u003csub\u003e2\u003c/sub\u003eRE phases possess a high\u0026ndash;melting\u0026ndash;point, making them challenging to dissolve within the matrix. Therefore, the elimination of the Mg\u0026ndash;RE eutectic structure is essential for improving mechanical properties. After the joint underwent T4 treatment at 500\u0026deg;C for 24 h, portion of the Mg\u0026ndash;RE eutectic structure remains undissolved, with a small limited amount of linear Al\u003csub\u003e2\u003c/sub\u003eRE phase precipitating within the grain. As the T4 temperature increases to 525\u0026deg;C, the Mg\u0026ndash;RE eutectic structure undergoes complete dissolution, accompanied by a growth in in both the quantity and size of the linear Al\u003csub\u003e2\u003c/sub\u003eRE phase. With a further rise in T4 temperature to 550\u0026deg;C, the linear Al\u003csub\u003e2\u003c/sub\u003eRE phase is entirely eliminated, however, the AGG phenomenon appears at the FZ edge. Based on these observations, a T4 treatment at 525\u0026deg;C for 24 h is determined to be optimal.\u003c/p\u003e\u003c/li\u003e\u003c/span\u003e\u003cspan\u003e\u003cli\u003e\u003cp\u003eThe duration of the aging treatment has a markedly influence on the type, morphology, and size of precipitates. When the T4\u0026ndash;2 sample undergoes aging at 200\u0026deg;C for 2 h, a limited quantity of granular β\u0026prime;\u0026prime; phase precipitates form in the FZ region, with an averaging about 5 nm in length. With an extension of aging time to 10 h, both the quantity and size of the precipitates increase, and the phase type transitions to a flake\u0026ndash;like β\u0026prime; phase with an average length of 10 nm. Further extending the aging time to 100 h, the precipitate type remained as flake\u0026ndash;like β' phase, however, the number of precipitates decreases while their size increases to 56 nm.\u003c/p\u003e\u003c/li\u003e\u003c/span\u003e\u003cspan\u003e\u003cli\u003e\u003cp\u003eAfter T4 treatment at 525\u0026deg;C for 24 h and subsequent aging at 200\u0026deg;C for 16 h, the repair welded joint displayed excellent overall mechanical properties. The UTS, YS, and EL reached 292 MPa, 215 MPa, and 5.8%, correspondingly. These values are comparable to the mechanical properties of the T6 alloy in its original, unrepaired state, demonstrating that the optimized heat treatment effectively restores the strength and ductility of the repair welded joint.\u003c/p\u003e\u003c/li\u003e\u003c/span\u003e\u003c/ol\u003e\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eAuthor contribution\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eZhengtao Liu\u003c/strong\u003e: Data curation, Investigation, Methodology, Writing \u0026ndash; original draft. \u003cstrong\u003eLei Wang\u003c/strong\u003e: Conceptualization, Funding acquisition, Writing \u0026ndash; review \u0026amp; editing. \u003cstrong\u003eHongyang Duan\u003c/strong\u003e: Investigation, Methodology, Validation. \u003cstrong\u003eSicong Zhao\u003c/strong\u003e: Data curation, Formal analysis, Funding acquisition.\u003cstrong\u003e\u0026nbsp;Erjun Guo\u003c/strong\u003e: Supervision, Validation. \u003cstrong\u003eYicheng Feng\u003c/strong\u003e: Data curation.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eFunding\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThis work was supported by the Heilongjiang Province Postdoctoral Science Foundation (LBH-Z24308).\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eData availability\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eData in the present work can be obtained from the corresponding author on request.\u003cstrong\u003e\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eEthics approval\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe manuscript was approved by all authors for publication.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eConsent to participate:\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eNot applicable.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eConsent for publication:\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eNot applicable.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCompeting interests\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\u003cli\u003e\u003cspan\u003eYang J, Zhu Z, Han S, Gu Y (2024) Evolution, limitations, advantages, and future challenges of magnesium alloys as materials for aerospace applications. 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Mater Sci Eng A 677(20):411\u0026ndash;420. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.msea.2016.09.044\u003c/span\u003e\u003cspan address=\"10.1016/j.msea.2016.09.044\" 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":"Magnesium alloy, Repair welding, Heat treatment, Mechanical properties, Grain thermal stability","lastPublishedDoi":"10.21203/rs.3.rs-7018798/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-7018798/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eMg\u0026ndash;RE alloys refined with Al demonstrate excellent grain thermal stability, enhancing the mechanical properties of repair welded joints after heat treatment. This research examines the impact of heat treatment on the microstructure and mechanical properties of repair welded joints composed of Mg\u0026ndash;4.02Y\u0026ndash;3.06Nd\u0026ndash;1.56Al alloy. The findings indicate that the microstructure of the repaired joint comprises primarily of an equiaxed α\u0026ndash;Mg matrix, an Mg\u0026ndash;RE eutectic structure, particle and needle\u0026ndash;like Al\u0026ndash;RE phases. The particle and needle\u0026ndash;like Al\u0026ndash;RE phases have a high\u0026ndash;melting\u0026ndash;point and do not readily dissolve during the heat treatment process. Eliminating the Mg\u0026ndash;RE eutectic structure is crucial for improving the mechanical properties of the repair welded joint. If the solution temperature is too low (500\u0026deg;C), an excessive amount of the residual Mg\u0026ndash;RE eutectic structure will hinder the improvement of mechanical properties. At a solution temperature of 550\u0026deg;C, although the pinning effect of Al\u003csub\u003e2\u003c/sub\u003eRE on the grain boundary inhibits abnormal grain growth (AGG), localized AGG still occurs at the FZ edge. At 525\u0026deg;C, the eutectic structure fully dissolved, and no AGG occurred, achieving an optimal solid solution influence. Subsequently, after aging at 200\u0026deg;C for 16 h, a high density β' strengthening phase uniformly precipitated in the matrix, markedly enhancing the alloy's strength. As a result, the joint's UTS (292 MPa), YS (215 MPa), and EL (5.8%) were comparable to those of the original, unrepaired alloy.\u003c/p\u003e","manuscriptTitle":"Strengthening and toughening mechanism of Mg–RE–Al alloy repair welded joints by suppressing abnormal grain growth during heat treatment","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-07-28 18:17:01","doi":"10.21203/rs.3.rs-7018798/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":"e1957c8c-a47a-4615-906f-6b832fc75a30","owner":[],"postedDate":"July 28th, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"posted","subjectAreas":[],"tags":[],"updatedAt":"2025-09-03T11:08:01+00:00","versionOfRecord":[],"versionCreatedAt":"2025-07-28 18:17:01","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-7018798","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-7018798","identity":"rs-7018798","version":["v1"]},"buildId":"8U1c8b4HqxoKbykW_rLl7","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}