Exceptional High-temperature Strength in an Additively Manufactured Al-based Superalloy with Stable Nanoscale Eutectic Cellular Network

Exceptional High-temperature Strength in an Additively Manufactured Al-based Superalloy with Stable Nanoscale Eutectic Cellular Network | 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 Article Exceptional High-temperature Strength in an Additively Manufactured Al-based Superalloy with Stable Nanoscale Eutectic Cellular Network Zhe Chen, Siming Ma, Haixing Fang, Gang Ji, Mingliang Wang, Yuchi Cui, and 7 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-5840317/v1 This work is licensed under a CC BY 4.0 License Status: Published Journal Publication published 22 Nov, 2025 Read the published version in Nature Communications → Version 1 posted You are reading this latest preprint version Abstract Metallic materials typically experience significant strength degradation at elevated temperatures. Traditional strengthening methods, which rely on thermally stable particle dispersion, exhibit limited effectiveness owing to the challenges in suppressing thermally activated dislocation motion. This work introduces a novel strategy for achieving exceptional high-temperature strength through a thermally stable nanoscale eutectic cellular network (ECN) enabled by additive manufacturing. A near-eutectic AlLaScZr alloy is developed for laser powder bed fusion, incorporating an Al-La nanoscale ECN and dense intracellular nanoprecipitates. This alloy demonstrates excellent printability and remarkable high-temperature yield strength above 0.6 T m (~ 250 MPa at 300°C), outperforming conventional aluminium alloys by 2–5 times with minimal degradation after prolonged annealing. Compared with the conventional configuration of particle dispersion, the nanoscale ECN architecture enhances load-bearing capacity and strengthens aluminium by caging dislocation motion within ultrafine cells (~ 200 nm), effectively mitigating intrinsic high-temperature softening. The proposed transformative approach paves the way for designing next-generation heat-resistant alloys, unlocking new possibilities for additive manufacturing in high-temperature applications. Physical sciences/Materials science/Structural materials/Metals and alloys Physical sciences/Materials science/Structural materials/Mechanical properties Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Introduction Softening owing to severe strength degradation is inevitable in metallic materials at elevated temperatures, especially above 0.5 T m . For instance, at 300°C, most Al alloys exhibit a low yield strength of only a few tens of MPa, typically 10–40% of that at ambient temperature 1 , 2 . Strategies to enhance high-temperature strength have traditionally focused on incorporating second-phase reinforcements, such as micro- and nanosized intermetallics, precipitates, or ceramic particles, to promote well-known load transfer and Orowan strengthening mechanisms 1 . In this context, efforts have been primarily aimed at improving the coarsening resistance and increasing the volume fraction of reinforcements. For example, a Sc-modified Al-Cu-Mg-Ag alloy contains high-density and thermally stable coherent nanoprecipitates V -(Al,Cu,Sc) 3 , exhibiting a superior yield strength of ~ 100 MPa at 400°C. In addition, a multiple precipitate/matrix interface engineering strategy was reported to stabilise the θ′-Al 2 Cu precipitates in an Al-Cu-Mg-Ag-Si-Sc alloy, retaining 97% yield strength after thermal exposure at 200°C 4 . In conventionally fabricated alloys, the abovementioned strategies rely on the uniform distribution of hard particles within a softer matrix. However, this architecture may limit the high-temperature strengthening effects. At elevated temperatures, thermally activated dislocation motions such as cross-slipping and climbing become dominant, allowing dislocations to bypass obstacles more easily and weakening the Orowan strengthening effect 5 , 6 . Additionally, dispersed particles exert limited geometrical constraints on dislocation movement in the continuous alloy matrix, as thermally activated dislocation motion can easily accommodate the strain incompatibility between the alloy matrix and particles 7 , reducing the load transfer effectiveness. Recently, additive manufacturing (AM) has demonstrated notable industrial application prospects in the rapid prototyping of complex metallic components with unprecedented freedom 8 , 9 , as well as the realisation of novel metastable and ultrafine microstructures through rapid solidification. Specifically, a notable 3D ultrafine cellular-like architecture, where one phase forms cells enclosed by a 3D interconnected network of another phase, has been reported in AM Al 10 , Fe 11 , Cu 12 , high-entropy alloys (HEAs) 13 , and other alloys. For example, recent studies on AM AlSi10Mg have demonstrated the intrinsic exceptional room-temperature (RT) mechanical properties, attributable to a nanoscale eutectic Si cellular network 14 , 15 . These materials achieve yield and fatigue strengths of ~ 470 MPa in defect-free micro-sized samples, surpassing conventionally fabricated counterparts 14 . During co-deformation, the ultrafine cellular network exhibits superior load transfer 7 , 16 , 17 , with stress in the Si phase at cell boundaries measured at 1.5–2 GPa 16 , 4–5 times that in conventional Si alloys 18 , 19 . Additionally, this network constrains dislocation motion within cells as ‘dislocation cages’, leading to strong work hardening and delayed damage under tensile and fatigue loads 14 , 15 . However, these benefits have been limited to RT applications owing to the low thermal stability of the Si cellular network 20 , which undergoes spheroidisation into dispersed Si particles after a short holding time at 300°C. The high-temperature strength is therefore significantly degraded 21 . Forming a thermally stable 3D cellular network architecture may address this limitation, paving the way for enhanced high-temperature strength. This paper introduces a novel architecture design strategy for developing high-temperature ‘Al-based superalloys’ using AM to achieve superior high-temperature strength and thermal stability. A novel AM AlLaScZr alloy was fabricated using laser powder bed fusion (LPBF), in which La undergoes eutectic alloying to form a thermally stable nanoscale Al-La eutectic cellular network (ECN). La was selected owing to its extremely low diffusivity and solubility in face-centred cubic (fcc)-Al among all alloying elements. Additional micro-alloying with Sc and Zr was performed to form coarsening-resistant L 12 coherent nanoprecipitates, providing synergistic strengthening. The LPBF AlLaScZr alloy exhibited extraordinary high-temperature yield strength (YS), ranking among the highest reported for Al alloys above 0.6 T m (~ 250 MPa at 300°C and ~ 110 MPa at 400°C). The superior high-temperature strength was attributable to the high load-bearing capacity of the Al-La ECN, combined with the high stress in the soft α-Al phase owing to the restricted dislocation free path within ultrafine-sized cells (~ 200 nm). Synergic strengthening was contributed by the intracellular nanoprecipitates as additional barriers of dislocation motion. Even after prolonged annealing and high-temperature testing, the alloy exhibited superior YS retention (~ 200 MPa at 300°C and ~ 100 MPa at 400°C) owing to the highly dense and well dispersed La-rich nanoparticles inherited from the ECN. The proposed strategy can inspire the design of other thermally stable nanostructures using state-of-the-art AM technologies, facilitating the development of alloys with enhanced high-temperature mechanical properties. Results Superior high-temperature strength We prepared an Al-9.8La-0.46Sc-0.26Zr (wt.%) alloy for LPBF using a microstructure refinement (heterogeneous nucleation on the primary nuclei) and eutectic solidification strategy (near-eutectic composition) to achieve the optimal printability based on the CALPHAD method 22 . The temperature–solidification fraction ( T – f s ) curve for this alloy exhibits an ‘L’-shaped profile (Fig. 1 a), indicating that solidification initiates with the precipitation of primary L 12 -Al 3 (Sc, Zr) dispersoids over a broad temperature range (740°C − 640°C), represented by a nearly vertical line. The horizontal segment of the curve corresponds to eutectic solidification of fcc-Al and Al 11 La 3 , extending from f s > 0.2 to the terminal stage at ~ 638°C with a near-zero solidification temperature range. Crack-free and nearly fully dense bulk samples (~ 99.7% relative density) were printed (inset of Fig. 1 b) using optimised LPBF parameters (190 W power, 1500 mm/s scanning speed) (Fig. S1 b). The pore sizes are mostly under 40 µm, as observed in 3D micro-computed tomography (CT) scans (Figs. S1e, f). The as-built samples were subjected to post-heat treatment at 325°C for 1 h to achieve peak hardening through L 12 -Al 3 (Sc, Zr) precipitation (Fig. S1 c). The LPBF AlLaScZr alloy exhibits a bimodal grain structure, with large columnar grains (10–50 µm) and small equiaxed grains (< 10 µm) without preferential texture (Fig. 1 b, Fig. S2). The refined equiaxed grains indicate the inoculation effect of the primary L 12 -Al 3 (Sc, Zr) phase 23 . At the subgrain scale, a ubiquitous nanoscale ECN structure is observed (Fig. 1 c), with La-rich eutectic phases forming bright-contrast cell walls around α-Al cells, as confirmed by the scanning transmission electron microscopy (STEM) high-angle annular dark-field (HAADF) image and corresponding energy-dispersive spectroscopy (EDS) maps (Fig. 1 d). The cellular structure exhibits slight heterogeneity, with coarser cellular dendrites at melt pool boundaries and finer equiaxed cells at melt pool interiors owing to cooling rate variations 24 (Fig. S3). A 3D view of the highly interconnected ECN structure is shown in the inset in Fig. 1 c, reconstructed from a series of 2D backscattered electron (BSE) images. The ECN-enclosed cells are slightly elongated along the building direction (BD) and feature an ultrafine cell size (203.3 ± 69.2 nm), finer than those reported in LPBF Al-Si alloys (0.5–1 µm) 15 , 24 and other LPBF materials such as Cu 12 , Ti 25 , 26 , and steels 27 , 28 (0.4–1 µm on average). The LPBF-processed and peak-aged AlLaScZr alloy with the nanoscale ECN structure is referred to as the AlLaScZr-ECN alloy hereafter. The AlLaScZr-ECN alloy demonstrates outstanding mechanical performance at both RT and elevated temperatures, achieving exceptionally high YS at and above 300°C (~ 0.6 T m of Al) (Fig. 1 e). At RT, the AlLaScZr-ECN alloy exhibits a YS of 418 MPa, an ultimate tensile strength (UTS) of 453 MPa, and an elongation of ~ 7%, comparable to those of the high-strength LPBF Scalmalloy®. Notably, at 300°C, the alloy reaches a remarkable YS of 246 MPa and UTS of 265 MPa, surpassing conventional Al alloys by fivefold and conventional heat-resistant Al alloys by twofold. Furthermore, it significantly outperforms commercial LPBF Al alloys, such as AlSi10Mg and AlMgScZr, by 2–4 times, as well as recently developed heat-resistant near-eutectic LPBF Al alloys, including Al-Ni, Al-Ce, and Al-Fe systems (Fig. 1 f). The outstanding high-temperature YS of the AlLaScZr-ECN alloy highlights its potential for lightweight structural applications at a medium temperature range (300–400°C), currently dominated by ferrous, titanium, and nickel-based alloys 1 . The specific YS of the AlLaScZr-ECN alloy at 300°C is 86 MPa/(kg·m − 3 ), nearly double that of commercial heat-resistant 2618 Al alloy and superior to several commercially pure titanium (CP-Ti) alloys, stainless steels, and Ni-based alloys (Fig. 1 g). Excellent thermal stability above 0.6 T m A common challenge in AM alloys is the spheroidisation and coarsening of fine cellular structures and precipitates at elevated temperatures, attributable to the thermodynamically large driving force to reduce the interface energy, as observed in LPBF Al-Si alloys 7 , 44 . In contrast, the LPBF AlLaScZr-ECN alloy demonstrates excellent microstructural thermal stability and high strength retention even after prolonged thermal exposure at and above 300°C, despite its large cell boundary (CB) area and dense nanoprecipitates (Fig. 2 ). The thermal stability is primarily attributable to the La-rich cell walls, which provide both thermodynamic and kinetic resistance to coarsening owing to La’s extremely low diffusivity and solubility in α-Al. At 300°C, La exhibits a diffusivity of 6.9 × 10 − 21 m 2 ·s − 1 and a maximum solid solubility of 0.01 at.%, both significantly lower than those of Si (2.6 × 10 − 16 m 2 ·s − 1 and 1.6 at.%, respectively) 45 , 46 . In addition, a high density (7.3 × 10 23 m − 3 ) of coarsening-resistant L 12 -Al 3 Sc nanoprecipitates are dispersed ubiquitously within α-Al cells, with an average size of 3.1 nm, as revealed by atomic-probe tomography (APT) (Fig. 2 e) and high-resolution transmission electron microscopy (HRTEM) with the fast Fourier transformation (FFT) pattern (Fig. 3 b, d). These nanoprecipitates form owing to Sc solute supersaturation under rapid solidification during the LPBF 23 . Additionally, a small amount of Al 11 La 3 nanoparticles (5–10 nm in size) within α-Al cells (Fig. 3 b), likely precipitated owing to La supersaturation due to LPBF, can be observed 47 . This unique combination of thermodynamically and kinetically stable phases endows the alloy with superior thermal stability and mechanical performance at elevated temperatures. Moreover, the CBs exhibit high coherency and low interfacial energy, rendering them more stable than the high-angle grain boundaries (HAGBs) commonly observed in ultrafine-grained or nano-grained (UFG/NG) metals. The La-rich cell wall of the ECN is identified as the Al 11 La 3 phase (orthorhombic, group: Immm) by X-ray diffraction (XRD) analysis (Fig. 3 a), HRTEM analysis with the FFT pattern (Fig. 3 a, c), and APT stoichiometric analysis (Fig. 2 e). The Al 11 La 3 cell walls (and intracellular Al 11 La 3 nanoparticles) present a specific orientation relationship (OR) with α-Al along the Al[001] zone axis: Al[001]//Al 11 La 3 [010], Al(200)//Al 11 La 3 (200) (Fig. 3 a). The interface of Al 11 La 3 /α-Al presents a remarkably low interfacial misfit of 0.5% (Fig. 3 e) and an interfacial energy (0.45 J·m − 2 ) lower than that (0.8–2.0 J·m − 2 ) of conventional high-energy HAGBs in UFG/NG alloys 48 , as supported by density functional theory (DFT) calculations (Table S3, Supplementary Text). Additionally, localised segregation of Sc on the Al 11 La 3 CBs (circled in Figs. 3 a, 3 b, and 2 e) further stabilises the CBs by accommodating the misfit of α-Al/Al 11 La 3 interfaces. HRTEM images reveal that Sc segregation results in the formation of Al 3 Sc at the Al 11 La 3 cell wall surface with an OR of Al 3 Sc[100]//Al 11 La 3 [010], Al 3 Sc(200)//Al 11 La 3 (200) (Fig. 3 a), similar to that between α-Al and Al 11 La 3 . Similar Sc segregation at phase boundaries has been observed in other eutectic Al alloys, such as Al-Ce and Al-Fe-Ni 49 , 50 , where the larger atomic size of Sc alleviates local lattice distortion. DFT calculations indicate that the lattice misfit between Al 3 Sc and Al 11 La 3 is -1.2%, while that between α-Al and Al 11 La 3 is + 0.5%. These outcomes suggest that Al 3 Sc segregation reduces the interface energy by accommodating the positive lattice misfit at the α-Al/Al 11 La 3 interface (Fig. 3 e, Table S2). Consequently, even after prolonged annealing at 300°C for 168 h, the ECN demonstrates remarkable resistance to coarsening, retaining its connectivity and stability (Fig. 2 b, Fig. S4a). The interconnectivity of the thermally stable ECN is disrupted upon annealing at a higher temperature of 325°C for an extended period of 168 h (Fig. 2 c, Fig. S4a). This process results in the complete spheroidisation of Al 11 La 3 cell walls into Al 11 La 3 nanoparticles (55.2 ± 18.2 nm), forming a structure referred to as AlLaScZr-NP, which is compared with the original AlLaScZr-ECN (Figs. 2 a, c). XRD patterns and atom stoichiometric ratios of the La-rich zone derived from APT mapping confirm that the phase transformation of Al 11 La 3 (Figs. 2 d, e) does not occur during the annealing process. Sc segregation along Al 11 La 3 persists. Zr segregates within Sc-rich regions (Fig. 2 e), indicating the formation of Al 3 (Sc, Zr) with a Zr-rich shell, which retards Al 3 Sc coarsening 51 . The intracellular Al 3 Sc particle size increases slightly from 3.1 to 4.6 nm, demonstrating strong resistance to coarsening, consistent with previous reports on Al alloys micro-alloyed with Sc (Zr) 23 , 52 . Additionally, grain coarsening is minimal after prolonged thermal exposure, owing to the abundant Al 11 La 3 nanoparticles pinning the grain boundaries (Fig. S2). In terms of mechanical properties after the spheroidisation annealing, the YS of the AlLaScZr-NP alloy decreases at both RT and 300°C (Fig. 1 e). Nevertheless, the alloy retains a high YS of ~ 200 MPa at 300°C. Moreover, at an elevated tensile testing temperature of 400°C, the YS difference between the AlLaScZr-ECN and AlLaScZr-NP alloys becomes minimal, as the ECN structure fully spheroidises into Al 11 La 3 nanoparticles during holding in the tensile test (Fig. S4b). Despite this structural change, both alloys maintain an exceptional YS of over 100 MPa at 400°C, which is at least double that of conventional heat-resistant 2618 Al alloys (Fig. 1 g). This result suggests that even after the transformation of Al 11 La 3 cell walls into nanoparticles, their high density, fine particle size (~ 55 nm), and uniform dispersion contribute to the retention of superior YS at elevated temperatures, similar to other alloys with high-density reinforcements 3 , 53 . ECN-induced high-temperature strengthening mechanisms To investigate the strengthening mechanisms of the LPBF AlLaScZr alloy, particularly at elevated temperatures, we conducted in situ synchrotron X-ray diffraction (SXRD) tensile tests at RT and 300°C on both AlLaScZr-ECN and AlLaScZr-NP samples, as shown in Fig. 4 (see Fig. S5, Methods, and Supplementary Text for details). First, the reinforced Al₁₁La₃ phase in the nanoscale ECN configuration exhibits superior load transfer ability at both RT and elevated temperatures compared with the dispersed Al 11 La 3 nanoparticles. Larger lattice strains are observed in different crystalline planes of Al 11 La 3 in the ECN configuration (Fig. 4 a, b), increasing by ~ 0.002 at RT and by ~ 0.001 at 300°C at the yield point. This is followed by plastic deformation in the (101), (130), and (132) planes, resulting in a 100–200 MPa increase in phase stress, which represents a 20–25% improvement compared to the dispersed particle configuration. (Fig. 4 c, d). Additionally, during plastic deformation, Al 11 La 3 in the ECN configuration shows prolonged load-bearing and delayed load shedding compared with the dispersed nanoparticles, resulting in enhanced work hardening and UTS. This effect is more pronounced at RT than at 300°C. Second, the α-Al phase in the ECN exhibits extraordinarily high phase stress even at 300°C. At RT, the α-Al phase stress in the AlLaScZr-ECN alloy is moderately higher than that in the AlLaScZr-NP alloy, with YS rising from 315 MPa to 360 MPa (only a ~ 15% increase). In contrast, at 300°C, the α-Al phase in the ECN configuration displays a marked improvement, with an average lattice strain of 0.003 at yield point and a YS of ~ 180 MPa, approximately 60 MPa higher than that in the AlLaScZr-NP sample (a ~ 50% increase) and 3–7 times that of conventional Al alloys at 300°C (25–50 MPa) 2 , 43 . These results demonstrate that the Al 11 La 3 reinforcing phase provides superior load-bearing ability in the ECN configuration compared with dispersed nanoparticles at both RT and elevated temperatures. The ECN configuration also synergistically promotes high flow stress in the α-Al phase, contributing to the alloy’s exceptional high-temperature strength. This superior performance of cellular structures in AM alloys has been well-documented at RT, as their strong cell walls and refined cell size offer substantial load-bearing advantages over conventional alloys 7 , 12 , 16 . For example, in LPBF AlSi10Mg alloys, nanoscale Si cellular networks bear 4–5 times the stress and are more fracture-resistant than the coarser Si networks and particles found in cast Al-Si alloys 16 . Moreover, Li et al. reported that a continuous Si network in LPBF AlSi10Mg enhances YS and work hardening compared with the annealed state with discrete Si particles 7 . In this study, direct comparisons of lattice strain and stress partitioning demonstrate the superior load-bearing capability of the Al 11 La 3 phase in the nanoscale ECN configuration not only at RT but also at elevated temperatures, provided thermal stability is maintained. From the perspective of dislocation behaviour, prior studies have shown that cellular structures in AM alloys facilitate co-deformation of the matrix and cell walls, primarily through dislocation accumulation at cell boundaries. For instance, in LPBF AlSi10Mg alloys, Si cell wall constraints lead to high-density geometrically necessary dislocations at Al/Si interfaces to manage misfit strain during deformation 7 . Similarly, in LPBF 316L stainless steel, dislocations have been observed to pin and tangle at cell walls 54 . Kwon et al. recently proposed a dislocation-based constitutive model for the AM-induced cellular structure in a LPBF Cu-Sn model alloy, describing that dislocations originate from the intercellular Frank–Read (F–R) sources and deposit from the cell interior onto the walls 12 . In contrast, this study reveals different dislocation behaviours during the co-deformation of α-Al and Al 11 La 3 cell walls at RT and 300°C. Dislocations are confined within individual cells and remain as single dislocation segments between cell walls even at a high strain level ( ε ≈ 7%) during tensile tests at RT (Fig. S6a) and 300°C (Figs. 5 a, b). Unlike in previous reports 7 , 12 , no substantial dislocation segregation or entanglement is observed at the cell boundaries. In terms of the microstructural differences from other reported AM-induced cellular structures, the ultrafine (~ 200 nm) cell size of the Al-La ECN is smaller than the typical cell sizes (0.5–1 µm) found in LPBF Al-Si, steel, and Cu alloys. This refined cell size likely activates the confined layer slip (CLS) mechanism, as proposed for nanolamellar architectures 55 . According to the CLS theory, ductile layers with dimensions below 200 nm deform through single dislocation segments forming Orowan-type loops between two parallel interfaces, rather than through dislocation arrays at a larger scale 56 , 57 . Given the size distribution of the Al-La ECN-enclosed cells (~ 200 nm on average), the CLS mechanism likely dominates (Fig. 5 g), with dislocations gliding as Orowan-type loops confined by the Al 11 La 3 cell walls. This behaviour is consistent with the isolated dislocation segments observed in the cells. The critical stress for Orowan bowing of dislocations confined between two interfaces is calculated using Eq. 1 58 : $$\:{{\sigma\:}}_{CLS}=\frac{MGb}{8\pi\:h}\left(\frac{4-\nu\:}{1-\nu\:}\right)\text{ln}\left(\frac{\alpha\:h}{b}\right)$$ 1 At 300°C, although the shear modulus of Al decreases from 26.2 GPa (at RT) to 22.5 GPa 6 , \(\:{{\sigma\:}}_{CLS}\) can still reach values of 100–180 MPa within the cell size range of 100–200 nm, corresponding to the layer thickness h (Fig. 5 g, Table S4). This supports the high stress observed in the α-Al phase enclosed by the ECN during SXRD analysis. Specifically, a short dislocation segment constrained by neighbouring cell walls (inset in Fig. 5 h) results in severe lattice distortion in α-Al and localised high strain, as indicated by the inverse fast Fourier transform (IFFT) pattern and geometrical phase analysis (GPA) (Fig. 5 i). In comparison, in the AlLaScZr-NP sample, longer dislocation segments, spanning several hundred nanometres, are observed within the continuous α-Al matrix among dispersed particles (Fig. 5 d) at a deformation of ε = 2% at 300°C. At the fracture strain ( ε ≈ 13.6%), clear interactions and segregation of dislocations are visible within the continuous α-Al matrix, similar to other alloys with dispersed second-phase particles 59 (Fig. 5 e). This suggests that in the continuous α-Al matrix without a cellular structure – or with a cellular structure at a larger cell size, as commonly found in LPBF alloys – the dislocation free path (DFP) increases, making the CLS mechanism less effective. Instead, the F–R mechanism becomes the predominant source of dislocations 60 . The critical stress (σₙ) required to activate an F–R source is inversely proportional to the DFP (or the dislocation segment length λ ), as indicated in Eq. 2 61 : $$\:{{\sigma\:}}_{F-R}=\frac{MGb}{\lambda\:}\frac{1}{2\pi\:(1-\nu\:)}[\left(1-\frac{3}{2}\nu\:\right)\text{ln}\left(\frac{\lambda\:}{b}\right)+\frac{\upsilon\:-2}{2}]$$ 2 The critical stresses for the CLS and F–R mechanisms are compared as a function of the cell size or dislocation segment length (Fig. 5 g). CLS requires significantly higher stress at ultrafine cell sizes below 200 nm, which explains the high flow stress in the ECN-confined α-Al phase of the AlLaScZr-ECN sample. In contrast, \(\:{{\sigma\:}}_{F-R}\) is significantly lower for larger DFPs, resulting in reduced flow stress for the α-Al phase at elevated temperatures in the AlLaScZr-NP sample. This behaviour is consistent with observations in conventional Al alloys and other LPBF alloys with coarser dispersed reinforcements or cellular structures 7 , 12 , 43 . The abovementioned results are validated by comparing the fracture modes (Figs. 5 c, f). Dislocation motion in the AlLaScZr-ECN sample is confined within α-Al cells. Load is transferred to the Al 11 La 3 cell walls from α-Al during their co-deformation. Micro-cracking initiates by fracturing the cell walls (Fig. 5 c), which requires higher stress. In contrast, dislocation segregation and interactions are more frequent within the α-Al matrix at the fracture strain, inducing local strain concentration and ductile failure characterised by micro-cracks distributed throughout the softer α-Al matrix (Fig. 5 f). Moreover, the advantages of the ECN in load transfer and dislocation confinement are confirmed by the minimal YS difference (< 10 MPa) between the ECN and NP samples under tensile testing at 400°C (Fig. 1 e, Fig. S8c). At this temperature, the ECN fully spheroidises into nanoparticles, similar to the NP sample, thereby losing its reinforcing benefits (Fig. S4b). Furthermore, densely distributed L1 2 -Al 3 (Sc,Zr) nanoprecipitates contribute to strengthening by acting as obstacles to dislocation motion. At RT, these nanoprecipitates significantly shorten the DFP within α-Al 6 , 62 , 63 . Similar YS values and dislocation segment lengths in the α-Al phase are observed between the ECN and NP samples at RT, as the nanoprecipitates are uniformly dispersed in both samples (Fig. 2 e, Fig. S6). However, at elevated temperatures (≥ 300°C), the obstacle effects weaken as cross-slip and climb motions enable dislocations to bypass the precipitates 64 , 65 . In this case, the ECN plays a crucial role in limiting the DFP, thereby maintaining a high yield stress in the α-Al phase. Without the ECN, the DFP increases within the continuous α-Al matrix (as observed in the longer segments in Fig. 5 d), resulting in a substantial drop in YS. Despite this, at elevated temperatures (300°C and 400°C), the highly dense intracellular L 12 -Al 3 (Sc, Zr) nanoprecipitates still provide synergistic high-temperature strengthening alongside the ECN through effective dislocation–precipitate interactions (inset in Fig. 5 a). This behaviour has been reported in Sc, Zr-modified Al alloys at elevated temperatures 5 . Micro-alloying of Sc, Zr further augments the YS by ~ 45 MPa at 300°C and ~ 40 MPa at 400°C, compared with LPBF Al-La alloy counterparts without Sc and Zr (Table S6 and Fig. S8). Consistent with the proposed strengthening mechanisms, the macroscopic tensile behaviour also indicates enhanced dislocation storage and reduced dislocation annihilation due to the ECN. This is further supported by quantitative analysis using the dislocation-based Kocks–Mecking (K–M) model 66 , 67 applied to the tensile stress–strain curves of AlLaScZr-ECN and AlLaScZr-NP alloys at both RT and 300°C (Fig. S7a, Supplementary Text). The results (Fig. S7, Table S5) suggest a higher dislocation storage rate ( \(\:{k}_{1}\) ) in the AlLaScZr-ECN sample than in the AlLaScZr-NP sample at RT, consistent with findings for LPBF AlSi10Mg 7 , 17 . Additionally, the high dislocation storage capacity is largely preserved at 300°C in the AlLaScZr-ECN sample, with \(\:{k}_{1}\) retaining 61% of its RT value, higher than that (17%) for the AlLaScZr-NP sample. Additionally, the ECN reduces the dynamic annihilation rate \(\:{k}_{2}\) at both RT and 300°C. This quantitative comparison using the K–M model further validates the role of the ECN in enhancing flow stress in α-Al at elevated temperatures, driven by its high dislocation storage capacity and ability to inhibit dislocation annihilation. In conclusion, the LPBF AlLaScZr alloy achieves significantly enhanced high-temperature YS above 0.6 T m of Al, while addressing the longstanding challenge of severe strength degradation in metals after prolonged thermal exposure, attributable to its thermally stable nanoscale ECN architecture and dense nanoprecipitate distribution. This alloy demonstrates substantial potential as an Al-based superalloy, extending the application range of lightweight Al alloys into the medium temperature range (300–400°C) currently dominated by ferrous, titanium, and nickel-based superalloys for structural applications. Moreover, this study highlights the distinct advantages of the 3D interconnected network architecture over the traditional uniform particle dispersion configuration, which has been the standard for high-temperature strengthening. Additionally, it provides critical insights into the high-temperature strengthening mechanisms associated with AM-induced cellular structures, particularly the role of ultrafine cell sizes in restricting dislocation motion through the CLS mechanism, enabling effective strengthening of metal matrix to counteract the intrinsic softening. Based on our findings, a novel strategy is proposed to achieve exceptional high-temperature strength by integrating thermally stable phases into a nanoscale 3D network architecture, leveraging low-diffusivity alloying elements and the rapid solidification process of AM. We propose that this strategy can be broadly applied to other alloy systems, enabling extraordinary high-temperature mechanical performance through advanced AM technologies. Methods LPBF process and heat treatment AlLaScZr powders were produced through gas atomisation from pre-alloyed ingots, resulting in spherical particles with diameters ranging from 15 to 53 µm (Fig. S1 a). LPBF was performed using a metal 3D printer (Truprint 1000, Germany) equipped with a 200 W fibre laser and a fixed spot size of 30 µm. The hatching distance and layer thickness were set as 100 µm and 30 µm, respectively, with a laser power of 190 W. The laser scanning direction was alternated by 90° between consecutive layers without preheating the base plate. The optimal scanning speed of 1500 mm/s was determined based on the relative density of fully dense bulk (theoretical density: 2.86 g/cm 3 ) and hardness of test cubes printed at various speeds (Fig. S1 b). The density was measured using the Archimedes method, and hardness tests were conducted using a Vickers hardness tester (EZ-mat CARAT 930, Germany) with a 10 kg load and 15 s dwell time. Results across nine points were averaged. The as-built samples were subjected to isothermal ageing at 325°C for 1 h, yielding peak-aged samples, referred to as AlLaScZr-ECN in the main text (Fig. S1 c). Additional samples were annealed for 168 h at 325°C to disrupt the ECN structure, producing AlLaScZr-NP samples (Fig. 2 c). Other heat treatment conditions are specified in the main text. Mechanical properties Dog-bone-shaped bulk samples were printed (inset in Fig. 1 b) and subsequently sliced to 1.6 mm thickness using electrical discharge machining. The samples were then polished into rectangular tensile samples. The sample gauge length and width were 15 mm and 2.8 mm, respectively. Tensile testing (at both RT and elevated temperatures) was conducted on the peak-aged samples using a tensile testing machine (Zwick/Roell Z005, Germany) with a laser extensometer, following ASTM E8 and E21 standards. The tensile direction was oriented perpendicular to the BD, and the strain rate was maintained at 1 × 10 − 3 s − 1 for all tests. For elevated temperature tests, samples were heated at 15°C/min and held for 30 min at the target temperature to achieve thermal equilibration prior to loading. Microstructure characterisation The chemical composition of the AlLaScZr powders and as-printed samples was determined using inductively coupled plasma–atomic emission spectrometry (ICP-AES, iCAP6300, Thermo Fisher Scientific, USA), as detailed in Table S1 . XRD was performed on as-printed and peak-aged samples using a D8 ADVANCE Da Vinci X-ray diffractometer (Bruker Corporation, Germany) with Cu Kα radiation (λ = 0.1542 nm), scanning from 10° to 90° at 2°/min. Microstructural analysis was performed using a Zeiss AxioVision optical microscope (Germany) and a MAIA3 field emission scanning electron microscope (SEM) (TESCAN, Czech Republic) equipped with EDS and electron backscatter diffraction (EBSD) detectors. The SEM instrument operated at 10 kV, and samples were electrolytically polished in a nitric acid–methanol solution (15 V, 10 s) for EBSD imaging. The EBSD scans used a 0.78 µm step size at 20 kV and were processed with HKL Channel 5 software. Samples for transmission electron microscopy (TEM) were prepared by mechanical grinding, followed by electro-polishing in a nitric acid–methanol solution (1:3 by volume) at -30°C and 15 V. TEM and corresponding EDS mapping were conducted using Talos 200X and Titan Themis microscopes (Thermo Fisher Scientific, USA) with a HAADF detector and SuperX EDS with four silicon drift detectors. ImageJ was used to measure the cellular structure dimensions. APT samples were prepared using a focused ion beam (FIB) and analysed using a LEAP-5000XR microscope (Colorado MicroDissect, Inc., USA), with data processed in Cameca IVAS (France). Additionally, 3D X-ray micro-CT was performed with an Xradia 510 Versa instrument (Zeiss, Germany) using a voxel size of 1.5 µm to image large-volume specimens (1.35 mm × 1.38 mm × 0.85 mm). Pore size distribution analysis was conducted using Dragonfly software (ORS, Canada) (Fig. S1 e, f), and 3D reconstruction of the ECN was achieved from 2D BSE images sliced with FIB (inset in Fig. 1 c), with an interlayer distance of 7.5 nm. Thermodynamic calculations The solidification behaviour and phase evolution of the AlLaScZr alloy were simulated using the CALPHAD method in Pandat™ software (CompuTherm LLC, USA), based on the PanAl75 commercial database. To simulate the rapid solidification during LPBF, Scheil mode simulations were conducted from 800°C to 600°C with a step of 1°C, based on the composition of the LPBF AlLaScZr alloy. In situ SXRD test In situ SXRD tensile testing was conducted at beamline ID11 of the European Synchrotron Radiation Facility. The beam energy was 65.3508 keV, and the wavelength was 0.01897 nm. The incident beam size was 0.2 mm × 0.2 mm. For elevated temperature tests, the samples were heated to the target temperature and held for 30 min prior to tensile loading. A Frelon36 detector (2048×2048 pixels and pixel size of 47 µm × 47 µm) was used to acquire the diffraction patterns. The cross-section of the samples was approximately 0.5 mm × 0.5 mm. The specimen-to-detector distance was 179.60 mm, calibrated using a CeO 2 powder standard. The AlLaScZr-ECN and AlLaScZr-NP samples were tested using a Nanox tensile stress rig 68 at both RT and 300°C. The Nanox voltage for loading was maintained through a constant ramp rate of 0.095 V/s at RT and 1.42 V/s at 300°C, corresponding to strain rates of 1.2 × 10 − 4 s − 1 and 1.4 × 10 − 3 s − 1 , respectively. At RT, a lower strain rate was used than that in the macro-tensile tests (1 × 10 − 3 s − 1 ) to capture sufficient data points, as the material’s mechanical properties were relatively insensitive to strain rate. At 300°C, the strain rate matched that of the macro-tensile test. Background noise from the loading equipment was subtracted from the 2D diffraction images before calibration and integration. Detailed calculations for lattice strain and stress partitioning of constitutive phases are demonstrated in Supplementary Materials. First-principles calculations based on DFT DFT calculations were performed using the Vienna Ab initio Simulation Package (VASP) with projector augmented wave pseudopotentials and a plane-wave basis set 69 . The Perdew–Burke–Ernzerhof exchange–correlation functional within the generalised gradient approximation was applied 70 – 72 . A plane-wave cut-off of 450 eV was set with an energy convergence parameter of 10 − 5 eV/atom for electronic self-consistency. The maximum forces on each relaxed atom converged to 0.02 eV/Å during structural relaxation through conjugate gradient minimisation. Gamma-centred 22 × 22 × 22 k-point meshes were used for the bulk unit cell of Al and Al 3 Sc, while 21 × 6 × 6 k-point meshes were used for Al 11 La 3 , with densities lower than 0.03 Å −1 . A vacuum layer with a thickness of 20 Å was added to the surface supercell to prevent unwanted interactions between the slab and its periodic arrangement. During the geometry optimisation, both the volume and atomic positions were relaxed for the initial interface systems, and only the atomic positions of the alloyed interface systems were relaxed. The lattice parameter for α-Al was a = 4.04 Å, while those for Al 11 La 3 were a = 4.44 Å, b = 10.14 Å, and c = 13.14 Å 73 . Detailed calculations are demonstrated in Supplementary Materials. Declarations Data availability All data supporting the findings of this study are available in the main text or the Supplementary Materials. Acknowledgements This work was supported by the National Natural Science Foundation of China 52304405, 52371034, and 52204393. The authors also thank the European Synchrotron Radiation Facility (ESRF) for providing synchrotron radiation facilities under proposal number ma6301 (doi.org/10.15151/ESRF-ES-1830139531). Professional English language editing support provided by AsiaEdit (asiaedit.com). Author contributions Conceptualisation: S.M., Z.C., M.W., and J. L. Methodology: S.M., H.F., G.J., Y.C., Y.L., Y.W., and Y.Z. Investigation: S.M., Z.C., Y.C., and H.C. Visualisation: S.M., H.F., G.J., Y.L, and S.Z. Supervision: Z.C., M.W., and S.N. Writing—original draft: S.M. Writing—review & editing: Z.C., G.J., M.W. and J. L. Competing interests The authors declare that they have no competing interests. References Michi RA, Plotkowski A, Shyam A, Dehoff RR, Babu SS (2022) Towards high-temperature applications of aluminium alloys enabled by additive manufacturing. 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Supplementary Files SupplementaryMaterialsforncommsubmit.docx Supplementary Materials for Exceptional High-temperature Strength in an Additively Manufactured Al-based Superalloy with Stable Nanoscale Eutectic Cellular Network Cite Share Download PDF Status: Published Journal Publication published 22 Nov, 2025 Read the published version in Nature Communications → Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. Our growing team is made up of researchers and industry professionals working together to solve the most critical problems facing scientific publishing. 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\u003cem\u003eT\u003c/em\u003e–\u003cem\u003ef\u003c/em\u003e\u003csub\u003es\u003c/sub\u003e curves under Scheil conditions. The corresponding phase evolution is shown in the inset; \u003cstrong\u003eb\u003c/strong\u003e Typical electron backscatter diffraction (EBSD) inverse pole figure (IPF) parallel to the BD. The inset shows the dog-bone-shaped bulk samples produced using LPBF; \u003cstrong\u003ec\u003c/strong\u003e Backscattered electron (BSE) image revealing the subgrain ECN, with the inset showing the 3D reconstructed image; \u003cstrong\u003ed\u003c/strong\u003e Scanning transmission electron microscopy high-angle annular dark-field (STEM-HAADF) image of the ECN with corresponding energy-dispersive spectroscopy (EDS) elemental mappings; \u003cstrong\u003ee\u003c/strong\u003e Uniaxial tensile stress–strain curves of the AlLaScZr alloys at RT and elevated temperatures; \u003cstrong\u003ef\u003c/strong\u003e Plot comparing the tensile yield strength (YS) at room temperature (RT) with that at 300 °C for the LPBF AlLaScZr alloy and other commercial and LPBF aluminium alloys (Note: the YS of the commercial Al alloys at 300 °C is taken by linearly interpolation of the YS at 260 °C (500 °F) and 316 °C (600 °F)); \u003cstrong\u003eg\u003c/strong\u003e Specific YS as function of temperature for the LPBF AlLaScZr alloy, compared with several typical Al, Ti, stainless steels, and Ni-based alloys.\u003c/p\u003e","description":"","filename":"Fig1.jpg","url":"https://assets-eu.researchsquare.com/files/rs-5840317/v1/14be9d010c2feab28611cda9.jpg"},{"id":75146060,"identity":"72b8630f-8e6f-45e6-970b-fcd32017ccf2","added_by":"auto","created_at":"2025-01-31 06:53:27","extension":"jpg","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":1065122,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eMicrostructural thermal stability of the LPBF AlLaScZr alloy.\u003c/strong\u003e Morphology of the ECN \u003cstrong\u003ea\u003c/strong\u003e in the peak-aged state, with the inset showing the cell size distribution; \u003cstrong\u003eb\u003c/strong\u003e after annealing at 300 °C for 168 h; \u003cstrong\u003ec\u003c/strong\u003e after annealing at 325 °C for 168 h, with the inset showing the distribution of spheroidised particle size; \u003cstrong\u003ed\u003c/strong\u003e X-ray diffraction (XRD) patterns of AlLaScZr-ECN and AlLaScZr-NP; \u003cstrong\u003ee\u003c/strong\u003e Atomic-probe tomography (APT) analysis of selected regions in AlLaScZr-ECN and AlLaScZr-NP samples, with localised linear scanning identifying the composition of La-rich regions.\u003c/p\u003e","description":"","filename":"Fig2.jpg","url":"https://assets-eu.researchsquare.com/files/rs-5840317/v1/eb97bc8eae3e3a6e83193591.jpg"},{"id":75146039,"identity":"1647ca9e-d6d8-4ac5-a916-2e0f6609b0da","added_by":"auto","created_at":"2025-01-31 06:53:25","extension":"jpg","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":1506159,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eInterfacial behaviours of phases and density functional theory (DFT) calculations. a \u003c/strong\u003ehigh-resolution transmission electron microscopy\u003cstrong\u003e (\u003c/strong\u003eHRTEM) image showing the interface between the cell wall and α-Al matrix; \u003cstrong\u003eb\u003c/strong\u003e intracellular nanoprecipitations with corresponding EDS elemental mappings; FFT patterns of \u003cstrong\u003ec\u003c/strong\u003e the image in \u003cstrong\u003ea\u003c/strong\u003e and \u003cstrong\u003ed\u003c/strong\u003e boxed regions in \u003cstrong\u003eb\u003c/strong\u003e; \u003cstrong\u003ee\u003c/strong\u003e DFT-calculated atomic models of the interfaces between Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e/α-Al and Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e/L\u003csub\u003e12\u003c/sub\u003e-Al\u003csub\u003e3\u003c/sub\u003eSc;\u003c/p\u003e","description":"","filename":"Fig3.jpg","url":"https://assets-eu.researchsquare.com/files/rs-5840317/v1/0b57f76ad6d21ba6232f84be.jpg"},{"id":75146066,"identity":"721df9aa-ec6c-4d69-b653-4955258dffe1","added_by":"auto","created_at":"2025-01-31 06:53:27","extension":"jpg","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":817119,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eIn situ synchrotron X-ray diffraction (SXRD) tensile tests at RT and elevated temperatures.\u003c/strong\u003e Lattice strain evolution of different crystalline planes of Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e and the average lattice strain of α-Al in the AlLaScZr-ECN and AlLaScZr-NP alloys at \u003cstrong\u003ea\u003c/strong\u003e RT and \u003cstrong\u003eb\u003c/strong\u003e 300 °C; Evaluated phase stress partitioning of the AlLaScZr-ECN and AlLaScZr-NP alloys during in situ tensile testing, plotted against the engineering stress–strain curves at \u003cstrong\u003ec\u003c/strong\u003e RT and \u003cstrong\u003ed\u003c/strong\u003e 300 °C.\u003c/p\u003e","description":"","filename":"Fig4.jpg","url":"https://assets-eu.researchsquare.com/files/rs-5840317/v1/bce495e11c86305798bba6b4.jpg"},{"id":75146044,"identity":"56f832cb-5880-4b38-9548-eac1e607263b","added_by":"auto","created_at":"2025-01-31 06:53:26","extension":"jpg","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":3619413,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eDeformation behaviour of the LPBF AlLaScZr alloy\u003c/strong\u003e \u003cstrong\u003eat 300 °C.\u003c/strong\u003e \u003cstrong\u003ea\u003c/strong\u003eScanning transmission electron microscopy–bright-field (STEM-BF) image showing dislocations within α-Al cells of the AlLaScZr-ECN sample under tensile strain of 2 %. The inset shows a magnified STEM-HAADF image illustrating the intracellular dislocations; \u003cstrong\u003eb\u003c/strong\u003e STEM-BF image showing the dislocations in α-Al cells at the fracture strain; \u003cstrong\u003ec\u003c/strong\u003e Micro-cracks owing to broken ECN near the fracture surface; \u003cstrong\u003ed\u003c/strong\u003e STEM-BF image showing dislocations in the AlLaScZr-NP sample subjected to tensile strain of 2 %. The inset shows a magnified STEM-HAADF image indicating longer dislocation segments in the continuous Al matrix; \u003cstrong\u003ee\u003c/strong\u003e STEM-BF image showing the dislocation behaviour in the annealed sample deformed at 300 °C at the fracture strain; \u003cstrong\u003ef\u003c/strong\u003eMicro-cracks in continuous Al matrix near the fracture surface; \u003cstrong\u003eg\u003c/strong\u003eCritical stress for Orowan bowing at 300 °C during confined layer slip (CLS) and Frank–Read (F–R) source activation as a function of the dislocation segment length, accompanied by schematic illustration of dislocation behaviour in the presence of ECN or dispersed particles; \u003cstrong\u003eh\u003c/strong\u003e HRTEM image depicting a local highly strained region near a dislocation line within Al matrix between neighbouring cell walls. \u003cstrong\u003ei\u003c/strong\u003e Inverse fast Fourier transform (IFFT) pattern and geometrical phase analysis (GPA) mappings of \u003cstrong\u003eh\u003c/strong\u003e.\u003c/p\u003e","description":"","filename":"Fig5.jpg","url":"https://assets-eu.researchsquare.com/files/rs-5840317/v1/bbb53a0a1b51919e926802b1.jpg"},{"id":98927991,"identity":"2206136b-832b-4474-b5b5-ddf401540b6a","added_by":"auto","created_at":"2025-12-24 08:10:13","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":10939648,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-5840317/v1/108651c8-cebb-4e4d-be5f-ce39e212219b.pdf"},{"id":75146029,"identity":"1acd3f46-44cb-45cc-89bb-b86ad1933799","added_by":"auto","created_at":"2025-01-31 06:53:25","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":6393244,"visible":true,"origin":"","legend":"Supplementary Materials for Exceptional High-temperature Strength in an Additively Manufactured Al-based Superalloy with Stable Nanoscale Eutectic Cellular Network","description":"","filename":"SupplementaryMaterialsforncommsubmit.docx","url":"https://assets-eu.researchsquare.com/files/rs-5840317/v1/f1b4fea62badf8d214dd9219.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"Exceptional High-temperature Strength in an Additively Manufactured Al-based Superalloy with Stable Nanoscale Eutectic Cellular Network","fulltext":[{"header":"Introduction","content":"\u003cp\u003eSoftening owing to severe strength degradation is inevitable in metallic materials at elevated temperatures, especially above 0.5\u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e. For instance, at 300\u0026deg;C, most Al alloys exhibit a low yield strength of only a few tens of MPa, typically 10\u0026ndash;40% of that at ambient temperature \u003csup\u003e\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e,\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e\u003c/sup\u003e. Strategies to enhance high-temperature strength have traditionally focused on incorporating second-phase reinforcements, such as micro- and nanosized intermetallics, precipitates, or ceramic particles, to promote well-known load transfer and Orowan strengthening mechanisms \u003csup\u003e\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e. In this context, efforts have been primarily aimed at improving the coarsening resistance and increasing the volume fraction of reinforcements. For example, a Sc-modified Al-Cu-Mg-Ag alloy contains high-density and thermally stable coherent nanoprecipitates \u003cem\u003eV\u003c/em\u003e-(Al,Cu,Sc) \u003csup\u003e\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e\u003c/sup\u003e, exhibiting a superior yield strength of ~\u0026thinsp;100 MPa at 400\u0026deg;C. In addition, a multiple precipitate/matrix interface engineering strategy was reported to stabilise the θ\u0026prime;-Al\u003csub\u003e2\u003c/sub\u003eCu precipitates in an Al-Cu-Mg-Ag-Si-Sc alloy, retaining 97% yield strength after thermal exposure at 200\u0026deg;C \u003csup\u003e\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e \u003cp\u003eIn conventionally fabricated alloys, the abovementioned strategies rely on the uniform distribution of hard particles within a softer matrix. However, this architecture may limit the high-temperature strengthening effects. At elevated temperatures, thermally activated dislocation motions such as cross-slipping and climbing become dominant, allowing dislocations to bypass obstacles more easily and weakening the Orowan strengthening effect \u003csup\u003e\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e,\u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e\u003c/sup\u003e. Additionally, dispersed particles exert limited geometrical constraints on dislocation movement in the continuous alloy matrix, as thermally activated dislocation motion can easily accommodate the strain incompatibility between the alloy matrix and particles \u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u003c/sup\u003e, reducing the load transfer effectiveness.\u003c/p\u003e \u003cp\u003eRecently, additive manufacturing (AM) has demonstrated notable industrial application prospects in the rapid prototyping of complex metallic components with unprecedented freedom \u003csup\u003e\u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e,\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e\u003c/sup\u003e, as well as the realisation of novel metastable and ultrafine microstructures through rapid solidification. Specifically, a notable 3D ultrafine cellular-like architecture, where one phase forms cells enclosed by a 3D interconnected network of another phase, has been reported in AM Al \u003csup\u003e\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e\u003c/sup\u003e, Fe \u003csup\u003e\u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e\u003c/sup\u003e, Cu \u003csup\u003e\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e\u003c/sup\u003e, high-entropy alloys (HEAs) \u003csup\u003e\u003cspan citationid=\"CR13\" class=\"CitationRef\"\u003e13\u003c/span\u003e\u003c/sup\u003e, and other alloys. For example, recent studies on AM AlSi10Mg have demonstrated the intrinsic exceptional room-temperature (RT) mechanical properties, attributable to a nanoscale eutectic Si cellular network \u003csup\u003e\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e,\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e\u003c/sup\u003e. These materials achieve yield and fatigue strengths of ~\u0026thinsp;470 MPa in defect-free micro-sized samples, surpassing conventionally fabricated counterparts \u003csup\u003e\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e\u003c/sup\u003e. During co-deformation, the ultrafine cellular network exhibits superior load transfer \u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e,\u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e,\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e\u003c/sup\u003e, with stress in the Si phase at cell boundaries measured at 1.5\u0026ndash;2 GPa \u003csup\u003e\u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e\u003c/sup\u003e, 4\u0026ndash;5 times that in conventional Si alloys \u003csup\u003e\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e,\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e\u003c/sup\u003e. Additionally, this network constrains dislocation motion within cells as \u0026lsquo;dislocation cages\u0026rsquo;, leading to strong work hardening and delayed damage under tensile and fatigue loads \u003csup\u003e\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e,\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e\u003c/sup\u003e. However, these benefits have been limited to RT applications owing to the low thermal stability of the Si cellular network \u003csup\u003e\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e\u003c/sup\u003e, which undergoes spheroidisation into dispersed Si particles after a short holding time at 300\u0026deg;C. The high-temperature strength is therefore significantly degraded \u003csup\u003e\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e\u003c/sup\u003e. Forming a thermally stable 3D cellular network architecture may address this limitation, paving the way for enhanced high-temperature strength.\u003c/p\u003e \u003cp\u003eThis paper introduces a novel architecture design strategy for developing high-temperature \u0026lsquo;Al-based superalloys\u0026rsquo; using AM to achieve superior high-temperature strength and thermal stability. A novel AM AlLaScZr alloy was fabricated using laser powder bed fusion (LPBF), in which La undergoes eutectic alloying to form a thermally stable nanoscale Al-La eutectic cellular network (ECN). La was selected owing to its extremely low diffusivity and solubility in face-centred cubic (fcc)-Al among all alloying elements. Additional micro-alloying with Sc and Zr was performed to form coarsening-resistant L\u003csub\u003e12\u003c/sub\u003e coherent nanoprecipitates, providing synergistic strengthening. The LPBF AlLaScZr alloy exhibited extraordinary high-temperature yield strength (YS), ranking among the highest reported for Al alloys above 0.6 \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e (~\u0026thinsp;250 MPa at 300\u0026deg;C and ~\u0026thinsp;110 MPa at 400\u0026deg;C). The superior high-temperature strength was attributable to the high load-bearing capacity of the Al-La ECN, combined with the high stress in the soft α-Al phase owing to the restricted dislocation free path within ultrafine-sized cells (~\u0026thinsp;200 nm). Synergic strengthening was contributed by the intracellular nanoprecipitates as additional barriers of dislocation motion. Even after prolonged annealing and high-temperature testing, the alloy exhibited superior YS retention (~\u0026thinsp;200 MPa at 300\u0026deg;C and ~\u0026thinsp;100 MPa at 400\u0026deg;C) owing to the highly dense and well dispersed La-rich nanoparticles inherited from the ECN. The proposed strategy can inspire the design of other thermally stable nanostructures using state-of-the-art AM technologies, facilitating the development of alloys with enhanced high-temperature mechanical properties.\u003c/p\u003e"},{"header":"Results","content":"\u003cdiv id=\"Sec3\" class=\"Section2\"\u003e \u003ch2\u003eSuperior high-temperature strength\u003c/h2\u003e \u003cp\u003eWe prepared an Al-9.8La-0.46Sc-0.26Zr (wt.%) alloy for LPBF using a microstructure refinement (heterogeneous nucleation on the primary nuclei) and eutectic solidification strategy (near-eutectic composition) to achieve the optimal printability based on the CALPHAD method \u003csup\u003e\u003cspan citationid=\"CR22\" class=\"CitationRef\"\u003e22\u003c/span\u003e\u003c/sup\u003e. The temperature\u0026ndash;solidification fraction (\u003cem\u003eT\u003c/em\u003e\u0026ndash;\u003cem\u003ef\u003c/em\u003e\u003csub\u003es\u003c/sub\u003e) curve for this alloy exhibits an \u0026lsquo;L\u0026rsquo;-shaped profile (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea), indicating that solidification initiates with the precipitation of primary L\u003csub\u003e12\u003c/sub\u003e-Al\u003csub\u003e3\u003c/sub\u003e(Sc, Zr) dispersoids over a broad temperature range (740\u0026deg;C \u0026minus;\u0026thinsp;640\u0026deg;C), represented by a nearly vertical line. The horizontal segment of the curve corresponds to eutectic solidification of fcc-Al and Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e, extending from \u003cem\u003ef\u003c/em\u003e\u003csub\u003es\u003c/sub\u003e \u0026gt; 0.2 to the terminal stage at ~\u0026thinsp;638\u0026deg;C with a near-zero solidification temperature range. Crack-free and nearly fully dense bulk samples (~\u0026thinsp;99.7% relative density) were printed (inset of Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb) using optimised LPBF parameters (190 W power, 1500 mm/s scanning speed) (Fig. \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003eb). The pore sizes are mostly under 40 \u0026micro;m, as observed in 3D micro-computed tomography (CT) scans (Figs. S1e, f). The as-built samples were subjected to post-heat treatment at 325\u0026deg;C for 1 h to achieve peak hardening through L\u003csub\u003e12\u003c/sub\u003e-Al\u003csub\u003e3\u003c/sub\u003e(Sc, Zr) precipitation (Fig. \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003ec). The LPBF AlLaScZr alloy exhibits a bimodal grain structure, with large columnar grains (10\u0026ndash;50 \u0026micro;m) and small equiaxed grains (\u0026lt;\u0026thinsp;10 \u0026micro;m) without preferential texture (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb, Fig. S2). The refined equiaxed grains indicate the inoculation effect of the primary L\u003csub\u003e12\u003c/sub\u003e-Al\u003csub\u003e3\u003c/sub\u003e(Sc, Zr) phase \u003csup\u003e\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e\u003c/sup\u003e. At the subgrain scale, a ubiquitous nanoscale ECN structure is observed (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ec), with La-rich eutectic phases forming bright-contrast cell walls around α-Al cells, as confirmed by the scanning transmission electron microscopy (STEM) high-angle annular dark-field (HAADF) image and corresponding energy-dispersive spectroscopy (EDS) maps (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ed). The cellular structure exhibits slight heterogeneity, with coarser cellular dendrites at melt pool boundaries and finer equiaxed cells at melt pool interiors owing to cooling rate variations \u003csup\u003e\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e\u003c/sup\u003e (Fig. S3). A 3D view of the highly interconnected ECN structure is shown in the inset in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ec, reconstructed from a series of 2D backscattered electron (BSE) images. The ECN-enclosed cells are slightly elongated along the building direction (BD) and feature an ultrafine cell size (203.3\u0026thinsp;\u0026plusmn;\u0026thinsp;69.2 nm), finer than those reported in LPBF Al-Si alloys (0.5\u0026ndash;1 \u0026micro;m) \u003csup\u003e\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e,\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e\u003c/sup\u003e and other LPBF materials such as Cu \u003csup\u003e\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e\u003c/sup\u003e, Ti \u003csup\u003e\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e,\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e\u003c/sup\u003e, and steels \u003csup\u003e\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e,\u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e\u003c/sup\u003e (0.4\u0026ndash;1 \u0026micro;m on average). The LPBF-processed and peak-aged AlLaScZr alloy with the nanoscale ECN structure is referred to as the AlLaScZr-ECN alloy hereafter.\u003c/p\u003e \u003cp\u003eThe AlLaScZr-ECN alloy demonstrates outstanding mechanical performance at both RT and elevated temperatures, achieving exceptionally high YS at and above 300\u0026deg;C (~\u0026thinsp;0.6 \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e of Al) (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ee). At RT, the AlLaScZr-ECN alloy exhibits a YS of 418 MPa, an ultimate tensile strength (UTS) of 453 MPa, and an elongation of ~\u0026thinsp;7%, comparable to those of the high-strength LPBF Scalmalloy\u0026reg;. Notably, at 300\u0026deg;C, the alloy reaches a remarkable YS of 246 MPa and UTS of 265 MPa, surpassing conventional Al alloys by fivefold and conventional heat-resistant Al alloys by twofold. Furthermore, it significantly outperforms commercial LPBF Al alloys, such as AlSi10Mg and AlMgScZr, by 2\u0026ndash;4 times, as well as recently developed heat-resistant near-eutectic LPBF Al alloys, including Al-Ni, Al-Ce, and Al-Fe systems (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ef). The outstanding high-temperature YS of the AlLaScZr-ECN alloy highlights its potential for lightweight structural applications at a medium temperature range (300\u0026ndash;400\u0026deg;C), currently dominated by ferrous, titanium, and nickel-based alloys \u003csup\u003e\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e. The specific YS of the AlLaScZr-ECN alloy at 300\u0026deg;C is 86 MPa/(kg\u0026middot;m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e), nearly double that of commercial heat-resistant 2618 Al alloy and superior to several commercially pure titanium (CP-Ti) alloys, stainless steels, and Ni-based alloys (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eg).\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003e \u003cb\u003eExcellent thermal stability above 0.6\u003c/b\u003e \u003cb\u003eT\u003c/b\u003e \u003csub\u003e \u003cb\u003em\u003c/b\u003e \u003c/sub\u003e \u003c/p\u003e \u003cp\u003eA common challenge in AM alloys is the spheroidisation and coarsening of fine cellular structures and precipitates at elevated temperatures, attributable to the thermodynamically large driving force to reduce the interface energy, as observed in LPBF Al-Si alloys \u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e,\u003cspan citationid=\"CR44\" class=\"CitationRef\"\u003e44\u003c/span\u003e\u003c/sup\u003e. In contrast, the LPBF AlLaScZr-ECN alloy demonstrates excellent microstructural thermal stability and high strength retention even after prolonged thermal exposure at and above 300\u0026deg;C, despite its large cell boundary (CB) area and dense nanoprecipitates (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e). The thermal stability is primarily attributable to the La-rich cell walls, which provide both thermodynamic and kinetic resistance to coarsening owing to La\u0026rsquo;s extremely low diffusivity and solubility in α-Al. At 300\u0026deg;C, La exhibits a diffusivity of 6.9 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;21\u003c/sup\u003e m\u003csup\u003e2\u003c/sup\u003e\u0026middot;s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003eand a maximum solid solubility of 0.01 at.%, both significantly lower than those of Si (2.6 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;16\u003c/sup\u003e m\u003csup\u003e2\u003c/sup\u003e\u0026middot;s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e and 1.6 at.%, respectively) \u003csup\u003e\u003cspan citationid=\"CR45\" class=\"CitationRef\"\u003e45\u003c/span\u003e,\u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e\u003c/sup\u003e. In addition, a high density (7.3 \u0026times; 10\u003csup\u003e23\u003c/sup\u003e m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e) of coarsening-resistant L\u003csub\u003e12\u003c/sub\u003e-Al\u003csub\u003e3\u003c/sub\u003eSc nanoprecipitates are dispersed ubiquitously within α-Al cells, with an average size of 3.1 nm, as revealed by atomic-probe tomography (APT) (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ee) and high-resolution transmission electron microscopy (HRTEM) with the fast Fourier transformation (FFT) pattern (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb, d). These nanoprecipitates form owing to Sc solute supersaturation under rapid solidification during the LPBF \u003csup\u003e\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e\u003c/sup\u003e. Additionally, a small amount of Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e nanoparticles (5\u0026ndash;10 nm in size) within α-Al cells (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb), likely precipitated owing to La supersaturation due to LPBF, can be observed \u003csup\u003e\u003cspan citationid=\"CR47\" class=\"CitationRef\"\u003e47\u003c/span\u003e\u003c/sup\u003e. This unique combination of thermodynamically and kinetically stable phases endows the alloy with superior thermal stability and mechanical performance at elevated temperatures.\u003c/p\u003e \u003cp\u003eMoreover, the CBs exhibit high coherency and low interfacial energy, rendering them more stable than the high-angle grain boundaries (HAGBs) commonly observed in ultrafine-grained or nano-grained (UFG/NG) metals. The La-rich cell wall of the ECN is identified as the Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e phase (orthorhombic, group: Immm) by X-ray diffraction (XRD) analysis (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea), HRTEM analysis with the FFT pattern (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea, c), and APT stoichiometric analysis (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ee). The Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e cell walls (and intracellular Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e nanoparticles) present a specific orientation relationship (OR) with α-Al along the Al[001] zone axis: Al[001]//Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e[010], Al(200)//Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e(200) (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea). The interface of Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e/α-Al presents a remarkably low interfacial misfit of 0.5% (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ee) and an interfacial energy (0.45 J\u0026middot;m\u003csup\u003e\u0026minus;\u0026thinsp;2\u003c/sup\u003e) lower than that (0.8\u0026ndash;2.0 J\u0026middot;m\u003csup\u003e\u0026minus;\u0026thinsp;2\u003c/sup\u003e) of conventional high-energy HAGBs in UFG/NG alloys \u003csup\u003e\u003cspan citationid=\"CR48\" class=\"CitationRef\"\u003e48\u003c/span\u003e\u003c/sup\u003e, as supported by density functional theory (DFT) calculations (Table S3, Supplementary Text). Additionally, localised segregation of Sc on the Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e CBs (circled in Figs.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea, \u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb, and \u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ee) further stabilises the CBs by accommodating the misfit of α-Al/Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e interfaces. HRTEM images reveal that Sc segregation results in the formation of Al\u003csub\u003e3\u003c/sub\u003eSc at the Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e cell wall surface with an OR of Al\u003csub\u003e3\u003c/sub\u003eSc[100]//Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e[010], Al\u003csub\u003e3\u003c/sub\u003eSc(200)//Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e(200) (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea), similar to that between α-Al and Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e. Similar Sc segregation at phase boundaries has been observed in other eutectic Al alloys, such as Al-Ce and Al-Fe-Ni \u003csup\u003e\u003cspan citationid=\"CR49\" class=\"CitationRef\"\u003e49\u003c/span\u003e,\u003cspan citationid=\"CR50\" class=\"CitationRef\"\u003e50\u003c/span\u003e\u003c/sup\u003e, where the larger atomic size of Sc alleviates local lattice distortion. DFT calculations indicate that the lattice misfit between Al\u003csub\u003e3\u003c/sub\u003eSc and Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e is -1.2%, while that between α-Al and Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e is +\u0026thinsp;0.5%. These outcomes suggest that Al\u003csub\u003e3\u003c/sub\u003eSc segregation reduces the interface energy by accommodating the positive lattice misfit at the α-Al/Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e interface (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ee, Table S2). Consequently, even after prolonged annealing at 300\u0026deg;C for 168 h, the ECN demonstrates remarkable resistance to coarsening, retaining its connectivity and stability (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003eb, Fig. S4a).\u003c/p\u003e \u003cp\u003eThe interconnectivity of the thermally stable ECN is disrupted upon annealing at a higher temperature of 325\u0026deg;C for an extended period of 168 h (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ec, Fig. S4a). This process results in the complete spheroidisation of Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e cell walls into Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e nanoparticles (55.2\u0026thinsp;\u0026plusmn;\u0026thinsp;18.2 nm), forming a structure referred to as AlLaScZr-NP, which is compared with the original AlLaScZr-ECN (Figs.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ea, c). XRD patterns and atom stoichiometric ratios of the La-rich zone derived from APT mapping confirm that the phase transformation of Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e (Figs.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ed, e) does not occur during the annealing process. Sc segregation along Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e persists. Zr segregates within Sc-rich regions (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ee), indicating the formation of Al\u003csub\u003e3\u003c/sub\u003e(Sc, Zr) with a Zr-rich shell, which retards Al\u003csub\u003e3\u003c/sub\u003eSc coarsening \u003csup\u003e\u003cspan citationid=\"CR51\" class=\"CitationRef\"\u003e51\u003c/span\u003e\u003c/sup\u003e. The intracellular Al\u003csub\u003e3\u003c/sub\u003eSc particle size increases slightly from 3.1 to 4.6 nm, demonstrating strong resistance to coarsening, consistent with previous reports on Al alloys micro-alloyed with Sc (Zr) \u003csup\u003e\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e,\u003cspan citationid=\"CR52\" class=\"CitationRef\"\u003e52\u003c/span\u003e\u003c/sup\u003e. Additionally, grain coarsening is minimal after prolonged thermal exposure, owing to the abundant Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e nanoparticles pinning the grain boundaries (Fig. S2).\u003c/p\u003e \u003cp\u003eIn terms of mechanical properties after the spheroidisation annealing, the YS of the AlLaScZr-NP alloy decreases at both RT and 300\u0026deg;C (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ee). Nevertheless, the alloy retains a high YS of ~\u0026thinsp;200 MPa at 300\u0026deg;C. Moreover, at an elevated tensile testing temperature of 400\u0026deg;C, the YS difference between the AlLaScZr-ECN and AlLaScZr-NP alloys becomes minimal, as the ECN structure fully spheroidises into Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e nanoparticles during holding in the tensile test (Fig. S4b). Despite this structural change, both alloys maintain an exceptional YS of over 100 MPa at 400\u0026deg;C, which is at least double that of conventional heat-resistant 2618 Al alloys (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eg). This result suggests that even after the transformation of Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e cell walls into nanoparticles, their high density, fine particle size (~\u0026thinsp;55 nm), and uniform dispersion contribute to the retention of superior YS at elevated temperatures, similar to other alloys with high-density reinforcements \u003csup\u003e\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e,\u003cspan citationid=\"CR53\" class=\"CitationRef\"\u003e53\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e\n\u003ch3\u003eECN-induced high-temperature strengthening mechanisms\u003c/h3\u003e\n\u003cp\u003eTo investigate the strengthening mechanisms of the LPBF AlLaScZr alloy, particularly at elevated temperatures, we conducted in situ synchrotron X-ray diffraction (SXRD) tensile tests at RT and 300\u0026deg;C on both AlLaScZr-ECN and AlLaScZr-NP samples, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e (see Fig. S5, Methods, and Supplementary Text for details). First, the reinforced Al₁₁La₃ phase in the nanoscale ECN configuration exhibits superior load transfer ability at both RT and elevated temperatures compared with the dispersed Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e nanoparticles. Larger lattice strains are observed in different crystalline planes of Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e in the ECN configuration (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea, b), increasing by ~\u0026thinsp;0.002 at RT and by ~\u0026thinsp;0.001 at 300\u0026deg;C at the yield point. This is followed by plastic deformation in the (101), (130), and (132) planes, resulting in a 100\u0026ndash;200 MPa increase in phase stress, which represents a 20\u0026ndash;25% improvement compared to the dispersed particle configuration. (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ec, d). Additionally, during plastic deformation, Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e in the ECN configuration shows prolonged load-bearing and delayed load shedding compared with the dispersed nanoparticles, resulting in enhanced work hardening and UTS. This effect is more pronounced at RT than at 300\u0026deg;C.\u003c/p\u003e \u003cp\u003eSecond, the α-Al phase in the ECN exhibits extraordinarily high phase stress even at 300\u0026deg;C. At RT, the α-Al phase stress in the AlLaScZr-ECN alloy is moderately higher than that in the AlLaScZr-NP alloy, with YS rising from 315 MPa to 360 MPa (only a\u0026thinsp;~\u0026thinsp;15% increase). In contrast, at 300\u0026deg;C, the α-Al phase in the ECN configuration displays a marked improvement, with an average lattice strain of 0.003 at yield point and a YS of ~\u0026thinsp;180 MPa, approximately 60 MPa higher than that in the AlLaScZr-NP sample (a\u0026thinsp;~\u0026thinsp;50% increase) and 3\u0026ndash;7 times that of conventional Al alloys at 300\u0026deg;C (25\u0026ndash;50 MPa) \u003csup\u003e\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e,\u003cspan citationid=\"CR43\" class=\"CitationRef\"\u003e43\u003c/span\u003e\u003c/sup\u003e. These results demonstrate that the Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e reinforcing phase provides superior load-bearing ability in the ECN configuration compared with dispersed nanoparticles at both RT and elevated temperatures. The ECN configuration also synergistically promotes high flow stress in the α-Al phase, contributing to the alloy\u0026rsquo;s exceptional high-temperature strength. This superior performance of cellular structures in AM alloys has been well-documented at RT, as their strong cell walls and refined cell size offer substantial load-bearing advantages over conventional alloys \u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e,\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e,\u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e\u003c/sup\u003e. For example, in LPBF AlSi10Mg alloys, nanoscale Si cellular networks bear 4\u0026ndash;5 times the stress and are more fracture-resistant than the coarser Si networks and particles found in cast Al-Si alloys\u003csup\u003e\u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e\u003c/sup\u003e. Moreover, Li et al. reported that a continuous Si network in LPBF AlSi10Mg enhances YS and work hardening compared with the annealed state with discrete Si particles \u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u003c/sup\u003e. In this study, direct comparisons of lattice strain and stress partitioning demonstrate the superior load-bearing capability of the Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e phase in the nanoscale ECN configuration not only at RT but also at elevated temperatures, provided thermal stability is maintained.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFrom the perspective of dislocation behaviour, prior studies have shown that cellular structures in AM alloys facilitate co-deformation of the matrix and cell walls, primarily through dislocation accumulation at cell boundaries. For instance, in LPBF AlSi10Mg alloys, Si cell wall constraints lead to high-density geometrically necessary dislocations at Al/Si interfaces to manage misfit strain during deformation \u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u003c/sup\u003e. Similarly, in LPBF 316L stainless steel, dislocations have been observed to pin and tangle at cell walls \u003csup\u003e\u003cspan citationid=\"CR54\" class=\"CitationRef\"\u003e54\u003c/span\u003e\u003c/sup\u003e. Kwon et al. recently proposed a dislocation-based constitutive model for the AM-induced cellular structure in a LPBF Cu-Sn model alloy, describing that dislocations originate from the intercellular Frank\u0026ndash;Read (F\u0026ndash;R) sources and deposit from the cell interior onto the walls \u003csup\u003e\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e\u003c/sup\u003e. In contrast, this study reveals different dislocation behaviours during the co-deformation of α-Al and Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e cell walls at RT and 300\u0026deg;C. Dislocations are confined within individual cells and remain as single dislocation segments between cell walls even at a high strain level (\u003cem\u003eε\u003c/em\u003e\u0026thinsp;\u0026asymp;\u0026thinsp;7%) during tensile tests at RT (Fig. S6a) and 300\u0026deg;C (Figs.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ea, b). Unlike in previous reports \u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e,\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e\u003c/sup\u003e, no substantial dislocation segregation or entanglement is observed at the cell boundaries. In terms of the microstructural differences from other reported AM-induced cellular structures, the ultrafine (~\u0026thinsp;200 nm) cell size of the Al-La ECN is smaller than the typical cell sizes (0.5\u0026ndash;1 \u0026micro;m) found in LPBF Al-Si, steel, and Cu alloys. This refined cell size likely activates the confined layer slip (CLS) mechanism, as proposed for nanolamellar architectures \u003csup\u003e\u003cspan citationid=\"CR55\" class=\"CitationRef\"\u003e55\u003c/span\u003e\u003c/sup\u003e. According to the CLS theory, ductile layers with dimensions below 200 nm deform through single dislocation segments forming Orowan-type loops between two parallel interfaces, rather than through dislocation arrays at a larger scale \u003csup\u003e\u003cspan citationid=\"CR56\" class=\"CitationRef\"\u003e56\u003c/span\u003e,\u003cspan citationid=\"CR57\" class=\"CitationRef\"\u003e57\u003c/span\u003e\u003c/sup\u003e. Given the size distribution of the Al-La ECN-enclosed cells (~\u0026thinsp;200 nm on average), the CLS mechanism likely dominates (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003eg), with dislocations gliding as Orowan-type loops confined by the Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e cell walls. This behaviour is consistent with the isolated dislocation segments observed in the cells. The critical stress for Orowan bowing of dislocations confined between two interfaces is calculated using Eq.\u0026nbsp;1 \u003csup\u003e58\u003c/sup\u003e:\u003cdiv id=\"Equ1\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equ1\" name=\"EquationSource\"\u003e\n$$\\:{{\\sigma\\:}}_{CLS}=\\frac{MGb}{8\\pi\\:h}\\left(\\frac{4-\\nu\\:}{1-\\nu\\:}\\right)\\text{ln}\\left(\\frac{\\alpha\\:h}{b}\\right)$$\u003c/div\u003e\u003cdiv class=\"EquationNumber\"\u003e1\u003c/div\u003e\u003c/div\u003e\u003c/p\u003e \u003cp\u003eAt 300\u0026deg;C, although the shear modulus of Al decreases from 26.2 GPa (at RT) to 22.5 GPa \u003csup\u003e\u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e\u003c/sup\u003e, \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:{{\\sigma\\:}}_{CLS}\\)\u003c/span\u003e\u003c/span\u003e can still reach values of 100\u0026ndash;180 MPa within the cell size range of 100\u0026ndash;200 nm, corresponding to the layer thickness \u003cem\u003eh\u003c/em\u003e (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003eg, Table S4). This supports the high stress observed in the α-Al phase enclosed by the ECN during SXRD analysis. Specifically, a short dislocation segment constrained by neighbouring cell walls (inset in Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003eh) results in severe lattice distortion in α-Al and localised high strain, as indicated by the inverse fast Fourier transform (IFFT) pattern and geometrical phase analysis (GPA) (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ei).\u003c/p\u003e \u003cp\u003eIn comparison, in the AlLaScZr-NP sample, longer dislocation segments, spanning several hundred nanometres, are observed within the continuous α-Al matrix among dispersed particles (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ed) at a deformation of \u003cem\u003eε\u003c/em\u003e\u0026thinsp;=\u0026thinsp;2% at 300\u0026deg;C. At the fracture strain (\u003cem\u003eε\u003c/em\u003e\u0026thinsp;\u0026asymp;\u0026thinsp;13.6%), clear interactions and segregation of dislocations are visible within the continuous α-Al matrix, similar to other alloys with dispersed second-phase particles \u003csup\u003e\u003cspan citationid=\"CR59\" class=\"CitationRef\"\u003e59\u003c/span\u003e\u003c/sup\u003e (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ee). This suggests that in the continuous α-Al matrix without a cellular structure \u0026ndash; or with a cellular structure at a larger cell size, as commonly found in LPBF alloys \u0026ndash; the dislocation free path (DFP) increases, making the CLS mechanism less effective. Instead, the F\u0026ndash;R mechanism becomes the predominant source of dislocations \u003csup\u003e\u003cspan citationid=\"CR60\" class=\"CitationRef\"\u003e60\u003c/span\u003e\u003c/sup\u003e. The critical stress (σₙ) required to activate an F\u0026ndash;R source is inversely proportional to the DFP (or the dislocation segment length \u003cem\u003eλ\u003c/em\u003e), as indicated in Eq.\u0026nbsp;2 \u003csup\u003e61\u003c/sup\u003e:\u003cdiv id=\"Equ2\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equ2\" name=\"EquationSource\"\u003e\n$$\\:{{\\sigma\\:}}_{F-R}=\\frac{MGb}{\\lambda\\:}\\frac{1}{2\\pi\\:(1-\\nu\\:)}[\\left(1-\\frac{3}{2}\\nu\\:\\right)\\text{ln}\\left(\\frac{\\lambda\\:}{b}\\right)+\\frac{\\upsilon\\:-2}{2}]$$\u003c/div\u003e\u003cdiv class=\"EquationNumber\"\u003e2\u003c/div\u003e\u003c/div\u003e\u003c/p\u003e \u003cp\u003eThe critical stresses for the CLS and F\u0026ndash;R mechanisms are compared as a function of the cell size or dislocation segment length (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003eg). CLS requires significantly higher stress at ultrafine cell sizes below 200 nm, which explains the high flow stress in the ECN-confined α-Al phase of the AlLaScZr-ECN sample. In contrast, \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:{{\\sigma\\:}}_{F-R}\\)\u003c/span\u003e\u003c/span\u003e is significantly lower for larger DFPs, resulting in reduced flow stress for the α-Al phase at elevated temperatures in the AlLaScZr-NP sample. This behaviour is consistent with observations in conventional Al alloys and other LPBF alloys with coarser dispersed reinforcements or cellular structures \u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e,\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e,\u003cspan citationid=\"CR43\" class=\"CitationRef\"\u003e43\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e \u003cp\u003eThe abovementioned results are validated by comparing the fracture modes (Figs.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ec, f). Dislocation motion in the AlLaScZr-ECN sample is confined within α-Al cells. Load is transferred to the Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e cell walls from α-Al during their co-deformation. Micro-cracking initiates by fracturing the cell walls (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ec), which requires higher stress. In contrast, dislocation segregation and interactions are more frequent within the α-Al matrix at the fracture strain, inducing local strain concentration and ductile failure characterised by micro-cracks distributed throughout the softer α-Al matrix (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ef). Moreover, the advantages of the ECN in load transfer and dislocation confinement are confirmed by the minimal YS difference (\u0026lt;\u0026thinsp;10 MPa) between the ECN and NP samples under tensile testing at 400\u0026deg;C (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ee, Fig. S8c). At this temperature, the ECN fully spheroidises into nanoparticles, similar to the NP sample, thereby losing its reinforcing benefits (Fig. S4b).\u003c/p\u003e \u003cp\u003eFurthermore, densely distributed L1\u003csub\u003e2\u003c/sub\u003e-Al\u003csub\u003e3\u003c/sub\u003e(Sc,Zr) nanoprecipitates contribute to strengthening by acting as obstacles to dislocation motion. At RT, these nanoprecipitates significantly shorten the DFP within α-Al \u003csup\u003e\u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e,\u003cspan citationid=\"CR62\" class=\"CitationRef\"\u003e62\u003c/span\u003e,\u003cspan citationid=\"CR63\" class=\"CitationRef\"\u003e63\u003c/span\u003e\u003c/sup\u003e. Similar YS values and dislocation segment lengths in the α-Al phase are observed between the ECN and NP samples at RT, as the nanoprecipitates are uniformly dispersed in both samples (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ee, Fig. S6). However, at elevated temperatures (\u0026ge;\u0026thinsp;300\u0026deg;C), the obstacle effects weaken as cross-slip and climb motions enable dislocations to bypass the precipitates \u003csup\u003e\u003cspan citationid=\"CR64\" class=\"CitationRef\"\u003e64\u003c/span\u003e,\u003cspan citationid=\"CR65\" class=\"CitationRef\"\u003e65\u003c/span\u003e\u003c/sup\u003e. In this case, the ECN plays a crucial role in limiting the DFP, thereby maintaining a high yield stress in the α-Al phase. Without the ECN, the DFP increases within the continuous α-Al matrix (as observed in the longer segments in Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ed), resulting in a substantial drop in YS. Despite this, at elevated temperatures (300\u0026deg;C and 400\u0026deg;C), the highly dense intracellular L\u003csub\u003e12\u003c/sub\u003e-Al\u003csub\u003e3\u003c/sub\u003e(Sc, Zr) nanoprecipitates still provide synergistic high-temperature strengthening alongside the ECN through effective dislocation\u0026ndash;precipitate interactions (inset in Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ea). This behaviour has been reported in Sc, Zr-modified Al alloys at elevated temperatures \u003csup\u003e\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e\u003c/sup\u003e. Micro-alloying of Sc, Zr further augments the YS by ~\u0026thinsp;45 MPa at 300\u0026deg;C and ~\u0026thinsp;40 MPa at 400\u0026deg;C, compared with LPBF Al-La alloy counterparts without Sc and Zr (Table S6 and Fig. S8).\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eConsistent with the proposed strengthening mechanisms, the macroscopic tensile behaviour also indicates enhanced dislocation storage and reduced dislocation annihilation due to the ECN. This is further supported by quantitative analysis using the dislocation-based Kocks\u0026ndash;Mecking (K\u0026ndash;M) model \u003csup\u003e\u003cspan citationid=\"CR66\" class=\"CitationRef\"\u003e66\u003c/span\u003e,\u003cspan citationid=\"CR67\" class=\"CitationRef\"\u003e67\u003c/span\u003e\u003c/sup\u003e applied to the tensile stress\u0026ndash;strain curves of AlLaScZr-ECN and AlLaScZr-NP alloys at both RT and 300\u0026deg;C (Fig. S7a, Supplementary Text). The results (Fig. S7, Table S5) suggest a higher dislocation storage rate (\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:{k}_{1}\\)\u003c/span\u003e\u003c/span\u003e) in the AlLaScZr-ECN sample than in the AlLaScZr-NP sample at RT, consistent with findings for LPBF AlSi10Mg \u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e,\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e\u003c/sup\u003e. Additionally, the high dislocation storage capacity is largely preserved at 300\u0026deg;C in the AlLaScZr-ECN sample, with \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:{k}_{1}\\)\u003c/span\u003e\u003c/span\u003e retaining 61% of its RT value, higher than that (17%) for the AlLaScZr-NP sample. Additionally, the ECN reduces the dynamic annihilation rate \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:{k}_{2}\\)\u003c/span\u003e\u003c/span\u003e at both RT and 300\u0026deg;C. This quantitative comparison using the K\u0026ndash;M model further validates the role of the ECN in enhancing flow stress in α-Al at elevated temperatures, driven by its high dislocation storage capacity and ability to inhibit dislocation annihilation.\u003c/p\u003e \u003cp\u003eIn conclusion, the LPBF AlLaScZr alloy achieves significantly enhanced high-temperature YS above 0.6\u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e of Al, while addressing the longstanding challenge of severe strength degradation in metals after prolonged thermal exposure, attributable to its thermally stable nanoscale ECN architecture and dense nanoprecipitate distribution. This alloy demonstrates substantial potential as an Al-based superalloy, extending the application range of lightweight Al alloys into the medium temperature range (300\u0026ndash;400\u0026deg;C) currently dominated by ferrous, titanium, and nickel-based superalloys for structural applications. Moreover, this study highlights the distinct advantages of the 3D interconnected network architecture over the traditional uniform particle dispersion configuration, which has been the standard for high-temperature strengthening. Additionally, it provides critical insights into the high-temperature strengthening mechanisms associated with AM-induced cellular structures, particularly the role of ultrafine cell sizes in restricting dislocation motion through the CLS mechanism, enabling effective strengthening of metal matrix to counteract the intrinsic softening. Based on our findings, a novel strategy is proposed to achieve exceptional high-temperature strength by integrating thermally stable phases into a nanoscale 3D network architecture, leveraging low-diffusivity alloying elements and the rapid solidification process of AM. We propose that this strategy can be broadly applied to other alloy systems, enabling extraordinary high-temperature mechanical performance through advanced AM technologies.\u003c/p\u003e"},{"header":"Methods","content":"\u003cdiv id=\"Sec6\" class=\"Section2\"\u003e \u003ch2\u003eLPBF process and heat treatment\u003c/h2\u003e \u003cp\u003eAlLaScZr powders were produced through gas atomisation from pre-alloyed ingots, resulting in spherical particles with diameters ranging from 15 to 53 \u0026micro;m (Fig. \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003ea). LPBF was performed using a metal 3D printer (Truprint 1000, Germany) equipped with a 200 W fibre laser and a fixed spot size of 30 \u0026micro;m. The hatching distance and layer thickness were set as 100 \u0026micro;m and 30 \u0026micro;m, respectively, with a laser power of 190 W. The laser scanning direction was alternated by 90\u0026deg; between consecutive layers without preheating the base plate. The optimal scanning speed of 1500 mm/s was determined based on the relative density of fully dense bulk (theoretical density: 2.86 g/cm\u003csup\u003e3\u003c/sup\u003e) and hardness of test cubes printed at various speeds (Fig. \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003eb). The density was measured using the Archimedes method, and hardness tests were conducted using a Vickers hardness tester (EZ-mat CARAT 930, Germany) with a 10 kg load and 15 s dwell time. Results across nine points were averaged. The as-built samples were subjected to isothermal ageing at 325\u0026deg;C for 1 h, yielding peak-aged samples, referred to as AlLaScZr-ECN in the main text (Fig. \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003ec). Additional samples were annealed for 168 h at 325\u0026deg;C to disrupt the ECN structure, producing AlLaScZr-NP samples (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ec). Other heat treatment conditions are specified in the main text.\u003c/p\u003e \u003c/div\u003e\n\u003ch3\u003eMechanical properties\u003c/h3\u003e\n\u003cp\u003eDog-bone-shaped bulk samples were printed (inset in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb) and subsequently sliced to 1.6 mm thickness using electrical discharge machining. The samples were then polished into rectangular tensile samples. The sample gauge length and width were 15 mm and 2.8 mm, respectively. Tensile testing (at both RT and elevated temperatures) was conducted on the peak-aged samples using a tensile testing machine (Zwick/Roell Z005, Germany) with a laser extensometer, following ASTM E8 and E21 standards. The tensile direction was oriented perpendicular to the BD, and the strain rate was maintained at 1 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e for all tests. For elevated temperature tests, samples were heated at 15\u0026deg;C/min and held for 30 min at the target temperature to achieve thermal equilibration prior to loading.\u003c/p\u003e \u003cdiv id=\"Sec8\" class=\"Section2\"\u003e \u003ch2\u003eMicrostructure characterisation\u003c/h2\u003e \u003cp\u003eThe chemical composition of the AlLaScZr powders and as-printed samples was determined using inductively coupled plasma\u0026ndash;atomic emission spectrometry (ICP-AES, iCAP6300, Thermo Fisher Scientific, USA), as detailed in Table \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003e. XRD was performed on as-printed and peak-aged samples using a D8 ADVANCE Da Vinci X-ray diffractometer (Bruker Corporation, Germany) with Cu Kα radiation (λ\u0026thinsp;=\u0026thinsp;0.1542 nm), scanning from 10\u0026deg; to 90\u0026deg; at 2\u0026deg;/min. Microstructural analysis was performed using a Zeiss AxioVision optical microscope (Germany) and a MAIA3 field emission scanning electron microscope (SEM) (TESCAN, Czech Republic) equipped with EDS and electron backscatter diffraction (EBSD) detectors. The SEM instrument operated at 10 kV, and samples were electrolytically polished in a nitric acid\u0026ndash;methanol solution (15 V, 10 s) for EBSD imaging. The EBSD scans used a 0.78 \u0026micro;m step size at 20 kV and were processed with HKL Channel 5 software. Samples for transmission electron microscopy (TEM) were prepared by mechanical grinding, followed by electro-polishing in a nitric acid\u0026ndash;methanol solution (1:3 by volume) at -30\u0026deg;C and 15 V. TEM and corresponding EDS mapping were conducted using Talos 200X and Titan Themis microscopes (Thermo Fisher Scientific, USA) with a HAADF detector and SuperX EDS with four silicon drift detectors. ImageJ was used to measure the cellular structure dimensions. APT samples were prepared using a focused ion beam (FIB) and analysed using a LEAP-5000XR microscope (Colorado MicroDissect, Inc., USA), with data processed in Cameca IVAS (France). Additionally, 3D X-ray micro-CT was performed with an Xradia 510 Versa instrument (Zeiss, Germany) using a voxel size of 1.5 \u0026micro;m to image large-volume specimens (1.35 mm \u0026times; 1.38 mm \u0026times; 0.85 mm). Pore size distribution analysis was conducted using Dragonfly software (ORS, Canada) (Fig. \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003ee, f), and 3D reconstruction of the ECN was achieved from 2D BSE images sliced with FIB (inset in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ec), with an interlayer distance of 7.5 nm.\u003c/p\u003e \u003c/div\u003e\n\u003ch3\u003eThermodynamic calculations\u003c/h3\u003e\n\u003cp\u003eThe solidification behaviour and phase evolution of the AlLaScZr alloy were simulated using the CALPHAD method in Pandat\u0026trade; software (CompuTherm LLC, USA), based on the PanAl75 commercial database. To simulate the rapid solidification during LPBF, Scheil mode simulations were conducted from 800\u0026deg;C to 600\u0026deg;C with a step of 1\u0026deg;C, based on the composition of the LPBF AlLaScZr alloy.\u003c/p\u003e\n\u003ch3\u003eIn situ SXRD test\u003c/h3\u003e\n\u003cp\u003eIn situ SXRD tensile testing was conducted at beamline ID11 of the European Synchrotron Radiation Facility. The beam energy was 65.3508 keV, and the wavelength was 0.01897 nm. The incident beam size was 0.2 mm \u0026times; 0.2 mm. For elevated temperature tests, the samples were heated to the target temperature and held for 30 min prior to tensile loading. A Frelon36 detector (2048\u0026times;2048 pixels and pixel size of 47 \u0026micro;m \u0026times; 47 \u0026micro;m) was used to acquire the diffraction patterns. The cross-section of the samples was approximately 0.5 mm \u0026times; 0.5 mm. The specimen-to-detector distance was 179.60 mm, calibrated using a CeO\u003csub\u003e2\u003c/sub\u003e powder standard. The AlLaScZr-ECN and AlLaScZr-NP samples were tested using a Nanox tensile stress rig \u003csup\u003e\u003cspan citationid=\"CR68\" class=\"CitationRef\"\u003e68\u003c/span\u003e\u003c/sup\u003e at both RT and 300\u0026deg;C. The Nanox voltage for loading was maintained through a constant ramp rate of 0.095 V/s at RT and 1.42 V/s at 300\u0026deg;C, corresponding to strain rates of 1.2 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;4\u003c/sup\u003e s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e and 1.4 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e, respectively. At RT, a lower strain rate was used than that in the macro-tensile tests (1 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e) to capture sufficient data points, as the material\u0026rsquo;s mechanical properties were relatively insensitive to strain rate. At 300\u0026deg;C, the strain rate matched that of the macro-tensile test. Background noise from the loading equipment was subtracted from the 2D diffraction images before calibration and integration. Detailed calculations for lattice strain and stress partitioning of constitutive phases are demonstrated in Supplementary Materials.\u003c/p\u003e \u003cdiv id=\"Sec11\" class=\"Section2\"\u003e \u003ch2\u003eFirst-principles calculations based on DFT\u003c/h2\u003e \u003cp\u003eDFT calculations were performed using the Vienna Ab initio Simulation Package (VASP) with projector augmented wave pseudopotentials and a plane-wave basis set \u003csup\u003e\u003cspan citationid=\"CR69\" class=\"CitationRef\"\u003e69\u003c/span\u003e\u003c/sup\u003e. The Perdew\u0026ndash;Burke\u0026ndash;Ernzerhof exchange\u0026ndash;correlation functional within the generalised gradient approximation was applied \u003csup\u003e\u003cspan additionalcitationids=\"CR71\" citationid=\"CR70\" class=\"CitationRef\"\u003e70\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR72\" class=\"CitationRef\"\u003e72\u003c/span\u003e\u003c/sup\u003e. A plane-wave cut-off of 450 eV was set with an energy convergence parameter of 10\u003csup\u003e\u0026minus;\u0026thinsp;5\u003c/sup\u003e\u0026nbsp;eV/atom for electronic self-consistency. The maximum forces on each relaxed atom converged to 0.02 eV/\u0026Aring; during structural relaxation through conjugate gradient minimisation. Gamma-centred 22 \u0026times; 22 \u0026times; 22 k-point meshes were used for the bulk unit cell of Al and Al\u003csub\u003e3\u003c/sub\u003eSc, while 21 \u0026times; 6 \u0026times; 6 k-point meshes were used for Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e, with densities lower than 0.03 \u0026Aring;\u003csup\u003e\u0026minus;1\u003c/sup\u003e. A vacuum layer with a thickness of 20 \u0026Aring; was added to the surface supercell to prevent unwanted interactions between the slab and its periodic arrangement. During the geometry optimisation, both the volume and atomic positions were relaxed for the initial interface systems, and only the atomic positions of the alloyed interface systems were relaxed. The lattice parameter for α-Al was \u003cem\u003ea\u003c/em\u003e\u0026thinsp;=\u0026thinsp;4.04 \u0026Aring;, while those for Al\u003csub\u003e11\u003c/sub\u003eLa\u003csub\u003e3\u003c/sub\u003e were \u003cem\u003ea\u003c/em\u003e\u0026thinsp;=\u0026thinsp;4.44 \u0026Aring;, \u003cem\u003eb\u003c/em\u003e\u0026thinsp;=\u0026thinsp;10.14 \u0026Aring;, and \u003cem\u003ec\u003c/em\u003e\u0026thinsp;=\u0026thinsp;13.14 \u0026Aring; \u003csup\u003e\u003cspan citationid=\"CR73\" class=\"CitationRef\"\u003e73\u003c/span\u003e\u003c/sup\u003e. Detailed calculations are demonstrated in Supplementary Materials.\u003c/p\u003e \u003c/div\u003e "},{"header":"Declarations","content":"\u003ch3\u003eData availability\u003c/h3\u003e\n\u003cp\u003eAll data supporting the findings of this study are available in the main text or the Supplementary Materials.\u003c/p\u003e\n\u003ch3\u003eAcknowledgements\u003c/h3\u003e\n\u003cp\u003eThis work was supported by the National Natural Science Foundation of China 52304405, 52371034, and 52204393. The authors also thank the European Synchrotron Radiation Facility (ESRF) for providing synchrotron radiation facilities under proposal number ma6301 (doi.org/10.15151/ESRF-ES-1830139531). Professional English language editing support provided by AsiaEdit (asiaedit.com).\u003c/p\u003e\n\u003ch3\u003eAuthor contributions\u003c/h3\u003e\n\u003cp\u003eConceptualisation: S.M., Z.C., M.W., and J. L.\u003c/p\u003e\n\u003cp\u003eMethodology: S.M., H.F., G.J., Y.C., Y.L., Y.W., and Y.Z.\u003c/p\u003e\n\u003cp\u003eInvestigation: S.M., Z.C., Y.C., and H.C.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eVisualisation: S.M., H.F., G.J., Y.L, and S.Z.\u003c/p\u003e\n\u003cp\u003eSupervision: Z.C., M.W., and S.N.\u003c/p\u003e\n\u003cp\u003eWriting\u0026mdash;original draft: S.M.\u003c/p\u003e\n\u003cp\u003eWriting\u0026mdash;review \u0026amp; editing: Z.C., G.J., M.W. and J. L.\u003c/p\u003e\n\u003ch3\u003eCompeting interests\u003c/h3\u003e\n\u003cp\u003eThe authors declare that they have no competing interests.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\u003cli\u003e\u003cspan\u003eMichi RA, Plotkowski A, Shyam A, Dehoff RR, Babu SS (2022) Towards high-temperature applications of aluminium alloys enabled by additive manufacturing. 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Mater Today Commun 24:101101\u003c/span\u003e\u003c/li\u003e\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":true,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":true,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"nature-portfolio","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"","title":"Nature Portfolio","twitterHandle":"","acdcEnabled":false,"dfaEnabled":false,"editorialSystem":"ejp","reportingPortfolio":"","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"","lastPublishedDoi":"10.21203/rs.3.rs-5840317/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-5840317/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eMetallic materials typically experience significant strength degradation at elevated temperatures. Traditional strengthening methods, which rely on thermally stable particle dispersion, exhibit limited effectiveness owing to the challenges in suppressing thermally activated dislocation motion. This work introduces a novel strategy for achieving exceptional high-temperature strength through a thermally stable nanoscale eutectic cellular network (ECN) enabled by additive manufacturing. A near-eutectic AlLaScZr alloy is developed for laser powder bed fusion, incorporating an Al-La nanoscale ECN and dense intracellular nanoprecipitates. This alloy demonstrates excellent printability and remarkable high-temperature yield strength above 0.6\u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e (~\u0026thinsp;250 MPa at 300\u0026deg;C), outperforming conventional aluminium alloys by 2\u0026ndash;5 times with minimal degradation after prolonged annealing. Compared with the conventional configuration of particle dispersion, the nanoscale ECN architecture enhances load-bearing capacity and strengthens aluminium by caging dislocation motion within ultrafine cells (~\u0026thinsp;200 nm), effectively mitigating intrinsic high-temperature softening. The proposed transformative approach paves the way for designing next-generation heat-resistant alloys, unlocking new possibilities for additive manufacturing in high-temperature applications.\u003c/p\u003e","manuscriptTitle":"Exceptional High-temperature Strength in an Additively Manufactured Al-based Superalloy with Stable Nanoscale Eutectic Cellular Network","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-01-31 06:53:07","doi":"10.21203/rs.3.rs-5840317/v1","editorialEvents":[],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"nature-communications","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"NCOMMS","sideBox":"Learn more about [Nature Communications](http://www.nature.com/ncomms/)","snPcode":"","submissionUrl":"https://mts-ncomms.nature.com/","title":"Nature Communications","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"ejp","reportingPortfolio":"Nature Communications","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"0ff7cc40-cbe8-400f-82a9-e61f0f2c99fd","owner":[],"postedDate":"January 31st, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"published-in-journal","subjectAreas":[{"id":43630741,"name":"Physical sciences/Materials science/Structural materials/Metals and alloys"},{"id":43630742,"name":"Physical sciences/Materials science/Structural materials/Mechanical properties"}],"tags":[],"updatedAt":"2025-12-24T08:10:02+00:00","versionOfRecord":{"articleIdentity":"rs-5840317","link":"https://doi.org/10.1038/s41467-025-66441-0","journal":{"identity":"nature-communications","isVorOnly":false,"title":"Nature Communications"},"publishedOn":"2025-11-22 05:00:00","publishedOnDateReadable":"November 22nd, 2025"},"versionCreatedAt":"2025-01-31 06:53:07","video":"","vorDoi":"10.1038/s41467-025-66441-0","vorDoiUrl":"https://doi.org/10.1038/s41467-025-66441-0","workflowStages":[]},"version":"v1","identity":"rs-5840317","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-5840317","identity":"rs-5840317","version":["v1"]},"buildId":"8U1c8b4HqxoKbykW_rLl7","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}