Nanoscale multi-gradient ordered architectures driven exceptional strength and ductility in low thermal expansion magnesium alloy

Domain: physics
Published: 2025-12-19
Editorial outcome: major_revision

Abstract
Alloys with low thermal expansion, high strength, and superior plasticity are crucial for critical applications in industries such as aerospace. Although the in-situ formation of low or negative thermal expansion particles presents a promising strategy, these materials generally suffer from limited ductility. Here, we demonstrate that the construction of nanoscale gradient ordered architectures can effectively addresses this limitation. By facilitating the diffusion and reaction of aluminum (Al) atoms into boronized manganese (Mn-B), we induce a gradient ordering architecture between rare-earthed magnesium (Mg) alloys and MnB phase, achieving an near-zero thermal expansion coefficient of 0.8 × 10-6·°C-1 within the temperature range of 280–320 °C. As fully transitioning from Mn-B to a multi-gradient ordered Mn-Al-B configuration results in a stabilized thermal expansion coefficient of 23 × 10-6·°C-1 across a broad temperature range (25–400 °C). This nanoscale architecture not only mitigates brittle interfacial fractures by maintaining the mechanical integrity but also enables the Mg alloy to reach an ultrahigh compressive strength of 507 MPa, with a 23.8% compressive strain. Our findings highlight the potential of designing gradient ordered architectures as a strategic approach to enhance the mechanical properties of lightweight alloys.

Introduction
Thermal expansion is a widely observed phenomenon with significant implications for materials science and engineering1,2. This behavior is quantitatively characterized by the thermal expansion coefficient (CTE)3,4. An excessively high CTE presents substantial challenges for metals and alloys, such as the risk of hot cracking during casting processes and dimensional instability caused by interface misalignments under service conditions5–9. These challenges are particularly pronounced in lightweight metals like magnesium (Mg). Owing to its relatively low atomic mass, Mg experiences substantial atomic vibrations at elevated temperatures, resulting in a higher CTE compared to conventional metals such as aluminum (Al), iron (Fe), copper (Cu), and titanium (Ti)10–13. Consequently, thermal cracking during casting or processing becomes a pronounced issue, seriously restricting its production ratio and structural applications in engineering fields, such as aerospace, automotive14–16.

A variety of strategies have been explored to reduce the CTE in metals, including alloying, composite design, and processing optimization17–21. Among these strategies, the incorporation of low or negative thermal expansion (NTE) materials, guided by composite theory, has proven to be the most effective approach22–26. Nevertheless, such strategies often compromise the ductility of metals due to undesirable insufficient interfacial bonding27,28. At high temperature, these issues are further compounded by detrimental chemical reactions between NTE phases and matrices, as well as the inherent limitations of NTE materials, which typically display an anomalous expansion behavior only within narrow temperature windows linked to specific phase transformations29,30. Therefore, the challenge of reducing the CTE of lightweight alloys at high temperatures while simultaneously preserving superior mechanical properties, particularly room-temperature plasticity, remains a critical and unresolved issue in the field of materials design.

In this study, we propose a strategy for the in-situ construction of a nanoscale multi-gradient ordered architecture to reduce the thermal expansion while enhancing the mechanical properties of a rare-earthed WE43 alloy (composed of Mg-4% Y-3.3% RE (Nd, Gd) −0.5% Zr by weight). The WE43 alloy was selected as a representative model system to address its inherent susceptibility to thermal cracking during processing, a key limitation hindering its broader application31–33. By optimizing the Al content in a boronized manganese (Mn-B) reinforced Mg matrix, we aim to achieve a serious transformation from Mn-B to various Mn-Al-B phases within multi-gradient configuration through Al diffusion and reaction. The in-situ thermal expansion strain analysis is used to elucidate the critical role of restructuring in alleviating internal stresses and strains during thermal exposure. This distinctive configuration enables the development of lightweight structural alloys with low thermal expansion and excellent mechanical properties.

Results
Microstructure characterization
The X-ray diffraction (XRD) analysis of the Al-modified rare-earthed Mg-Mn-B alloys show that the absence of distinctive diffraction peaks associated with Mn, MnB, and MgAl phases (Fig. 1a). Notably, the diffraction peaks indexed as (101), (002), and (100) of the Mg matrix exhibit a pronounced shift towards lower angles upon the incorporation of 5 wt.% Al. However, these peaks remain unchanged with an increase in Al content to 10 wt.%. This behavior indicates a release of lattice strain in the Mg matrix as the Al concentration increases to 10 wt.%. Although the formation of Al2Y phases is evident with 5 wt.% Al, no discernible diffraction signals from the 7 wt.% Mn-B constituents are observed. This absence suggests either the presence of nanocrystalline Mn-B that fall below the detection limit of conventional XRD or the existence of amorphous Mn-B nanoparticles.

Micromorphological analysis (Fig. 1b) indicates that the refinement of Mg grains remains minimal with the addition of Mn-B and Al34–36. We found that a significant reduction in pore density is observed as the Al content increases (Supplementary Fig. 1). The detailed Energy-dispersive X-ray spectroscopy (EDS) mapping shows that Al is predominantly located at the interfaces and partially diffuses into the Mn-B particles as increasing its contents (Supplementary Figs. 2, 3), significantly enhancing interfacial cohesion37. We also found that a diffusion of Al within the Mn-B particles (Fig. 1c), with complete diffusion occurring in particles smaller than 15 μm (Supplementary Fig. 4). Kernel Average Misorientation (KAM) analysis confirms strain accommodation in the 10 wt.% Al system, with substantially reduced strain localization at interfaces compared to conventional ceramic-reinforced Mg composites (e.g., SiCp/Mg and AlN/Mg)38,39 or Mg composites containing negative thermal expanded LaFe9.8Co1.1Al2.1 particles40, highlighting the unique interfacial compatibility achieved through this alloy design.

Interface observation
The high-resolution transmission electron microscopy (HRTEM) analysis (Fig. 2) elucidates atomic-scale interfacial structures and chemical distributions at the interface. Elemental mapping demonstrates complete Mg-Al-Mn intermixing at the interfaces (Fig. 2a–d), with notable Al diffusion into Mn-B particles. Atomic-scale imaging shows perfect lattice continuity across Mg/AlMn-B interfaces (Fig. 2e–g) with two distinct strain effects. One is a lattice compression in Mg adjacent to interfaces, and the other one is an expansion near Mn-B particles (Fig. 2i). This strain coupling effectively mitigates interfacial embrittlement through elastic energy dissipation. Most remarkably, we directly resolved a 4–10 nm gradient ordered architecture within AlMn-B, generated via Al-Mn-B reaction, exhibiting continuous lattice spacing evolution (Fig. 2j–m). This configuration coherently accommodates the lattice mismatch while maintaining perfect crystallographic registry, enabling a unique structural configuration achievement in Mg matrix composites.

The diffraction pattern reveals that the AlMn-B structure can be indexed to either the [110] zone axis of the Mn4Al11 phase or the [001] zone axis of the MnB phase (Fig. 2h). Partial overlap between the (001) diffraction spots of Mn4Al11 and the (010) spots of MnB indicates symmetry breaking along this orientation. This can be attributed to the Al/B defects during the formation of nanoscale gradient metastable AlMn-B. In the 5 wt.% Al sample, regions without interfacial cracks reveals that the interfacial formed Mn2Al3 phase preserves gradient-varying interfacial contact (Supplementary Fig. 5) suggesting the AlMn-B phase maintains strong cohesion with the Mg matrix and mitigates brittle interfacial fracture. This further supports our expectation that varying Al contents induce distinct phase formations.

In the Mg matrix adjacent to the interface, geometric phase analysis (GPA) and dislocation density comparisons between 0 wt.% and 10 wt.% Al samples reveal a stark contrast (Supplementary Fig. 6). The former displays pronounced interfacial strain concentrations and high dislocation density, whereas the latter exhibits a near dislocation-free interface with dramatically reduced strain. The absence of deformation dislocations in Mg matrix, variable lattice spacings, and asymmetric spot broadening in AlMn-B confirm that these chemically nanoscale ordered transition layers effectively relieve interfacial stresses. Real-space HRTEM and reciprocal-space SEAD data consistently demonstrate this Al diffusion and reaction drive the formation of nanoscale multi-gradient ordered architectures with progressive crystallographic symmetry variation. The above results suggest that the nanoscale gradient ordered interfacial architecture is established and have a good cohesion within Mg matrix.

Mechanical property and Thermal expansion behavior
The compressive mechanical properties and thermal expansion characteristics of the samples were systematically investigated (Fig. 3, Supplementary Figs. 7, 8, Supplementary Tables 1, 2). With increasing Al content, both the ultimate compressive strength and compressive strain of the alloys demonstrated a progressive enhancement. Notably, at an Al concentration of 10 wt.%, the alloy attained a compressive strength of 507 ± 8 MPa accompanied by an elongation of 23.8 ± 0.6%, indicative of significant plasticity improvements in the Mg alloy (Fig. 3a, Supplementary Fig. 9). The tensile performance improvement is less significant than the compressive performance (Supplementary Fig. 10), which reflects the typical tension-compression asymmetry of magnesium alloys due to their hexagonal close-packed crystal structure. Nevertheless, compared with the WE43 alloy containing only 7 wt.% Mn-B, the addition of 10 wt.% Al raises the yield tensile strength from 157 ± 4 MPa to 204 ± 2 MPa, while maintain an elongation of 8 ± 0.5%, which is a notable enhancement compared with the nearly complete loss of ductility observed without Al addition (Supplementary Fig. 11).

Interestingly, thermal strain analysis revealed an near-zero expansion regime in samples with 0 wt.% and 5 wt.% Al (Fig. 3b), analogous to behaviors observed in alloys solely addition of pure MnB phase (Supplementary Fig. 12). This regime coincides with the phase transition temperature of MnB, approximately 280–300 °C41,42. The results suggest the potential formation of the MnB phase during alloy preparation, as corroborated by EDS point analyses (Supplementary Fig. 13). In comparison to magnesium alloys previously reported in the literature for their low thermal expansion and robust plasticity7,21,43–52, the alloy synthesized through our approach exhibits superior mechanical strength and plasticity (Fig. 3c). This highlights the critical role of interface control through alloying in achieving outstanding thermal expansion and mechanicalload transfer properties.

The formation of the MnB phase typical requires high pressure, indicating significant thermal stress during synthesis. Interestingly, alloys modified with 5 wt.% Al display an expanded near-zero thermal strain range compared to those influenced predominantly by MnB alone. We propose that this anomalous thermal stability arises from a synergistic effect among Al-induced deformation recovery, the presence of porosity in the Mg matrix, and interactions between MnB and Mg, collectively broadening the thermal stability range40. This hypothesis is supported by XRD analysis, which indicates significant residual strain in the sample with 5 wt.% Al. In contrast, the sample with 10 wt.% Al shows reduced residual strain, as evidenced by TEM and GPA characterization of the Mg matrix (Supplementary Fig. 6), indicating strain relaxation and a stable CTE of 23 × 10-6 °C-1. Remarkably, this CTE aligns with that of many Al alloys, yet the material maintains a substantially lower density of 1.83 g/cm³ compared to 2.7 g/cm³ of conventional Al alloys10,21. This research advances our understanding of macroscopic thermal strain, phase transitions, and residual strain during processing by examining variations in the CTE across different temperature ranges, providing a scientific basis for optimizing manufacturing processability and operational reliability of lightweight alloys.

Discussion
Formation mechanism of the gradient ordered architecture
Three-dimensional atom probe tomography (3D-APT) and proximity histogram analysis across the Mg/AlMn–B interface reveal a nanoscale intermixing region (2–5 nm wide) exhibiting near-perfect spatial overlap of B, Mn, and Al elemental distributions (Fig. 4a, b). This colocalization confirms that B atoms dissolve atomically into the Mn-Al lattice, forming a compositionally graded metastable phase rather than discrete MnB precipitates at the interface. To accommodate local strain from solute B atoms or Al deficiencies, Mn atoms undergo coordinated lattice adjustments, consistent with stress-activated structural and electronic adaptations7,53,54. The segregation of solute B atoms generates strain accommodating regions through atomic-scale structural modulation every few lattice periods55,56, which contribute to the coherent interfacial architecture observed in Fig. 2.

To further elucidate the formation mechanism of AlMn-B, we conducted first-principles calculations of temperature-dependent free energies for MnB, Mn2Al3, Mn4Al11, MnAl-B, and Mn4Al11-B phases. Our calculations (Fig. 4c) indicate that Mn4Al11-B possesses the lowest free energy at 520 °C, elucidating its preferential formation in Al-rich interfacial regions. In contrast, at room temperature, Mn2Al3 and Mn4Al11 are thermodynamically more stable, suggesting that the metastable AlMn-B phase undergoes decomposition or phase separation during cooling57–59. The final phase assemblage depends strongly on the initial Al content, which governs both the primary phase formed and the composition of precipitates generated upon cooling.

We thus proposed the following formation pathway to elucidate the evolution of this distinctive compositional gradient ordered architecture and its effect on Mg matrix (inset of Fig. 4c, Supplementary Fig. 14). During hot-pressing at 520 °C, interfacial Al atoms rapidly diffuse into Mn-B, leading to the formation of a B-supersaturated AlMn-B phase. Upon cooling, this phase undergoes decomposition via B precipitation, while spatial variations in Al content and localized stress promote the development of a compositional gradient through coordinated B segregation. In the absence of Al, direct formation of MnB particles at the Mg/Mn–B interface induces high residual strain due to mechanical and thermal mismatch. With limited Al addition, although the stable Mn2Al3 enhances interfacial bonding, the mismatch-induced localized stress concentration still promotes MnB precipitation and residual strain accumulation in the Mg matrix. At higher Al contents, however, the high-temperature Mn4Al11-B phase decomposes under cooling through a defect-mediated restructuring process. Enrichment of solute B atoms effectively suppresses MnB formation and facilitates discontinuous transformation within the Mn4Al11, dissipating interfacial energy and minimizing microcracking. Consequently, in the 10 wt.% Al alloy, MnB is eliminated, strain concentration is absent, and no zero-thermal-expansion regime is observed. This collectively demonstrate the role of the nanoscale gradient structure in achieving interfacial stability and thermodynamic equilibrium.

Mechanism of strength and ductility synergy
Microstructural analysis of the 10 wt.% Al composites deformed to 24% strain reveal the emergence of Mn5Al8 phase (Fig. 5a–d). GPA shows pronounced strain localization within the Mg matrix adjacent to the Mn5Al8 regions, while the original Mn4Al11-B phases maintain substantially lower strain levels (Supplementary Fig. 15). Moreover, dislocations observed within Mn5Al8 further confirm plastic activity in the secondary particles. First-principles calculations demonstrate that applied strain drives the transition from metastable Mn4Al11-B to Mn5Al8 (Fig. 5e, Supplementary Fig. 16), revealing a stress-induced phase transformation mechanism.

Based on the observed microstructural evolution, we propose the following strengthening and toughening mechanism (Fig. 5f). Stress induced phase transformation and plastic activity in the secondary particle enable high density dislocation storage for strengthening and work hardening60, while simultaneously improving ductility through enhanced dislocation mobility. The mitigation of brittle interfacial fractures achieved by the nanoscale gradient ordered architecture, along with reduced residual matrix strain, facilitates mechanically synergistic particle-matrix interactions. This process is accompanied by various phase transformations mediated volumetric adjustments that reinforce interfacial cohesion and delay fracture initiation. Furthermore, dislocation activity within the newly formed phases at higher strain levels provides additional enhancement of both strength and ductility27,61,62. This nanoscale compositional gradient architecture effectively converts conventional weak interfaces into active energy dissipation zones, ultimately enabling the exceptional strength-ductility synergy observed in this system.

In conclusion, our study demonstrates the formation of a nanoscale gradient ordered Mn-Al-B architecture in Mg alloys. This unique structural feature effectively mitigates stress concentration at the interface, resulting in an exceptional combination of high strength (507 MPa) and ductility (23.8% compressive strain) in boronized manganese particle reinforced Mg matrices, surpassing the performance of conventional alloys reported to date. The atomic level gradient ordered architecture engineering strategy facilitates the development of lightweight structural alloys with low thermal expansion and tunable residual thermal strain. The remarkable combination of low thermal expansion and favorable elastic properties positions these alloys for applications in high-precision devices requiring stable elastic rigidities, such as mechanical chronometers for space missions. This work introduces a design strategy for composites, with the nanoscale gradient architecture formed via reaction and segregation gradients laying the groundwork for future composite material studies.

Methods
Sample preparation
The Mn powder and B powder were mixed in a planetary ball mill at an atomic ratio of one to one for 60 min, and the rotation speed was 180 r/min. Then the vacuum sealing tube was placed in a crucible filled with graphite powder. The crucible was placed in an insulating brick and heated in a microwave oven for 5–10 minutes to obtain Mn and B mixed powder, denoted as boronized manganese (Mn-B). Different mass ratios of Mn-B powder, Al powder and WE43 powder were mixed in a planetary ball mill at a speed of 180 r/min for 2.5 h. Then it was sintered in a vacuum hot-press sintering furnace with a vacuum degree of 1 × 10−5 Pa, a pressure of 20 MPa, a temperature of 520 °C, and a holding time of 1 h. The sample is then cooled to room temperature in the furnace.

Microstructure characterization
The X-ray diffraction (XRD) pattern of samples was recorded by the D/Max 2500pc X-ray diffractometer with monochromatic Cu-Kα radiation (λ = 1.5418 Å). The microstructure was examined by a JEOL-7800F type scanning electron microscopy (SEM) and FEI Talos F200X type field emission transmission electron microscopy (TEM) and FEI 200. The Fast Fourier transform pattern was processed using DigitalMicrograph software (version 3.43.3212.0) with a Wiener filter applied for noise reduction. The APT measurements were conducted on a CAMECA LEAP 5000 XR. Electron backscatter diffraction (EBSD) mapping was conducted utilizing an Oxford Instruments AZtec system (version 2.1.2), with step sizes varying between 0.26 and 0.46 µm. The specific step size and resultant indexing rate for each scan are provided in Supplementary Table 3. Subsequent analysis of the EBSD data was performed using the MTEX toolbox (version 5.8.0) operating within MATLAB R2024A. For grain reconstruction, a misorientation threshold of 15° was applied to delineate grain boundaries, which then underwent smoothing with a parameter of 5. Final sample preparation was performed by argon ion polishing, and the shooting instrument was JEOL JSM-7800F.

Mechanical properties and thermal expansion performance test
The dimension of the specimens for CTE measurement was Ø 3 mm × 3 mm. The linear thermal expansion data (dL/L) of samples were measured using a NETZSCH 402C Dilatometer under a heating rate of 10 °C/min and temperature ranging from room temperature to 200 °C. The room temperature compressive mechanical properties were measured by a CMT-5305-300KN universal testing machine. The size of the compressed sample was Ø 4 mm × 8 mm.

First-principles Calculations
Theoretical investigations were conducted to understand the mechanisms behind the structure evolution and its effect on the mechanical properties. The calculations of phonon dispersion spectra and related thermodynamic properties were conducted using density functional perturbation theory (DFPT)63. All first-principles calculations were performed employing the Vienna Ab initio Simulation Package (VASP, version 6.3.2)64,65,66]. The projector augmented-wave (PAW) pseudopotentials67 and the Perdew-Burke-Ernzerhof (PBE)68 generalized gradient approximation functional were used to describe the electron exchange-correlation interactions69.

We began with the geometric optimization of the pristine MnAl unit cell, using a Brillouin zone k-point mesh of 3 × 3 × 4. The energy convergence criterion for the electronic self-consistent cycle was set to 10−8 eV, and the ionic relaxation was considered converged when the forces on atoms were below 0.001 eV/Å. Upon obtaining the optimized host structure, we constructed a supercell, introduced B atom doping, and built a supercell model containing a total of 64 atoms. The energy cut-off of the plane-wave basis sets was 450 eV.

Based on this doped supercell, the first-order force constants required for the dynamical matrix were computed directly using the DFPT method implemented in VASP. Since this 64-atom supercell is sufficiently large to capture the relevant atomic interactions, it was treated as a 1 × 1 × 1 primitive cell for the phonon calculations. A q-point mesh of 7 × 7 × 4 was used to sample its Brillouin zone. The force constants obtained from the DFPT calculations were then processed using the Phonopy software to obtain the phonon dispersion relations and the phonon density of states (PDOS). Finally, within the harmonic approximation, the thermodynamic properties of the system, including the Helmholtz free energy (F), entropy (S), and constant-volume heat capacity (Cv), were derived by integrating the PDOS and are presented as functions of temperature.

The atomic coordinates of all optimized computational models are provided in the Source Data file.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The authors show that incorporating Mg-B particles and regulating interfacial diffusion in Mg alloys results in the formation of\na nanoscale gradient-ordered Mn-Al-B structure. This architecture effectively alleviates stress concentration, thereby\nachieving an exceptional strength-ductility combination (500 MPa, 24% compressive strain) and low thermal expansion.\nThis work pioneers a novel \"atomic-level gradient-ordered engineering\" strategy for high-precision lightweight composites.\nHowever, the manuscript exhibits several significant issues that necessitate major revisions to enhance its scientific rigor\nand persuasiveness.\n(1) APT characterization is essential to conclusively demonstrate the claimed interfacial elemental gradient distribution. The\nauthors attribute significant left-shift of XRD peaks solely to lattice strain but provide no compelling evidence. Neutron\ndiffraction is required to validate the proposed lattice strain mechanism for the XRD peak shift.\n(2) The manuscript fails to specify which alloy composition (Al5 vs. Al10) exhibits near-zero CTE at 280-320°C. Crucially, it\ndoes not clarify whether the reported 500 MPa strength and 24% ductility correspond to the low-CTE composition or\ndemonstrate a simultaneous achievement.\n(3) The fundamental mechanism by which the nanoscale gradient structure synergistically enhances both strength and\nductility remains unexplained. Thermodynamic/kinetic analysis of Al diffusion-driven gradient formation is lacking.\n(4) The selection of WE43 alloy as the model material lacks justification. Its high rare-earth content (Y, Nd) contradicts the\nlightweight objective, and alternative Mg-RE alloys with potentially superior strength-to-weight ratios exist but were not\nconsidered. Furthermore, tensile properties of the modified WE43 alloy are absent, limiting assessment of practical\nrelevance.\n(5) The critique of existing CTE-reduction strategies (alloying, processing) causing weak interfaces is undermined, as the\nauthors' own solution relies fundamentally on composition control (Al addition) and process optimization (diffusion control).\nThe claimed \"gradient structure\" appears largely observational rather than a distinct design strategy.\n(6) Claims that the Mn-Al-B gradient \"effectively releases interfacial stress\" lack direct experimental validation (e.g., interface\nstress mapping). Characterization of as-sintered interface stress states is also missing.\n(7) Figure 2's structural analysis must explicitly state the Al content (5Al or 10Al?). Characterization of the 5Al alloy\n(reportedly better CTE) is needed to determine if the gradient structure exists and why its mechanical properties differ.", "reviewer_2": "In the manuscript, the “gradient-ordered architecture” is more a descriptive term than a rigorously defined phenomenon. The\nauthors should provide a schematic or mechanistic diagram clearly outlining the Al diffusion process, MnB → Mn₄Al₁₁\ntransition, and how this sequence leads to internal stress accommodation and thermal expansion suppression. Without this, the structural\ninterpretation lacks the clarity required for broader generalization.\n1.Lines 123-124\n“Notably, at an Al concentration of 10 wt.%, the alloy attained a compressive strength 123\nof 500 MPa accompanied by an elongation of 24%” The error values should be presented.\n2.Lines 93-95\nGeometric phase analysis could effectively reveal the strain distribution across the Mg/AlMn-B interface.\n3.Lines 149-150\nThe analysis of the influence of gradient-ordered architecture on compressive strength and ductility remains superficial.\nAdditional experiments or simulations are necessary to elucidate the mechanisms. For instance, microstructure\ncharacterization of the tensile sample with different strain to reveal strengthening effects of the gradient-ordered architecture\nduring compressive deformation.\n4.The discussion of the formation mechanism of the gradient ordered architecture requires further validation for its rigor and\npersuasiveness. The authors only investigated the interface structure of 10 Al-7(Mn-B)/WE43 alloy. The interface structure of\nother sample with different contents of Al should be characterized to further confirm the formation process of nanoscale multi-\ngradient ordered architecture.\n5.The manuscript only reports compressive properties, without any tensile data. Given that magnesium alloys typically show\nstrong tension–compression asymmetry, and that the introduced Mn-B particles and Al-induced phase transformation (MnB\n→ Mn₄Al₁₁) likely create brittle interfaces prone to cracking under tension, the absence of tensile results is a major concern.\nTensile performance is essential for structural applications and should be presented or at least properly discussed."}, "author_response": "Point by point response to reviewers’ comments\nDear reviewers,\nWe sincerely appreciate your valuable and constructive feedback on our manuscript. We have\nthoroughly revised the manuscript to address all comments. Additionally, new figures have been\nincluded in the Supplementary Materials to further clarify and support the revisions. The\namendments have been highlighted in both the revised manuscript and the Supplementary Materials\n(with the order of figures in the Supplementary Materials adjusted to align with the changes in the\nmain text).\nA detailed, point-by-point response to the reviewers’ comments is enclosed below. We hope that the\nrevised version of the manuscript is now suitable for publication in Nature Communications.\nYours Sincerely\nSujuan Wu on behalf of all authors\nReviewer #1 (Remarks to the Author):\nThe authors show that incorporating Mn-B particles and regulating interfacial diffusion in Mg\nalloys results in the formation of a nanoscale gradient-ordered Mn-Al-B structure. This architecture\neffectively alleviates stress concentration, thereby achieving an exceptional strength-ductility\ncombination (500 MPa, 24% compressive strain) and low thermal expansion. This work pioneers a\nnovel \"atomic-level gradient-ordered engineering\" strategy for high-precision lightweight\ncomposites. However, the manuscript exhibits several significant issues that necessitate major\nrevisions to enhance its scientific rigor and persuasiveness.\nAuthor reply:\nWe sincerely appreciate the reviewer’s valuable comments on our work. In response, we have\nsubstantially revised the manuscript to improve its scientific rigor and persuasiveness. Specifically,\nwe performed additional microstructural characterizations across different Al contents (0, 5, and 10\nwt.% Al alloys), complemented by thermodynamic calculations, to systematically validate the\nproposed formation mechanism of the nanoscale gradient ordered Mn-Al-B structure. We also\nexpanded the discussion on residual strain formation and its influence on mechanical properties,\nproviding clearer mechanistic insights into how the gradient architecture alleviates stress\nconcentration and enables strength and ductility synergy. Furthermore, new figures and\nsupplementary data have been incorporated to support these analyses. We believe these\ncomprehensive revisions have significantly strengthened the manuscript and addressed the\nreviewer’s concerns.\n(1) APT characterization is essential to conclusively demonstrate the claimed interfacial elemental\ngradient distribution. The authors attribute significant left-shift of XRD peaks solely to lattice strain\nbut provide no compelling evidence. Neutron diffraction is required to validate the proposed lattice\nstrain mechanism for the XRD peak shift.\nAuthor reply:\nAPT Characterization:\nWe thank the reviewer for this constructive suggestion. We have conducted additional APT\nexperiments to investigate the interfacial elemental gradient distribution. The new APT results\n(Figure R1.1), combined with the atomic-scale evidence for the interfacial gradient structure of Al\nin Mn-B observed by STEM-HAADF imaging and corresponding EDS mapping (Figure 2 in the\nmain text), provide direct evidence for B segregation in the meta stabled Mn Al -B particles and\n4 11\nthe formation of nanoscale gradient architecture at the interface.\nSpecifically, the three-dimensional elemental reconstruction (Figure R1.1 A) and the corresponding\nproximity histogram (Figure R1.1 B) across the Mg/Mn-Al interface reveal a nanoscale (2-5 nm\nwide) region where B, Mn, and Al elements exhibit a perfect overlap in their distribution. This\nconfined overlap is a definitive signature of the formation of a nanoscale gradient structure, rather\nthan the presence of separate boride particles at interfaces. We deduce that B atoms are dissolved\nwithin the Mn-Al lattice to form a metastable Mn Al -B phase with a composition gradient. This\n4 11\natomic-scale evidence for the interfacial gradient structure is consistent with and strongly supported\nby our STEM-HAADF imaging and EDS mapping results presented in Figure 2 of the main text.\nThe formation of this metastable Mn Al -B gradient structure is crucial, as it can transform into\n4 11\nstable phases under applied stress. This transformation stabilizes the evolving interface throughout\nthe phase transition and mitigates brittle interfacial fractures by preserving mechanical integrity.\nThank you once again for this comment, which has significantly strengthened the evidence for the\nnanoscale gradient structure in our manuscript. This new APT analysis has been included in the\nrevised manuscript: Three-dimensional atom probe tomography (3D-APT) and proximity histogram\nanalysis across the Mg/AlMn-B interface reveal a nanoscale intermixing region (2-5 nm wide)\nexhibiting near-perfect spatial overlap of B, Mn, and Al elemental distributions (Figure 4A). This\ncolocalization confirms that B atoms dissolve atomically into the Mn Al lattice, forming a\n4 11\ncompositionally graded metastable phase rather than discrete MnB precipitates at the interface. To\naccommodate local strain from interstitial B atoms or Al deficiencies, Mn atoms undergo\ncoordinated lattice adjustments, consistent with stress activated structural and electronic\nadaptations7,43,44. The segregation of B atoms results in the formation of strain accommodating\nregions with atomic scale structural modulation every few lattice periods45,46.\nFigure R1.1. APT characterization of the nanoscale interfacial gradient. (A) 3D atom map\nreconstruction of the Mg/Mn-Al interface, showing the distribution of Mg, Al, Mn and B atoms. (B) One-\ndimensional compositional profile across the Mg/Mn-Al interface, illustrating the nanoscale (2-5 nm)\nsegregation and perfect overlap of B, Mn, and Al elements. This overlap confirms the formation of a B-\ncontaining AlMn-B gradient structure (e.g., metastable Mn Al -B) at the interface.\n4 11\nValidation lattice strain\nFor the lattice strain, we agree that definitive evidence is required to attribute the XRD peak shift to\nlattice strain. We performed neutron diffraction experiments to directly response the suggestion.\nHowever, as shown in the new Figure R1.2, due to the strong neutron absorption by elements present\nin the WE43 matrix (particularly Gd) and the MnB additions, no usable diffraction peaks were\nobtained despite a prolonged irradiation time of 18 hours. Given this technical limitation, we\npursued complementary multi-scale characterization to validate the lattice strain.\nFigure R1.2. Neutron diffraction pattern for strain observation. Neutron diffraction pattern\nobtained from the WE43 sample. The strong neutron absorption by Gd resulted in a very weak signal\nto noise ratio, preventing the acquisition of usable diffraction peaks for strain analysis, even after\nan extended irradiation time of 18 hours.\nWe used HRTEM to verify the existence of strain in the Mg matrix. As shown in Figure R1.3, a\nsubstantially higher density of dislocations and associated lattice strain is observed in the matrix of\nthe 5 wt.% Al sample, in contrast to the 10 wt.% Al sample which shows fewer defects. This inverse\nrelationship, where a higher Al contents results in fewer matrix defects, deviates from conventional\nwisdom. To understand this abnormal microstructure and their corresponding property, we designed\nthe experiment as presented in our manuscript and verified the interfacial stabilization effect\nproposed in our work.\nFigure R1.3. Comparative analysis of lattice strain and dislocation density in 5 wt.% Al and 10 wt.%\nAl alloys. HADDF-STEM and HRTEM images revealing the dislocation density and lattice strain in the\nmatrix of (A-D) 5 wt.% Al and (E-H) 10 wt.% Al samples. A significantly higher density of dislocations\n(examples indicated by white “T” symbols) and associated strain contrast is evident in the 5 wt.% Al\nsample, providing direct microstructural evidence for the residual lattice strain.\nTo further confirm the existence of macroscopic strain, we employ CTE measurements for\nillustration. It can be observed that the CTE of the 5 wt.% Al sample in the low-temperature region\nis significantly lower than that of the pure WE43 sample and also lower than the sample without Al\naddition. The reduction in the CTE is attributed to the presence of strain, i.e., a large number of\ninternal defects (such as dislocations and solute atoms). During the heating process, strain recovery\n(also known as residual strain recovery) occurs. The volume contraction caused by this strain\nrecovery—mainly the reduction in volume due to the disappearance of defects—partially\ncompensates for the volume expansion induced by thermal expansion. Consequently, the net\nvolumetric change is reduced, resulting in a smaller apparent CTE. The greater the reduction in CTE,\nthe larger the amount of residual strain retained in the material. Therefore, the observed suppression\nof CTE in the 5 wt.% Al sample before any phase transformation occurs strongly indicates the\npresence of substantial residual strain. This provides compelling macroscopic evidence supporting\nour interpretation of the XRD peak shifts. For a more comprehensive discussion of the correlation\nbetween strain and its effect on thermal expansion, we refer to our prior publication in Applied\nPhysics Letters (doi: 10.1063/5.0284514) as well as ongoing, unpublished work.\nCollectively, the evidence from HRTEM (direct observation of defects) and CTE (macroscopic\nstrain signature) robustly confirms that the significant XRD peak left-shift in the 5 wt.% Al sample\nis indeed primarily caused by lattice strain. We sincerely thank the reviewer for guiding us toward\nthis important improvement.\n(2) The manuscript fails to specify which alloy composition (Al5 vs. Al10) exhibits near-zero CTE\nat 280-320°C. Crucially, it does not clarify whether the reported 500 MPa strength and 24%\nductility correspond to the low-CTE composition or demonstrate a simultaneous achievement.\nAuthor reply:\nWe thank the reviewer for pointing out this important clarification. In the revised manuscript, we\nhave explicitly specified the composition and property correspondence.\nAs shown in Figure 3 and summarized in Table 1 and Figure R1.4, the 5 wt.% Al sample exhibits\nnear-zero thermal expansion (CTE ≈ 0.8×10-6·°C⁻¹) in the 280-320 °C range while maintaining a\ncompressive strength of 295 ± 3 MPa and elongation of 18.5 ± 0.4 % at room temperature. In\ncomparison, the 10 wt.% Al sample demonstrates a low CTE of 23×10-6·°C⁻¹ with superior\nmechanical properties, achieving 507 ± 8 MPa compressive strength and 23.8 ± 0.6 % elongation.\nThese findings are further supported by the repeat mechanical property test and the following\nrevision has been highlighted in the revised manuscript.\nFigure R1.4. Summarize the mechanical properties. Room temperature compressive strength and\nelongation to failure of the WE43, 7 wt.% Mn-B/WE43 (0 wt.% Al+7*), 5 wt.% Al+7 wt.% Mn-B/WE43\n(5 wt.% Al+7*) and 10 wt.% Al + 7 wt.% Mn-B/WE43 (10 wt.% Al+7*) alloys.\nThis distinction has now been clearly articulated in the main text: Notably, at an Al concentration\nof 10 wt.%, the alloy attained a compressive strength of 507 ± 8 MPa accompanied by an elongation\nof 23.8 ± 0.6%, indicative of significant plasticity improvements in the Mg alloy.\n(3) The fundamental mechanism by which the nanoscale gradient structure synergistically enhances\nboth strength and ductility remains unexplained. Thermodynamic/kinetic analysis of Al diffusion-\ndriven gradient formation is lacking.\nAuthor reply:\nWe thank the reviewer for this insightful comment. In the revised manuscript, we have provided\nboth mechanistic and thermodynamic/kinetic analyses to clarify how the nanoscale gradient\nstructure synergistically enhances strength and ductility, and how Al diffusion drives the gradient\nformation.\nMechanism of strength and ductility synergy\nThe mechanism by which nanoscale gradient structures simultaneously enhance strength and\nductility has been illustrated in Figure R1.5, and consistent with prior studies (Science 2020, 369,\n427-432; Acta Mater. 2023, 259, 119238). Generally, applied stress in particle-reinforced alloys\nusually accumulates at particle-matrix interfaces, often causing interface failure and crack\npropagation. In our system, however, the accumulated stress is partially released through stress-\ninduced phase transformation, such as the meta stabled Mn Al -B to Mn Al . This phase\n4 11 5 8\ntransformation, accompanied by local volumetric changes, simultaneously strengthen the interface\nby delaying failure and increase strain through transformation-enabled strain redistribution.\nFigure R1.5. Schematic mechanism of the strength-ductility synergy mediated by stress induced\nphase transformation and further strain redistribution. (Ⅰ) The dislocations generated by applied\nstress pile up at the nanoscale metastable Mn Al -B interface, leading to stress concentration. (Ⅱ) Once\n4 11\nthe stress concentration exceeds the threshold value, it triggers the structural transformation of meta\nstable Mn Al -B to stable Mn Al , which releases a portion of dislocations, relieves the stress\n4 11 5 8\nconcentration, and redistributes it to the adjacent regions. (Ⅲ) With increasing stress, the transformed\nregion cooperates with untransformed areas to accommodate further deformation, such as active\ndislocations in the particles, effectively mitigating damage nucleation and enhancing strain hardening\nWe have conducted additional experimental and computational analyses to provide a clear and\nfundamental explanation, which has been incorporated into the revised manuscript (Figure 5).\nIn the 10 wt.% Al alloy subjected to 24% strain, a new Mn Al phase is observed (Figure R1.6). Free-\n5 8\nenergy calculations show that Mn Al with higher free energy than both Mn Al -B and Mn Al at\n5 8 4 11 4 11\nroom temperature (Figure R1.7 A). However, applied strain drives Mn Al -B above the energy\n4 11\nlevel of Mn Al at ~2% strain (Figure R1.7 B). This energy crossover provides the driving force for\n5 8\nthe stress-induced phase transformation from Mn Al -B to Mn Al during mechanical loading,\n4 11 5 8\ncorroborating the experimental observations.\nFigure R1.6. Microstructural evidence of stress induced phase transformation and dislocation\nactivity. TEM image and diffraction patterns of the 10 wt.% Al alloy after 24% compressive strain,\nconfirming the formation of the new Mn Al phase and the presence of dislocations within the\n5 8\ntransformed particle, indicating plastic activity in the particles.\nFigure R1.7. Calculated free energy of Mn Al -B, Mn Al , and Mn Al phases: (A) Equilibrium\n4 11 4 11 5 8\nstates at 15℃-40℃; (B) Under applied strain at 20 °C. At 0% strain (equilibrium), Mn Al possess the\n5 8\nhighest energy than Mn Al -B, Mn Al , and cannot stable under equilibrium condition. Under applied\n4 11 4 11\nstrain (~2%), the energy of Mn Al -B rises above that of Mn Al , providing the thermodynamic driving\n4 11 5 8\nforce for the stress induced phase transformation observed experimentally.\nAt the gradient architecture, the phase transformation does not occur uniformly throughout a particle\nbut initiates in localized nanoregions where the stress threshold is first exceeded. This initial\ntransformation relieves the local stress concentration, effectively redistributing the stress to adjacent\nregions. As the global stress increases, these previously transformed regions then cooperate with\nneighboring areas to accommodate further deformation (schematic in Figure R1.5). This sequential,\nnanoscale process enables a continuous and efficient mitigation of stress concentrations across the\ninterface, effectively preventing damage nucleation.\nMoreover, we observed unusual dislocations in the transformed particle regions under 24% strain.\nThis process, recently highlighted in Science (2024, 385, 422-427), significantly extends the\neffective dislocation glide path on the Mg matrix and enables additional plastic deformation within\nthe second phase itself—a phenomenon rarely observed in conventional composites. The combined\neffect of phase transformation-induced strain accommodation and dislocation activity in\nparticle leads to simultaneous gains in strength and ductility.\nCollectively, the synergy is therefore attributed to mitigating stress concentration and inhibiting\ndamage at the nanoscale gradient interfaces and particle deformation, enhancing strain hardening\nand extending ductility. To clarify this mechanism, we have added a dedicated section entitled\n“Mechanism of strength and ductility synergy” in the revised manuscript, accompanied by new\nexperimental evidence and a schematic illustration (Figure 5).\nFormation mechanism of gradient architecture\nOur thermodynamic/kinetic analysis further clarifies how Al diffusion and reaction govern gradient\nstructure formation (Figure R1.8). Free-energy calculations reveal that Mn Al -B is more stable at\n4 11\n520 °C (-248 KJ·mol-1) than other phases, such as Mn Al (-237 KJ·mol-1), Mn Al (-212 KJ·mol-\n4 11 2 3\n1) and Mn Al -B (-232 KJ·mol-1) explaining its preferential formation in Al-rich interfacial regions.\n2 3\nDuring hot pressing, Al atoms diffuse into Mn-B domains and react with Mn to form B-\nsupersaturated Mn Al -B. Upon cooling, B partially precipitates and segregate at Mn Al to form\n4 11 4 11\nthis metastable gradient architecture. While varying Al contents and applied pressure, different Mn-\nAl phases and gradient AlMn-B architecture can be created at the interface (Materials\nCharacterization 140 (2018) 1–8). In contrast, Al-free systems under high pressure benefit for MnB\nformation, which induces substantial residual strain in Mg matrix during cooling due to the\nsignificant modulus mismatch and thermal expansion discrepancy.\nFigure R1.8. Thermodynamic and kinetic pathways for gradient interface formation. Free energy\ndiagram comparing Mn Al , Mn Al -B, Mn Al , Mn Al -B and MnB from 0 °C to 520 °C. The inset\n4 11 4 11 2 3 2 3\nschematically illustrates the Al interdiffusion and reaction front progression during hot pressing, leading\nto B supersaturation and compositional grading.\nBeyond conventional interfacial reaction, the strong electronic correlations within Mn lattices drive\nlocal lattice distortions and magnetic domain formation during B partially precipitates and segregate.\nThese discontinuous distortions coupled with second-order magnetic transitions (Acta Materialia\n174 (2019) 227236; Nanomaterials 11 (2021) 896), stabilize gradient structures with varying atomic\nspacings while preserving lattice continuity. Such lattice-magnetic-structure correlations further\ncontribute to stress accommodation.\nWe have revised the manuscript accordingly and added a schematic illustration to clarify this\nformation mechanism: During hot-pressing at 520 °C, interfacial Al atoms rapidly diffuse into Mn-\nB, leading to the formation of a B-supersaturated AlMn-B phase. Upon cooling, this phase\nundergoes decomposition via B precipitation, while spatial variations in Al content and applied\nstress promote the development of a compositional gradient through coordinated B atomic\nsegregation. In the absence of Al, direct formation of MnB particles at the Mg/Mn-B interface\ninduces high residual strain due to mechanical and thermal mismatch. With limited Al addition,\nalthough stabled Mn Al enhances interfacial bonding, the induced localized stress concentration\n2 3\nstill promotes MnB precipitation and residual strain accumulation in the Mg matrix. At higher Al\ncontents, however, the high-temperature Mn Al -B phase decomposes under cooling through a\n4 11\ndefect-mediated restructuring process. Enrichment of solute B atoms effectively suppresses MnB\nnucleation and facilitates discontinuous transformation within the Mn Al matrix, dissipating\n4 11\ninterfacial energy and minimizing microcracking. Consequently, in the 10 wt.% Al alloy, MnB is\neliminated, strain concentration is absent, and no zero-thermal-expansion regime is observed. These\nresults collectively underscore the role of the nanoscale gradient structure in achieving interfacial\nstability and thermodynamic equilibrium.\nIn above, our work demonstrates that nanoscale gradient structures release interfacial stress through\nstress-induced phase transformations and further dislocation activity in the secondary phase\nparticles. This dual pathway not only clarifies the origin of the strength-ductility synergy but also\nhighlights the novelty of gradient interface design as a strategy for damage mitigation in Mg alloys.\nWe appreciate the reviewer’s constructive feedback, which has help us to deepen our mechanistic\nanalysis and to present these findings with the rigor and clarity required for broader generalization.\n(4) The selection of WE43 alloy as the model material lacks justification. Its high rare-earth content\n(Y, Nd) contradicts the lightweight objective, and alternative Mg-RE alloys with potentially superior\nstrength-to-weight ratios exist but were not considered. Furthermore, tensile properties of the\nmodified WE43 alloy are absent, limiting assessment of practical relevance.\nAuthor reply:\nWe thank the reviewer for raising these important points regarding the choice of WE43 and tensile\nproperties.\nJustification for WE43 Selection\nThe WE43 alloy was selected based on its industrial applicability and processing advantages rather\nthan solely lightweight performance. Although it contains rare-earth elements (Y, Nd), which\nincrease density relative to pure Mg, the choice is justified by three practical considerations:\nProcessing Safety: WE43 alloy powder offers significantly higher ignition resistance and\n\nan elevated melting point compared to pure Mg, greatly enhancing safety during high-\ntemperature processing such as hot pressing, sintering and additive manufacturing.\nIndustrial Relevance: WE43 alloy is a commercially mature alloy widely used in aerospace\n\nand additive manufacturing, making it a representative system for studying and mitigating\nthermally induced deformation and cracking through thermal expansion and mechanical\nproperty control.\nResearch Focus: This work aims to demonstrate a novel approach for concurrently\n\nimproving thermal expansion and mechanical performance in a commercially viable Mg\nalloy, rather than screening alternative Mg–RE systems.\nWhile the addition of rare-earth elements slightly increases density (1.84 g/ cm3 vs. 1.74 g/ cm3 for\npure Mg), WE43 still retains the lightweight advantage intrinsic to Mg alloys. To clarify this point,\nwe have revised the Introduction to emphasize that optimizing the thermos and mechanical stability\nof WE43 is crucial for mitigating cracking during manufacturing and service.\nAs suggested, we have supplemented relevant background information in the introduction: The\nWE43 alloy was selected as a representative model system to address its inherent susceptibility to\nthermal cracking during processing, a key limitation hindering its broader application31–33.\nTensile Properties\nRegarding the tensile properties, we have added the tensile data of the samples, as shown in Figure\nR1.9. The tensile performance improvement is less significant than the compressive performance,\nwhich reflects the typical tension-compression asymmetry of magnesium alloys due to their hcp\ncrystal structure. Nevertheless, compared with the pure WE43 alloy, the tensile strength of the alloy\nwith 10 wt.% Al addition increased to 250 ± 5 MPa, and the elongation decreased slightly to 8 ±\n0.5 %. These data provide a more comprehensive mechanical characterization and confirm the\npractical relevance of the material under different loading conditions.\nWe have revised the manuscript accordingly and added tensile properties to confirm the practical\nrelevance: The tensile performance improvement is less significant than the compressive\nperformance (Figure S10), which reflects the typical tension-compression asymmetry of magnesium\nalloys due to their hexagonal close-packed crystal structure. Nevertheless, compared with the WE43\nalloy containing only 7 wt.% Mn-B, the addition of 10 wt.% Al raises the yield tensile strength from\n157 ± 4 MPa to 204 ± 2 MPa, while maintain an elongation of 8 ± 0.5 %, which is a notable\nenhancement compared with the nearly complete loss of ductility observed without Al addition\n(Figure S11).\nFigure R1.9. Tensile properties of the WE43 based composites. (A) Engineering stress-strain curves\nof the WE43 alloy, and the composite with different Al content in 7 (Mn-B)/WE43. (B) Comparison of\nkey tensile properties (ultimate tensile strength and elongation to failure) between these materials.\n(5) The critique of existing CTE-reduction strategies (alloying, processing) causing weak interfaces\nis undermined, as the authors' own solution relies fundamentally on composition control (Al\naddition) and process optimization (diffusion control). The claimed \"gradient structure\" appears\nlargely observational rather than a distinct design strategy.\nAuthor reply:\nWe sincerely thank the reviewer for this insightful critique, which provides us with the opportunity\nto clarify both the conceptual foundation and the distinctiveness of our strategy. We agree that\ncomposition and processing are indispensable tools in materials design; however, our approach\ntranscends conventional alloying or processing routes by harnessing intrinsic material physics to\ndrive self-organization of a functional meta stabled gradient interface, rather than merely achieving\nequilibrium phases or homogeneous structures.\nOur initial aim was to employ in-situ synthesis of Mn-B compounds with Al addition to regulate\nstress distribution at the Mg/Mn interface, thereby improving structural compatibility. The choice\nof Al was intentional, owing to its high solubility and mobility in both Mg and Mn, which enhances\ninterfacial wettability. Unexpectedly, we found that Al preferentially diffuses into Mn and triggers\nthe formation of Mn-Al intermetallics. This reaction is facilitated by B segregation and stabilized\nby Mn’s magnetic structure, which promotes local ordering. The outcome is not a homogeneous\ninterfacial reaction layer, but a chemically and structurally graded interface.\nThis serendipitous finding prompted us to systematically explore Al content (5 wt.% vs. 10 wt.%)\nand optimize hot-pressing conditions, leading to reproducible gradient formation. Importantly, these\ngradients were not mere observational artifacts but were confirmed to correlate with macroscopic\nproperties—including phase evolution, thermal expansion suppression, and enhanced mechanical\ncompatibility. In this regard, our contribution is not simply composition/process optimization, but\nthe identification, mechanistic understanding, and deliberate exploitation of defect-diffusion-\nmagnetism interplay to enable self-organized gradient architectures.\nWe acknowledge that this strategy was initially uncovered through serendipity, but the novelty lies\nin recognizing and systematizing the underlying physics into a new design principle: guiding self-\norganization to achieve functional gradient interfaces. This principle advances beyond conventional\nCTE-reduction approaches (e.g., random alloying, particle dispersion, or processing-induced\ntextures) by establishing a pathway where interfacial compatibility is achieved in situ through\ncontrolled local transformations. We have expanded the revised manuscript to explicitly highlight\nthis distinction, and we are grateful to the reviewer for encouraging this clarification.\n(6) Claims that the Mn-Al-B gradient \"effectively releases interfacial stress\" lack direct\nexperimental validation (e.g., interface stress mapping). Characterization of as-sintered interface\nstress states is also missing.\nAuthor reply:\nWe thank the reviewer for this important comment and fully agree that interfacial stress\ncharacterization is crucial. While direct measurement of interface stresses is challenging, we have\nadopted a multi-technique approach that provides consistent and compelling indirect evidence that\nthe Mn-Al-B gradient architecture effectively mitigates interfacial stress.\nFirst, XRD measurements reveal only negligible peak shifts in the 10 wt.% Al sample, concomitant\nwith its less reduced CTE, indicating a lower level of residual strain. This observation suggests that\nthermal stresses are effectively relaxed during the formation of the AlMn-B gradient architecture.\nSecond, dislocation-related features provide further insight. TEM characterization shows a notable\nreduction of dislocation arrays in the 10 wt.% Al alloy (Figure R1.3). Geometric phase analysis\n(GPA, Figure R1.10) further confirms that while the 0 wt.% Al alloy exhibits pronounced local strain\nfields and interfacial dislocations, the 10 wt.% Al alloy exhibits no discernible interfacial strain and\nessentially no dislocations at the particle-matrix boundary. These findings strongly imply that the\ngradient transition alleviates stress concentrations.\nFigure R1.10. Nanoscale strain analysis at the particle-matrix interface via Geometric Phase\nAnalysis (GPA). HRTEM images of the particle-matrix interface in the alloys without Al (A) and with\n10 wt.% Al (D). According to corresponding GPA E strain maps, A pronounced strain field is present\nXX\nat the sharp interface in the alloy without Al (B), while the 10 wt.% Al alloy shows a significantly reduced\nstrain field (E). Inverse FFT analysis reveals a higher density of dislocations at the interface in the 0 wt.%\nAl sample (C) compared with the nearly dislocation-free interface in the 10 wt.% Al sample (F).\nThird, we evaluated our system against other negative thermal expansion (NTE) particles. EBSD-\nbased kernel average misorientation (KAM) analysis (Figure R1.11) demonstrates that composites\nreinforced with LaFe Co Al particles, which is known to negative thermal expansion, show\n9.8 1.1 2.1\ndistinct strain localization (Appl. Phys. Lett. 127 (2025) 081901). In sharp contrast, the AlMn-B\ngradient architecture exhibit significantly suppressed strain localization and a lower density of\ngeometrically necessary dislocations (GNDs).\n[Redacted]\nFigure R1.11. EBSD based analysis of strain localization at particle interfaces. (A, B) Inverse pole\nfigure (IPF) map and kernel average misorientation (KAM) map for this work. (C, D) Inverse pole figure\n(IPF) map and kernel average misorientation (KAM) map for a reference Mg composite containing\nLaFe Co Al particles (Applied Physics Letters 127 (2025) 081901). The KAM maps, which reflect\n9.8 1.1 2.1\nthe density of geometrically necessary dislocations (GNDs), show strong strain localization around\nLaFe Co Al particles but significantly weaker localization in the AlMn-B gradient composites,\n9.8 1.1 2.1\nindicating effective interfacial stress release.\nFinally, we note that characterization of “as-sintered” interface stress is not meaningful in this\npowder-consolidated system: the relevant stress state emerges only after interfacial bonding during\nsintering.\nTogether, these complementary results (XRD, TEM/GPA, EBSD/KAM) consistently demonstrate\nthat the AlMn-B gradient architecture effectively releases interfacial stress. We have revised the\nmanuscript to incorporate these data and references, and we thank the reviewer for prompting us to\nstrengthen this point.\n(7) Figure 2's structural analysis must explicitly state the Al content (5 wt.% Al or 10 wt.% Al?).\nCharacterization of the 5 wt.% Al alloy (reportedly better CTE) is needed to determine if the\ngradient structure exists and why its mechanical properties differ.\nAuthor reply:\nWe thank the reviewer for highlighting the need for clarification. In the revised manuscript, we have\nexplicitly stated that the structural analysis shown in Figure 2 corresponds to the 10 wt.% Al alloy.\nTo further address this point, we first conducted comparative SEM-EDS and TEM-EDS mapping\non both the 5 wt.% Al and 10 wt.% Al alloys (Figure R1.12). The results demonstrate that Al\ndiffusion into Mn-B particles is highly sensitive to overall Al content: in the 10 wt.% Al alloy, Al\npenetrates more deeply into Mn-B particles, producing a well-developed gradient structure, whereas\nin the 5 wt.% Al alloy, the diffusion depth is relatively shallow. The particle size and local stress\nstate strongly influence diffusion heterogeneity, suggesting that both thermal dynamic driving force\nand stress-strain fields govern gradient formation. Moreover, a significant reduction in pore density\non the Mg matrix is observed as the Al content increases to 10 wt.%.\nFigure R1.12. Comparative EDS analysis of Al diffusion depth. SEM images and corresponding\nAl elemental maps of the 5 wt.% Al (A) and 10 wt.% Al (C). TEM images and corresponding Al\nelemental maps of the 5 wt.% Al (B) and 10 wt.% Al (D). Line profiles of Al concentration across\nthe particle-matrix interface, demonstrating deeper Al penetration and a more pronounced\nconcentration gradient in the 10 wt.% Al alloy.\nTEM analysis of the 5 wt.% Al alloy reveals the presence of Mn Al interfacial phases (Figure\n2 3\nR1.13). While these phases indicate the onset of gradient formation, the depth of continuous gradient\narchitecture is longer and no formation of meta stable AlMn-B, such as Mn Al -B, as compared\n4 11\nwith the 10 wt.% Al alloy. This distinction accounts for the differences in thermal and mechanical\nperformance. In the 5 wt.% Al alloy, limited Al diffusion leaves residual strain in the Mg matrix,\nresults in fewer AlMn-B phases and mainly formation of MnB phase. Consequently, stress\nredistribution via interfacial phase transformation and crack-arresting toughening are less effective,\nleading to lower strength and ductility. By contrast, the 10 wt.% Al alloy develops abundant meta\nstable Mn Al -B phases that not only enhance load transfer but also undergo stress-assisted\n4 11\ntransformations, thereby improving both thermal stability and mechanical performance.\nFigure R1.13. Interfacial microstructure of the 5 wt.% Al alloy. (A) Low-magnification TEM image\nof the interface. (B) HRTEM image of the region marked in (A), showing the interface between the Mg\nmatrix and a Mn Al -type phase. (C) Corresponding lattice and FFT patterns of the Mn Al phase. (D-\n2 3 2 3\nF) Lattice spacing profile across the interface, revealing a gradient structure in 5 wt.% Al alloy.\nWe have revised the manuscript accordingly to explicitly indicate the alloy compositions in Figure\n2, and incorporate these comparative results and interpretations for the 5 wt.% Al alloy: In the 5 wt.%\nAl sample, regions without interfacial cracks reveals that the interfacial formed Mn Al phase\n2 3\npreserves gradient-varying interfacial contact, suggesting the AlMn-B phase maintains strong\ncohesion with the Mg matrix and mitigates brittle interfacial fracture. This further supports our\nexpectation that varying Al contents induce distinct phase formations.\nWe sincerely thank the reviewer for encouraging us to provide this essential data, which\nsubstantially strengthens the mechanistic depth of our work.\nReviewer #2 (Remarks to the Author):\nIn the manuscript, the “gradient-ordered architecture” is more a descriptive term than a rigorously\ndefined phenomenon. The authors should provide a schematic or mechanistic diagram clearly\noutlining the Al diffusion process, MnB → Mn Al transition, and how this sequence leads to\n4 11\ninternal stress accommodation and thermal expansion suppression. Without this, the structural\ninterpretation lacks the clarity required for broader generalization.\nAuthor reply:\nWe sincerely thank the reviewer for this valuable suggestion. We agree that a clearer visual and\nmechanistic representation of the gradient ordered architecture and its formation process is essential\nto convey the novelty and generality of our approach.\nWe have now included a new schematic diagram (Figure R2.1) in the revised Supplementary\nInformation, which explicitly illustrates the key stages of the structural evolution and the underlying\nmechanism. Specifically, we analyzed the structural evolution in both matrix and phase evolution\nby comparation the 7 wt.% Mn-B added alloys containing 0 wt.% Al, 5 wt.% Al, and 10 wt.% Al,\ndetailed descriptions are provided in our response to the fourth comment. For the 0 wt.% Al alloy,\nMn-B at the interface tends to transform into MnB during hot pressing. MnB is a hard phase, and\nupon cooling, the mismatch between the rigid interface and the contracting Mg matrix generates\nsignificant residual strain. In the 5 wt.% Al alloy, limited Al diffusion results in partial formation of\nAlMn-B alongside stable Mn Al and MnB, creating a heterogeneous interface with moderate stress\n2 3\naccommodation and significant residual strain. A nascent gradient exists, but less effective in stress\naccommodation due to the absent of meta stabled AlMn-B phase. With sufficient Al in the 10 wt.%\nAl sample, extensive diffusion promotes AlMn-B formation at high temperature. Upon cooling, B\nparticipation and segregation in AlMn-B phase could balance the local thermal stress, suppressing\nMnB formation and enabling a nanoscale, compositionally graded meta stabled architecture. This\ngradient structure accommodates thermal and lattice mismatch through spatially varying elastic\nfields and B distributions, allow sequential stress-induced transformations and localized atomic\nmigration, thereby alleviating mismatch strain at the particle-matrix interface.\nFigure R2.1. Schematic illustration of Al diffusion and reaction induced interfacial phase evolution.\n(A) Comparison of alloys containing 0 wt.% Al, 5 wt.% Al, and 10 wt.% Al. In the 0 wt.% Al alloy, Mn-\nB transforms into rigid MnB, producing a sharp Mg/MnB interface with large residual strain. In the 5\nwt.% Al alloy, partial Al reaction yields Mn Al and MnB phases and residual strain in the Mg matrix.\n2 3\nIn the 10 wt.% Al alloy, extensive Al stabilizes Mn Al -B and Mn Al , while suppressing MnB,\n4 11 4 11\nresulting in a continuous gradient-ordered architecture due to the segregation of B atoms and the strain\naround. (B) Schematic representation of the formation process of multi-gradient ordered architecture\nthrough the modulation of Al content.\nThe above content combined with free energy calculation results for multiple phases have been\nschematic summarized in Figure R2.2 and integrated into the main text. For completeness and to\navoid redundancy, detailed descriptions are provided in our response to Reviewer 1’s third comment,\nwhile this section emphasizes the mechanistic linkage between phase evolution, strain\naccommodation, and thermal expansion behavior.\nHerein, we specifically investigated the CTE changes of these materials from room temperature to\nMnB phase transition temperature, as well as the evolution of phase transition temperature ranges.\nIt can be observed that the CTE of the 5 wt.% Al sample in the low-temperature region is\nsignificantly lower than that of the pure WE43 sample and also lower than the sample without Al\naddition. This CTE reduction is attributed to the presence of strain in the Mg matrix, i.e., a large\nnumber of internal defects (such as dislocations or solute atoms). During the heating process, strain\nrecovery (also known as residual strain recovery) occurs. The volume contraction caused by this\nstrain\nrecovery—mainly the reduction in volume due to the disappearance of defects—partially\ncompensates for the volume expansion induced by thermal expansion. Consequently, the net\nvolumetric change is reduced, resulting in a smaller apparent CTE. The greater the reduction in CTE,\nthe larger the amount of residual strain retained in the material. Therefore, the observed suppression\nof CTE in the 5 wt.% Al sample before MnB phase transformation can be attributed to the presence\nof substantial residual strain in the Mg matrix. In the 0 wt.% Al alloy, residual strain recovery in Mg\nmatrix and the negative thermal expansion of MnB partially offset positive CTE, producing a near-\nzero expansion within the transition window. In the 5 wt.% Al alloy, higher MnB content and greater\nresidual strain broaden the near-zero CTE window due to interactions between soft and rigid phases.\nThe well-bonded interface enables the strain generated during particle phase transition to induce Mg\nmatrix deformation, which further extends the temperature window via deformation recovery. For\na more comprehensive discussion on the correlation between strain and its effect on thermal\nexpansion, please refer to our previous study (Appl. Phys. Lett. doi: 10.1063/5.0284514) and our\nongoing unpublished work. In the 10 wt.% Al alloy, reduced residual strain and suppressed MnB\nformation limit compensatory effects, resulting in a less pronounced CTE reduction and continuous\nthermal deformation.\nThese results, together with the new schematic diagrams (Figure R2.1) and thermodynamic and\nkinetic calculations (Figure R2.2), provide a coherent mechanistic framework that explicitly links\nAl diffusion → interfacial phase transition (Mn-B → Mn Al -B) → stress accommodation →\n4 11\nthermal expansion suppression. We have included these analyses in the revised manuscript and\nSupporting Information.\nWe appreciate the reviewer’s constructive feedback, which has enabled us to significantly\nstrengthen the mechanistic clarity and general relevance of our work.\nFigure R2.2. Formation mechanism of the compositional gradient ordered architecture. (A) 3D\natom map reconstruction of the Mg/Mn-Al interface, showing the distribution of Mg, Al, Mn and B atoms.\n(B) One-dimensional compositional profile across the Mg/Mn-Al interface, illustrating the nanoscale (2-\n5 nm) co-segregation and perfect overlap of B, Mn, and Al elements. (C) Free energy profiles of MnB,\nMn Al , Mn Al , Mn Al -B and Mn Al -B phases range of temperatures, illustrating their\n4 11 2 3 2 3 4 11\nthermodynamic stability and interactions within the system. The inset schematically illustrates the Al\ninterdiffusion and reaction front progression during hot pressing, leading to B supersaturation and\ncompositional grading.\n1.Lines 123-124\n“Notably, at an Al concentration of 10 wt.%, the alloy attained a compressive strength 123 of 500\nMPa accompanied by an elongation of 24%” The error values should be presented.\nAuthor reply:\nWe sincerely thank the reviewer for this valuable suggestion. The error values for both compressive\nstrength and elongation have now been statistically analyzed (Figure R2.3) and are included in the\nrevised manuscript. Error bars have been added to the relevant figures (the newly provided\nSupplementary Figure S5). The original data from the same batch of samples showed a compressive\nstrength of 507 ± 8 MPa and an elongation of 23.8 ± 0.6%.\nThe description in the main text has been revised accordingly to: Notably, at an Al concentration of\n10 wt.%, the alloy attained a compressive strength of 507 ± 8 MPa accompanied by an elongation\nof 23.8 ± 0.6%, indicative of significant plasticity improvements in the Mg alloy.\nFigure R2.3. Compressive properties of the samples. Compressive strength and elongation of\nalloys with different Al contents. The 10 wt.% Al alloy exhibits the highest performance, with a\ncompressive strength of 507 ± 8 MPa and an elongation of 23.8 ± 0.6%.\n2.Lines 93-95\nGeometric phase analysis could effectively reveal the strain distribution across the Mg/AlMn-B\ninterface.\nAuthor reply:\nWe thank the reviewer for this valuable suggestion regarding the use of Geometric Phase Analysis\n(GPA) to probe strain distribution across the Mg/AlMn-B interface.\nWe performed GPA on high-resolution TEM (HRTEM) images acquired across the Mg/Mn-B\ninterface in both the 0 wt.% Al and 10 wt.% Al alloys. The results, now presented as Figure R2.4,\nprovide direct evidence supporting our mechanistic interpretation. The GPA strain mapping (Figure\nR2.4 B, E) reveals that the strain field in the Mg matrix adjacent to the interface in the 10 wt.% Al\nalloy is significantly lower than that observed in the 0 wt.% Al alloy. Furthermore, a detailed\nexamination of the HRTEM images and corresponding inverse Fast Fourier Transform (IFFT)\nanalyses (Figure R2.4 C, F) shows a high density of dislocations in the Mg matrix at the interface\nin the 0 wt.% Al alloy, whereas no dislocations are observed at the interface in the 10 wt.% Al alloy.\nThis key observation aligns perfectly with our hypothesis that the dissolution of Al and the in-situ\nformation of a nano scale compositionally graded meta stabled phases at the interface effectively\naccommodates and releases the localized strain that would otherwise be concentrated at interfaces,\ncausing the deformation in the soft matrix or the interface failure and crack propagation. While in\nthe 0 wt.% Al alloy, the large thermal and structural mismatch of the soft Mg matrix and hard, high-\nmodulus MnB/Mn-B phases, generates substantial residual thermal and back stresses upon cooling,\nleading to the pronounced strain accumulation and dislocation generation in the Mg matrix, as\nconfirmed by our GPA and defect analysis. Therefore, this comparation indicates the nanoscale\ngradient interface architecture prevents strain accumulation.\nFigure R2.4. Strain analysis at the interface via Geometric Phase Analysis (GPA). HRTEM images\nof the particle-matrix interface in the alloys without Al (A) and with 10 wt.% Al (D). According to\ncorresponding GPA E strain maps (denoted as yellow box in A and red box in D), a strain field is\nXX\npresent at the Mg matrix in the alloy without Al (B), while the 10 wt.% Al alloy shows a significantly\nreduced strain field (E). Inverse FFT analysis reveals a higher density of dislocations in the 0 wt.% Al\nsample (C) compared with the nearly dislocation-free in the 10 wt.% Al sample (F).\nWe have incorporated this new GPA analysis and the associated discussion into the revised\nmanuscript, emphasizing that the nanoscale gradient interface is pivotal for mitigating strain\nconcentration. “In the Mg matrix adjacent to the interface, geometric phase analysis (GPA) and\ndislocation density comparisons between 0 wt.% and 10 wt.% Al samples reveal a stark contrast\n(Figure S6). The former displays pronounced interfacial strain concentrations and high dislocation\ndensity, whereas the latter exhibits a near dislocation-free interface with dramatically reduced strain.”\nWe appreciate the reviewer’s constructive suggestions, which has helped us strengthen the strain\ndistribution on the interface.\n3.Lines 149-150\nThe analysis of the influence of gradient-ordered architecture on compressive strength and ductility\nremains superficial. Additional experiments or simulations are necessary to elucidate the\nmechanisms. For instance, microstructure characterization of the tensile sample with different\nstrain to reveal strengthening effects of the gradient-ordered architecture during compressive\ndeformation.\nAuthor reply：\nWe thank the reviewer for this important comment. We have conducted additional microstructural\ncharacterization on deformed samples and supplemented our analysis with free-energy calculations\nto clarify the strengthening and toughening mechanisms of the gradient-ordered architecture.\nWe conducted microstructural characterization of the 10 wt.% Al sample subjected to 24% strain\nand identified the presence of Mn Al and Mn Al phases in different regions. GPA analysis further\n5 8 2 3\nrevealed that the strain in the Mg matrix surrounding Mn Al particles was significantly higher than\n5 8\nthat around Mn Al -B particles (Figure R2.5). Thermodynamic calculations confirm this stress-\n4 11\ninduced phase transformation mechanism (Figure R2.6). While Mn Al has higher free energy than\n5 8\nMn Al -B and Mn Al at equilibrium (0% strain), its energy becomes lower than Mn Al -B under\n4 11 4 11 4 11\n~2% applied strain, providing thermodynamic impetus for this stress-driven transition. Moreover,\nwe observed unusual dislocations in the transformed particle regions under 24% strain (Figure R2.7\nA-D). This process, recently highlighted in Science (2024, 385, 422-427), significantly extends the\neffective dislocation glide path on the Mg matrix and enables additional plastic deformation within\nthe second phase itself. We thus deduced that the combined effect of newly formed phase\ntransformation induced strain accommodation and dislocation activity in particle leads to\nsimultaneous gains in strength and ductility. The schematic mechanism of the strength-ductility\nsynergy is shown in Figure. R2.7 F and its detailed descriptions are provided in our response to\nReviewer 1’s third comment.\nFigure R2.5 Evolution of the gradient-ordered architecture under compressive strain. Under\nthe condition of 24 % strain, different MnAl structural phases were observed in different strain\nregions, (A-D) the larger strain region of Mg base, and Mn Al structure was observed. (E-H) a\n5 8\nsmall strain region of Mg matrix, and Mn Al structure is observed. (I-L) the Mn Al -Bphase\n2 3 4 11\nobserved without strain.\nNotably, compositional gradients at newly formed AlMn-B phases, arising from Al content\nvariations and B segregation, establish multiple potential transformation pathways. For instance,\nregions with lower Al content may undergo transformation to phases such as Mn Al . The\n2 3\nsegregation and diffusion of B atoms facilitate these continuous transitions, benefit for preserving\ninterfacial integrity. Besides, residual strain manifests distinct characteristics across different alloy\nsystems. In the 0 wt.% Al alloy, direct Mg/MnB interfaces exhibit strain localization, leading to\npremature failure. In the 5 wt.% Al alloy, formation of Mn Al and MnB improves interfacial\n2 3\nbonding and CTE, but leaves significant residual strains in the Mg matrix, accelerating crack\ninitiation. Only in the 10 wt.% Al alloy does the metastable Mn Al -B phase enable dynamic\n4 11\ninterfacial transformations that accommodate strain, resulting in a relatively strain free Mg matrix.\nIt is precisely this synergy that the relaxation of residual strain in the Mg matrix coupled with the\ndynamic response of the metastable Mn Al -B interface that enables the 10 wt.% Al alloy to achieve\n4 11\nits exceptional strength-ductility combination.\nFigure R2.6. Calculated free-energy pathway under applied strain. The free energies of Mn Al -B,\n4 11\nMn Al , Mn Al , and Mn Al under different strain conditions. At 0% strain (equilibrium), Mn Al\n4 11 2 3 5 8 5 8\npossess the highest energy than Mn Al -B, Mn Al , and cannot stable under equilibrium condition.\n4 11 4 11\nUnder applied strain (~2%), the energy of Mn Al -B rises above that of Mn Al , providing the\n4 11 5 8\nthermodynamic driving force for the stress-induced phase transformation observed experimentally.\nFigure R2.7. Stress-induced phase transformation and its role in strength and ductility synergy.\n(A-C) Under the condition of 24 % strain, the larger strain region of Mg matrix, and Mn Al structure\n5 8\nwas observed. (D) Inverse fast Fourier transform image shows the dislocations in the Mn Al particles,\n5 8\nsuggesting the deformation of the particles. (E) Calculated free-energy under applied strain. Mn Al\n5 8\npossess the highest energy than Mn Al -B, Mn Al , and cannot stable under equilibrium condition.\n4 11 4 11\nUnder applied strain (~2%), the energy of Mn Al -B rises above that of Mn Al , providing the\n4 11 5 8\nthermodynamic driving force for the stress-induced phase transformation (F) Schematic mechanism of\nthe strength-ductility synergy mediated by stress induced phase transformation and plastic activity in\ngradient architecture.\nConsequently, the compositional gradient architecture functions not merely as a passive stress\ndistributor but as an active energy dissipater. It undergoes progressive transformation under stress\nto absorb energy, delay crack initiation, and homogenize strain, as illustrated in Figure R2.7.\nWe have incorporated these new findings, including the high-strain microstructural characterization\nand thermodynamic calculations, into the revised manuscript:\nMicrostructural analysis of the 10 wt.% Al composites deformed to 24% strain reveals the\nemergence of Mn Al phase (Figure 5A-D). GPA shows pronounced strain localization within the\n5 8\nMg matrix adjacent to the Mn Al regions, while the original Mn Al -B phases maintain\n5 8 4 11\nsubstantially lower strain levels. (Figure S15). Moreover, dislocations observed within Mn Al\n5 8\nfurther confirm plastic activity in the secondary particles. First-principles calculations demonstrate\nthat applied strain drives the transition from metastable Mn Al -B to Mn Al (Figure 5E, Figure\n4 11 5 8\nS16), revealing a stress-induced phase transformation mechanism.\nBased on the observed microstructural evolution, we propose the following strengthening and\ntoughening mechanism (Figure 5F). Stress induced phase transformation and plastic activity in the\nsecondary particle enable high density dislocation storage for strengthening and work hardening50,\nwhile simultaneously improving ductility through enhanced dislocation mobility. The mitigation of\nbrittle interfacial fractures achieved by the nanoscale gradient ordered architecture, along with\nreduced residual matrix strain, facilitates mechanically synergistic particle-matrix interactions. This\nprocess is accompanied by various phase transformations mediated volumetric adjustments that\nreinforce interfacial cohesion and delay fracture initiation. Furthermore, dislocation activity within\nthe newly formed phases at higher strain levels provides additional enhancement of both strength\nand ductility27,51,52. This nanoscale compositional gradient architecture effectively converts\nconventional weak interfaces into active energy dissipation zones, ultimately enabling the\nexceptional strength-ductility synergy observed in this system.\nWe believe that this comprehensive analysis provides a deeper, mechanistic understanding of how\nthe gradient-ordered architecture enables the observed exceptional mechanical performance. We\nsincerely thank the reviewer for guiding us toward this important improvement.\n4.The discussion of the formation mechanism of the gradient ordered architecture requires further\nvalidation for its rigor and persuasiveness. The authors only investigated the interface structure of\nshould be characterized to further confirm the formation process of nanoscale multi-gradient\nordered architecture.\nAuthor reply:\nWe appreciate the reviewer’s insightful comment. To strengthen the validation of the formation\nmechanism of the nanoscale gradient ordered architecture, we further characterized the interfacial\nstructures of alloys with different Al contents (0, 5, 10 wt.% Al) sample (Figure R2.8).\nFor the 0 wt.% Al alloy, Mn-B and MnB are the dominant phases, the interfaces mainly contain Mn-\nrich regions without gradient structure. In the 5 wt.% Al alloy, limited Al diffusion into Mn-B\nparticles occurs, leading to the formation of Mn Al (Figure R2.9). The Al:Mn ratio in Mn Al is\n2 3 2 3\nlower than that in Mn Al , and the diffusion depth is significantly shallower than in the 10 wt.%\n4 11\nAl alloy. These results demonstrate that Al reaction within Mn-B particles strongly depends on the\noverall Al content, and diffusion heterogeneity is additionally influenced by particle size and local\nstrain environment, indicating that both chemical driving force and stress-strain state govern\ngradient formation.\nFigure R2.8. Comparation the interface microstructure of the 0 wt.% Al (A), 5 wt.% Al (B),\nand 10 wt.% Al (C) alloy. It can be found that the Al diffusion and reaction induced the formation\nof gradient-ordered architecture.\nFigure R2.9. Interfacial microstructure of the 5 wt.% Al alloy. (A) Low-magnification TEM image\nof the interface. (B) HRTEM image of the region marked in (A), showing the interface between the Mg\nmatrix and a Mn Al -type phase. (C) Corresponding lattice and FFT patterns of the Mn Al phase. (D-\n2 3 2 3\nF) Lattice spacing profile across the interface, revealing a gradient structure in 5 wt.% Al alloy.\nMoreover, free-energy calculations further confirm these experimental findings in Figure R2.6.\nMn Al is energetically more stable than Mn Al at low temperature, while Mn Al is favored at\n2 3 4 11 4 11\nelevated temperature. This implies that Mn Al tends to form during cooling, whereas the Mn Al -\n2 3 4 11\nB gradient structure observed at room temperature in the 10 wt.% Al alloy corresponds to a\nmetastable state. Similar to the role of C-enriched retained austenite in TRIP steels, this metastable\ngradient-ordered Mn Al -B structure accommodates thermal stresses during cooling, suppressing\n4 11\nthe transformation to MnB down to room temperature. Consequently, the metastable nanoscale\ngradient structure can be preserved by B segregation and subsequently undergo stress-induced\ntransformation during deformation, which provides a mechanistic basis for its strengthening and\ntoughening effects.\nAccordingly, we propose a formation mechanism to elucidate the development of this distinctive\narchitecture. In the absence of Al, MnB secondary particles form in-situ at the Mg/Mn-B interface\nunder the accumulated thermal stress during the hot pressing of Mn-B with WE43. Introducing a\nsmall amount of Al to improve the interfacial connection, these Al atoms predominantly concentrate\nat the interface and have a completely reaction to forming stable Mn Al phase, meanwhile, result\n2 3\nin the stress around the Mg matrix and the left unreacted Mn-B particles, benefit for the formation\nof MnB phase. During cooling, the unmatched new phases create severe lattice and thermal\nmismatch, leaving high residual strain in the softer Mg matrix. At higher Al levels, a high-Al\nMn Al -B phase forms at the interface under elevated temperatures; upon cooling, phase separation\n4 11\noccurs—driven by ambient stress, Al content, and cooling rate—yielding diverse structural phases.\nNotably, inhomogeneous B distribution leads to discontinuous precipitation with compositional\ngradients, which, coupled with local stress, results in a nanoscale gradient-ordered metastable phase\nstructure at the particle/matrix interface. Such metastable phases readily balance cooling-induced\nstress via interstitial atoms or deficiencies, reducing MnB formation and preventing stress\naccumulation in the Mg matrix, thereby dissipating interfacial energy and mitigating thermal strain\nbuildup. Consequently, MnB formation is eliminated, residual strain and interface crack is\nminimized, and no zero-thermal-expansion region is detected in the 10 wt. % Al sample. This\nindicates that the AlMn-B gradient architecture effectively releases interfacial stress, as\ncorroborated by observations that no strain concentration occurs near the Al-modified Mn-B\nparticles in KAM analysis.\nWe have added this systematic discussion, the new comparative figures (Figure R19 and\nsupplementary data), and the thermodynamic analysis to the revised manuscript: We thus proposed\nthe following formation pathway to elucidate the evolution of this distinctive compositional\ngradient–ordered architecture and its effect on Mg matrix (inset of Figure 4A, Figure S14). With\nlimited Al addition, although stabled Mn Al enhances interfacial bonding, the induced localized\n2 3\nstress concentration still promotes MnB precipitation and residual strain accumulation in the Mg\nmatrix. At higher Al contents, however, the high-temperature Mn Al -B phase decomposes under\n4 11\ncooling through a defect-mediated restructuring process. Enrichment of solute B atoms effectively\nsuppresses MnB nucleation and facilitates discontinuous transformation within the Mn Al matrix,\n4 11\ndissipating interfacial energy and minimizing microcracking. Consequently, in the 10 wt.% Al alloy,\nMnB is eliminated, strain concentration is absent, and no zero-thermal-expansion regime is\nobserved. These results collectively underscore the role of the nanoscale gradient structure in\nachieving interfacial stability and thermodynamic equilibrium.\nWe believe these additions have significantly strengthened the validation and persuasiveness of our\nproposed formation mechanism. Thank you again for this excellent suggestion.\n5. The manuscript only reports compressive properties, without any tensile data. Given that\nmagnesium alloys typically show strong tension-compression asymmetry, and that the introduced\nMn-B particles and Al-induced phase transformation (MnB → Mn Al ) likely create brittle\n4 11\ninterfaces prone to cracking under tension, the absence of tensile results is a major concern. Tensile\nperformance is essential for structural applications and should be presented or at least properly\ndiscussed.\nAuthor reply：\nWe thank the reviewer for this highly professional suggestion. We fully agree that Mg alloys exhibit\npronounced tension-compression asymmetry, with compressive properties generally superior to\ntensile ones.\nTo address this concern, we have performed tensile tests on the Mn-B containing alloys with and\nwithout Al addition. As shown in Figure R2.10, the Mn-B reinforced alloy exhibits a tensile strength\nof 157 ± 13 MPa and an elongation of 1.3 ± 0.4 %. With additional 5 wt.% Al, the tensile elongation\nis significantly improved, reaching 9.7 ± 0.2 %, the tensile strength is also increased to 221 ± 2.4\nMPa. When the Al content is added to 10 wt.%, the tensile strength is further improved to 250 ± 5\nMPa, and the elongation is 8 ± 0.5 %. This enhancement in tensile ductility provides strong\nexperimental evidence that the in-situ formed gradient-ordered Mn Al -B interface is far less brittle\n4 11\nthan the original sharp Mn-B/Mg interface. We attribute this improvement to two key mechanisms\nfacilitated by the gradient architecture: (1) The reaction of Al and Mn induces good interface\nintegration with the Mg matrix, which inhibits crack initiation. (2) the metastable nature of the\nMn Al -B phase allow for free residual strain in the Mg matrix and benefit for accommodate the\n4 11\nstrain and delay fracture during loading by various phase transitions and the plastic activity in the\nnewly formed secondary particles.\nFigure R2.10. Tensile properties of the WE43-based composites. (A) Engineering stress-strain curves\nof the WE43 alloy, and the composite with different Al content in 7(Mn-B)/WE43. (B) Comparison of\nkey tensile properties (ultimate tensile strength and elongation to failure) between these samples.\nWe have added a dedicated section discussing the tensile properties, the asymmetry with\ncompression, and the proposed mechanisms by which the gradient structure mitigates interfacial\nbrittleness under tensile loading: The tensile performance improvement is less significant than the\ncompressive performance, which reflects the typical tension-compression asymmetry of magnesium\nalloys due to their hcp crystal structure. Nevertheless, compared with the WE43 alloy with merely\naddition of 7 wt.% Mn-B, the yield tensile strength of the alloy with 10 wt.% Al addition increased\nfrom 157 MPa to 204 ± 2 MPa, and the elongation maintained at 8 ± 0.5 % as than almost loss\nwithout Al addition.\nWe sincerely thank the reviewer for encouraging us to provide this essential data, which\nsubstantially strengthens both the practical relevance and mechanistic depth of our work."}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors show that incorporating Mg-B particles and regulating interfacial diffusion in Mg\nalloys results in the formation of a nanoscale gradient-ordered Mn-Al-B structure. This\narchitecture effectively alleviates stress concentration, thereby achieving an exceptional strength-\nductility combination (500 MPa, 24% compressive strain) and low thermal expansion. This work\npioneers a novel \"atomic-level gradient-ordered engineering\" strategy for high-precision\nlightweight composites.\nThe author has meticulously addressed each of the concerns raised, leading to a substantial\nimprovement in the overall quality of the manuscript. The dataset is complete and rigorously\nvalidated, and the reasoning is logically sound. I therefore recommend that the manuscript be\naccepted for publication.", "reviewer_2": "Reviewer #1 (Remarks to the Author):\nThe authors show that incorporating Mg-B particles and regulating interfacial diffusion in Mg\nalloys results in the formation of a nanoscale gradient-ordered Mn-Al-B structure. This\narchitecture effectively alleviates stress concentration, thereby achieving an exceptional strength-\nductility combination (500 MPa, 24% compressive strain) and low thermal expansion. This work\npioneers a novel \"atomic-level gradient-ordered engineering\" strategy for high-precision\nlightweight composites.\nThe author has meticulously addressed each of the concerns raised, leading to a substantial\nimprovement in the overall quality of the manuscript. The dataset is complete and rigorously\nvalidated, and the reasoning is logically sound. I therefore recommend that the manuscript be\naccepted for publication."}, "author_response": "Authors reply:\nWe thank the reviewer for the insightful comments provided during the review process. We\nappreciate the reviewer’s acknowledgement of our efforts in thoroughly addressing all comments\nand strengthening the manuscript accordingly.\nWe sincerely thank the reviewer for recommending acceptance of the manuscript."}]

Decision letter