Breaking immiscibility barriers: ultrafast sintering of interlocked Cu-Fe-based composites

Domain: chemistry
Published: 2026-01-06
Editorial outcome: major_revision

Abstract
With increasingly stringent application environments, there is a growing demand for designing and fabricating of promising structural materials. Conventional metallurgical methods typically struggle to produce positive mixing enthalpy alloys with uniform microstructures and consistent properties due to mutual immiscibility and uncontrollable segregation during solidification. To this end, we employ a bottom-up approach to fabricate Cu-50 vol% Fe55Cr25Mo16B2C2 positive mixing enthalpy composites, utilizing ultrafast high-temperature sintering and instantaneous quenching. The method effectively prevents liquid segregation during the molten state and allows precise control over the size and distribution of the two phases. In addition, incorporated metallic glass particles as the hard phase additives, which can quickly form the localized multi-phase nanocrystals and establish a robust interlocking structure with pure copper during rapid sintering. Upon rapid quenching, the two phases solidify with a tight and conformal interface, while the elemental cross-diffusion and phase separation are mitigated. The mechanical tensile strength of the resulting composite is approximately 8 times greater than that of pure Cu. The wear resistance improves by 40-50 times, and the Vickers hardness increases by a factor of 10. The innovative sintering method provides a feasible pathway for developing and manufacturing of positive mixing enthalpy composite materials.

Introduction
For thousands of years, traditional alloy design and development have primarily focused on systems with negative mixing enthalpy1,2. Expanding the compositional range of positive mixing enthalpy alloy systems and controlling the microstructural arrangement of different phases could lead to the development of materials with characteristics previously unseen in alloy materials, offering immense potential3,4. Unlike negative mixing enthalpy alloy systems, which can easily achieve highly uniform solid solution structures5–7, positive mixing enthalpy systems face inherent challenges8,9. Due to the immiscibility of metals in the liquid phase and the prolonged heat treatment durations of conventional metallurgical methods, such as suction casting10, vacuum hot pressing11, high-pressure torsion12, or mechanical repetitive rolling13, the resulting positive mixing enthalpy systems tend to have severe phase separation and poor mechanical properties. Moreover, the temperature-dependent thermodynamic driving forces induce variations in elemental segregation kinetics, leading to uncontrolled heterogeneous microstructures of immiscible phases due to the instability of phase separation dynamics14,15. Despite decades of research, precisely controlling elemental segregation to achieve a uniformly distributed biphasic structure remains an unsolved challenge16,17.

To address this issue, introducing intermediate elements has become a widely adopted strategy for inducing stable solid solutions18,19. For instance, adding a third component (e.g., Ni or Co) can reduce the overall mixing enthalpy or stabilize metastable phases, leading to greater elemental mixing and reduced segregation20,21. Nevertheless, the strategy necessitates precise regulation of alloying element stoichiometry and spatial distribution, concurrent maintenance of thermodynamic equilibrium, and interfacial stability across multiphase systems. Alternatively, multiscale structural design and hierarchical heat treatment have also been employed in developing metals with positive mixing enthalpy22. Recent studies, however, have been limited to adding small amounts of positive mixing enthalpy elements (<0.5%) to either promote the precipitation of coherent nanophases (typically confined to phase diagram edges5,6) or form limited solid solutions23,24. While hypoeutectic doping leverages metastable phase diagrams to confine positive-enthalpy elements at grain boundaries, such compositional constraints inherently limit bulk-scale property engineering25,26.

Given these constraints, we herein employ a bottom-up fabrication strategy for positive mixing enthalpy alloys, incorporating precise microstructural regulation to suppress high-temperature phase segregation. Based on the ultrafast high-temperature sintering (UHS) technology27,28, we have introduced an ultrafast cooling system to enable rapid quenching during the final stage of sintering, achieving ultrahigh-temperature rapid sintering and quenching (UHSQ) (Supplementary Fig. 1a, b and Supplementary Movie 1). Combining ultra-high heating and quenching rates ensures a stable and controllable liquid-phase window in positive mixing enthalpy systems (Fig. 1a). This approach retains the high heating rate (500–1000 K/s) of the UHS process while providing an enhanced cooling rate (up to 1000 K/s at the sample surface; see Fig. 1b; Supplementary Fig. 2 and Supplementary Movie 2). Specifically, the precursor compact was wrapped in zirconia (ZrO₂) fiber cloth for electrical insulation (Supplementary Fig. 1c), positioned between two carbon felts in an argon atmosphere, rapidly sintered for 10 seconds, and then immediately quenched in high thermal conductivity silicone oil at 243 K (Fig. 1c). Employing the UHSQ technique, we successfully sintered a Cu-50 vol% Fe55Cr25Mo16B2C2 positive mixing enthalpy composite, denoted as Cu-Fe-based composite. Concurrent ultrafast quenching (~1 s) induces three critical microstructural evolutions: atomic-scale heterogeneity redistribution in metallic glass; dislocation substructure modification in copper; and nanoscale rivet-like interlocking at interfaces arising from differential thermal contraction coefficients.

Results
The exceptional hardness of metallic glass is a significant advantage, but the smooth surface poses challenges for interface attachment32 (Fig. 2a and Supplementary Fig. 4). Mechanical milling has been proven to be an effective method for enhancing the attachment between phases33. To verify this, we employed high-energy ball milling and observed notable changes in the covering degree between phases as a function of milling time (Fig. 2b, c and Supplementary Fig. 5a). It was evident that Fe55Cr25Mo16B2C2 metallic glass precursors subjected to high-energy ball milling exhibited significantly improved interfacial attachment with Cu particles. The continuous milling process not only introduces surface roughening into the Fe55Cr25Mo16B2C2 metallic glass itself but also ensures uniform coverage of Cu on the surface of metallic glass particles34, facilitating subsequent connection with the Cu phase during sintering (Supplementary Fig. 5b). Notably, the particle size of Fe55Cr25Mo16B2C2 metallic glass remains unchanged during high-energy ball milling. Additionally, the starting Fe55Cr25Mo16B2C2 metallic glass particles are generally large, highlighting a potential direction for future optimization by reducing the size of metallic glass particles (Fig. 2d).

To demonstrate the universality of our UHSQ technique, we fabricated Cu-Fe55Cr25Mo16B2C2 composite materials with a 1:1 volume ratio, where the Fe content reached 27.5%, far exceeding that of conventional Cu-Fe alloys. As a comparison, we also employed various sintering techniques to prepare samples. The XRD patterns of samples prepared via UHS and traditional furnace sintering were similar, displaying pure Cu phases (Fig. 2e left). Notably, in UHSQ specimens quenched within 1 second, weak crystalline peaks emerged (Fig. 2e right), suggesting partial devitrification of the amorphous matrix. In contrast, specimens subjected to a 10-second dwell time prior to quenching achieved near-complete crystallization, restoring the equilibrium Fe–Cr-Mo phase assemblage.

Optical and scanning electron microscopy images (Fig. 2f) demonstrate that conventional sintering (annealing at 1073 K for 2 h in Ar, followed by annealing at 573 K for 6 h) fails to mitigate elemental segregation in Cu-Fe₅₅Cr₂₅Mo₁₆B₂C₂ positive mixing enthalpy alloys, resulting in phase-separated Cu-rich and Fe-rich coarsened domains. This phenomenon primarily originates from thermocapillary-driven Marangoni convection during liquid-liquid phase separation, as described by the Cahn-Hilliard model35,36. In contrast, UHS leverages ultrafast sintering rates to suppress elemental segregation and control the crystallization of the metallic glass reinforcements. Upon softening, the viscosity of the metallic glass temporarily behaves like an adhesive, tightly bonding the two phases and effectively extending the processing window. UHS is ultimately limited by the intrinsic immiscibility of positive-enthalpy systems and subcritical cooling rates, permitting interfacial phase separation that generates porous weak interfaces (Supplementary Fig. 6a–c and Supplementary Fig. 7). This interfacial failure is directly linked to the sintering duration; prolonging the high-temperature exposure allows for the severe phase segregation demonstrated in Supplementary Fig. 8d. Minimizing the sintering time is therefore critical to kinetically trap the system and suppress elemental segregation (Supplementary Fig. 8a–c). Remarkably, UHSQ, with its ultrafast cooling rates, refines grain structures while retaining the advantages of UHS. Moreover, UHSQ precisely regulates the devitrification pathway of the amorphous matrix, generating hierarchical microstructures that integrate nanocrystals with high-density lattice defects (Fig. 2f right and Supplementary Fig. 6d). The extreme quenching rate triggers distinct nanoscale architectural interlocking through stress-induced transformation. The microscopic rivet-like architecture at the nanoscale firmly bonds the Cu-Fe interface while constraining the distribution of individual components, resulting in a homogeneous interlocked Cu-Fe-based composite structure (Fig. 3a, b).

The rapid quenching process induces the formation of a multiscale structure within the Fe-based metallic glass (Supplementary Fig. 9). Bright-field TEM imaging reveals a regular and uniform distribution of nanoscale crystal particles within the metallic glass (Fig. 3c). Figure 3b collectively demonstrates that the Fe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass devitrifies into a multiphase nanocrystal, in which nanocrystalline Fe–Cr ferrite forms the continuous skeleton of the microstructure. The FCC Fe–Cr matrix surrounds and mechanically pins two secondary constituents, Mo-rich nanocrystals and Fe-rich areas. Bcc Mo-rich nanocrystals populate junctions and grain corners in a quasi-periodic manner, occupying ~14 vol% of the microstructure (Supplementary Figs. 10 and 11a, b), surrounding each Mo grain is a shell of bcc Fe ferrite. Notably, the Fe-rich areas contain a mixture of pure Fe, partially retained amorphous phase (Supplementary Fig. 11c, e), and ordered B2 intermetallic phases that contribute to the overall intricate structure (Supplementary Fig. 11d). The complex arrangement indicates an incomplete devitrification process, resulting in a multiphase nanocrystalline structure with retained amorphous pockets. The overall microstructure exhibits a high density of interfaces and structural complexity. This intricate structural architecture, arising from the controlled devitrification of the initial metallic glass, significantly influences the material’s properties. These special nanocrystal structural features are critical for coherent Cu-Fe interfaces with periodic misfit dislocations, which geometrically resemble traditional mortise-tenon interlocking structures. The interlocking of different rivet structures creates a tight interface, with a strong attachment between them (Fig. 3a). Such microstructural control further contributes to the macroscopic uniformity of the composite. Furthermore, at different interface angles, no elemental cross-diffusion is observed, providing clear evidence that this technique effectively controls elemental segregation at the microscopic level (Fig. 3b right). Such unprecedented control over both structural hierarchy and elemental distribution redefines manufacturability limits for immiscible systems.

We systematically compared the mechanical properties of materials processed under different sintering conditions. The optimized UHSQ-processed composites demonstrate a remarkable and comprehensive enhancement of mechanical properties, spanning bulk behavior (tensile and flexural strength) and surface performance (hardness and wear resistance). Using a composite with 50% Fe₅₅Cr₂₅Mo₁₆B₂C₂ as an example, the ultimate tensile yield strength reached ~685 MPa, which is 6.5 times higher than pure Cu (Fig. 3d). Remarkably, this strength advantage persists at elevated temperatures, with the composite maintaining 290 MPa yield strength at 923 K (Fig. 3e). For comparison, we prepared control samples (Supplementary Fig. 12) with several key variations: reinforcement content, reinforcement type (pure Fe; Fe-based alloy), and high-temperature performance. All comparisons demonstrated significant improvements in mechanical performance originating from the unique, interlocking Cu-Fe interface, where continuous Cu-rich and Fe-rich phases form interpenetrating networks resembling rivet joints, a structure forged by the UHSQ method. These results emphasize the substantial performance enhancement achieved by UHSQ-processed high-Fe-content positive mixing enthalpy composites, including in high-temperature applications (Supplementary Fig. 13). The improvement is primarily attributed to the high strength of multiphase nanocrystals37.

SEM characterization reveals distinct morphological evolution between the samples prepared by UHS and UHSQ protocols (Fig. 3f, g). UHS-processed specimens exhibited compromised interfacial cohesion (average interfacial gap length: 2.1 ± 0.3 μm), accounting for their inferior mechanical performance (Supplementary Fig. 14). Suboptimal interfacial integrity originates from prolonged thermal relaxation during UHS processing, where thermal expansion mismatch induced interfacial decohesion. In contrast, UHSQ samples displayed tight interfacial bonding. The main reason is that the ultrafast cooling rate prevents the system from reaching equilibrium during the two-phase shrinkage process, enabling a strong attachment of different phases at high temperatures to be maintained at room temperature. The multiphase nanocrystals formed strong interfaces with Cu, acting as rivets during tensile testing. The high interfacial bonding strength provided by the rivet-like structures is best demonstrated by the material’s failure mechanism. Instead of debonding at the interface, failure occurs via brittle trans-particle fracture, where cracks propagate directly through the reinforcement phase itself (Supplementary Fig. 15). The behavior shows that the interface is stronger than the reinforcement material. The interfacial integrity is also reflected in the bulk mechanical properties shown in Supplementary Fig. 16. While the difference in elastic modulus between the UHS and UHSQ samples contributes to the overall stress-strain response, the strength is predominantly attributed to a multiphase synergistic strengthening mechanism forged by the UHSQ process. The enhancement is not due to a change in intrinsic material stiffness but is a direct reflection of the strengthening effect from the improved interface bonding achieved by rapid quenching.

Fractography of UHSQ specimens revealed a homogeneous distribution of equiaxed dimples in the Cu matrix (Fig. 3g), indicative of good ductility. This structural uniformity is confirmed by energy-dispersive X-ray spectroscopy (EDS) mapping of the same fracture regions, which shows no evidence of large-scale elemental segregation (Supplementary Fig. 17). The suppression of elemental diffusion is further corroborated at higher resolution by corresponding SEM and TEM-based EDS line scans performed on cross-sections (Supplementary Fig. 18). Overall, we believe the enhancement in mechanical strength of this structure is manifested at different levels. Macroscopically, the Cu-rich and Fe-rich phases form two distinct interwoven network structures, similar to a rivet configuration, imparting excellent strength and ductility to the sample. For the composite with a 1:1 volume ratio of Cu to Fe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass, the composite achieves 8.91% tensile ductility, while enhancing the material’s overall strength and performance. Also, UHSQ allows for the adjustment of the Cu-Fe ratio to fine-tune the overall strength and ductility of the composite material. At the microscopic level, a small amount of Fe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass is dispersed within the Cu phase, creating a dispersion-strengthening effect that further improves mechanical performance (Fig. 3h). Furthermore, at the interface, the Cu-Fe phases bond seamlessly with no gaps, and the distribution of phases is uniform. Additionally, crystallization of the Fe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass leads to the formation of multi-phase nanocrystals, which exhibit remarkable strength to further enhance the material’s overall performance. The synergistic combination of rivet-like reinforcement and dispersion strengthening offers a clear pathway for further performance optimization, which can be achieved by tailoring particle size and distribution. This multiscale design strategy, which combines macro-architectural control with nanoscale precipitation engineering, provides a viable method for developing high-performance hybrid composites with enhanced properties.

As shown in Fig. 4a, composites processed via UHSQ exhibit a significantly higher Vickers hardness than those prepared by UHS or pure Cu, achieving a high and consistent average hardness of 900 HV. The value is approximately three times higher than that of UHS-processed samples and eight times higher than pure Cu (Supplementary Fig. 19). The hardness significantly exceeds the prediction from a simple linear rule of mixtures, indicating powerful synergistic strengthening effects. The origin of this synergy can be understood from the microhardness distribution shown in Fig. 4b. The Fe-based reinforcement contains in-situ formed nanocrystalline intermetallic compounds and possesses a high intrinsic hardness of ~1100 HV, which provides a strong foundation for the composite’s overall performance. Furthermore, the co-continuous, interlocking network architecture imposes a strong mechanical constraint on the softer Cu phase (which has a hardness of only ~90 HV). The constraint is enabled by a robust interface, which exhibits a high average hardness of ~500 HV and ensures efficient load transfer during indentation (Supplementary Figs. 20 and 21). Significant work hardening occurs in the ductile Cu matrix during deformation, further increasing its contribution to the overall hardness. It is the combination of these effects that results in the composite’s hardness.

We summarized the hardness and electrical conductivity of Cu alloys, ceramic-metal composites, and metallic glasses reported in the literature22,40–55 (Fig. 4c). The positive mixing enthalpy composites synthesized via UHSQ exhibit superior hardness compared to conventional Cu alloys, approaching the strength of metallic glasses and ceramic-metal composites. The electrical conductivity ranges between 15 and 20% IACS, which is generally superior to metallic glasses or ceramic-metal composites. This property holds significant potential in certain specialized materials, such as the conductive rails of electromagnetic launchers and the tracks of electromagnetic catapult systems, which require conductive materials that possess ultrahigh hardness, heat resistance, and erosion resistance38,39. Moreover, the mechanical properties can be tuned by adjusting the content of Fe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass. Increasing the proportion improves strength, whereas decreasing it enhances ductility. This tunability highlights the potential for future optimization and broader applications.

We conducted wear tests on UHS and UHSQ-processed samples, which are compared to pure Cu. The tests applied a 10 MPa load, rotated at 500 r/min, with mass loss measured every 5 minutes. Fitted data curves show that UHSQ-processed composites demonstrate 40–50 times higher wear resistance than pure Cu (Fig. 4d). Samples with 25% and 50% Fe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass clearly exhibited strong Fe-based networks, contributing to greater macroscopic strength (Fig. 4e). Further SEM characterization of UHSQ-processed samples before and after 15 minutes and 60 minutes of wear testing revealed surface changes (Fig. 4f). The multi-phase nanocrystals showed minimal structural degradation, demonstrating their contribution to the enhancement of wear resistance. Strong interfaces and robust attachment also play a critical role in improving tribological performance. Photographic evidence of wear damage under identical test conditions highlights the superior wear resistance of UHSQ-processed composites compared to sintered pure Cu (Supplementary Fig. 22). This study demonstrates an application view of positive mixing enthalpy composite materials. The seamless interlocking architecture with strong interfaces opens substantial possibilities for next-generation sliding electrical contacts, high-load tribological components, and thermally stable conductor substrates. It addresses critical needs in sustainable energy systems, heavy machinery, and aerospace electronics, where concurrent electrical and mechanical robustness is essential.

Our method provides valuable guidance for designing advanced composite materials and highlights the potential of positive mixing enthalpy composites for structural and functional applications. The ability to tailor mechanical properties through compositional adjustments underscore the flexibility and scalability of this approach. In terms of scalability, the process’s fast speed enables high-throughput manufacturing of large-area sheets and coatings, with current capabilities extending to near-centimeter thick bulk components; the primary focus for future scaling is addressing the engineering challenge of uniform quenching in even larger volumes. At its core, the method’s success lies in the microscale rivet-like structure induced by rapid quenching, which produces an unprecedented strengthening effect at the interface. The approach produces homogeneous positive mixing enthalpy composites with a well-balanced structure, successfully controlling phase segregation while preserving both toughness and hardness. The universality of this approach suggests that leveraging its advantages in future material design will significantly expand the boundaries of composite materials synthesis and development.

We have developed a bottom-up approach using the UHSQ process to fabricate positive mixing enthalpy composites that uniquely combine the excellent ductility of pure Cu with the ultrahigh strength of hard alloys. The process leverages controlled partial crystallization during rapid heating to accelerate densification, while subsequent rapid quenching is then critical to preserve a key portion of the amorphous phase in the final microstructure. The strength of the composite originates from the robust, rivet-like interlocking interface forged between the Cu matrix and the multi-phase nanocrystals. Our interlocked Cu-Fe-based composite overcomes the traditional strength-ductility trade-off, achieving a good combination of electrical conductivity (15–20% IACS) and ultrahigh hardness (900 HV). This work demonstrates the tremendous potential of positive mixing enthalpy composites for advanced structural and functional applications. Our approach facilitates the development of previously inaccessible immiscible alloy systems. The good balance of properties makes the material highly promising for demanding applications, including advanced mechanical components, aerospace systems, and particularly for electromagnetic propulsion and high-performance tribological components. Ultimately, the UHSQ process preserves the unique properties of the constituent phases, significantly expanding the synthesis boundaries for advanced metallic composites.

Methods
Synthesis
The Cu-coated Fe55Cr25Mo16B2C2 metallic glass precursor was prepared by mixing Cu (99%, Aladdin) and metallic glass (99% Yijinxc) in a 1:9 mass ratio using ZrO₂ as the grinding medium. High-energy ball milling was performed for 20 hours. The prepared coated precursor was then mixed with Cu (99%, Aladdin) in a 1:1 volume ratio. After further ball milling for 0.5 hours with ZrO₂ balls, the precursor was retrieved and shaped into a bulk form. Cold isostatic pressing at 250 MPa was employed to prepare a composite billet. Subsequently, the UHSQ technique was applied. The composite billet was wrapped in ZrO₂ fiber cloth, which serves as a high-temperature electrical insulator to prevent short-circuiting between the conductive sample and the carbon heaters. The high temperature of the carbon heater was monitored using a thermal infrared imager (628 CH-L25).

Mechanical property characterization
The stress-strain curves at room temperature of the samples were measured with a Shimadzu AGS-X 10Kg tensile machine operated at a strain rate of 1×10−3 s−1. An extensometer (RVX-112B) was used to ensure accurate strain reading. The stress-strain curves of the samples at high temperature were measured with an Instron 45582 tensile machine operated at a strain rate of 1×10−3 s−1. Specific sample dimensions are provided in Supplementary Fig. 23b. The hardness testing was performed using a HengYi MH-31 tester. The nanoindentation tests were performed on an Oxford Instruments Femto-Tools nanoindenter, and the specific test locations were selected based on the tensile fracture samples divided into regions at different distances (Supplementary Fig. 23a).

Materials characterization
The metallographic photographs were taken after polishing and grinding using an optical microscope of ZEISS AXIO Imager (A1m), to observe the distribution of metallic glass and Cu. The scanning electron microscope (SEM) images were taken using a cold field emission scanning electron microscope (SU8220), Zeiss G500 SEM, and Guoyi Quantum SEM 5000X to observe the morphology of the sample. The crystal phases were examined by Ultima IV (Rigaku Corporation, Japan) XRD and scanned between 10° and 80° using Cu Kα radiation. A differential scanning calorimeter (DSC-Q200) from TA Instruments was used to monitor and measure the glass transition temperature (Tg) of the Fe55Cr25Mo16B2C2 metallic glass. Furthermore, the detailed microstructures of the interfaces and the corresponding elemental distributions were analyzed using TEM and EDS on an FEI Talos F200X G2 system.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "Comments to NCOMMS-25-21591\nTitle: Breaking Immiscibility Barriers: Ultrafast Sintering of Interlocked Cu-Fe-Based Composites\nThe authors show a novel method (UHSQ) to fabricate Cu/metallic glass composites and present interesting microstructural\nfeatures. The results indicate that UHSQ is indeed a reliable technique for processing metallic glass composites. However,\nthe performance and related mechanism analysis presented in the manuscript are not particularly compelling, and the\nselected material system does not fully demonstrate the advantages of the UHSQ technique. Specific comments are as\nfollows:\n1. What is the relative density of the UHSQ-sintered samples? Is there any evidence of residual porosity?\n2. Page 10: “The optimized UHSQ-processed composites demonstrate remarkable mechanical enhancement across\nmultiple length scales.” What exactly is meant by \"multiple length scales\" in this context?\n3. The comparison “30 times higher than pure copper under same conditions” is questionable, given that the composite\ncontains only 50% Cu. This comparison is not meaningful. A more relevant benchmark would be W–Cu alloys.\n4. The term “bimodal architecture” is introduced without any prior mention or explanation in the microstructure section.\n5. What are the specific phases referred to as “nanocrystalline intermetallic compounds”?\n6. How was the “average debonding length of 2.1 ± 0.3 μm” measured?\n7. “ultimately fracturing along crystalline boundaries”. It lacks direct evidence. In composites with 50% reinforcement phase,\nsuch fracture behavior is unlikely. Cracks are more likely to initiate and propagate along interfaces.\n8. “UHS and UHSQ samples exhibited similar yield strains, but UHSQ samples demonstrated higher yield strength (Figure\nS12), reflecting the strengthening effect of rapid quenching and improved interface bonding”. The difference in yield strain\ncould be attributed to differences in elastic modulus, rather than interface bonding.\n9. What is the “fine Fe phase”? In Figure 3h, the region is labeled as metallic glass.\n10. “At room temperature, UHSQ-treated composites achieve an average hardness of 900 HV” requires further clarification.\nGiven the two-phase nature (Cu and metallic glass), it is difficult to avoid indentation into different phases. Why is the\nhardness variation so small?\n11. As shown in Figure 4b, there is a significant hardness difference between the two phases. Error bars should be added in\nFigure 4a. Also, considering that Cu accounts for 50% of the composite, what causes the overall hardness to far exceed the\nvalue predicted by the rule of mixtures?\n12. There is a typographical error in Figure 4d: “10 MP”.", "reviewer_2": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature\nCommunications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career\nResearchers who co-review manuscripts.", "reviewer_3": "Review of manuscript submitted to Nature Communications\nTitle: Breaking Immiscibility Barriers: Ultrafast Sintering of Interlocked Cu-Fe-Based Composites\nThis manuscript shows the view on the possibility of densifying positive mixing enthalpy alloys with a homogenous\nmicrostructure and achieve a high tensile strength, wear resistance and hardness. Authors propose ultrafast high-\ntemperature sintering followed by quenching as the processing approach. Of paramount importance is to demonstrate and\ndiscuss the link between processing, the corresponding microstructure and the subsequent mechanical properties. In fact,\nauthors are still required to discuss the critical role of ultrafast high-temperature sintering. In addition to contribution of\nquenching, it is necessary to clarify in detail whether the unique microstructure and mechanical properties are obtained as a\nresult of a) the high sintering temperature, b) short sintering duration.\nThe following comments must also be considered when revising the manuscript.\n1. The Abstract is ambiguous and does not highlight the material(s) of study. Without providing the composition of the\n‘composite’, authors mention ‘The mechanical tensile strength of the resulting composite is approximately 8 times greater\nthan that of pure Cu.’ This ambiguity also continues to Introduction section.\nIn Introduction\n2. There is a sharp contrast between values of heating rate (500-100 ºC/s) and that is represented by authors for this study.\n3. Figure 1b and 1c, I wonder how temperature profile was obtained.\n4. The term ‘current activated sintering’ is typically used for techniques like flash sintering, or spark plasma sintering when\nelectric current contributes to sintering process. However, as noted in Ref. 28, in ultrafast high-temperature sintering, there is\nalmost no contribution of electric current during sintering.\nHigh-energy ball milling pre-treatment to enhance attachment\n5. General comment, the title of the section and context are in contradiction.\n6. I do not understand the correlation between ‘hardness of metallic glass’, ‘challenges for interface attachment’ and Figures\n2a and S1.\n7. It is difficult to realize what the authors mean by coarsening in ‘In contrast, UHS leverages ultra-fast sintering rates to\nsuppress coarsening’.\nGood Mechanical Properties\n8. The title of the section could be Mechanical Properties. Please leave the decision for the reader to decide whether\nmechanical properties are good or not. Same also holds for the next section.\n9. Energy-dispersive X-ray spectroscopy (EDS) line scan analysis (most likely) in Scanning Electron Microscopy (SEM) is\nemployed to demonstrate absence of cross diffusion (Figure S13). However, the analysis cannot be considered as the most\nreliable approach, considering the limitations of EDS in SEM. It is recommended to perform the analysis using Transmission\nElectron Microscopy, having a nano scale resolution.\nGood Hardness and Wear Resistance\n10. ‘The Fe55Cr25Mo16B2C2 metallic glass maintains a high hardness of approximately 1100 HV.’ Metallic glass or the\ncrystallized phase?\nFurther concerns and comments\nUnits of measurement are not kept consistent throughout the manuscript!\nAbbreviations are used without defining them. Please do not use abbreviations in the conclusion section.\nScale bar in some Figures is missing, for most it is provided on the right bottom side; however, for one part of Figure 2f it is\nshown on the left bottom side.\nFurther advanced analysis could be used to support the statement of porous weak interface (Figure S3 a-c).\nDescriptions for Figures’ captions still need improvement, and revision.\nFigure S11, the corresponding element labels are missing.\nFigure S13, the corresponding micrographs are not provided, if any. In addition, in the manuscript it noted ‘mapping’.\nHowever, Figure S13 is line scan analysis.", "reviewer_4": "The study presents a novel ultrafast high-temperature sintering and quenching (UHSQ) method for fabricating high-strength\npositive mixing enthalpy Cu-Fe-based composites, and it provides strong evidence supporting the claimed benefits of this\nmethod. The resulting composite materials exhibit significantly enhanced mechanical properties compared to pure\ncopper and materials processed by conventional methods. Overall, the work is of very high quality and innovative, and can be\nconsidered for publication in Nature Communications after addressing the comments as below:\n(1) The extent and nature of the \"partial devitrification\" should be further elucidated in this work. The manuscript mentions\nweak crystalline peaks in UHSQ specimens, suggesting the formation of \"nanoscale intermetallic compounds\". While this\ncan contribute to the enhanced properties, the phase characterization, including crystal structure identification, phase\nvolume fraction, size distribution, and, perhaps even thermal stability, should be discussed in more detail. If these\nnanocrystals are not stable at the intended application temperatures or if their formation is inconsistent, it could lead to\ndegraded properties over time. Therefore, a more detailed characterization of these fine intermetallic compounds is\nsuggested, as it will help understand the fundamental mechanisms at play. As a start, an isothermal treatment to grow the\nphase bigger will help identify the crystal structure and phase composition for these intermetallic compounds. TEM or atom\nprobe should have a good chance of assisting in this analysis. Otherwise, authors should justify with a convincing statement.\n(2) It is also helpful to discuss the scale-up potential of such a technique, as the tensile test samples are very limited in terms\nof dimension compared to ASTM standards. Note that there is no need to be convinced that this technique can produce\neverything on a large scale, but it is important to mention related limitations or any specific requirements to scale up."}, "author_response": "Point-by-point response to the reviewer’s remarks\n(Black italic: Reviewer’s remarks; Blue: Our response;\nFig. RX appears only in the response.)\nFirst, we appreciate the reviewers' time and efforts spent reviewing our manuscript. We very much\nappreciate the reviewers' thoughtful comments and suggestions, which have significantly\nimproved the quality and clarity of our manuscript. We have performed additional experiments\nand revised the manuscript to address all comments and concerns. The changes are highlighted in\nyellow in the revised manuscript. Our point-by-point responses are presented below.\nReviewer #1:\nThe authors show a novel method (UHSQ) to fabricate Cu/metallic glass composites and present\ninteresting microstructural features. The results indicate that UHSQ is indeed a reliable technique\nfor processing metallic glass composites. However, the performance and related mechanism\nanalysis presented in the manuscript are not particularly compelling, and the selected material\nsystem does not fully demonstrate the advantages of the UHSQ technique. Specific comments are\nas follows:\nOur response: Thank you for your valuable time and incisive comments, and we especially\nappreciate your recognition of the novelty and reliability of our Ultrafast High-temperature\nSintering and Quenching (UHSQ) method for metallic-glass composites. Your incisive critiques\nof the mechanistic interpretation and material selection have guided our comprehensive revision,\nincorporating new experiments and analyses. In response, we have performed substantial revisions\nto address these points.\nRegarding the justification for the selected material system and its ability to highlight the\nadvantages of UHSQ: We acknowledge that the strategic rationale for selecting the Cu-Fe-based\nimmiscible system was underarticulated in the original draft. In the revised manuscript, we state\nin the Introduction and Discussion that our goal is to overcome the processing bottleneck of\nimmiscible alloys with large positive mixing enthalpy, exemplified by Cu–Fe. This makes it an\nideal platform to demonstrate the unique capabilities of our method. To further highlight the\nindispensable advantages of UHSQ, we have reinforced our comparison against other sintering\ntechniques, demonstrating that the ultra-fast quenching step of UHSQ can produce the defect-free,\nstrongly-bonded interface, thereby preventing the interfacial decohesion seen in UHS or the gross\nphase segregation in conventional methods. Finally, we have expanded our discussion to include\nthe potential universality of this method for other immiscible systems, broadening the impact of\nour findings.\nThank you again for your valuable feedback. We are confident that these substantial revisions have\nsignificantly improved the scientific rigor and the strength of the conclusions drawn in the\nmanuscript. We will now address your specific comments on a point-by-point basis.\n1\n1. What is the relative density of the UHSQ-sintered samples? Is there any evidence of residual\nporosity?\nOur response: We thank the reviewer for this important question. In response, we have now\nincorporated a more comprehensive analysis of density and porosity into the revised manuscript.\nTo accurately characterize the densification of the Cu-50wt%FeMG composite prepared by\nUHSQ, we employed two complementary methods. Firstly, the relative density measured by the\nArchimedes method is 97.9% (measured: 8.30 g/cm³; theoretical: 8.48 g/cm³). Secondly, to\ndirectly assess the internal porosity, we performed statistical analysis on SEM images of polished\ncross-sections, which yielded a porosity of approximately 1.2%, corresponding to a relative density\nof 98.8%. These two datasets converge, confirming that the sample achieved a high relative density\nof approximately 98-99%.\nThe characteristics of the minor residual porosity (~1-2%) are now presented in Supplementary\nFigures S5. As shown in the SEM images, the pores in the UHSQ sample are small, isolated, and\nwell-contained within the Cu matrix (It may be caused by corrosion from metallographic polishing\nsolution), crucially away from the primary load-bearing Cu/FeMG interfaces (Almost one hundred\npercent). This is in stark contrast to the UHS sample (without quenching in the Figures S4), which\nexhibits significant porosity concentrated at these interfaces. The favorable pore distribution,\ncombined with a robust rivet-like interlocking structure at the interface, prevents the propagation\nof voids under strain. This comparison unequivocally demonstrates that the ultra-fast quenching\nstep in the UHSQ process is essential for eliminating such detrimental interfacial defects, thereby\ncreating the robust microstructure that underpins high mechanical performance of the material.\nconventional UHS method. (a-b) SEM micrographs providing a general overview of the UHS-\nsintered control sample. These images demonstrate that interfacial defects between the\nFe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass particles and the Cu matrix are widespread throughout the\nmaterial. (c) A higher-magnification micrograph of a single FeCrMo nanocrystal, presenting a\n a\nclear, representative example of poor interfacial bonding.\nfabricated by UHSQ. (a-d) SEM micrographs at increasing magnifications reveal a dense, two-\nphase composite where multi-phase nanocrystals are uniformly embedded within a continuous Cu\nmatrix with clean, sharp interfaces. (e,f) Corresponding EDS elemental maps confirm the distinct\nchemical separation between the MG particles (rich in Fe, Cr, Mo) and the Cu matrix.\n2. Page 10: “The optimized UHSQ-processed composites demonstrate remarkable mechanical\nenhancement across multiple length scales.” What exactly is meant by \"multiple length\nscales\" in this context?\nOur response: We thank the reviewer for this insightful comment and agree that our original\nphrasing was ambiguous. “Multiple length scales” distinguishes bulk mechanics from localized\nsurface behavior. Our analysis encompasses:\n1. Bulk properties, such as the tensile strength and ductility evaluated across the entire sample\nvolume, as presented in the stress-strain curves (Figure 3d, e).\n2. Surface-dominant properties, such as the microhardness and wear resistance, which probe the\nmaterial's response in localized regions near the surface (Figure 4a, d).\nTo eliminate this ambiguity and provide better clarity, we have revised the sentence in the\nmanuscript as follow:\n\"The optimized UHSQ-processed composites demonstrate a remarkable and comprehensive\nenhancement of mechanical properties, spanning bulk behavior (tensile and flexural strength) and\nsurface performance (hardness and wear resistance).\"\n55 25 16 2 2\nmetal matrix composites. (a) TEM image showing the clean, well-bonded interface between the\nCu matrix and the Fe-based reinforcement. (b) Higher-magnification view revealing the nanoscale\ninterlocking structure, with corresponding EDS maps confirming the distinct phases. (c) HAADF-\nSTEM analysis of the devitrified reinforcement, revealing a nano-architecture of an Fe-Cr matrix,\nMo-rich precipitates, and Fe-rich inclusions, confirmed by EDS mapping. The plot below\nquantifies the size distribution of these nanocrystals. (d) Room-temperature tensile engineering\nstress-strain curves for the composites processed by UHS and UHSQ, compared with a monolithic\nCu control sample. (e) Elevated-temperature (650 °C) tensile engineering stress-strain curves for\nthe UHS and UHSQ composites. (f-g) SEM fracture images of UHS sample tensile fracture at\nroom temperature and SEM fracture images of UHSQ sample tensile fracture at room temperature.\n(h) SEM images showing the distribution of small-sized metallic glass particles in the composite\nmaterial.\ncomposites. (a) Microhardness profiles measured from the sample surface inwards for the UHSQ\ncomposite, UHS composite, and monolithic pure Cu. The shaded bands represent the standard\ndeviation of the measurements. (b) Representative microhardness line scan across the Cu matrix,\nthe Cu/FeCrMo interface, and multi-phase nanocrystals. (c) The plot of microhardness versus\nelectrical conductivity (% IACS), benchmarking the UHSQ composite against various material\nclasses reported in the literature 43-60. (d) Wear test results showing the wear loss rate as a function\nof friction cycles under a load of 10 MPa and a sliding speed of 500 r/min. (e) Optical micrographs\nof the worn surfaces of UHSQ composites with 25 vol% and 50 vol% Fe-MG reinforcement,\nshowing uniform wear tracks. (f) High-magnification SEM images tracking the morphology of a\nsingle multi-phase nanocrystals on the wear surface after 15 and 60 minutes of sliding.\n6\n3. The comparison “30 times higher than pure copper under same conditions” is questionable,\ngiven that the composite contains only 50% Cu. This comparison is not meaningful. A more\nrelevant benchmark would be W–Cu alloys.\nOur response: We sincerely thank the reviewer for this constructive suggestion. We agree that\ncomparing our 50% composite to pure copper is not the most relevant benchmark and could be\nmisleading. While our initial intent was simply to use copper as a baseline to show thermal stability\nimprovement, we acknowledge that a comparison with state-of-the-art materials is more\nappropriate.\nAs suggested, we have considered W-Cu alloys. However, a direct comparison is not\nstraightforward, as most high-performance systems in the literature do not have a comparably high\n(~50%) reinforcement content1-5. This is also where our technological advantage lies, as it can\nachieve the ultra-large component homogeneous synthesis of positive mixed alloy systems while\nmaintaining the excellent characteristics of both.\nTo address this concern and improve the manuscript, we have removed the direct comparison to\npure copper from the text. The sentence has been revised to focus on the material's intrinsic high-\ntemperature performance.\nWe have incorporated the aforementioned revisions into the manuscript, which is attached for your\nconsideration.\n\"Remarkably, this strength advantage persists at elevated temperatures, with the composite\nmaintaining 290 MPa yield strength at 650 °C (Figure 3e).\"\nReferences\n1. Han, T. et al. W–Cu composites with excellent comprehensive properties. Composites Part B:\nEngineering 233, 109664 (2022).\n2. Duan, Q. et al. Synergistic Enhancement of Mechanical Properties and Electrical Conductivity\nof Immiscible Bimetal: A Case Study on W–Cu. Engineering 46, 224–235 (2025).\n3. Liu, J.-K., Wang, K.-F., Chou, K.-C. & Zhang, G.-H. Fabrication of ultrafine W-Cu composite\npowders and its sintering behavior. Journal of Materials Research and Technology 9, 2154–\n2163 (2020).\n4. Hou, C., Cao, L., Li, Y., Tang, F. & Song, X. Hierarchical nanostructured W-Cu composite\nwith outstanding hardness and wear resistance. Nanotechnology 31, 084003 (2020).\n5. Xiaoyong, Z., Jun, Z., Junling, C. & Yucheng, W. Structure and Properties of W-Cu/AlN\nComposites Prepared via a Hot Press-Sintering Method. Rare Metal Materials and\nEngineering 44, 2661–2664 (2015).\n7\n4. The term “bimodal architecture” is introduced without any prior mention or explanation in\nthe microstructure section.\nOur response: We thank the reviewer for this critical and insightful comment. We concede that\n“bimodal architecture” was ill-defined and potentially misleading, and this has been rectified.\nHonestly, we aimed not to imply a bimodal size distribution, but to emphasize the unique 3D\ninterlocked Cu/FeCrMo interface. This key microstructural feature, which results from our novel\nUHSQ process, is clearly visualized in Figure 3b.\nAs the reviewer correctly inferred, this complex interfacial morphology is far more significant than\na simple boundary. It functions as a microscopic \"mortise and tenon\" joint, creating a powerful\nmechanical lock between the ductile matrix and the hard reinforcement. This architecture\nunderpins the composite’s exceptional performance, furnishing seamless load transfer and\nsuppressing interfacial decohesion to simultaneously boost strength and toughness (Fig. 3d, g).\nTo rectify this, we have performed the following specific revisions throughout the manuscript:\n\"All comparisons demonstrated significant improvements in mechanical performance originating\nfrom the unique, interlocking Cu-Fe interface where continuous Cu and Fe phases form\ninterpenetrating networks resembling dovetail joints, a structure forged by the UHSQ method.\"\n5. What are the specific phases referred to as “nanocrystalline intermetallic compounds”?\nOur response: We thank the reviewer for this comment. We agree that the term “nanocrystalline\nintermetallic compounds” was not precise enough. Our analysis reveals that the material is a\ncomplex multiphase nanocrystalline composite rather than a simple mixture of intermetallic.\nWe have revised the manuscript to explicitly define each phase as detailed below, supported by the\nnew data in the Supplementary Information. We have performed detailed characterization of these\nphases and will provide a thorough explanation below.\na) Nanocrystalline Fe-Cr Ferrite Matrix\nThe foundational structure of our material is an Fe-Cr ferrite matrix. As shown in the HAADF\nimages (Figure S7 a, c), this phase constitutes the large, gray-contrast grains labeled \"Fe-Cr.\" The\ncorresponding EDS elemental maps (Figure S7 b) confirm the homogeneous co-distribution of Fe\nand Cr throughout these grains, indicating the formation of a solid solution. Crucially, the high-\nresolution TEM image and its Fast Fourier Transform (FFT) pattern (Figure S7 d, center panel)\nunambiguously identify its crystal structure as Face-Centered Cubic (FCC). This Fe-Cr ferrite\nphase serves as the primary framework for the entire nanocrystalline ensemble.\nb) Mo-rich Precipitate Phase\nEmbedded within the ferrite matrix are nanoscale, Mo-rich precipitates. These are clearly\ndistinguished in the HAADF images (Figure S7 a, c) as bright particles, a direct result of the higher\natomic number of molybdenum (Z=42) providing strong Z-contrast. The EDS map for Mo (Figure\nS7 b) confirms the significant enrichment of molybdenum in these localized regions. It should be\nnoted that the use of a Mo grid may slightly enhance the overall Mo signal, but the distinct,\n8\nlocalized enrichment in the precipitates is unequivocal. Structural analysis via HRTEM (Figure S7\nd, right panel) reveals that this Mo-rich phase possesses a BCC structure. These precipitates are a\nkey strengthening component within the overall nanocrystalline structure.\nc) Complex Fe-rich Regions\nMost notably, we have identified complex regions labeled \"Fe-rich Area\" which are themselves\nnanocomposites. Detailed HRTEM analysis (Figure S8 a) reveals these regions are not a single\nphase but are comprised of at least three distinct constituents:\nFirst is ordered b2 intermetallic phase: Within these Fe-rich areas, we identified an ordered b2\nintermetallic phase. The HRTEM image and corresponding FFT pattern (Figure S8 b) show clear\n{010} superlattice reflections in addition to the fundamental BCC reflections. The presence of\nthese superlattice spots is definitive proof of b2 ordering.\nThe second is disordered nanocrystals regions: Coexisting with the ordered b2 phase are pockets\nof untransformed FeCrMo amorphous phase (metallic glass), as labeled \"FeCrMo MG\". This is\nconfirmed by HRTEM images showing a characteristic disordered, maze-like atomic arrangement\nand a corresponding diffuse halo in the FFT pattern (Figure S8 c).\nThe last is Fe nanocrystals: Pure Fe nanocrystals with a BCC structure are also present within these\ncomplex regions, as identified by HRTEM (Figure S-d, left panel).\nIn summary, “Nanocrystalline intermetallic compounds” denotes the entire multiphase ensemble:\nFe–Cr ferrite, Mo-rich precipitates, and Fe-rich domains (ordered b2 intermetallic phase,\ndisordered nanocrystals, and pure Fe nanocrystals). The interplay between these intimately mixed,\nnanoscale phases govern the material's overall properties. We revise the manuscript to explicitly\ndefine this term and clarify the multiphase nature of the material. The details are as follows.\n\"The rapid quenching process induces the formation of a multiscale structure within the Fe-based\nmetallic glass (Figure S5). Bright-field TEM imaging reveals a regular and uniform distribution of\nnanoscale crystal particles within the metallic glass (Figure 3c). Figure 3b collectively demonstrate\nthat the Fe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass devitrifies into a multi-phase nanocrystal, in which\nnanocrystalline Fe–Cr ferrite forms the continuous skeleton of the microstructure (Figure S6). The\nfcc Fe–Cr matrix surrounds and mechanically pins two secondary constituents, Mo-rich\nnanocrystals and Fe-rich areas. Bcc Mo-rich nanocrystals that populate junctions and grain corners\nin a quasi-periodic manner, occupying ~14 vol % of the microstructure (Figure S7), surrounding\neach Mo grain is a shell of bcc Fe ferrite. Notably, in the Fe-rich areas, containing both pure Fe\nand disordered nanocrystals regions (Figure S8a). Furthermore, these Fe-rich areas contain ordered\nb2 intermetallic phases, contributing to the overall intricate structure (Figure S8b). The complex\narrangement indicates an incomplete devitrification process, resulting in a multiphase\nnanocrystalline structure with retained amorphous pockets and the formation of ordered\nintermetallic compounds (Figure S8c). The overall microstructure exhibits a high density of\ninterfaces and structural complexity. This intricate structural architecture, arising from the\ncontrolled devitrification of the initial metallic glass, significantly influences the material's\nproperties.\"\n(a-b) Analysis from two different regions of the reinforcement phase, demonstrating the consistent\nformation of nano-architecture. Each panel contains a HAADF-STEM image, a corresponding BF-\nS image, and EDS elemental maps for Fe, Cr, and Mo. The HAADF images use Z-contrast to\ndistinguish the phases, which are then chemically identified by the EDS maps. The analysis\nconfirms the structure consists of a continuous Fe-Cr matrix, discrete Mo-rich precipitates, and\nFe-rich areas occupying the interstitial spaces.\nreinforcement phase. (a-b) HAADF-STEM image revealing a complex nano-architecture\nconsisting of multiple phases distinguished, corresponding EDS elemental maps of the region\nboxed in (a), confirming the chemical distribution of Fe, Cr, and Mo. The analysis identifies a\nprimary Fe-Cr matrix, localized Mo-rich precipitates, and Fe-rich inclusions. (c) The higher-\nmagnification HAADF-STEM image providing a detailed view of the phase morphology and\ndistribution. (d) Atomic-resolution HRTEM images with corresponding Fast Fourier Transform\n(FFT) patterns for the Fe-rich, Fe-Cr, and Mo-rich regions, respectively.\nTEM image of the Fe-rich areas, revealing a nanocomposite structure. (b) HRTEM image of a\ncrystalline domain; inset FFT indicates a primary bcc lattice with superlattice reflections consistent\nwith a coexisting b2 phase. (c) HRTEM image showing a region of untransformed amorphous\nphase; inset FFT displays a diffuse halo. The Fe-rich reinforcement comprises bcc-Fe, b2\nintermediate phase, and disordered nanocrystals regions.\n6. How was the “average debonding length of 2.1 ± 0.3 μm” measured?\nOur response: We sincerely thank the reviewer for this important question. We acknowledge that\n“debonding length” was a misnomer for the observed feature. We agree that this term was not\noptimal, as it implies a post-fracture event. We have revised the manuscript to use the more\naccurate term \"interfacial gap length\" to describe the pre-existing microstructural defects in the as-\nsintered UHS control samples. The following details this corrected analysis.\nThe pre-existing interfacial gaps in as-sintered UHS controls have been quantified. These gaps\nare inherent defects at the Cu-FeCrMo multi-phase nanocrystals interfacial gap length interface\nresulting from the non-quenched cooling process and are the primary source of weak bonding. The\ndetailed statistical analysis is presented in Figure R1. The process involved analyzing SEM images\nof polished cross-sections from three independent UHS samples using ImageJ software. The\ncollected data for each sample are shown as histograms in Figure R3a-c, yielding highly consistent\naverage gap lengths of D = 1.78 ± 0.07 μm (Sample 1), D = 1.93 ± 0.05 μm (Sample 2), and D =\n2.11 ± 0.19 μm (Sample 3). The summary plot in Figure R3d confirms the excellent reproducibility\nof these measurements.\nThis analysis provides a quantitative measure of the poor as-sintered integrity of the UHS control\nsamples. This stands in stark contrast to the UHSQ composites, which exhibit clean, tightly-\n12\nbonded, gap-free interfaces. This comparison unequivocally highlights that the ultra-fast\nquenching step in the UHSQ process is critical for preventing the formation of these detrimental\ninterfacial defects, thus achieving a robust metallurgical bond.\nFigure R1. Statistical analysis of the pre-existing interfacial gap length in as-sintered UHS control\nsamples. The measurements were performed on scanning electron microscopy (SEM) images of\npolished cross-sections to quantify inherent defects at the Cu-FeCrMo multi-phase nanocrystals\ninterfaces. (a-c) Histograms showing the statistical distribution of measured gap lengths for three\nindependent samples (Sample 1, 2, and 3). A log-normal distribution curve (red line) is fitted to\neach dataset, and the calculated average gap length with its standard deviation is displayed. (d)\nThe summary plot illustrating the average gap length with standard deviation for each of the three\nsamples.\n13\n7. “ultimately fracturing along crystalline boundaries”. It lacks direct evidence. In composites\nwith 50% reinforcement phase, such fracture behavior is unlikely. Cracks are more likely to\ninitiate and propagate along interfaces.\nOur response: We thank the reviewer for highlighting this ambiguity. The reviewer makes a valid\npoint; our use of the phrase 'fracturing along crystalline boundaries' lacked the necessary\nspecificity. We also agree that in many composites, interfacial debonding is a common failure\nmode. However, our evidence indicates a different mechanism: brittle trans-particle fracture of the\nreinforcement phase, as detailed in Figure S12.\nFirst, Energy-Dispersive X-ray Spectroscopy (EDS) mapping confirms the chemical identity of\nthe features on the fracture surface. Figure S12d shows Fe-Cr-Mo-rich fragments amid Cu-rich\ndimples, confirming reinforcement fracture within the ductile matrix. Second, high-magnification\nscanning electron microscopy (SEM) images provide direct visual evidence of the failure mode.\nThe images in Figure S12a-c show no signs of interfacial debonding, such as particle pull-out or\nempty sockets left behind in the matrix. Instead, the reinforcement particles have fractured cleanly\nthrough their cross-section, displaying flat surfaces with features characteristic of brittle cleavage.\nThis trans-particle fracture strongly indicates that the bond at the Cu-FeCrMo multi-phase\nnanocrystals interface is exceptionally robust—stronger than the intrinsic fracture strength of the\nreinforcement phase itself. The crack, therefore, follows the path of least resistance by cleaving\nthe particle rather than separating the interface.\nPers the reviewers' suggestions and the emphasis on this mechanism, the revised manuscript is as\nfollows.\n\"The high interfacial bonding strength provided by the rivet-like structures is best demonstrated\nby the material's failure mechanism. Instead of debonding at the interface, failure occurs via brittle\ntrans-particle fracture, where cracks propagate directly through the reinforcement phase itself\n(Figure S12)\".\ntesting, providing evidence for robust interfacial bonding. (a-c) High-magnification SEM images\nshowing the characteristic failure modes. The Cu matrix exhibits classic ductile fracture with fine,\nequiaxed dimples. Critically, the Fe-based reinforcement particle fails via brittle trans-particle\nfracture, with the crack propagating directly through the particle rather than along the interface.\n(d) Corresponding EDS elemental mapping confirms the chemical identity of the phases, showing\nthe fractured particle is rich in Fe, Cr, and Mo, surrounded by the Cu-rich matrix. The fracture\nbehavior demonstrates that the interfacial bond strength exceeds the intrinsic fracture strength of\nthe reinforcement phase itself.\n8. “UHS and UHSQ samples exhibited similar yield strains, but UHSQ samples demonstrated\nhigher yield strength (Figure S12), reflecting the strengthening effect of rapid quenching and\nimproved interface bonding”. The difference in yield strain could be attributed to differences\nin elastic modulus, rather than interface bonding.\nOur response: We thank the reviewer for raising this important point regarding the interplay\nbetween modulus, strength, and strain. We recognize that this complex relationship warrants a\nmore detailed explanation than we initially provided. A a clearer, data-driven account: direct\nflexural stress–strain analysis in Figure S14 has been offered.\nThe difference in the effective elastic modulus between the UHS and UHSQ samples will indeed\ninfluence the stress-strain behavior, and we agree that our original statement did not capture the\nfull picture. While a change in elastic modulus certainly contributes, we propose that the\nsubstantial increase in yield strength is predominantly rooted in the unique and complex\nmicrostructure enabled by our rapid quenching process. The UHSQ technique fosters a\nhierarchical, multiphase nanocomposite—characterized by a nanocrystalline matrix, hard Mo-rich\nprecipitates and ordered b2 intermetallic, and disordered nanocrystals regions—where the\nresulting high density of interfaces establishes a complex, three-dimensional network of barriers\n15\nto plastic deformation. Consequently, the remarkable enhancement in yield strength is a direct\nresult of a multiphase synergistic strengthening mechanism, where the interplay between the\ndifferent constituents provides a strengthening effect far exceeding that attributable to modulus\nchanges alone. This intricate, synergistic architecture is a direct and unique advantage of our\nUHSQ processing strategy.\nAccordingly, we have revised the relevant passage in the manuscript to correct the inaccuracies\nand reflect our more precise, data-driven analysis.\n\"While the difference in elastic modulus between the UHS and UHSQ samples contributes to the\noverall stress-strain response, the strength is predominantly attributed to a multiphase synergistic\nstrengthening mechanism forged by the UHSQ process. The enhancement is not due to a change\nin intrinsic material stiffness but is a direct reflection of the strengthening effect from the improved\ninterface bonding achieved by rapid quenching.\"\nFigure S14. Comparison of flexural stress-strain curves for composites processed via UHSQ and\nUHS at room temperature. The data shows that the UHSQ composite exhibits a significantly higher\nflexural strength (approaching 900 MPa) compared to the control sample processed without\nquenching (UHS).\n16\n9. What is the “fine Fe phase”? In Figure 3h, the region is labeled as metallic glass.\nOur response: We would like to thank the reviewers for raising this critical issue. Which we have\nrevised “fine Fe phase” to \"Fe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass\" in the original manuscript.\n10. “At room temperature, UHSQ-treated composites achieve an average hardness of 900 HV”\nrequires further clarification. Given the two-phase nature (Cu and metallic glass), it is\ndifficult to avoid indentation into different phases. Why is the hardness variation so small?\nOur response: We thank the reviewer for this excellent question, which allows us to clarify a key\nfeature of our material. The consistency of the Vickers hardness values is a direct consequence of\nthe measurement scale relative to the composite's unique, homogeneous microstructure. The\nmicrohardness indentation size is significantly larger than the scale of the uniform microstructure.\nConsequently, each measurement samples a statistically representative volume, resulting in data\nwith low variance and high statistical significance.\nThe Vickers test creates an indentation (typically tens of micrometers) that is significantly larger\nthan the individual phase domains in our composite. Therefore, each measurement assesses the\ncollective mechanical response of numerous copper and metallic glass regions, inherently\naveraging their properties rather than probing the intrinsic hardness of a single phase (approx. 90\nHV for the Cu phase vs. 1100 HV for the FeCrMo nanocrystal). This averaging is effective because\nthe UHSQ process produces a highly uniform, co-continuous interlocked network of the two\nphases, ensuring that any macroscopic indentation samples a statistically consistent volume\nfraction of both hard and soft components. Mechanically, this cooperative deformation is enabled\nby the exceptionally strong interfacial bonding, which is confirmed by TEM analysis showing\nclean, gap-free interfaces and further evidenced by high interfacial microhardness values of ~500\nHV. The robust interface loads the hard network through the ductile matrix, causing monolithic\nresponse. The low hardness spread is thus a design outcome, not an anomaly.\n11. As shown in Figure 4b, there is a significant hardness difference between the two phases.\nError bars should be added in Figure 4a. Also, considering that Cu accounts for 50% of the\ncomposite, what causes the overall hardness to far exceed the value predicted by the rule of\nmixtures?\nOur response: We are grateful to the reviewer for these perceptive comments. Per the request,\nFigure 4a now displays error bands that clearly convey the tight standard deviation in hardness\n(error bands as a percentage of the mean)\nRegarding the second point, the reviewer is correct that the composite's average hardness (~900\nHV) significantly exceeds the value predicted by the simple linear rule of mixtures (ROM), which\nwould be approximately 600 HV. This positive deviation is a hallmark of our material's design and\nis attributed to several powerful synergistic strengthening mechanisms not captured by the ROM.\nFirst, the Fe-based reinforcement is itself an in-situ formed nanocrystalline intermetallic\ncompounds with exceptionally high intrinsic hardness (~1100 HV) due to the Hall-Petch effect.\nsofter Cu phase, while the robust interface-confirmed by TEM and high interfacial microhardness\n-ensures efficient load transfer. Finally, the extensive plastic deformation of the Cu matrix prior to\nfailure, evidenced by dense dimples on the fracture surface, leads to significant work hardening,\nincreasing its in-situ contribution to the overall hardness. We have incorporated this detailed\nexplanation into the revised manuscript to clarify these crucial strengthening mechanisms.\nTo ensure that this detailed explanation of the synergistic strengthening mechanisms is clear to the\nreader, we have revised the relevant paragraph in the 'Hardness and Wear Behavior' section. For\nclarity, the revision is as follows:\n\"The Fe-based reinforcement is itself an in-situ formed nanocrystalline intermetallic compounds\nand possesses a high intrinsic hardness of approximately 1100 HV. Which provides a strong\nfoundation for the composite's overall performance. Furthermore, the co-continuous, interlocking\nnetwork architecture imposes a strong mechanical constraint on the softer Cu phase (which has a\nhardness of only ~90 HV). The constraint is enabled by a robust interface, which exhibits a high\naverage hardness of ~500 HV and ensures efficient load transfer during indentation. Significant\nwork hardening occurs in the ductile Cu matrix during deformation, further increasing its in-situ\ncontribution to the overall hardness. It is the combination of these effects that results in the\ncomposite's hardness.\"\n55 25 16 2 2\nmetal matrix composites. (a) TEM image showing the clean, well-bonded interface between the\nCu matrix and the Fe-based reinforcement. (b) Higher-magnification view revealing the nanoscale\ninterlocking structure, with corresponding EDS maps confirming the distinct phases. (c) HAADF-\nSTEM analysis of the devitrified reinforcement, revealing a nano-architecture of an Fe-Cr matrix,\nMo-rich precipitates, and Fe-rich inclusions, confirmed by EDS mapping. The plot below\nquantifies the size distribution of these nanocrystals. (d) Room-temperature tensile engineering\nstress-strain curves for the composites processed by UHS and UHSQ, compared with a monolithic\nCu control sample. (e) Elevated-temperature (650 °C) tensile engineering stress-strain curves for\nthe UHS and UHSQ composites. (f-g) SEM fracture images of UHS sample tensile fracture at\nroom temperature and SEM fracture images of UHSQ sample tensile fracture at room temperature.\n(h) SEM images showing the distribution of small-sized metallic glass particles in the composite\nmaterial.\ncomposites. (a) Microhardness profiles measured from the sample surface inwards for the UHSQ\ncomposite, UHS composite, and monolithic pure Cu. The shaded bands represent the standard\ndeviation of the measurements. (b) Representative microhardness line scan across the Cu matrix,\nthe Cu-FeCrMo interface, and multi-phase nanocrystals. (c) The plot of microhardness versus\nelectrical conductivity (% IACS), benchmarking the UHSQ composite against various material\nclasses reported in the literature 43-60. (d) Wear test results showing the wear loss rate as a function\nof friction cycles under a load of 10 MPa and a sliding speed of 500 r/min. (e) Optical micrographs\nof the worn surfaces of UHSQ composites with 25 vol% and 50 vol% Fe-MG reinforcement,\nshowing uniform wear tracks. (f) High-magnification SEM images tracking the morphology of a\nsingle multi-phase nanocrystals on the wear surface after 15 and 60 minutes of sliding.\n12. There is a typographical error in Figure 4d: “10 MP”.\nOur response: We thank the reviewer for pointing out this error. The unit has been corrected to\n\"MPa\" in the revised Figure 4d.\nThis manuscript shows the view on the possibility of densifying positive mixing enthalpy alloys\nwith a homogenous microstructure and achieve a high tensile strength, wear resistance and\nhardness. Authors propose ultrafast high-temperature sintering followed by quenching as the\nprocessing approach. Of paramount importance is to demonstrate and discuss the link between\nprocessing, the corresponding microstructure and the subsequent mechanical properties. In fact,\nauthors are still required to discuss the critical role of ultrafast high-temperature sintering. In\naddition to contribution of quenching, it is necessary to clarify in detail whether the unique\nmicrostructure and mechanical properties are obtained as a result of a) the high sintering\ntemperature, b) short sintering duration.\nOur response: We thank the reviewer for this insightful suggestion. It prompted us to more\nclearly articulate the fundamental principles of our UHSQ process. We agree that a detailed\ndiscussion on the distinct roles of high sintering temperature and short sintering duration is\ncrucial and have revised the manuscript accordingly.\nOur central hypothesis is that for this challenging positive-mixing-enthalpy system, there is a\nnarrow \"kinetic window\" for achieving full densification while suppressing undesirable\nthermodynamic evolution (e.g., phase segregation and grain coarsening). To experimentally\ndecouple the effects of temperature and time, we performed sintering experiments at a lower\ntemperature (600 °C) for different durations, with the results presented in a new Figure R6.\nAs shown in Figure R2a, b, at 600 °C for 1 minute, the material is severely under-sintered with\nonly initial particle contact. Extending the time to 10 minutes (Figure R2c, d) results in only\nmarginal improvement in necking between Cu particles, while the metallic glass particles remain\nlargely inert and the structure remains highly porous. When combined with our previous finding\nthat high temperatures with long durations (i.e., conventional sintering) lead to detrimental phase\nsegregation, it becomes clear that it is the unique combination of high temperature (to enable\ndensification) and short duration (to preserve the microstructure) that is essential. Our UHSQ\nmethod is designed to precisely navigate this kinetic window, which is the key to its success.\nsintering (600 °C). (a, b) After 1 minute, the particles show only initial point contact and\nsignificant porosity remains. (c, d) Extending the time to 10 minutes leads to only slight necking\nin the Cu matrix, while the metallic glass particles remain un-sintered and the overall structure is\nfar from dense. This confirms that high temperature is a prerequisite for successful densification\nin this system.\n1. The Abstract is ambiguous and does not highlight the material(s) of study.\nWithout providing the composition of the ‘composite’, authors mention ‘The mechanical tensile\nstrength of the resulting composite is approximately 8 times greater than that of pure Cu.’\nThis ambiguity also continues to Introduction section.\nOur response: We sincerely thank the reviewer for this crucial feedback. We have accordingly\nforegrounded the material composition early in the manuscript. To address this ambiguity, we have\nadjusted the Abstract to introduce the Cu-50 vol% Fe₅₅Cr₂₅Mo₁₆B₂C₂ system more prominently\nbefore the property claims are made. Similarly, we have revised the Introduction to state the\nspecific composite material immediately after introducing our general fabrication strategy. This\nimmediately grounds the subsequent discussion in the context of our chosen system. We believe\nthese revisions make the focus of our study clear from the very beginning.\nFor clarity, the specific revisions to both sections are shown below:\n\" To this end, we employ a bottom-up approach to fabricate Cu-50 vol% Fe Cr Mo B C\n55 25 16 2 2\n22\npositive mixing enthalpy composites, utilizing ultrafast high-temperature sintering and\ninstantaneous quenching.\"\n\" Employing the UHSQ technique, we successfully sintered a Cu-50 vol% positive mixing\nenthalpy composites \"\n2. There is a sharp contrast between values of heating rate (500-1000 ºC/s) and that is\nrepresented by authors for this study.\nOur response: We thank the reviewer for pointing out this discrepancy. The reviewer correctly\nnotes that the schematic diagram did not accurately represent the high heating rates (500-1000\n°C/s) achieved in our experiments. This was an oversight in the illustration, and the figure has been\ncorrected to match the experimental data.\nWe have redrawn the UHSQ schematic (Figure R3) to visualize the high heating rates. Besides, to\nbetter illustrate the characteristics of my heating and cooling, we added the infrared curve of the\nactual sintering temperature to better illustrate the rapid change of the reaction temperature.\nFurthermore, as the reviewer's next comment also touches upon the feasibility of such rapid\nthermal cycling, we have prepared a more detailed explanation of our experimental setup. As we\nwill elaborate in our response to Comment #3.\nthe overall UHSQ apparatus, showing the upper heating stage (UHS) and the lower quenching\nbath. (b) Schematic illustration of the heating principle, where a sample is insulated by ZrO₂ fibers\nand heated by surrounding carbon felts. (c) Photographs of the UHS process in action (left) and\nthe ZrO₂ fiber insulation material (right). (d) A representative thermal image captured by an\ninfrared pyrometer during sintering, showing a uniform temperature distribution. (e) The\nmeasured temperature-time profile for a typical UHSQ experiment, detailing the ultrafast heating,\nsintering, and quenching stages. (f) A temperature profile demonstrating the system's capability\nfor rapid pulsed heating, validating its high thermal responsiveness.\n3. Figure 1b and 1c, I wonder how temperature profile was obtained.\nOur response: We thank the reviewer for this important question, which let us to elaborate on the\nexperimental setup that enables our Ultrafast High-temperature Sintering and Quenching (UHSQ)\nprocess. The original manuscript lacked these crucial details, and we have now added them to\nprovide full transparency.\nOur methodology is adapted from the well-established UHS technique reported by Wang et al. in\nScience (2020, 368, 521-526), UHS technology has also become very mature and widely used in\nrecent years1-6. Which we have customized for our specific alloy system. The temperature profile\n24\nwas achieved and recorded as follows:\nThe temperature-time profile was not simulated or inferred but measured directly and in-situ\nduring the experiment. As illustrated in Figure 1b and 1c, the sample is heated via Joule heating\nof the surrounding carbon felts. We used a high-speed infrared pyrometer (model 628CH-L25) to\nlive-monitor the surface temperature of the sample assembly in real-time. Figure 1d shows a\nrepresentative thermal image captured by the pyrometer, which visually confirms the uniform\ntemperature distribution in the heating zone, reaching ~1800 K. The quantitative plot in Figure 1e\nis the temperature data extracted directly from these real-time pyrometer measurements over a\ncomplete heating and quenching cycle. To further validate the system's high performance and rapid\nthermal response, we have included Figure 1f, which shows a pulsed heating profile and confirms\nour setup can achieve the ultrafast heating/cooling rates reported.\nReferences\n1. Chen, L. et al. A nanocrystal garnet skeleton-derived high-performance composite solid-\nstate electrolyte membrane. Energy Storage Materials 65, 103140 (2024).\n2. Guo, M. et al. Ultrafast high-temperature sintering to avoid metal loss toward high-\nperformance and scalable cermets. Matter 5, 594–604 (2022).\n3. Shan, Y. et al. Programmable and rapid fabrication of complex-shape ceramics. Nat\nCommun 15, 9973 (2024).\n4. Kong, X. et al. Maximizing interface stability in all-solid-state lithium batteries through\nentropy stabilization and fast kinetics. Nat Commun 15, 7247 (2024).\n5. Dong, Q. et al. Depolymerization of plastics by means of electrified spatiotemporal heating.\nNature 616, 488–494 (2023).\n6. Zeng, K. et al. Finely tailoring the local ensembles in heterostructured high entropy alloy\ncatalysts through pulsed annealing. Nat Commun 16, 3403 (2025).\n4. The term ‘current activated sintering’ is typically used for techniques like flash sintering, or\nspark plasma sintering when electric current contributes to sintering process. However, as\nnoted in Ref. 28, in ultrafast high-temperature sintering, there is almost no contribution of\nelectric current during sintering.\nOur response: We thank the reviewer for this insightful observation. We thank the reviewer for\ndistinguishing our UHS method from true “current-activated” techniques. The revised manuscript\nnow clarifies that our process electrically insulates the sample.\nIn contrast to techniques like Spark Plasma Sintering (SPS), where electric current passes directly\nthrough the sample, our UHS setup heats the sample indirectly. As explicitly illustrated in Figure\n1b and 1c, the current passes exclusively through two external graphite carbon felts, which act as\nresistive heating elements. The sample is heated purely by thermal radiation and conduction from\nthese incandescent felts. This is achieved by sandwiching the sample between two layers of ZrO₂\nfiber felt, an excellent electrical insulator. This crucial layer prevents any significant electrical\ncurrent from passing through our highly conductive composite. The uniform temperature field\nobserved in the thermal image (Figure 1d) is characteristic of this radiation-dominated heating,\nfurther confirming that the process is governed by the external elements and not by current flow\n25\nwithin the sample.\nAs the reviewer notes, the technique is a radiation/conduction-based UHS process, not a current-\nactivated one. We have revised the manuscript in \"radiative joule-heating sintering\" to state this\ndistinction clearly.\n\"The sequential stages of UHSQ processing – including precursor assembly, radiative joule-\nheating sintering (UHS), and forced-convection quenching – are visually documented in (Figure\n1e).\"\n5. High-energy ball milling pre-treatment to enhance attachment\nGeneral comment, the title of the section and context are in contradiction.\nOur response: We sincerely thank the reviewer for this crucial feedback. The reviewer is entirely\ncorrect that the original section title was too narrow and created a contradiction with the section's\nbroader content. The discussion in this section goes far beyond the HEBM pre-treatment step\nalone.\nTo resolve this contradiction and better reflect the complete scientific narrative, we have renamed\nthe section to a more comprehensive title: \"Fabrication and Microstructural Characterization of\nCu-MG Composites\".\nThis new title accurately reflects the full scope of the section, which, as we also clarified in our\nresponse to Comment #6, presents a complete story from the initial material challenge to the final\nsuccessful outcome. The revised title now fully captures the entire fabrication pathway. It\nexplicitly incorporates the complete sequence that begins with the High-Energy Ball Milling\n(HEBM) pre-treatment illustrated in Fig. 2b–c, a step specifically implemented to overcome the\nsurface-bonding issues inherent to the starting constituents. The title further extends to the\nsubsequent Ultrafast High-temperature Sintering and Quenching (UHSQ) consolidation stage\nexamined in Fig. 2e–f, a rapid thermal route that is indispensable for attaining the final, fully dense\ncomposite. The Comprehensive Microstructural Characterization: This covers the analysis of the\nraw materials that define the initial challenge (Fig. 2a), the intermediate processed powders, and\nthe final interfaces, phase structures, and phase coarsening behavior of the composites (Fig. 2e-f).\nmicrographs showing the morphology of Fe Cr Mo B C metallic glass particles. (b-c) The\n55 25 16 2 2\neffect of high-energy ball milling for 1 h and 10 h, respectively, showing the progressive coating\nof the hard MG particles with ductile Cu to create an ideal precursor. (d) Statistical size distribution\nof ball-milled amorphous alloy particles. (e) XRD patterns and magnified 2θ regions of Cu-Fe\ncomposites fabricated via distinct sintering routes. (f) Phase coarsening behavior in Cu–Fe\ncomposites processed by different sintering techniques. UHS and UHSQ methods suppress\nelemental segregation, with UHSQ further enabling a defect-free interfacial morphology.\n6. I do not understand the correlation between ‘hardness of metallic glass’, ‘challenges for\ninterface attachment’ and Figures 2a and S1.\nOur response: We thank the reviewer for their comment, which has prompted us to explicitly\nstate the connection between the raw material properties and the need for our specific processing\nstrategy.\nThe core challenge arises from two intrinsic properties of the as-received metallic glass (MG)\npowders (originally shown in our manuscript's Fig. 2a and S1): 1) Exceptional Hardness, which\nprevents the particles from deforming to create mechanical interlocking, and 2) Smooth Surface\nThe direct result of simply mixing and sintering these raw powders is a weak and unreliable\ninterface. We now provide clear visual evidence of this problem in Figure R4. These SEM images\nof a fractured composite prepared without our pre-processing step explicitly show a \"Poor\ninterface,\" characterized by widespread interfacial debonding and clean particle pull-out. This is\nthe \"challenge for interface attachment\" that our work overcomes.\nOur solution is a two-part synergistic strategy. First, we introduce a crucial pre-processing step:\nhigh-energy ball milling (HEBM). By enveloping the rigid MG spheres in a pliant copper sheath,\nthe process converts the intrinsically challenging MG–Cu interface into a readily sintered Cu–Cu\ncontact. Second, our UHSQ process capitalizes on this improved initial contact. The ultrafast high-\ntemperature sintering provides the energy for rapid and full densification, while the subsequent\nquenching locks the components into a robust, interlocked structure. In summary, the properties\nof the raw MG powder pose the challenge, and our combined HEBM + UHSQ process provides\nthe comprehensive solution. The revised manuscript presents this entire logical pathway with full\nclarity.\nAccordingly, we have revised the relevant section in the manuscript to address this point. For\nclarity, the revision is shown below:\n\"The exceptional hardness of metallic glass is a significant advantage but the smooth surface poses\nchallenges for interface attachment32(Figure 2a and Figure S1). Mechanical milling has been\nproven to be an effective method for enhancing the attachment between phases33\"\nFigure R4. Fracture surface morphology of the Cu-MG composite prepared without HEBM pre-\ntreatment. (a) Low- and (b) high-magnification SEM images show significant interfacial\ndebonding, characterized by the clean pull-out of MG particles (highlighted by dashed circles)\nfrom the Cu matrix.\n28\n7. It is difficult to realize what the authors mean by coarsening in ‘In contrast, UHS leverages\nultra-fast sintering rates to suppress coarsening’.\nOur response: We thank the reviewer for this insightful comment. The reviewer is correct that\nour use of the term \"coarsening\" was imprecise and confusing in this context, and we apologize\nfor the lack of clarity. We intended to describe two distinct phenomena related to the immiscible\nCu-Fe system, and we have revised the manuscript to use more accurate terminology.\nOur intended meaning, now clarified in the manuscript, is that the UHSQ process uses kinetic\ncontrol to bypass thermodynamic limitations. Conventional slow sintering allows the positive Cu–\nFe mixing enthalpy to induce extensive elemental segregation, yielding the phase-separated\nmicrostructure evident in our furnace-sintered samples. Concurrently, slow heating would lead to\nuncontrolled grain growth during the crystallization of the metallic glass reinforcement. In\ncontrast, the ultra-fast heating rate of our UHSQ process kinetically suppresses both of these\ndetrimental effects. The timescale is too short for significant long-range atomic diffusion, thus\npreventing elemental segregation. Simultaneously, the process confines the crystallization of the\nmetallic glass to the nanoscale, creating a desirable and stable in-situ nanostructure (Figure S13).\nTherefore, we agree \"coarsening\" was a poor choice. We have revised the manuscript to state,\nmore accurately, that the process is designed to\n\"In contrast, UHS leverages ultra-fast sintering rates to suppress elemental segregation and control\nthe crystallization of the metallic glass reinforcements.\"\ntesting, providing evidence for robust interfacial bonding. (a-c) High-magnification SEM images\nshowing the characteristic failure modes. The Cu matrix exhibits classic ductile fracture with fine,\nequiaxed dimples. Critically, the Fe-based reinforcement particle fails via brittle trans-particle\nfracture, with the crack propagating directly through the particle rather than along the interface.\n(d) Corresponding EDS elemental mapping confirms the chemical identity of the phases, showing\nthe fractured particle is rich in Fe, Cr, and Mo, surrounded by the Cu-rich matrix. The fracture\nbehavior demonstrates that the interfacial bond strength exceeds the intrinsic fracture strength of\nthe reinforcement phase itself.\n8. The title of the section could be Mechanical Properties. Please leave the decision for the\nreader to decide whether mechanical properties are good or not. Same also holds for the next\nsection.\nOur response: We thank the reviewer for this valuable guidance on improving the scientific rigor\nand objectivity of our manuscript. We completely agree that section titles should be neutral and\ndescriptive, allowing the presented data to speak for itself. Follow this excellent suggestion, we\nhave revised the section titles in the manuscript accordingly:\nThe section title \"Good Mechanical Properties\" has been changed to \"Multi-scale Strengthening\nMechanisms\". Apply the same principle, the title of the subsequent section has also been changed\nfrom \"Excellent Wear Performance\" to the more objective \"Hardness and Wear Behavior\"\nWe appreciate the reviewer's effort in helping us enhance the professionalism of our paper.\n9. Energy-dispersive X-ray spectroscopy (EDS) line scan analysis (most likely) in Scanning\nElectron Microscopy (SEM) is employed to demonstrate absence of cross diffusion (Figure\nS13). However, the analysis cannot be considered as the most reliable approach, considering\n30\nthe limitations of EDS in SEM. It is recommended to perform the analysis using Transmission\nElectron Microscopy, having a nano scale resolution.\nOur response: We are very grateful to the reviewer for this critical and expert comment. We\ncompletely agree that SEM-EDS is not sufficient to definitively prove the absence of diffusion at\nthe nanoscale and that TEM-EDS is the most rigorous approach. We apologize for the confusion;\nthis high-resolution TEM-EDS data was included in our original supplementary information, but\nwe failed to present it clearly and highlight its significance.\nTo directly resolve this issue, we have generated an updated, fully integrated Supplementary\nFigure S15 that delivers a hierarchical, multi-resolution characterization of the interfacial structure.\nAs shown in Figure S15a, the SEM-EDS line scan confirms a sharp elemental transition at the\nmicrometer scale, showing no large-scale segregation. More importantly, as the reviewer correctly\nrecommended, Figure S15b presents the high-resolution TEM-EDS analysis. This nanoscale\nanalysis provides definitive proof of a clean, abrupt bonding interface with an extremely sharp\nchemical gradient. The elemental transition zone between the Cu and Fe-based phases is confined\nto a width of less than 10 nm, which effectively demonstrates that our UHSQ process suppresses\natomic diffusion across the interface. We thank the reviewer for pushing us to present this crucial\ndata more clearly.\nFigure S15. Multi-scale analysis of the Cu-FeCrMo nanocrystal interface. (a) SEM-EDS line scan\nshowing a sharp elemental transition at the microscale. (b) High-resolution TEM-EDS line scan\nconfirming a chemically abrupt interface, which demonstrates the effectiveness of the UHSQ\nprocess in suppressing diffusion.\n31\n10. ‘The Fe55Cr25Mo16B2C2 metallic glass maintains a high hardness of approximately 1100\nHV.’ Metallic glass or the crystallized phase?\nOur response: We thank the reviewer for this insightful question. The reviewer correctly\nidentifies an apparent contradiction, which we acknowledge stemmed from our imprecise use of\nterminology. We have revised the text to resolve this ambiguity. We have thoroughly re-examined\nthis aspect and revised the manuscript to provide a much clearer and more accurate description.\nThe core of the clarification is this: the high hardness (~1100 HV) is indeed a property of a\ncrystalline structure (Figure S6), which is formed in-situ from the initial metallic glass powders\nduring our sintering process. To provide definitive evidence, we conducted a detailed\nmicrostructural investigation using TEM in the Figure S7, and the results are now presented in the\nrevised manuscript. Our analysis reveals the following:\nWe began with feedstock powders of the Fe Cr Mo B C composition, which were in an\n55 25 16 2 2\namorphous (metallic glass) state. This is the origin of the name used in the manuscript. During the\nUHSQ process, these amorphous particles underwent a partial in-situ crystallization. The final\nreinforcement phase is not amorphous but is, in fact, a complex nanocrystalline composite. As\nshown in the HAADF-STEM images and HRTEM analysis, this composite is composed of an Fe–\nCr ferrite, Mo-rich precipitates, and Fe-rich domains (ordered b2 intermetallic phase, retained\namorphous pockets, and pure Fe nanocrystals) (Figure S8).\nTherefore, the exceptional hardness is a direct result of this sophisticated, multi-phase\nnanocrystalline architecture, which benefits from both grain refinement and dispersion\nstrengthening. Based on this critical insight, prompted by the reviewer's comment, we have\nmeticulously revised the manuscript. We have corrected all instances of this ambiguity and now\nconsistently refer to the reinforcement phase as the \"multi-phase nanocrystals\" derived from the\noriginal metallic glass powders.\n(a-b) Analysis from two different regions of the reinforcement phase, demonstrating the consistent\nformation of nano-architecture. Each panel contains a HAADF-STEM image, a corresponding BF-\nS image, and EDS elemental maps for Fe, Cr, and Mo. The HAADF images use Z-contrast to\ndistinguish the phases, which are then chemically identified by the EDS maps. The analysis\nconfirms the structure consists of a continuous Fe-Cr matrix, discrete Mo-rich precipitates, and\nFe-rich areas occupying the interstitial spaces.\nreinforcement phase. (a-b) HAADF-STEM image revealing a complex nano-architecture\nconsisting of multiple phases distinguished, corresponding EDS elemental maps of the region\nboxed in (a), confirming the chemical distribution of Fe, Cr, and Mo. The analysis identifies a\nprimary Fe-Cr matrix, localized Mo-rich precipitates, and Fe-rich inclusions. (c) The higher-\nmagnification HAADF-STEM image providing a detailed view of the phase morphology and\ndistribution. (d) Atomic-resolution HRTEM images with corresponding Fast Fourier Transform\n(FFT) patterns for the Fe-rich, Fe-Cr, and Mo-rich regions, respectively.\nTEM image of the Fe-rich areas, revealing a nanocomposite structure. (b) HRTEM image of a\ncrystalline domain; inset FFT indicates a primary bcc lattice with superlattice reflections consistent\nwith a coexisting b2 phase. (c) HRTEM image showing a region of untransformed amorphous\nphase; inset FFT displays a diffuse halo. The Fe-rich reinforcement comprises bcc-Fe, b2\nintermediate phase, and disordered nanocrystals regions.\nReply to the further concerns and comments (Reviewer #3)\nUnits of measurement are not kept consistent throughout the manuscript!\nOur response: We thank the reviewer for pointing out this critical inconsistency. We have\nperformed a thorough review of the entire manuscript and have now standardized all units of\nmeasurement to ensure they are consistent throughout the text, figures, and tables.\nAbbreviations are used without defining them. Please do not use abbreviations in the conclusion\nsection\nOur response: We appreciate this important reminder. We have carefully revised the manuscript\nto ensure that all abbreviations are explicitly defined upon their first use. In accordance with the\nreviewer's suggestion, we have also removed all abbreviations from the Conclusion section to\nimprove its readability and impact.\nScale bar in some Figures is missing, for most it is provided on the right bottom side; however,\nfor one part of Figure 2f it is shown on a different side.\nOur response: To ensure clarity and consistency, we have reviewed all figures in both the main\nmanuscript and the supplementary information. We have now added the missing scale bars and\n35\nstandardized their position to the bottom right corner for all images. We apologize for this previous\ninconsistency.\nFurther advanced analysis could be used to support the statement of porous weak interface (Figure\nS3 a-c).\nOur response: This is a good suggestion to strengthen our claim. To provide more direct evidence\nfor the \"porous weak interface,\" we have conducted additional analysis and included new, high-\nmagnification SEM images in the revised manuscript (now designated as Figure S4 in the\nSupplementary Information). These new images clearly reveal the microporosity and areas of\nincomplete bonding at the original particle interfaces, thus providing strong visual support for our\nstatement. The manuscript text has been updated to reference this new evidence.\nDescriptions for Figures’ captions still need improvement, and revision.\nOur response: Following the reviewer's advice, we have undertaken a comprehensive revision of\nall figure captions throughout the manuscript and supplementary information. They have been\nrewritten to be more descriptive, precise, and self-contained, ensuring they provide all necessary\ninformation for a clear understanding of each figure.\nFigure S11, the corresponding element labels are missing.\nOur response: We thank the reviewer for pointing out this oversight. The element labels were\nmissing from the original figure. We have revised new Figure S13 to include all the corresponding\nelement labels for clarity and apologize for the omission.\nFigure S13, the corresponding micrographs are not provided, if any. In addition, in the manuscript\nit noted ‘mapping’.\nHowever, Figure S13 is line scan analysis.\nOur response: We thank the reviewer for identifying this error. We have revised new Figure S15\nby correcting the terminology to 'line scan analysis', adding corresponding SEM/TEM micrographs\nto show the scan location, and redrawing the figure for clarity.\nconventional UHS method. (a-b) SEM micrographs providing a general overview of the UHS-\nsintered control sample. These images demonstrate that interfacial defects between the FeCrMo\nnanocrystal, particles and the Cu matrix are widespread throughout the material. (c) A higher-\nmagnification micrograph of a single FeCrMo nanocrystal, presenting a clear, representative\nexample of poor interfacial bonding.\n(UHSQ) method for fabricating high-strength positive mixing enthalpy Cu-Fe-based composites,\nand it provides strong evidence supporting the claimed benefits of this method. The resulting\ncomposite materials exhibit significantly enhanced mechanical properties compared to pure\ncopper and materials processed by conventional methods. Overall, the work is of very high quality\nand innovative, and can be considered for publication in Nature Communications after addressing\nthe comments as below:\n1. The extent and nature of the \"partial devitrification\" should be further elucidated in this work.\nThe manuscript mentions weak crystalline peaks in UHSQ specimens, suggesting the\nformation of \"nanoscale intermetallic compounds\". While this can contribute to the enhanced\nproperties, the phase characterization, including crystal structure identification, phase\nvolume fraction, size distribution, and perhaps even thermal stability, should be discussed in\nmore detail. If these nanocrystals are not stable at the intended application temperatures or\nif their formation is inconsistent, it could lead to degraded properties over time. Therefore, a\nmore detailed characterization of these fine intermetallic compounds is suggested, as it will\nhelp understand the fundamental mechanisms at play. As a start, an isothermal treatment to\ngrow the phase bigger will help identify the crystal structure and phase composition for these\nintermetallic compounds. TEM or atom probe should have a good chance of assisting in this\nanalysis. Otherwise, authors should justify with a convincing statement.\nOur response: We are very grateful to the reviewer for this good and highly constructive\ncomment. The reviewer is correct: a rigorous, quantitative characterization of the partially\ndevitrified phases is indispensable for elucidating the operative strengthening mechanisms and\nfor evaluating the material’s long-term structural stability. We fully acknowledge that the\noriginal manuscript lacked the requisite depth and clarity in this regard. To remedy this\ndeficiency, we have undertaken an exhaustive microstructural re-analysis and have\ncomprehensively revised the text, now underpinned by a suite of newly acquired, high-resolution\nfigures and supplementary data.\nOur detailed characterization, now presented in Figure S6, reveals that the \"partial\ndevitrification\" of the Fe-based reinforcement results in a desirable, in-situ formed, multi-phase\nnanocrystals composite.\na) Phase Characterization and Distribution: Using High-Angle Annular Dark-Field\nScanning Transmission Electron Microscopy (HAADF-STEM), we identified three\ndistinct phase regions based on their Z-contrast (Figure S7 a,c). The corresponding\nEnergy-Dispersive X-ray Spectroscopy (EDS) elemental maps (Figure A7 b) confirm\ntheir composition: a) continuous Fe-Cr ferrite matrix, b) discrete, faceted Mo-rich\nprecipitates, and c) complex, multi-component Fe-rich domains.\nb) Crystal Structure Identification: As the reviewer correctly recommended, we performed\ndetailed crystallographic analysis. The atomic-resolution HRTEM imaging coupled with\nFast Fourier Transform (FFT) analysis, presented in Figure S7d, definitively identifies\nMo-rich phases as having a Body-Centered Cubic (BCC) crystal structure, albeit with\ndifferent orientations. This directly confirms the nature of the nanocrystalline ensemble.\nthe Mo-rich precipitates are Body-Centered Cubic (BCC) (Figure R2d), the Fe-rich\ndomains are themselves a nanocomposite. Within these domains, we have definitively\nidentified (Figure S8): a) an ordered b2 intermetallic phase, confirmed by its\ncharacteristic {010} superlattice reflections in the FFT patterns; b) disordered\nnanocrystals regions, evidenced by diffuse halos in FFTs; and c) pure Fe nanocrystals\nwith a BCC structure. This multi-phase nature, containing crystalline, intermetallic, and\ndisordered nanocrystals regions.\nc) Size Distribution and Volume Fraction: The size distribution of the constituent phases,\ndetermined from statistical TEM analysis, is shown in Figure 3c. The Mo-rich\nprecipitates exhibit a narrow distribution centered at approximately 100 nm. The Fe-rich\nnanocrystals have a broader distribution peaking around 180 nm, while the Fe-Cr matrix\ngrains are the largest, with a mean size of approximately 250 nm. Separate, area-based\nimage analysis indicates the Mo-rich precipitates occupy approximately 14 vol% of the\nreinforcement phase.\nd) Advanced Strengthening Mechanisms: Beyond simple phase identification, our HRTEM\nanalysis reveals significant lattice distortion and pronounced coherency-strain fields.\nThese \"strain walls\" act as powerful obstacles to dislocation motion. Crucially, the\nnewly identified phases introduce even more potent strengthening effects: the hard,\nordered b2 intermetallic phase provides classic compound strengthening, the disordered\nnanocrystals regions serve as glide-plane-free barriers, and the exceptionally high\ndensity of diverse inter-phase boundaries (e.g., crystalline-amorphous, crystalline-b2)\ncreates a formidable, multi-tiered network of obstacles. This provides a synergistic\nstrengthening effect that supplements traditional Hall-Petch grain boundary\nstrengthening.\ne) Thermal Stability: Regarding the crucial point of thermal stability, our existing high-\ntemperature mechanical test data provides strong evidence. The ability of the composite\nto maintain a high yield strength of 290 MPa at 650°C demonstrates that this complex\nnanocrystalline architecture, now understood to be bolstered by the intrinsic thermal\nstability of the ordered b2 phase and sluggish diffusion within the amorphous remnants,\nis stable and does not degrade significantly at this application temperature.\nIn summary, we have now provided a detailed, multi-faceted characterization of the devitrified\nphases, confirming their composition, crystal structures (BCC, B2, and amorphous), size, and\nvolume fraction, while also elucidating their complex, synergistic contribution to strengthening\nand their operational stability. We have incorporated this comprehensive analysis and the new\nfigures into the revised manuscript. We thank the reviewer again for pushing us to so\nsignificantly strengthen this core aspect of our work.\n\"The rapid quenching process induces the formation of a multiscale structure within the Fe-based\nmetallic glass (Figure S5). Bright-field TEM imaging reveals a regular and uniform distribution of\nnanoscale crystal particles within the metallic glass (Figure 3c). Figure 3b collectively demonstrate\nthat the Fe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass devitrifies into a multi-phase nanocrystal, in which\n39\nnanocrystalline Fe–Cr ferrite forms the continuous skeleton of the microstructure (Figure S6). The\nfcc Fe–Cr matrix surrounds and mechanically pins two secondary constituents, Mo-rich\nnanocrystals and Fe-rich areas. Bcc Mo-rich nanocrystals that populate junctions and grain corners\nin a quasi-periodic manner, occupying ~14 vol % of the microstructure (Figure S7), surrounding\neach Mo grain is a shell of bcc Fe ferrite. Notably, in the Fe-rich areas, containing both pure Fe\nand disordered nanocrystals regions (Figure S8a). Furthermore, these Fe-rich areas contain ordered\nb2 intermetallic phases, contributing to the overall intricate structure (Figure S8b). The complex\narrangement indicates an incomplete devitrification process, resulting in a multiphase\nnanocrystalline structure with retained amorphous pockets and the formation of ordered\nintermetallic compounds (Figure S8c). The overall microstructure exhibits a high density of\ninterfaces and structural complexity. This intricate structural architecture, arising from the\ncontrolled devitrification of the initial metallic glass, significantly influences the material's\nproperties.\"\n(a-b) Analysis from two different regions of the reinforcement phase, demonstrating the consistent\nformation of nano-architecture. Each panel contains a HAADF-STEM image, a corresponding BF-\nS image, and EDS elemental maps for Fe, Cr, and Mo. The HAADF images use Z-contrast to\ndistinguish the phases, which are then chemically identified by the EDS maps. The analysis\nconfirms the structure consists of a continuous Fe-Cr matrix, discrete Mo-rich precipitates, and\nFe-rich areas occupying the interstitial spaces.\nreinforcement phase. (a-b) HAADF-STEM image revealing a complex nano-architecture\nconsisting of multiple phases distinguished, corresponding EDS elemental maps of the region\nboxed in (a), confirming the chemical distribution of Fe, Cr, and Mo. The analysis identifies a\nprimary Fe-Cr matrix, localized Mo-rich precipitates, and Fe-rich inclusions. (c) The higher-\nmagnification HAADF-STEM image providing a detailed view of the phase morphology and\ndistribution. (d) Atomic-resolution HRTEM images with corresponding Fast Fourier Transform\n(FFT) patterns for the Fe-rich, Fe-Cr, and Mo-rich regions, respectively.\nTEM image of the Fe-rich areas, revealing a nanocomposite structure. (b) HRTEM image of a\ncrystalline domain; inset FFT indicates a primary bcc lattice with superlattice reflections consistent\nwith a coexisting b2 phase. (c) HRTEM image showing a region of untransformed amorphous\nphase; inset FFT displays a diffuse halo. The Fe-rich reinforcement comprises bcc-Fe, b2\nintermediate phase, and disordered nanocrystals regions.\ncomposites. (a) Microhardness profiles measured from the sample surface inwards for the UHSQ\ncomposite, UHS composite, and monolithic pure Cu. The shaded bands represent the standard\ndard\ndeviation of the measurements. (b) Representative microhardness line scan across the Cu matrix,\nthe Cu-FeCrMo interface, and multi-phase nanocrystals. (c) The plot of microhardness versus\nelectrical conductivity (% IACS), benchmarking the UHSQ composite against various material\nclasses reported in the literature 43-60. (d) Wear test results showing the wear loss rate as a function\nof friction cycles under a load of 10 MPa and a sliding speed of 500 r/min. (e) Optical micrographs\nof the worn surfaces of UHSQ composites with 25 vol% and 50 vol% Fe-MG reinforcement,\nshowing uniform wear tracks. (f) High-magnification SEM images tracking the morphology of a\nsingle multi-phase nanocrystals on the wear surface after 15 and 60 minutes of sliding.\n44\n2. It is also helpful to discuss the scale-up potential of such a technique, as the tensile test\nsamples are very limited in terms of dimension compared to ASTM standards. Note that there\nis no need to be convinced that this technique can produce everything on a large scale, but it\nis important to mention related limitations or any specific requirements to scale up.\nOur response: We thank the reviewer for this important and practical question regarding the\nscalability of our UHSQ technique. We agree that a discussion on this topic is a valuable addition\nto the manuscript.\nScaling up our UHSQ process for industrial application is highly feasible. The underlying principle\nof ultrafast Joule heating has already been successfully implemented in high-throughput processes\nlike roll-to-roll manufacturing1-5. This provides a strong precedent for scaling the 2D dimensions\n(length and width) of our composites to very large areas, which is a key strength of our approach.\nThe primary consideration for scaling is maintaining microstructural uniformity, which is\ngoverned by heat transfer during the rapid heating and quenching cycles. While this presents a\nchallenge for the thickness dimension, our current process has demonstrated the capability to\neffectively treat components with thicknesses approaching the centimeter scale. This level of\nperformance is already sufficient for a wide range of demanding industrial applications, including\nhigh-performance plates, wear-resistant surface layers, and structural components.\nTherefore, our UHSQ technique is immediately promising for a significant class of industrial\nproducts. While scaling to produce much thicker bulk components would be a subject for future\nengineering innovation, the current demonstrated capability is already of great practical relevance.\nWe have added this balanced discussion, emphasizing both the substantial current potential and\nfuture considerations, to our revised manuscript.\n\"In terms of scalability, the process's good speed enables high-throughput manufacturing of large-\narea sheets and coatings, with current capabilities extending to near-centimeter thick bulk\ncomponents; the primary focus for future scaling is addressing the engineering challenge of\nuniform quenching in even larger volumes. At its core, the method's success lies in the microscale\ndovetail-like structure induced by rapid quenching, which produces an unprecedented\nstrengthening effect at the interface.\"\nReferences\n1. Deng, B. et al. Rare earth elements from waste. Sci. Adv. 8, eabm3132 (2022).\n2. Advincula, P. A. et al. Replacement of Concrete Aggregates with Coal-Derived Flash\nGraphene. ACS Appl. Mater. Interfaces 16, 1474–1481 (2024).\n3. Du, P. et al. Roll-to-Roll Flash Joule Heating to Stabilize Electrocatalysts onto Meter-Scale Ni\nFoam for Advanced Water Splitting. ACS Nano 19, 1327–1339 (2025).\n4. Cheng, Y. et al. Rapid electrothermal rejuvenation of spent lithium cobalt oxide cathode.\nEnergy Environ. Sci. 18, 6085–6093 (2025).\n5. Luong, D. X. et al. Gram-scale bottom-up flash graphene synthesis. Nature 577, 647–651\n(2020).\n(Black italic: Reviewer’s remarks; Blue: Our response;\nFig. RX appears only in the response.)\nFirst, we appreciate the reviewers' time and efforts spent reviewing our manuscript. We very much\nappreciate the reviewers' thoughtful comments and suggestions, which have significantly\nimproved the quality and clarity of our manuscript. We have performed additional experiments\nand revised the manuscript to address all comments and concerns. The changes are highlighted in\nyellow in the revised manuscript. Our point-by-point responses are presented below.\nReviewer #1:\n1. Please check the full text and unify the symbol “-”, for example, the “Cu-Fe” in Figure 2 (f)\nof the figure title.\nOur response: We sincerely thank the reviewer for the careful reading and for pointing out this\ninconsistency. We have reviewed the entire manuscript, including the main text, all figures, and\ntheir captions. The use of the hyphen symbol (“-”) has now been unified throughout the document\nto ensure consistency.\n2. The Vickers hardness values obtained in this work are questionable. The average hardness\nvalue is as high as 931HV, while the tensile yield strength is only 684MPa, which is hard to\naccept. It is recommended to use nanoindentation techniques to determine the hardness of\nthe interface, matrix, and reinforcement. SEM images of the indentations should be provided\nas supporting evidence.\nOur response: We sincerely thank the reviewer for this critical and constructive comment. We\nagree that elucidating the relationship between macroscopic Vickers hardness and bulk tensile\nyield strength in complex composite materials necessitates a more detailed, multi-scale\ninvestigation.\nFollowing the reviewer's suggestion, we have performed in-situ nanoindentation tests to precisely\nmeasure the mechanical properties of the individual microstructural constituents.\nWe utilized an in-situ nanoindenter (FemtoTools) integrated within a scanning electron\nmicroscope (SEM), as shown in Fig. S18a, which enabled the precise placement of indents on\nselected phases. Fig. S18b, d provide representative SEM images of an array of nano-indents\nsystematically placed across the Cu matrix, the FeCrMo nanocrystalline reinforcement, and the\ninterfacial regions. The quantitative hardness map, presented in Fig. S18c, reveals the distinct\nproperties of each phase.\nTo directly compare these values with the macroscopic Vickers hardness, we use the approximate\nconversion Hv=H /9.81×c, where H is the indentation hardness in GPa and c is a conversion\nIT IT\nfactor often close to 1. For simplicity in direct comparison, we approximate 1 GPa=100 HV.\nCombined with Fig. S19 the results are as follows:\n1\n FeCrMo Nanocrystals (Reinforcement): Exhibit an ultra-high hardness of approximately 12-\n Cu-FeCrMo Interface: Shows a high transitional hardness value of 5-8 GPa, corresponding to\n~500-800 HV, indicating strong interfacial bonding and effective load transfer.\n Cu Matrix: Possesses a hardness of 2-3 GPa, corresponding to ~200-300 HV. This value is\nsignificantly higher than that of pure annealed copper (~90 HV), which we attribute to grain\nrefinement into the sub-micron range and particularly small Fe₅₅Cr₂₅Mo₁₆B₂C₂ particles (as\nshown in Fig. 3a, b) are distributed diffusely during the UHSQ process.\nWe think that high macroscopic Vickers hardness (avg. 931 HV) is a composite \"rule of mixtures\"\nproperty, heavily influenced by the volume fraction (50 vol%) of the extremely hard FeCrMo\nreinforcement (~1200–1400 HV). In contrast, the bulk tensile yield strength (684 MPa) is governed\nby plastic deformation of the continuous, ductile Cu matrix.\nprocessed composite. (a) Photograph of the Femto-Tools nanoindenter setup used for the in-situ\nmeasurements inside a Scanning Electron Microscope (SEM). (b) SEM image showing a precise\narray of indentations placed across the Cu matrix, FeCrMo nanocrystals, and their interfaces. (c)\nCorresponding quantitative hardness map (in GPa) of the indented area. (d) Detailed close-up SEM\nmicrograph of representative indentations, visually confirming the vast hardness difference\nquantified in panel (c).\nFigure S19. Nanoindentation load-displacement curves of the constituent phases.\ncomposites. (a) Microhardness profiles measured from the sample surface inwards for the UHSQ\ncomposite, UHS composite, and monolithic pure Cu. The shaded bands represent the standard\ndeviation of the measurements. (b) Representative microhardness line scan across the Cu matrix,\nthe Cu/FeCrMo interface, and multi-phase nanocrystals. (c) The plot of microhardness versus\nelectrical conductivity (% IACS), benchmarking the UHSQ composite against various material\nclasses reported in the literature 43-60. (d) Wear test results showing the wear loss rate as a function\nof friction cycles under a load of 10 MPa and a sliding speed of 500 r/min. (e) Optical micrographs\nof the worn surfaces of UHSQ composites with 25 vol% and 50 vol% Fe-MG reinforcement,\nshowing uniform wear tracks. (f) High-magnification SEM images tracking the morphology of a\nsingle multi-phase nanocrystals on the wear surface after 15 and 60 minutes of sliding.\n4\n3. The \"Multi-scale Strengthening Mechanisms\" are emphasized in the manuscript, but the\nanalysis of them is vague. How do the multi-scale strengthening phases contribute to the\nstrength?\nOur response: We thank the reviewer for this insightful comment. We have revised the \"Multi-\nscale Strengthening Mechanisms\" section to explicitly delineate the contributions at different\nlength scales.\n Co-continuous interlocking network: Macroscopically, the Cu and the Fe-based reinforcement\nphases form two independent, interpenetrating 3D networks. The continuous, high-strength\nFe-based network acts as the primary load-bearing skeleton, while the continuous, ductile Cu\nmatrix imparts toughness and prevents catastrophic brittle failure.\n Strong interfacial bonding and mechanical interlocking: The UHSQ process creates a pristine\ninterface between the Cu matrix and the Fe-based reinforcement, characterized by nanoscale\nrivet-like interlocking structures (as shown in Fig. 3b). The strong physical bond is critical\nfor effective load transfer from the softer Cu matrix to the much harder reinforcement phase.\n Intrinsic strength of the multi-phase nanocrystalline reinforcement: The original\nFe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass devitrifies during the UHSQ process into a complex, multi-\nphase nanocrystalline structure. This structure consists of an Fe-Cr ferrite matrix, Mo-rich\nnanocrystals, and ordered B2 intermetallic phases. Such a hierarchical nanocrystalline\ncomposite possesses exceptionally high intrinsic hardness and strength. The phase acts as the\nprimary source of the composite's high hardness and wear resistance.\n Dispersion strengthening within the Cu matrix: As observed in Fig. 3h, smaller\nFe₅₅Cr₂₅Mo₁₆B₂C₂ particles are dispersed within the Cu matrix. These hard, undeformable\nparticles act as obstacles to dislocation motion within the ductile copper, effectively increasing\nthe yield strength of the matrix itself via a dispersion strengthening mechanism (akin to\nOrowan strengthening).\nWe contend overall strength is not attributable to a single factor but is a product of these integrated,\nmulti-scale mechanisms. The macro-scale network establishes a robust framework, the strong\ninterface enable this framework to behave as a single unit, and the high intrinsic strength of both\nthe reinforcement phase and the dispersion-strengthened matrix provides the high stress-bearing\ncapacity.\n55 25 16 2 2\nmetal matrix composites. (a) TEM image showing the clean, well-bonded interface between the\nCu matrix and the Fe-based reinforcement. (b) Higher-magnification view revealing the nanoscale\ninterlocking structure, with corresponding EDS maps confirming the distinct phases. (c) HAADF-\nSTEM analysis of the devitrified reinforcement, revealing a nano-architecture of an Fe-Cr matrix,\nMo-rich precipitates, and Fe-rich inclusions, confirmed by EDS mapping. The plot below\nquantifies the size distribution of these nanocrystals. (d) Room-temperature tensile engineering\nstress-strain curves for the composites processed by UHS and UHSQ, compared with a monolithic\nCu control sample. (e) Elevated-temperature (650 °C) tensile engineering stress-strain curves for\nthe UHS and UHSQ composites. (f) SEM image of the tensile fracture surface of UHS sample at\nroom temperature and (g) SEM image of the tensile fracture surface of UHSQ sample at room\ntemperature. (h) SEM image showing the distribution of small-sized particles in the Cu matrix.\n6\n4. The analysis of mechanical and electrical properties in this work is not sufficiently in-depth.\nThe remarkable performance is only qualitatively analyzed. Please quantify the properties,\nespecially the mechanical properties.\nOur response: We thank the reviewer for the valuable feedback regarding the quantitative analysis\nof the material's properties. We agree that a more explicit and in-depth discussion of the\nquantitative data will strengthen our manuscript. While the mechanical properties data is presented\nin Fig. 3, Fig. 4, Fig. S9, Fig. S14, Fig. S18 and Fig. S19.\n Quantification of Mechanical Properties:\nTensile Strength: The optimized UHSQ-processed composite exhibits an ultimate tensile strength\nof approximately 685 MPa at room temperature, representing a 6.5-fold enhancement over the\npure Cu reference..\nHigh-Temperature Strength: The material demonstrates remarkable thermal stability, maintaining\na yield strength of 290 MPa at an elevated temperature of 650°C.\nHardness: The average Vickers hardness of the UHSQ composite is 931 HV, with a maximum\nvalue reaching 1297 HV. This represents a nearly 10-fold increase compared to pure Cu.\nWear Resistance: Under a 10 MPa load at 500 rpm, the UHSQ composite shows a 40–50 times\nimprovement in wear resistance over the pure Cu sample.\n Discussion of Electrical Properties:\nRegarding the electrical properties, we measured conductivity of 15-20% IACS. We acknowledge\nthis is a moderate value for a copper-based composite and have expanded our discussion to provide\na clear physical explanation. The limited conductivity is an inherent consequence of the\ncomposite's unique microstructure, which is essential for its mechanical performance. The high\nintrinsic electrical resistivity of the Fe₅₅Cr₂₅Mo₁₆B₂C₂ multi-phase nanocrystalline reinforcement\nis the primary factor. Since the hard phase forms a co-continuous, interlocking network with the\ncopper, it inevitably disrupts the electron transport pathways within the conductive Cu matrix,\nthereby limiting the bulk conductivity.\nWe reiterate that the central innovation of this work is the development of the UHSQ methodology\nwhich enables the synthesis of bulk, structurally-sound positive mixing enthalpy composites—a\nlong-standing challenge due to liquid-phase immiscibility. The primary objective of this study was\nto establish the feasibility of the UHSQ route and to document the resulting improvements in\nmechanical and tribological performance. Future efforts targeting architectural optimization and\nleveraging advances in metallic-glass-powder synthesis to engineer finer conductive networks are\nexpected to yield substantial gains in electrical performance.\nThese quantitative performance metrics are also summarized in the abstract of manuscript.\n\" The mechanical tensile strength of the resulting composite is approximately 8 times greater than\nthat of pure Cu. The wear resistance improves by 40-50 times, and the Vickers hardness increases\nby a factor of 10.\"\n55 25 16 2 2\nprepared by different sintering methods. Bar chart showing ultimate tensile strength of composites\nwith varying Cu content (25-50 vol%) processed by UHS and UHSQ at room temperature and\n650°C. Inset shows representative engineering stress-strain curves highlighting great mechanical\nproperties improvement in UHSQ samples. Inset SEM fractography of room temperature tensile\nUHSQ processed samples.\nFigure S14. Comparison of flexural stress-strain curves for composites processed via UHSQ and\nUHS at room temperature. The data shows that the UHSQ composite exhibits a significantly higher\nflexural strength (approaching 900 MPa) compared to the control sample processed without\nquenching (UHS).\nprocessed composite. (a) Photograph of the Femto-Tools nanoindenter setup used for the in-situ\nmeasurements inside a Scanning Electron Microscope (SEM). (b) SEM image showing a precise\narray of indentations placed across the Cu matrix, FeCrMo nanocrystals, and their interfaces. (c)\nCorresponding quantitative hardness map (in GPa) of the indented area. (d) Detailed close-up SEM\nmicrograph of representative indentations, visually confirming the vast hardness difference\nquantified in panel (c).\nFigure S19. Nanoindentation load-displacement curves of the constituent phases.\n1. It is necessary to clarify in detail whether the unique microstructure and mechanical\nproperties are obtained as a result of a) the high sintering temperature, b) short sintering\nduration. It is not clear why the authors have done experiments at a temperature of 600 °C,\nat which mass transfer mechanisms are hardly active.\nOur response: We thank the reviewer for this insightful comment. We agree with the reviewer\nthat both high temperature and short duration are critical parameters. The ultra-short duration is\ncrucial for kinetically trapping the system in a non-equilibrium, homogeneously intermixed state,\nthereby suppressing the thermodynamically driven phase segregation.\nWe have conducted a new set of comparative experiments at a high temperature of 1500 °C,\nvarying the holding time. The results are presented in Fig. R1 below.\n Sintering at 1500 °C for 120 seconds: As shown in the SEM micrograph and corresponding\nEDS maps (Fig. R1a-b), a prolonged duration at high temperature leads to severe phase\nsegregation. The Cu and Fe-rich phases have coarsened into large, distinct domains, which is\nthe expected outcome for a positive mixing enthalpy system approaching thermodynamic\nequilibrium.\n Sintering at 1500 °C for 10 seconds: In stark contrast, when the sintering duration is limited\nto just 10 seconds, the composite retains its fine, interlocked microstructure (Fig. R1a-b). The\nFeCrMo nanocrystals remain uniformly distributed within the Cu structure, and large-scale\nelemental segregation is effectively prevented.\nThese results clearly demonstrate that the ultra-short sintering duration is the dominant factor in\npreserving the desired microstructure. While a high temperature provides the driving force for\nsintering, it also promotes segregation. Our UHSQ method leverages this by using a timeframe so\nbrief that the atoms do not have sufficient time to diffuse over long distances and cause phase\nseparation, thus achieving a well-bonded and homogeneous composite. To resolve the ambiguity\narising from the 600 °C data, we provide additional results obtained at 1500 °C that directly answer\nthe reviewer’s valid question.\nof Cu-Fe₅₅Cr₂₅Mo₁₆B₂C₂ composites. (a) SEM micrograph and (b) corresponding EDS elemental\nmaps of a sample sintered for 120 seconds. Severe phase segregation and coarsening of the Cu-\nrich and Fe-rich domains are evident. (c) SEM micrograph and (d) corresponding EDS elemental\nmaps of a sample sintered for only 10 seconds, showing a homogeneous and finely dispersed\ninterlocked microstructure.\n2. Figure R3 and Figure 1 from the manuscript are in contrast. The authors are required to\nclarify how the experiments were done. To promote open science practices, I highly\nrecommend the authors to share a video of the experimental procedure using Ultrafast High-\ntemperature Sintering and quenching. From Figure R3, it appears the specimen was placed\nbetween zirconia fiber insulators. However, this is not demonstrated in the manuscript. Why\nwere zirconia insulators used? Figure R3 (a) is really difficult to understand and not clear at\nall. Figure R3(d), is there any sample between the insulators? Figure R3 (e) and (f), are these\nprofiles for the sample, carbon felt, or zirconia insulators?\nOur response: We sincerely thank the reviewer for these detailed questions and the suggestion to\nenhance the clarity of our experimental procedure. We have provided new supplementary materials\nto clarify our custom-built UHSQ system.\n Clarification of the Experimental Process and Video Evidence\nThe apparent contrast between Fig. 1 (in the manuscript) and the last Fig. R3 arises from the\ndifferent imaging conditions used. In Fig. 1e, we used a Neutral Density (ND) filter to reduce the\nintense brightness of the incandescent sample, which allowed us to clearly capture the physical\nstate of the sample and its surroundings during the sintering process. In contrast, the photographs\nin the last Fig. R3c are presented without a filter to convey the true intensity of the rapid, high-\n11\ntemperature heating.\nTo fully address the reviewer's request for clarity and to promote open science, we have prepared\ntwo new supplementary videos (Video S1 and Video S2). Video S1 shows the entire UHSQ\nprocess in real-time. Video S2 presents the sinter process filmed through an ND filter.\nFurthermore, we have created Fig. S23 to be much clearer, explicitly delineating the upper \"UHS\"\n(sintering) stage and the lower \"Quench\" stage of the apparatus.\n Purpose of the Zirconia (ZrO₂) Fiber Insulators\nThe reviewer's observation is correct; we use ZrO₂ fiber cloth as an insulating layer, which is a\ncritical component of our setup. UHSQ system relies on Joule heating, where a high current is\npassed through the resistive carbon felts to generate heat. As a Cu-based composite, the sample is\nhighly electrically conductive; direct contact with the carbon felt would establish a low-resistance\nshunt path, severely diminishing Joule heating efficiency and compromising temperature control.\nThe ZrO₂ fiber cloth, a commercially available and robust ceramic material, is used as a high-\ntemperature electrical insulator to separate the conductive sample from the carbon felt heaters. As\nshown in the Fig. S23b, the fiber cloth undergoes a pre-sintering step to remove binders, which\ncauses it to expand slightly, ensuring excellent and uniform insulation.\n Specific Clarifications for last Fig. R3 Sub-panels:\nLast Fig. R3d: We confirm that a sample is indeed present between the two insulating layers in\nthis thermal image. The image was captured by an infrared pyrometer, whose spatial resolution is\ninherently limited.\nLast Fig. R3e: The temperature profile represents the temperature measured at the surface of the\nsample itself. It shows the complete cycle of ultrafast heating, a short holding plateau for sintering,\nand subsequent ultrafast quenching.\nLast Fig. R3f: The profile demonstrates the rapid responsiveness of the carbon felt heaters. The\npulsed profile shows that the heating elements can be turned on and off almost instantaneously,\nhighlighting the precise temporal control of our system.\nTo further verify uniform high-temperature attainment across the sample, we have included a\nnew image sequence in Fig. S23c (captured with a strong ND filter). which shows that immediately\nafter the power is shut off, the sample itself continues to glow brightly and uniformly, confirming\nthat it has been effectively heated throughout. We want to show the UHSQ method first ensures\neffective and rapid sintering, and our introduction of the instantaneous quenching step further\nenables new possibilities in microstructural design and material property enhancement.\nRegard to fiber cloth, we emphasize the addition in the method of the manuscript:\n\"The UHSQ technique with ZrO fiber was subsequently applied, with sintering for 10 seconds\n2\nfollowed by a 5-second cooling period, to obtain a dense composite sheet.\"\nPhotograph of the overall UHSQ apparatus, showing the upper heating stage (UHS) and the lower\nquenching bath. (b) Schematic illustration of the heating principle, where a sample is insulated by\nZrO₂ fibers and heated by surrounding carbon felts. (c) Photographs of the UHS process in action\n(left) and the ZrO₂ fiber insulation material (right). (d) A representative thermal image captured\nby an infrared pyrometer during sintering, showing a uniform temperature distribution. (e) The\nmeasured temperature-time profile for a typical UHSQ experiment, detailing the ultrafast heating,\nsintering, and quenching stages. (f) A temperature profile demonstrating the system's capability\nfor rapid pulsed heating, validating its high thermal responsiveness.\n(a) The photograph of the overall experimental setup, illustrating the distinct upper UHS (Ultrafast\nHigh-temperature Sintering) stage and the lower Quenching stage containing a silicone oil bath.\n(b) A close-up view comparing the ZrO₂ fiber insulation layers before and after a pre-sintering\ntreatment, which is performed to remove binders and ensure effective electrical insulation. (c) The\nphotographic sequence demonstrating the principle of sample heating with ND filter. Immediately\nafter the power to the carbon heaters is turned off, the heaters dim while the sample itself remains\nincandescent.\nmatrix composites. (a) Schematic illustration of the \"interface riveting\" mechanism. An initially\nweak contact between immiscible Cu and Fe Cr Mo B C metallic glass particles is transformed\n55 25 16 2 2\ninto a strong, mechanically interlocked physical bond at the interface. (b) A typical temperature-\ntime profile of the UHSQ process, characterized by ultrafast heating (500~1000 K/s), a short\nsintering plateau at high temperature, and subsequent rapid quenching. (c) Schematic of the\nexperimental procedure, showing the sample being rapidly heated between two carbon electrodes\nand then quickly submerged into a silicone oil bath for quenching. (d) TEM images reveal rivet-\nlike connections formed by physical deformation between Cu and FeCrMo nanocrystal. (e) A\nphotographic sequence capturing the key moments of the actual UHSQ process. (f) The Cu-Fe\ncomposite pellet with positive mixing-enthalpy prepared by UHSQ sintering.\n15\n3. If zirconia insulator is used, the terminology of 'radiative joule-heating sintering' must be\ntreated carefully. Correspondingly, reconsideration is required. In addition, it is 'Joule'!\nFurthermore, in Spark Plasma Sintering, the amount of electric current that passes through\nthe sample depends the conductivity of the material.\nOur response: We thank the reviewer for this precise and important question.\nOur use of the term \"radiative Joule heating\" is a deliberate and accurate description of the physics\nat play. The heating process can be broken down as follows:\n Joule heating constitutes the sole heat source: a high direct current is passed through two\ncarbon felts that sandwich the sample, and their electrical resistance converts the current into\nthermal energy. To ensure that the felts—rather than the highly conductive Cu-based\ncomposite—act as the primary resistive heaters, a thin ZrO layer is interposed as an\n2\nelectrical insulator; this prevents shunting of current through the sample and thereby\npreserves the intended Joule-heating pathway in the carbon felts.\n The Source of Heat is Joule Heating: The core of our technique is indeed Joule heating. A\nhigh electrical current is passed directly through the two highly resistive carbon felts that\nsandwich the sample assembly. It is these carbon felts that act as the primary heating elements,\nconverting electrical energy into thermal energy via the Joule effect.\n The Mechanism of Heat Transfer to the Sample: The sample itself is not intended to be part\nof the electrical circuit. At operational temperatures of up to 1500 °C, the dominant mode of\nheat transfer from the incandescent carbon felts to the sample is thermal radiation. The hot\nfelts act as efficient radiators, rapidly and uniformly heating the sample's surface. Heat is\nalso transferred via conduction through the thin ZrO₂ fiber layers. Therefore, the term\n\"radiative\" correctly describes the primary mechanism by which the sample is heated.\n The Critical Role of the ZrO₂ Insulator: As the reviewer correctly infers, the ZrO₂ insulator\nis essential. Its purpose is to function as an electrical insulator, not as a thermal barrier.\nGiven that our Cu-based composite is highly conductive, direct contact with the carbon felts\nwould create a short circuit. This would divert the current through the low-resistance\nsample, preventing the carbon felts from heating up effectively. The insulator is therefore a\nnecessary measure to enable the Joule heating of the carbon felts, ensuring that they, and not\nthe sample, serve as the primary resistive heaters.\n Distinction from Conventional SPS: In many SPS configurations, the current is designed to\npass through the conductive sample and die, making them active components of the heating\ncircuit. In our method, we deliberately prevent this by using the insulator. This configuration\nenables rapid, controlled heating of highly conductive materials whose low resistance\nprecludes efficient Joule heating in conventional SPS.\n(Black italic: Reviewer’s remarks; Blue: Our response;\nFig. RX appears only in the response.)\nFirst, we appreciate the reviewers' time and efforts spent reviewing our manuscript. We deeply\nappreciate the reviewers' thoughtful comments and suggestions, which have significantly\nimproved the quality and clarity of our manuscript. We have performed additional experiments\nand revised the manuscript to address all comments and concerns. The changes are highlighted in\nyellow in the revised manuscript. Our point-by-point responses are presented below.\nReviewer #3:\n1. Units of measurement are not kept consistent throughout the manuscript!\nOur response: We thank the reviewer for identifying the inconsistencies in our units of\nmeasurement. We have now thoroughly reviewed the entire manuscript and standardized all units\nfor consistency. For example, all temperatures related to the sintering process are now uniformly\npresented in Kelvin (K) throughout the main text, figures, and all supplementary materials.\n\"This approach retains the high heating rate (500-1000 K /s) of the UHS process while providing\nan enhanced cooling rate (up to 1000 K /s at the sample surface; see Figure 1b; Figure S2 and\nVideo 2).\"\nIn addition, we have standardized key terminology for clarity. For instance, various descriptions\nof the material are now uniformly referred to as \"Cu-Fe-based composite\". We also consolidated\nthe Supplementary Information to eliminate redundancy; for example, the original Figures S11\nand S12 have been merged into the revised Figure S11. We have highlighted the changes in the\nmain text and supplementary files.\n2. The discussion on the importance of the short sintering duration is not comprehensive yet.\nFurther experiments at different time intervals and discussion must be performed and\nintegrated into the manuscript?\nOur response: We thank the reviewer for raising this important point. We have performed\nexperiments with varying sintering durations to investigate this effect, and the results are presented\nin Figure S8. These experiments confirm that a short sintering duration is critical to kinetically\nsuppress the severe phase segregation that occurs with prolonged heating.\nAs suggested, we have now integrated this discussion into the main manuscript:\n\"This interfacial failure is directly linked to the sintering duration; prolonging the high-temperature\nexposure allows for the severe phase segregation demonstrated in Figure S8d. Minimizing the\nsintering time is therefore critical to kinetically trap the system and suppress elemental segregation\n(Figure S8a-c).\"\nSEM micrographs and corresponding EDS elemental maps for samples sintered at 1773 K for (a)\n5 s, (b) 10 s, (c) 15 s, and (d) 20 s. While shorter durations (5-10 s) enable the samples to maintain\na relatively fine, intermixed structure, prolonged heating (15-20 s) leads to severe phase\nsegregation and coarsening of the Cu and Fe-rich phases.\n3. Explanation on the use of zirconia insulator must be also included in the manuscript?\nOur response: We thank the reviewer for this constructive suggestion. The role of the zirconia\nfiber is key to the reproducibility of our experiment.\nAs per the reviewer's recommendation, we have clarified its purpose in the manuscript. We have\nnow added this explanation in the main text and in the \"Methods\" section to enhance the clarity\nand completeness of our experimental description.\n\"Specifically, the precursor compact was wrapped in zirconia (ZrO₂) fiber cloth for electrical\ninsulation (Figure S1c), positioned between two carbon felts in an argon atmosphere, rapidly\nsintered for 10 seconds, and than immediately quenched in high thermal conductivity silicone oil\nat 243 K (Figure 1c).\"\n\"Subsequently, the UHSQ technique was applied. The composite billet was wrapped in ZrO₂ fiber\ncloth, which serves as a high-temperature electrical insulator to prevent short-circuiting between\nthe conductive sample and the carbon heaters.\"\n2\n4. There is still concern about Figure 1b and c. Which figure is reporting the real-time\ntemperature of the sample? Figure 1b or Figure R3e? Figure 1c must be consistent with the\ndeveloped procedure?\nOur response: We thank the reviewer for their feedback on the figures.\nIn response to the query, Figure 1b is a schematic profile designed to illustrate the ideal stages of\nthe UHSQ process. The representative, experimentally measured temperature data is provided in\nthe Supplementary Information as Figure S2 (Due to the setting limitations of the thermal infrared\nimager, the recorded temperature range for this measurement was 485 K to 2000 K); both are\nmutually unified. We also confirm that the process schematic in Figure 1c is consistent with our\ndescribed experimental procedure.\nWe have revised the manuscript to clarify it. The updated text now reads as follows:\n\"This approach retains the high heating rate (500-1000 K /s) of the UHS process while providing\nan enhanced cooling rate (up to 1000 K /s at the sample surface; see Figure 1b; Figure S2 and\nVideo 2).\"\nFigure S2. The measured temperature-time profile for a typical UHSQ experiment, detailing the\nultrafast heating, sintering, and quenching stages (Note: Due to the setting limitations of the\nthermal infrared imager, the recorded temperature range for this measurement was 485 K to 2000\nK).\nFe-based composite. (a) Schematic illustration of the \"interface riveting\" mechanism. An initially\nweak contact between immiscible Cu and Fe Cr Mo B C metallic glass particles is transformed\n55 25 16 2 2\ninto a strong, mechanically interlocked physical bond at the interface. (b) The schematic\ntemperature-time profile of the UHSQ process, characterized by ultrafast heating, a short sintering\nplateau at high temperature, and subsequent rapid quenching. (c) Schematic of the experimental\nprocedure, showing the sample, housed in a ZrO₂ die, being rapidly heated between two carbon\nelectrodes and then quickly submerged into a silicone oil bath for quenching. (d) Transmission\nelectron microscope (TEM) images reveal rivet-like connections formed by physical deformation\nbetween Cu and FeCrMo nanocrystal. (e) A photographic sequence capturing the key moments of\nthe actual UHSQ process. (f) The Cu-Fe-based composite pellet with positive mixing-enthalpy\nprepared by UHSQ sintering.\n4\n5. It is surprising to see that Figures numbering have not been updated in the revised\nmanuscript!!\nOur response: We sincerely apologize for this serious oversight regarding the figure numbering.\nThis was an unintentional error that occurred during our revision process, and we are grateful to\nthe reviewer for pointing it out.\nWe have now performed a thorough and careful review of the entire manuscript. All figure\nnumbers, along with their corresponding citations within the text, have been corrected to ensure\nthat they are sequential, accurate, and consistent throughout the document.\n6. What does the terminology of ‘radiative joule-heating’ mean?! This terminology must be\nrevised. Joule not joule (typo)!\nOur response: We thank the reviewer for this valuable feedback. We agree that the term \"radiative\njoule-heating\" was non-standard and have revised it to the standard term \"Joule heating\"\nthroughout the manuscript for improved clarity. For instance, the text has been corrected as\nfollows:\n\"The sequential stages of UHSQ processing – including precursor assembly, Joule heating\nsintering andquenching – are visually documented in Figure 1e and Figure S3.\"\n7. It is surprising to see the fire in the ‘Video - Real-time Video of the UHSQ Process’ during\nthe presintering step! Where does the oxygen come from?\nOur response: We thank the reviewer for their keen observation. The flame in the video confirms\nit was recorded in air. This video is a demonstration, filmed in an open environment to clearly\nvisualize the sintering process, which is difficult to be captured within the glovebox. All\nexperimental data in this manuscript were collected under a protective argon atmosphere, as shown\nwith the apparatus in Figure S1. To ensure this distinction is clear, we have added explanatory\nnotes in the video title.\nthe overall UHSQ apparatus, showing the heating stage (upper) and the quenching bath (lower).\n(b) Schematic illustration of the heating principle, where a sample is insulated by ZrO₂ fibers\nand heated by surrounding carbon felts. (c) Photographs of the UHS process in action (left) and\nthe ZrO₂ fiber insulation material (right).\n8. What does Ultrafast high temperature mean in ‘Ultrafast high temperature rapidly densifies\nthe sample’ in the ‘Video - Real-time Video of the UHSQ Process’?!\nOur response: We thank the reviewer for highlighting the imprecise language in the video\nannotation. We have since revised the annotation in the video to \"Joule heat rapid sintering\". This\nterm accurately describes the heating mechanism responsible for the rapid densification observed.\nFigure R1. Screenshot of updated annotations in the UHSQ demonstration video.\n6\n9. What does Joule heat radiation mean in the ‘Video - UHS Heating Process Visualized with\nan ND Filter’?!\nOur response: We thank the reviewer for their feedback on the video's terminology. Our intent\nwas to describe the heat transfer process from the heating element to the sample. We have revised\nthe video annotation to \"Joule heat from carbon felt to sample\" to accurately describe the physical\nprocess.\nFigure R2. Updated annotation in the demonstration video of UHS heating process.\n(Black italic: Reviewer’s remarks; Blue: Our response;\nFig. RX appears only in the response.)\nFirst, we appreciate the reviewers' time and efforts spent reviewing our manuscript. We deeply\nappreciate the reviewers' thoughtful comments and suggestions, which have significantly\nimproved the quality and clarity of our manuscript. We have performed additional experiments\nand revised the manuscript to address all comments and concerns. The changes are highlighted in\nyellow in the revised manuscript. Our point-by-point responses are presented below.\nReviewer #3:\n1. Manuscript has typos (than immediately quenched, andquenching, for example).\nOur response: We apologize for the carelessness in the previous draft. We have corrected the\nspecific typos highlighted by the reviewer, ensuring the final presentation is polished and error-\nfree.\n2. The terminology of Joule heating sintering is not appropriate yet. Please use 'ultra-high-\ntemperature rapid sintering', in the manuscript, Video 1 and 2. Please remove 'Joule heat\nrapid sintering', and 'Joule heat from carbon felt to sample' form both videos.\nOur response: We sincerely thank the reviewer for this valuable suggestion regarding the\nterminology. We have fully adopted the term \"ultra-high-temperature rapid sintering\" throughout\nthe manuscript. Correspondingly, we have carefully updated the subtitles in Video 1 and Video 2\nby removing the phrases \"Joule heat rapid sintering\" and \"Joule heat from carbon felt to sample\"\nto ensure perfect consistency with the text.\n3. Please add a small note within Video 1 in a corner to notify the video was recorded in ambient\natmosphere for demonstration.\nOur response: We have followed this suggestion. A note has been added to the corner of Video\n1 to confirm that the recording took place in an ambient atmosphere for demonstration.\n1"}, {"round": 2, "reviewer_comments": {"reviewer_1": "Comments to NCOMMS-25-21591A\nTitle: Breaking Immiscibility Barriers: Ultrafast Sintering of Interlocked Cu-Fe-Based Composites\nBased on the Reviewers' comments, the manuscript has been improved, but the following issues still require further revision:\n1. Please check the full text and unify the symbol “-”, for example, the “Cu-Fe” in Figure 2 (f) of the figure title.\n2. The Vickers hardness values obtained in this work are questionable. The average hardness value is as high as 931HV,\nwhile the tensile yield strength is only 684MPa, which is hard to accept. It is recommended to use nanoindentation\ntechniques to determine the hardness of the interface, matrix, and reinforcement. SEM images of the indentations should be\nprovided as supporting evidence.\n3. The \"Multi-scale Strengthening Mechanisms\" are emphasized in the manuscript, but the analysis of them is vague. How\ndo the multi-scale strengthening phases contribute to the strength?\n4. The analysis of mechanical and electrical properties in this work is not sufficiently in-depth. The remarkable performance\nis only qualitatively analyzed. Please quantify the properties, especially the mechanical properties.", "reviewer_2": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature\nCommunications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career\nResearchers who co-review manuscripts.", "reviewer_3": "Most of the concerns have been addressed in the revised manuscript. However, the followings still remain open.\n1) It is necessary to clarify in detail whether the unique microstructure and mechanical properties are obtained as a result of\na) the high sintering temperature, b) short sintering duration. It is not clear why the authors have done experiments at a\ntemperature of 600 °C, at which mass transfer mechanisms are hardly active.\n2) Figure R3 and Figure 1 from the manuscript are in contrast. The authors are required to clarify how the experiments were the\nexperiments were done. To promote open science practices, I highly recommend the authors to share a video of the\nexperimental procedure using Ultrafast High-temperature Sintering and quenching. From Figure R3, it appears the specimen was placed between\nzirconia fiber insulators. However, this is not demonstrated in the manuscript. Why were zirconia insulators used? Figure R3\n(a) is really difficult to understand and not clear at all. Figure R3(d), is there any sample between the insulators? Figure R3\n(e) and (f), are these profiles for the sample, carbon felt, or zirconia insulators?\n3) If zirconia insulator is used, the terminology of 'radiative joule-heating sintering' must be treated carefully.\nCorrespondingly, reconsideration is required. In addition, it is 'Joule'! Furthermore, in Spark Plasma Sintering, the amount of electric current that passes through the sample depends the conductivity of the material.", "reviewer_4": "The authors have satisfactorily addressed all of my previous concerns in their revision. I have no further questions, and I\nrecommend the manuscript for acceptance."}, "author_response": "Point-by-point response to the reviewer’s remarks\n(Black italic: Reviewer’s remarks; Blue: Our response;\nFig. RX appears only in the response.)\nFirst, we appreciate the reviewers' time and efforts spent reviewing our manuscript. We very much\nappreciate the reviewers' thoughtful comments and suggestions, which have significantly\nimproved the quality and clarity of our manuscript. We have performed additional experiments\nand revised the manuscript to address all comments and concerns. The changes are highlighted in\nyellow in the revised manuscript. Our point-by-point responses are presented below.\nReviewer #1:\n1. Please check the full text and unify the symbol “-”, for example, the “Cu-Fe” in Figure 2 (f)\nof the figure title.\nOur response: We sincerely thank the reviewer for the careful reading and for pointing out this\ninconsistency. We have reviewed the entire manuscript, including the main text, all figures, and\ntheir captions. The use of the hyphen symbol (“-”) has now been unified throughout the document\nto ensure consistency.\n2. The Vickers hardness values obtained in this work are questionable. The average hardness\nvalue is as high as 931HV, while the tensile yield strength is only 684MPa, which is hard to\naccept. It is recommended to use nanoindentation techniques to determine the hardness of\nthe interface, matrix, and reinforcement. SEM images of the indentations should be provided\nas supporting evidence.\nOur response: We sincerely thank the reviewer for this critical and constructive comment. We\nagree that elucidating the relationship between macroscopic Vickers hardness and bulk tensile\nyield strength in complex composite materials necessitates a more detailed, multi-scale\ninvestigation.\nFollowing the reviewer's suggestion, we have performed in-situ nanoindentation tests to precisely\nmeasure the mechanical properties of the individual microstructural constituents.\nWe utilized an in-situ nanoindenter (FemtoTools) integrated within a scanning electron\nmicroscope (SEM), as shown in Fig. S18a, which enabled the precise placement of indents on\nselected phases. Fig. S18b, d provide representative SEM images of an array of nano-indents\nsystematically placed across the Cu matrix, the FeCrMo nanocrystalline reinforcement, and the\ninterfacial regions. The quantitative hardness map, presented in Fig. S18c, reveals the distinct\nproperties of each phase.\nTo directly compare these values with the macroscopic Vickers hardness, we use the approximate\nconversion Hv=H /9.81×c, where H is the indentation hardness in GPa and c is a conversion\nIT IT\nfactor often close to 1. For simplicity in direct comparison, we approximate 1 GPa=100 HV.\nCombined with Fig. S19 the results are as follows:\n1\n FeCrMo Nanocrystals (Reinforcement): Exhibit an ultra-high hardness of approximately 12-\n Cu-FeCrMo Interface: Shows a high transitional hardness value of 5-8 GPa, corresponding to\n~500-800 HV, indicating strong interfacial bonding and effective load transfer.\n Cu Matrix: Possesses a hardness of 2-3 GPa, corresponding to ~200-300 HV. This value is\nsignificantly higher than that of pure annealed copper (~90 HV), which we attribute to grain\nrefinement into the sub-micron range and particularly small Fe₅₅Cr₂₅Mo₁₆B₂C₂ particles (as\nshown in Fig. 3a, b) are distributed diffusely during the UHSQ process.\nWe think that high macroscopic Vickers hardness (avg. 931 HV) is a composite \"rule of mixtures\"\nproperty, heavily influenced by the volume fraction (50 vol%) of the extremely hard FeCrMo\nreinforcement (~1200–1400 HV). In contrast, the bulk tensile yield strength (684 MPa) is governed\nby plastic deformation of the continuous, ductile Cu matrix.\nprocessed composite. (a) Photograph of the Femto-Tools nanoindenter setup used for the in-situ\nmeasurements inside a Scanning Electron Microscope (SEM). (b) SEM image showing a precise\narray of indentations placed across the Cu matrix, FeCrMo nanocrystals, and their interfaces. (c)\nCorresponding quantitative hardness map (in GPa) of the indented area. (d) Detailed close-up SEM\nmicrograph of representative indentations, visually confirming the vast hardness difference\nquantified in panel (c).\nFigure S19. Nanoindentation load-displacement curves of the constituent phases.\ncomposites. (a) Microhardness profiles measured from the sample surface inwards for the UHSQ\ncomposite, UHS composite, and monolithic pure Cu. The shaded bands represent the standard\ndeviation of the measurements. (b) Representative microhardness line scan across the Cu matrix,\nthe Cu/FeCrMo interface, and multi-phase nanocrystals. (c) The plot of microhardness versus\nelectrical conductivity (% IACS), benchmarking the UHSQ composite against various material\nclasses reported in the literature 43-60. (d) Wear test results showing the wear loss rate as a function\nof friction cycles under a load of 10 MPa and a sliding speed of 500 r/min. (e) Optical micrographs\nof the worn surfaces of UHSQ composites with 25 vol% and 50 vol% Fe-MG reinforcement,\nshowing uniform wear tracks. (f) High-magnification SEM images tracking the morphology of a\nsingle multi-phase nanocrystals on the wear surface after 15 and 60 minutes of sliding.\n4\n3. The \"Multi-scale Strengthening Mechanisms\" are emphasized in the manuscript, but the\nanalysis of them is vague. How do the multi-scale strengthening phases contribute to the\nstrength?\nOur response: We thank the reviewer for this insightful comment. We have revised the \"Multi-\nscale Strengthening Mechanisms\" section to explicitly delineate the contributions at different\nlength scales.\n Co-continuous interlocking network: Macroscopically, the Cu and the Fe-based reinforcement\nphases form two independent, interpenetrating 3D networks. The continuous, high-strength\nFe-based network acts as the primary load-bearing skeleton, while the continuous, ductile Cu\nmatrix imparts toughness and prevents catastrophic brittle failure.\n Strong interfacial bonding and mechanical interlocking: The UHSQ process creates a pristine\ninterface between the Cu matrix and the Fe-based reinforcement, characterized by nanoscale\nrivet-like interlocking structures (as shown in Fig. 3b). The strong physical bond is critical\nfor effective load transfer from the softer Cu matrix to the much harder reinforcement phase.\n Intrinsic strength of the multi-phase nanocrystalline reinforcement: The original\nFe₅₅Cr₂₅Mo₁₆B₂C₂ metallic glass devitrifies during the UHSQ process into a complex, multi-\nphase nanocrystalline structure. This structure consists of an Fe-Cr ferrite matrix, Mo-rich\nnanocrystals, and ordered B2 intermetallic phases. Such a hierarchical nanocrystalline\ncomposite possesses exceptionally high intrinsic hardness and strength. The phase acts as the\nprimary source of the composite's high hardness and wear resistance.\n Dispersion strengthening within the Cu matrix: As observed in Fig. 3h, smaller\nFe₅₅Cr₂₅Mo₁₆B₂C₂ particles are dispersed within the Cu matrix. These hard, undeformable\nparticles act as obstacles to dislocation motion within the ductile copper, effectively increasing\nthe yield strength of the matrix itself via a dispersion strengthening mechanism (akin to\nOrowan strengthening).\nWe contend overall strength is not attributable to a single factor but is a product of these integrated,\nmulti-scale mechanisms. The macro-scale network establishes a robust framework, the strong\ninterface enable this framework to behave as a single unit, and the high intrinsic strength of both\nthe reinforcement phase and the dispersion-strengthened matrix provides the high stress-bearing\ncapacity.\n55 25 16 2 2\nmetal matrix composites. (a) TEM image showing the clean, well-bonded interface between the\nCu matrix and the Fe-based reinforcement. (b) Higher-magnification view revealing the nanoscale\ninterlocking structure, with corresponding EDS maps confirming the distinct phases. (c) HAADF-\nSTEM analysis of the devitrified reinforcement, revealing a nano-architecture of an Fe-Cr matrix,\nMo-rich precipitates, and Fe-rich inclusions, confirmed by EDS mapping. The plot below\nquantifies the size distribution of these nanocrystals. (d) Room-temperature tensile engineering\nstress-strain curves for the composites processed by UHS and UHSQ, compared with a monolithic\nCu control sample. (e) Elevated-temperature (650 °C) tensile engineering stress-strain curves for\nthe UHS and UHSQ composites. (f) SEM image of the tensile fracture surface of UHS sample at\nroom temperature and (g) SEM image of the tensile fracture surface of UHSQ sample at room\ntemperature. (h) SEM image showing the distribution of small-sized particles in the Cu matrix.\n6\n4. The analysis of mechanical and electrical properties in this work is not sufficiently in-depth.\nThe remarkable performance is only qualitatively analyzed. Please quantify the properties,\nespecially the mechanical properties.\nOur response: We thank the reviewer for the valuable feedback regarding the quantitative analysis\nof the material's properties. We agree that a more explicit and in-depth discussion of the\nquantitative data will strengthen our manuscript. While the mechanical properties data is presented\nin Fig. 3, Fig. 4, Fig. S9, Fig. S14, Fig. S18 and Fig. S19.\n Quantification of Mechanical Properties:\nTensile Strength: The optimized UHSQ-processed composite exhibits an ultimate tensile strength\nof approximately 685 MPa at room temperature, representing a 6.5-fold enhancement over the\npure Cu reference..\nHigh-Temperature Strength: The material demonstrates remarkable thermal stability, maintaining\na yield strength of 290 MPa at an elevated temperature of 650°C.\nHardness: The average Vickers hardness of the UHSQ composite is 931 HV, with a maximum\nvalue reaching 1297 HV. This represents a nearly 10-fold increase compared to pure Cu.\nWear Resistance: Under a 10 MPa load at 500 rpm, the UHSQ composite shows a 40–50 times\nimprovement in wear resistance over the pure Cu sample.\n Discussion of Electrical Properties:\nRegarding the electrical properties, we measured conductivity of 15-20% IACS. We acknowledge\nthis is a moderate value for a copper-based composite and have expanded our discussion to provide\na clear physical explanation. The limited conductivity is an inherent consequence of the\ncomposite's unique microstructure, which is essential for its mechanical performance. The high\nintrinsic electrical resistivity of the Fe₅₅Cr₂₅Mo₁₆B₂C₂ multi-phase nanocrystalline reinforcement\nis the primary factor. Since the hard phase forms a co-continuous, interlocking network with the\ncopper, it inevitably disrupts the electron transport pathways within the conductive Cu matrix,\nthereby limiting the bulk conductivity.\nWe reiterate that the central innovation of this work is the development of the UHSQ methodology\nwhich enables the synthesis of bulk, structurally-sound positive mixing enthalpy composites—a\nlong-standing challenge due to liquid-phase immiscibility. The primary objective of this study was\nto establish the feasibility of the UHSQ route and to document the resulting improvements in\nmechanical and tribological performance. Future efforts targeting architectural optimization and\nleveraging advances in metallic-glass-powder synthesis to engineer finer conductive networks are\nexpected to yield substantial gains in electrical performance.\nThese quantitative performance metrics are also summarized in the abstract of manuscript.\n\" The mechanical tensile strength of the resulting composite is approximately 8 times greater than\nthat of pure Cu. The wear resistance improves by 40-50 times, and the Vickers hardness increases\nby a factor of 10.\"\n55 25 16 2 2\nprepared by different sintering methods. Bar chart showing ultimate tensile strength of composites\nwith varying Cu content (25-50 vol%) processed by UHS and UHSQ at room temperature and\n650°C. Inset shows representative engineering stress-strain curves highlighting great mechanical\nproperties improvement in UHSQ samples. Inset SEM fractography of room temperature tensile\nUHSQ processed samples.\nFigure S14. Comparison of flexural stress-strain curves for composites processed via UHSQ and\nUHS at room temperature. The data shows that the UHSQ composite exhibits a significantly higher\nflexural strength (approaching 900 MPa) compared to the control sample processed without\nquenching (UHS).\nprocessed composite. (a) Photograph of the Femto-Tools nanoindenter setup used for the in-situ\nmeasurements inside a Scanning Electron Microscope (SEM). (b) SEM image showing a precise\narray of indentations placed across the Cu matrix, FeCrMo nanocrystals, and their interfaces. (c)\nCorresponding quantitative hardness map (in GPa) of the indented area. (d) Detailed close-up SEM\nmicrograph of representative indentations, visually confirming the vast hardness difference\nquantified in panel (c).\nFigure S19. Nanoindentation load-displacement curves of the constituent phases.\n1. It is necessary to clarify in detail whether the unique microstructure and mechanical\nproperties are obtained as a result of a) the high sintering temperature, b) short sintering\nduration. It is not clear why the authors have done experiments at a temperature of 600 °C,\nat which mass transfer mechanisms are hardly active.\nOur response: We thank the reviewer for this insightful comment. We agree with the reviewer\nthat both high temperature and short duration are critical parameters. The ultra-short duration is\ncrucial for kinetically trapping the system in a non-equilibrium, homogeneously intermixed state,\nthereby suppressing the thermodynamically driven phase segregation.\nWe have conducted a new set of comparative experiments at a high temperature of 1500 °C,\nvarying the holding time. The results are presented in Fig. R1 below.\n Sintering at 1500 °C for 120 seconds: As shown in the SEM micrograph and corresponding\nEDS maps (Fig. R1a-b), a prolonged duration at high temperature leads to severe phase\nsegregation. The Cu and Fe-rich phases have coarsened into large, distinct domains, which is\nthe expected outcome for a positive mixing enthalpy system approaching thermodynamic\nequilibrium.\n Sintering at 1500 °C for 10 seconds: In stark contrast, when the sintering duration is limited\nto just 10 seconds, the composite retains its fine, interlocked microstructure (Fig. R1a-b). The\nFeCrMo nanocrystals remain uniformly distributed within the Cu structure, and large-scale\nelemental segregation is effectively prevented.\nThese results clearly demonstrate that the ultra-short sintering duration is the dominant factor in\npreserving the desired microstructure. While a high temperature provides the driving force for\nsintering, it also promotes segregation. Our UHSQ method leverages this by using a timeframe so\nbrief that the atoms do not have sufficient time to diffuse over long distances and cause phase\nseparation, thus achieving a well-bonded and homogeneous composite. To resolve the ambiguity\narising from the 600 °C data, we provide additional results obtained at 1500 °C that directly answer\nthe reviewer’s valid question.\nof Cu-Fe₅₅Cr₂₅Mo₁₆B₂C₂ composites. (a) SEM micrograph and (b) corresponding EDS elemental\nmaps of a sample sintered for 120 seconds. Severe phase segregation and coarsening of the Cu-\nrich and Fe-rich domains are evident. (c) SEM micrograph and (d) corresponding EDS elemental\nmaps of a sample sintered for only 10 seconds, showing a homogeneous and finely dispersed\ninterlocked microstructure.\n2. Figure R3 and Figure 1 from the manuscript are in contrast. The authors are required to\nclarify how the experiments were done. To promote open science practices, I highly\nrecommend the authors to share a video of the experimental procedure using Ultrafast High-\ntemperature Sintering and quenching. From Figure R3, it appears the specimen was placed\nbetween zirconia fiber insulators. However, this is not demonstrated in the manuscript. Why\nwere zirconia insulators used? Figure R3 (a) is really difficult to understand and not clear at\nall. Figure R3(d), is there any sample between the insulators? Figure R3 (e) and (f), are these\nprofiles for the sample, carbon felt, or zirconia insulators?\nOur response: We sincerely thank the reviewer for these detailed questions and the suggestion to\nenhance the clarity of our experimental procedure. We have provided new supplementary materials\nto clarify our custom-built UHSQ system.\n Clarification of the Experimental Process and Video Evidence\nThe apparent contrast between Fig. 1 (in the manuscript) and the last Fig. R3 arises from the\ndifferent imaging conditions used. In Fig. 1e, we used a Neutral Density (ND) filter to reduce the\nintense brightness of the incandescent sample, which allowed us to clearly capture the physical\nstate of the sample and its surroundings during the sintering process. In contrast, the photographs\nin the last Fig. R3c are presented without a filter to convey the true intensity of the rapid, high-\n11\ntemperature heating.\nTo fully address the reviewer's request for clarity and to promote open science, we have prepared\ntwo new supplementary videos (Video S1 and Video S2). Video S1 shows the entire UHSQ\nprocess in real-time. Video S2 presents the sinter process filmed through an ND filter.\nFurthermore, we have created Fig. S23 to be much clearer, explicitly delineating the upper \"UHS\"\n(sintering) stage and the lower \"Quench\" stage of the apparatus.\n Purpose of the Zirconia (ZrO₂) Fiber Insulators\nThe reviewer's observation is correct; we use ZrO₂ fiber cloth as an insulating layer, which is a\ncritical component of our setup. UHSQ system relies on Joule heating, where a high current is\npassed through the resistive carbon felts to generate heat. As a Cu-based composite, the sample is\nhighly electrically conductive; direct contact with the carbon felt would establish a low-resistance\nshunt path, severely diminishing Joule heating efficiency and compromising temperature control.\nThe ZrO₂ fiber cloth, a commercially available and robust ceramic material, is used as a high-\ntemperature electrical insulator to separate the conductive sample from the carbon felt heaters. As\nshown in the Fig. S23b, the fiber cloth undergoes a pre-sintering step to remove binders, which\ncauses it to expand slightly, ensuring excellent and uniform insulation.\n Specific Clarifications for last Fig. R3 Sub-panels:\nLast Fig. R3d: We confirm that a sample is indeed present between the two insulating layers in\nthis thermal image. The image was captured by an infrared pyrometer, whose spatial resolution is\ninherently limited.\nLast Fig. R3e: The temperature profile represents the temperature measured at the surface of the\nsample itself. It shows the complete cycle of ultrafast heating, a short holding plateau for sintering,\nand subsequent ultrafast quenching.\nLast Fig. R3f: The profile demonstrates the rapid responsiveness of the carbon felt heaters. The\npulsed profile shows that the heating elements can be turned on and off almost instantaneously,\nhighlighting the precise temporal control of our system.\nTo further verify uniform high-temperature attainment across the sample, we have included a\nnew image sequence in Fig. S23c (captured with a strong ND filter). which shows that immediately\nafter the power is shut off, the sample itself continues to glow brightly and uniformly, confirming\nthat it has been effectively heated throughout. We want to show the UHSQ method first ensures\neffective and rapid sintering, and our introduction of the instantaneous quenching step further\nenables new possibilities in microstructural design and material property enhancement.\nRegard to fiber cloth, we emphasize the addition in the method of the manuscript:\n\"The UHSQ technique with ZrO fiber was subsequently applied, with sintering for 10 seconds\n2\nfollowed by a 5-second cooling period, to obtain a dense composite sheet.\"\nPhotograph of the overall UHSQ apparatus, showing the upper heating stage (UHS) and the lower\nquenching bath. (b) Schematic illustration of the heating principle, where a sample is insulated by\nZrO₂ fibers and heated by surrounding carbon felts. (c) Photographs of the UHS process in action\n(left) and the ZrO₂ fiber insulation material (right). (d) A representative thermal image captured\nby an infrared pyrometer during sintering, showing a uniform temperature distribution. (e) The\nmeasured temperature-time profile for a typical UHSQ experiment, detailing the ultrafast heating,\nsintering, and quenching stages. (f) A temperature profile demonstrating the system's capability\nfor rapid pulsed heating, validating its high thermal responsiveness.\n(a) The photograph of the overall experimental setup, illustrating the distinct upper UHS (Ultrafast\nHigh-temperature Sintering) stage and the lower Quenching stage containing a silicone oil bath.\n(b) A close-up view comparing the ZrO₂ fiber insulation layers before and after a pre-sintering\ntreatment, which is performed to remove binders and ensure effective electrical insulation. (c) The\nphotographic sequence demonstrating the principle of sample heating with ND filter. Immediately\nafter the power to the carbon heaters is turned off, the heaters dim while the sample itself remains\nincandescent.\nmatrix composites. (a) Schematic illustration of the \"interface riveting\" mechanism. An initially\nweak contact between immiscible Cu and Fe Cr Mo B C metallic glass particles is transformed\n55 25 16 2 2\ninto a strong, mechanically interlocked physical bond at the interface. (b) A typical temperature-\ntime profile of the UHSQ process, characterized by ultrafast heating (500~1000 K/s), a short\nsintering plateau at high temperature, and subsequent rapid quenching. (c) Schematic of the\nexperimental procedure, showing the sample being rapidly heated between two carbon electrodes\nand then quickly submerged into a silicone oil bath for quenching. (d) TEM images reveal rivet-\nlike connections formed by physical deformation between Cu and FeCrMo nanocrystal. (e) A\nphotographic sequence capturing the key moments of the actual UHSQ process. (f) The Cu-Fe\ncomposite pellet with positive mixing-enthalpy prepared by UHSQ sintering.\n15\n3. If zirconia insulator is used, the terminology of 'radiative joule-heating sintering' must be\ntreated carefully. Correspondingly, reconsideration is required. In addition, it is 'Joule'!\nFurthermore, in Spark Plasma Sintering, the amount of electric current that passes through\nthe sample depends the conductivity of the material.\nOur response: We thank the reviewer for this precise and important question.\nOur use of the term \"radiative Joule heating\" is a deliberate and accurate description of the physics\nat play. The heating process can be broken down as follows:\n Joule heating constitutes the sole heat source: a high direct current is passed through two\ncarbon felts that sandwich the sample, and their electrical resistance converts the current into\nthermal energy. To ensure that the felts—rather than the highly conductive Cu-based\ncomposite—act as the primary resistive heaters, a thin ZrO layer is interposed as an\n2\nelectrical insulator; this prevents shunting of current through the sample and thereby\npreserves the intended Joule-heating pathway in the carbon felts.\n The Source of Heat is Joule Heating: The core of our technique is indeed Joule heating. A\nhigh electrical current is passed directly through the two highly resistive carbon felts that\nsandwich the sample assembly. It is these carbon felts that act as the primary heating elements,\nconverting electrical energy into thermal energy via the Joule effect.\n The Mechanism of Heat Transfer to the Sample: The sample itself is not intended to be part\nof the electrical circuit. At operational temperatures of up to 1500 °C, the dominant mode of\nheat transfer from the incandescent carbon felts to the sample is thermal radiation. The hot\nfelts act as efficient radiators, rapidly and uniformly heating the sample's surface. Heat is\nalso transferred via conduction through the thin ZrO₂ fiber layers. Therefore, the term\n\"radiative\" correctly describes the primary mechanism by which the sample is heated.\n The Critical Role of the ZrO₂ Insulator: As the reviewer correctly infers, the ZrO₂ insulator\nis essential. Its purpose is to function as an electrical insulator, not as a thermal barrier.\nGiven that our Cu-based composite is highly conductive, direct contact with the carbon felts\nwould create a short circuit. This would divert the current through the low-resistance\nsample, preventing the carbon felts from heating up effectively. The insulator is therefore a\nnecessary measure to enable the Joule heating of the carbon felts, ensuring that they, and not\nthe sample, serve as the primary resistive heaters.\n Distinction from Conventional SPS: In many SPS configurations, the current is designed to\npass through the conductive sample and die, making them active components of the heating\ncircuit. In our method, we deliberately prevent this by using the insulator. This configuration\nenables rapid, controlled heating of highly conductive materials whose low resistance\nprecludes efficient Joule heating in conventional SPS.\n(Black italic: Reviewer’s remarks; Blue: Our response;\nFig. RX appears only in the response.)\nFirst, we appreciate the reviewers' time and efforts spent reviewing our manuscript. We deeply\nappreciate the reviewers' thoughtful comments and suggestions, which have significantly\nimproved the quality and clarity of our manuscript. We have performed additional experiments\nand revised the manuscript to address all comments and concerns. The changes are highlighted in\nyellow in the revised manuscript. Our point-by-point responses are presented below.\nReviewer #3:\n1. Units of measurement are not kept consistent throughout the manuscript!\nOur response: We thank the reviewer for identifying the inconsistencies in our units of\nmeasurement. We have now thoroughly reviewed the entire manuscript and standardized all units\nfor consistency. For example, all temperatures related to the sintering process are now uniformly\npresented in Kelvin (K) throughout the main text, figures, and all supplementary materials.\n\"This approach retains the high heating rate (500-1000 K /s) of the UHS process while providing\nan enhanced cooling rate (up to 1000 K /s at the sample surface; see Figure 1b; Figure S2 and\nVideo 2).\"\nIn addition, we have standardized key terminology for clarity. For instance, various descriptions\nof the material are now uniformly referred to as \"Cu-Fe-based composite\". We also consolidated\nthe Supplementary Information to eliminate redundancy; for example, the original Figures S11\nand S12 have been merged into the revised Figure S11. We have highlighted the changes in the\nmain text and supplementary files.\n2. The discussion on the importance of the short sintering duration is not comprehensive yet.\nFurther experiments at different time intervals and discussion must be performed and\nintegrated into the manuscript?\nOur response: We thank the reviewer for raising this important point. We have performed\nexperiments with varying sintering durations to investigate this effect, and the results are presented\nin Figure S8. These experiments confirm that a short sintering duration is critical to kinetically\nsuppress the severe phase segregation that occurs with prolonged heating.\nAs suggested, we have now integrated this discussion into the main manuscript:\n\"This interfacial failure is directly linked to the sintering duration; prolonging the high-temperature\nexposure allows for the severe phase segregation demonstrated in Figure S8d. Minimizing the\nsintering time is therefore critical to kinetically trap the system and suppress elemental segregation\n(Figure S8a-c).\"\nSEM micrographs and corresponding EDS elemental maps for samples sintered at 1773 K for (a)\n5 s, (b) 10 s, (c) 15 s, and (d) 20 s. While shorter durations (5-10 s) enable the samples to maintain\na relatively fine, intermixed structure, prolonged heating (15-20 s) leads to severe phase\nsegregation and coarsening of the Cu and Fe-rich phases.\n3. Explanation on the use of zirconia insulator must be also included in the manuscript?\nOur response: We thank the reviewer for this constructive suggestion. The role of the zirconia\nfiber is key to the reproducibility of our experiment.\nAs per the reviewer's recommendation, we have clarified its purpose in the manuscript. We have\nnow added this explanation in the main text and in the \"Methods\" section to enhance the clarity\nand completeness of our experimental description.\n\"Specifically, the precursor compact was wrapped in zirconia (ZrO₂) fiber cloth for electrical\ninsulation (Figure S1c), positioned between two carbon felts in an argon atmosphere, rapidly\nsintered for 10 seconds, and than immediately quenched in high thermal conductivity silicone oil\nat 243 K (Figure 1c).\"\n\"Subsequently, the UHSQ technique was applied. The composite billet was wrapped in ZrO₂ fiber\ncloth, which serves as a high-temperature electrical insulator to prevent short-circuiting between\nthe conductive sample and the carbon heaters.\"\n2\n4. There is still concern about Figure 1b and c. Which figure is reporting the real-time\ntemperature of the sample? Figure 1b or Figure R3e? Figure 1c must be consistent with the\ndeveloped procedure?\nOur response: We thank the reviewer for their feedback on the figures.\nIn response to the query, Figure 1b is a schematic profile designed to illustrate the ideal stages of\nthe UHSQ process. The representative, experimentally measured temperature data is provided in\nthe Supplementary Information as Figure S2 (Due to the setting limitations of the thermal infrared\nimager, the recorded temperature range for this measurement was 485 K to 2000 K); both are\nmutually unified. We also confirm that the process schematic in Figure 1c is consistent with our\ndescribed experimental procedure.\nWe have revised the manuscript to clarify it. The updated text now reads as follows:\n\"This approach retains the high heating rate (500-1000 K /s) of the UHS process while providing\nan enhanced cooling rate (up to 1000 K /s at the sample surface; see Figure 1b; Figure S2 and\nVideo 2).\"\nFigure S2. The measured temperature-time profile for a typical UHSQ experiment, detailing the\nultrafast heating, sintering, and quenching stages (Note: Due to the setting limitations of the\nthermal infrared imager, the recorded temperature range for this measurement was 485 K to 2000\nK).\nFe-based composite. (a) Schematic illustration of the \"interface riveting\" mechanism. An initially\nweak contact between immiscible Cu and Fe Cr Mo B C metallic glass particles is transformed\n55 25 16 2 2\ninto a strong, mechanically interlocked physical bond at the interface. (b) The schematic\ntemperature-time profile of the UHSQ process, characterized by ultrafast heating, a short sintering\nplateau at high temperature, and subsequent rapid quenching. (c) Schematic of the experimental\nprocedure, showing the sample, housed in a ZrO₂ die, being rapidly heated between two carbon\nelectrodes and then quickly submerged into a silicone oil bath for quenching. (d) Transmission\nelectron microscope (TEM) images reveal rivet-like connections formed by physical deformation\nbetween Cu and FeCrMo nanocrystal. (e) A photographic sequence capturing the key moments of\nthe actual UHSQ process. (f) The Cu-Fe-based composite pellet with positive mixing-enthalpy\nprepared by UHSQ sintering.\n4\n5. It is surprising to see that Figures numbering have not been updated in the revised\nmanuscript!!\nOur response: We sincerely apologize for this serious oversight regarding the figure numbering.\nThis was an unintentional error that occurred during our revision process, and we are grateful to\nthe reviewer for pointing it out.\nWe have now performed a thorough and careful review of the entire manuscript. All figure\nnumbers, along with their corresponding citations within the text, have been corrected to ensure\nthat they are sequential, accurate, and consistent throughout the document.\n6. What does the terminology of ‘radiative joule-heating’ mean?! This terminology must be\nrevised. Joule not joule (typo)!\nOur response: We thank the reviewer for this valuable feedback. We agree that the term \"radiative\njoule-heating\" was non-standard and have revised it to the standard term \"Joule heating\"\nthroughout the manuscript for improved clarity. For instance, the text has been corrected as\nfollows:\n\"The sequential stages of UHSQ processing – including precursor assembly, Joule heating\nsintering andquenching – are visually documented in Figure 1e and Figure S3.\"\n7. It is surprising to see the fire in the ‘Video - Real-time Video of the UHSQ Process’ during\nthe presintering step! Where does the oxygen come from?\nOur response: We thank the reviewer for their keen observation. The flame in the video confirms\nit was recorded in air. This video is a demonstration, filmed in an open environment to clearly\nvisualize the sintering process, which is difficult to be captured within the glovebox. All\nexperimental data in this manuscript were collected under a protective argon atmosphere, as shown\nwith the apparatus in Figure S1. To ensure this distinction is clear, we have added explanatory\nnotes in the video title.\nthe overall UHSQ apparatus, showing the heating stage (upper) and the quenching bath (lower).\n(b) Schematic illustration of the heating principle, where a sample is insulated by ZrO₂ fibers\nand heated by surrounding carbon felts. (c) Photographs of the UHS process in action (left) and\nthe ZrO₂ fiber insulation material (right).\n8. What does Ultrafast high temperature mean in ‘Ultrafast high temperature rapidly densifies\nthe sample’ in the ‘Video - Real-time Video of the UHSQ Process’?!\nOur response: We thank the reviewer for highlighting the imprecise language in the video\nannotation. We have since revised the annotation in the video to \"Joule heat rapid sintering\". This\nterm accurately describes the heating mechanism responsible for the rapid densification observed.\nFigure R1. Screenshot of updated annotations in the UHSQ demonstration video.\n6\n9. What does Joule heat radiation mean in the ‘Video - UHS Heating Process Visualized with\nan ND Filter’?!\nOur response: We thank the reviewer for their feedback on the video's terminology. Our intent\nwas to describe the heat transfer process from the heating element to the sample. We have revised\nthe video annotation to \"Joule heat from carbon felt to sample\" to accurately describe the physical\nprocess.\nFigure R2. Updated annotation in the demonstration video of UHS heating process.\n(Black italic: Reviewer’s remarks; Blue: Our response;\nFig. RX appears only in the response.)\nFirst, we appreciate the reviewers' time and efforts spent reviewing our manuscript. We deeply\nappreciate the reviewers' thoughtful comments and suggestions, which have significantly\nimproved the quality and clarity of our manuscript. We have performed additional experiments\nand revised the manuscript to address all comments and concerns. The changes are highlighted in\nyellow in the revised manuscript. Our point-by-point responses are presented below.\nReviewer #3:\n1. Manuscript has typos (than immediately quenched, andquenching, for example).\nOur response: We apologize for the carelessness in the previous draft. We have corrected the\nspecific typos highlighted by the reviewer, ensuring the final presentation is polished and error-\nfree.\n2. The terminology of Joule heating sintering is not appropriate yet. Please use 'ultra-high-\ntemperature rapid sintering', in the manuscript, Video 1 and 2. Please remove 'Joule heat\nrapid sintering', and 'Joule heat from carbon felt to sample' form both videos.\nOur response: We sincerely thank the reviewer for this valuable suggestion regarding the\nterminology. We have fully adopted the term \"ultra-high-temperature rapid sintering\" throughout\nthe manuscript. Correspondingly, we have carefully updated the subtitles in Video 1 and Video 2\nby removing the phrases \"Joule heat rapid sintering\" and \"Joule heat from carbon felt to sample\"\nto ensure perfect consistency with the text.\n3. Please add a small note within Video 1 in a corner to notify the video was recorded in ambient\natmosphere for demonstration.\nOur response: We have followed this suggestion. A note has been added to the corner of Video\n1 to confirm that the recording took place in an ambient atmosphere for demonstration.\n1"}, {"round": 3, "reviewer_comments": {"reviewer_1": "Comments to NCOMMS-25-21591B\nTitle: Breaking Immiscibility Barriers: Ultrafast Sintering of Interlocked Cu-Fe-Based Composites\nThe authors have addressed all of my and my co-reviewer previous concerns in the revised version. We have no further\nquestions, and we recommend the manuscript for acceptance.", "reviewer_2": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature\nCommunications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career\nResearchers who co-review manuscripts.", "reviewer_3": "Changes in the manuscript were not highlighted in yellow in most instances! Manuscript has not improved substantially.\nFurther corrections and improvements are still required. Following are the comments.\n1. Units of measurement are not kept consistent throughout the manuscript!\n2. The discussion on the importance of the short sintering duration is not comprehensive yet. Further experiments at different\ntime intervals and discussion must be performed and integrated into the manuscript.\n3. Explanation on the use of zirconia insulator must be also included in the manuscript.\n4. There is still concern about Figure 1b and c. Which figure is reporting the real-time temperature of the sample? Figure 1b\nor Figure R3e? Figure 1c must be consistent with the developed procedure.\n5. It is surprising to see that Figures numbering have not been updated in the revised manuscript!!\n6. What does the terminology of ‘radiative joule-heating’ mean?! This terminology must be revised. Joule not joule (typo)!\n7. It is surprising to see the fire in the ‘Video - Real-time Video of the UHSQ Process’ during the presintering step! Where\ndoes the oxygen come from?\n8. What does Ultrafast high temperature mean in ‘Ultrafast high temperature rapidly densifies the sample’ in the ‘Video -\nReal-time Video of the UHSQ Process’?!\n9. What does Joule heat radiation mean in the ‘Video - UHS Heating Process Visualized with an ND Filter’?!"}, "author_response": "Point-by-point response to the reviewer’s remarks\n(Black italic: Reviewer’s remarks; Blue: Our response;\nFig. RX appears only in the response.)\nFirst, we appreciate the reviewers' time and efforts spent reviewing our manuscript. We deeply\nappreciate the reviewers' thoughtful comments and suggestions, which have significantly\nimproved the quality and clarity of our manuscript. We have performed additional experiments\nand revised the manuscript to address all comments and concerns. The changes are highlighted in\nyellow in the revised manuscript. Our point-by-point responses are presented below.\nReviewer #3:\n1. Manuscript has typos (than immediately quenched, andquenching, for example).\nOur response: We apologize for the carelessness in the previous draft. We have corrected the\nspecific typos highlighted by the reviewer, ensuring the final presentation is polished and error-\nfree.\n2. The terminology of Joule heating sintering is not appropriate yet. Please use 'ultra-high-\ntemperature rapid sintering', in the manuscript, Video 1 and 2. Please remove 'Joule heat\nrapid sintering', and 'Joule heat from carbon felt to sample' form both videos.\nOur response: We sincerely thank the reviewer for this valuable suggestion regarding the\nterminology. We have fully adopted the term \"ultra-high-temperature rapid sintering\" throughout\nthe manuscript. Correspondingly, we have carefully updated the subtitles in Video 1 and Video 2\nby removing the phrases \"Joule heat rapid sintering\" and \"Joule heat from carbon felt to sample\"\nto ensure perfect consistency with the text.\n3. Please add a small note within Video 1 in a corner to notify the video was recorded in ambient\natmosphere for demonstration.\nOur response: We have followed this suggestion. A note has been added to the corner of Video\n1 to confirm that the recording took place in an ambient atmosphere for demonstration.\n1"}, {"round": 4, "reviewer_comments": {"reviewer_3": "Most of the concerns have been addressed. However, minor improvements are required.\n1. Manuscript has typos (than immediately quenched, andquenching, for example).\n2. The terminology of Joule heating sintering is not appropriate yet. Please use 'ultra-high-temperature rapid sintering', in the\nmanuscript, Video 1 and 2. Please remove 'Joule heat rapid sintering', and 'Joule heat from carbon felt to sample' form both\nvideos.\n3. Please add a small note within Video 1 in a corner to notify the video was recorded in ambient atmosphere for\ndemonstration."}, "author_response": "Point-by-point response to the reviewer’s remarks\n(Black italic: Reviewer’s remarks; Blue: Our response;\nFig. RX appears only in the response.)\nFirst, we appreciate the reviewers' time and efforts spent reviewing our manuscript. We deeply\nappreciate the reviewers' thoughtful comments and suggestions, which have significantly\nimproved the quality and clarity of our manuscript. We have performed additional experiments\nand revised the manuscript to address all comments and concerns. The changes are highlighted in\nyellow in the revised manuscript. Our point-by-point responses are presented below.\nReviewer #3:\n1. Manuscript has typos (than immediately quenched, andquenching, for example).\nOur response: We apologize for the carelessness in the previous draft. We have corrected the\nspecific typos highlighted by the reviewer, ensuring the final presentation is polished and error-\nfree.\n2. The terminology of Joule heating sintering is not appropriate yet. Please use 'ultra-high-\ntemperature rapid sintering', in the manuscript, Video 1 and 2. Please remove 'Joule heat\nrapid sintering', and 'Joule heat from carbon felt to sample' form both videos.\nOur response: We sincerely thank the reviewer for this valuable suggestion regarding the\nterminology. We have fully adopted the term \"ultra-high-temperature rapid sintering\" throughout\nthe manuscript. Correspondingly, we have carefully updated the subtitles in Video 1 and Video 2\nby removing the phrases \"Joule heat rapid sintering\" and \"Joule heat from carbon felt to sample\"\nto ensure perfect consistency with the text.\n3. Please add a small note within Video 1 in a corner to notify the video was recorded in ambient\natmosphere for demonstration.\nOur response: We have followed this suggestion. A note has been added to the corner of Video\n1 to confirm that the recording took place in an ambient atmosphere for demonstration.\n1"}]

Decision letter
[Review submitted as PDF attachment - text not found in document]