Increasing fatigue resistance in ordered intermetallic alloys with multi-element symbiosis

Domain: biology
Published: 2026-03-18
Editorial outcome: minor_revision

Abstract
Intermetallic alloys, recognized for the long-range atomic ordering and resultant impressive mechanical properties, are highly sought after in various advanced fields, including aerospace, automotive, and nuclear energy. However, their widespread application is still hindered seriously due to the poor fatigue endurance. Here, we design a new-type L1 -structured multi-element symbiotic intermetallic alloy (MSIMA) and 2 achieve a fatigue limit of ~1,100 MPa that remarkably surpasses its yield strength by 1.1 times, which is superior to other structural alloys currently in use. The complex sublattice occupation strengthens the alloy by increasing the antiphase boundary energy of the superlattice, thereby suppressing the fatigue-induced lattice defects. Concurrently, the multi-element symbiosis enables the modulation of local chemistries and the architecting of the disordered interfacial nanolayer (DINL) near grain boundaries, thereby shifting the fatigue fracture mode from intergranular to transgranular cracking. Furthermore, serving as the ductilizing sources, these DINLs facilitate the unusual anti-fatigue mechanisms—mechanical faulting and twinning— that are rarely observed in ordered alloys at room temperature. This deformation behavior effectively alleviates the strain localization and blunts the crack propagation, thereby enhancing their fatigue resistance.

Introduction
Intermetallic alloys with long-range ordered superlattice structures bridge the gap between metals and ceramics, offering a unique combination of properties from both 1-6. In general, compared to disordered solid solution alloys, intermetallic alloys have stronger chemical bonding and lower atomic mobility, making them promising candidates for critical components in aerospace, power plants, and marine applications 1-3. Until now, however, most polycrystalline intermetallic alloys suffer from serious intergranular rupture under cyclic loading at room temperature due to the grain boundary (GB) embrittlement 4-7. Typically, under the uniaxial fatigue with a stress ratio (minimum stress/ maximum stress) of R=0.1, traditional chemically simple intermetallic alloys exhibit a fatigue limit (σ, i.e., the maximum f stress at a cycling lifetime ≥107) 8 below 90% of their yield strengths 6,7. Exacerbated by their insufficient yield strengths, the reported fatigue limits of intermetallic alloys are below 400 MPa 6,7. Such poor fatigue endurance, responsible for the majority of service failures in metallic components 9,10, has long been an Achilles’ heel of intermetallic alloys, greatly compromising their otherwise exceptional properties 4-7.

During the past decades, extensive efforts have been devoted to mitigating the GB brittleness through compositional modifications, including both the microalloying (usually in the parts-per-million range) to control grain-boundary composition and macroalloying additions to control ordered crystal structures 11-13. One classic example is the boron doping (0.1 atomic percent, at.%) in the Ni3Al alloy, which strengthens the GB cohesion via segregation and boosts a fracture elongation as high as 44% 11. However, despite the decent improvement in ductility, such chemical modifications have shown limited effectiveness in improving their fatigue resistance 6,7,14-16. For instance, the (Ni3Al)B alloy exhibits a fatigue limit of only 213 MPa, far below those of many other engineering alloys 7. In general, fatigue-induced stress concentration near GBs is much higher than that induced by the uniaxial tension 17,18. In the disordered alloys, stress concentrations could be released by emitting dislocations 19,20, twinning 21, nucleating second phases 22, and even diffusion 21, thereby preventing the GB cracking. Unfortunately, however, in the ordered intermetallic systems, stress relaxation is largely inhibited due to the presence of antiphase boundary (APB) energy 3,4. As a result, upon the cyclic loading, progressive stress localization ultimately exceeds the bearing capacity of GBs and results in intergranular fatigue failure 7.

To suppress fatigue-induced intergranular cracking, we devise a design strategy that couples chemical and structural tuning in multicomponent intermetallic alloys. We deviate from the stoichiometric composition 23 and exploit the cooperative interactions among various elements 3,24 to design a multi-element symbiotic intermetallic alloy (MSIMA) with a nominal composition of (Co47Ni30Ti4Al10Ta4V5)99B1. Such a design drives the segregation of face-centered cubic (FCC)-stabilizing elements to GBs, yielding a core-shell structure comprising a L1 -type ordered superlattice grain (OSG) encapsulated by an FCC-type 2 disordered interfacial nanolayer (DINL) (Fig. 1a). Such an order/disorder composite structure markedly enhances the fatigue resistance of intermetallic alloys, resulting in gigapascal-level fatigue limits (at R=0.1) that have not been reported in previous studies.

Results
Microstructure of the MSIMA

By controllably incorporating multiple elements, we synthesized the MSIMA with a nominal composition of (Co47Ni30Ti4Al10Ta4V5)99B1 (at.%) through arc melting. By adjusting the thermomechanical processing (see Methods), ultrafine-grained (UFG) and coarse-grained (CG) MSIMA samples were prepared, with average grain sizes of 0.8 ± 0.36 and 10.6 ± 5.70 μm, respectively (Fig. 1 and Suppl. Figs. 1-2). Both the as-prepared URFG and CG samples feature equiaxed grains with random orientation (Figs. 1b-c and Suppl. Figs. 2a-b) and contain a tiny amount (≤1.5 vol.%) of boride (Suppl. Figs. 1-3). High-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) analysis revealed that both samples exhibit a similar core-shell structure (Fig. 1d and Suppl. Fig. 4a). Specifically, the OSG has an L1 - 2 type ordered structure (Suppl. Fig. 5), while the DINL displays FCC structure with a thickness of ~2 nm and is highly coherent with the host OSG (Fig. 1d and Suppl. Fig. 4a). The corresponding volume fraction of DINL is 0.11 % in CG-MSIMA and 1.5 % in UFG-MSIMA, respectively (Suppl. Note 1). The ordered and disordered structures are identified based on the presence or absence of additional diffraction spots in the corresponding fast Fourier transform (FFT) images of selected areas (Fig. 1d and Suppl. Fig. 4a). Atomic resolution energy-dispersive X-ray spectroscopy (EDX) mapping (Fig. 1e and Suppl. Fig. 4b ) further revealed that the L1 phase corresponds to a (Co, Ni) (Ti, Al, Ta, V)-type compositionally complex 2 3 ordered superlattice, i.e., Co and Ni atoms occupy the face centers (A-sublattice), whereas Ti, Al, Ta, and V atoms occupy the vertices (B-sublattice) of the L1 unit cell. Moreover, EDX results highlighted the 2 strong tendency of Co segregation and the depletion of Ni, Ti, Al, Ta, and V along the GBs (Suppl. Figs. 4c-d and 6). Since the EDX approach has limitations in accurately quantifying the distribution of light elements like B, we further employed the three-dimensional atom probe tomography (3D-APT) for a more precise compositional analysis. As shown in Fig. 1f, the grain interior is compositionally homogeneous without elemental clustering. In contrast, obvious co-enrichment of Co and B is observed near the GB, along with the depletion of Ni, Ti, Al, Ta, and V. The one-dimensional compositional profile across the GB (Fig. 1g) shows that the peak concentrations of Co and B elements inside the DINL reach 55.8 and 1.5 at.%, respectively—approximately 1.2 and 2.1 times those of the OSG. The pronounced segregation of Co at GBs promotes the FCC-type disordering by reducing the alloy’s ordering energy, thereby facilitating the formation of the DINL3.

Mechanical properties

Leveraging the OSG/DINL composite structure, we achieve an enhanced synergy of strength, ductility, and fatigue resistance in the MSIMAs at room temperature (Fig. 2). The CG-MSIMA has a yield strength (σ ) of 698 ± 9.2 MPa, an ultimate tensile strength (σ ) of 1,561 ± 1.4 MPa, and a uniform elongation (ε ) y u u of 35 ± 2.7% (Fig. 2a). The UFG sample demonstrates superior performance, with a σ y of 1,020 ± 11.7 MPa and an σ of 1,680 ± 8.6 MPa, while maintaining a comparable ε of 33 ± 2.3%. Such exceptional strength- u u ductility combinations outperform those of conventional chemically simple intermetallic alloys without the DINL. For instance, the (Co Ti ) B alloy with an average grain size of 11.8 ± 3.63 μm (Suppl. Fig. 7) 76 24 99 1 exhibits a σ of 374 ± 5.7 MPa and an ε of 7 ± 0.1 % (Fig. 2a). We attribute the superior yield strength of y u our MSIMAs primarily to their high APB energy of 328 ± 32 mJ/m² of the complex L1 -type superlattice 2 (calculated by Density functional theory, DFT, see Methods), which is larger than those of binary counterparts—282 mJ/m2 for Co Ti 25 and 190 mJ/m2 for Ni Al 26. The higher APB energy enhances the 3 3 ordering strengthening effect by increasing the critical stress for superdislocation generation 3,27. Additionally, the high concentration of solute atoms further strengthens the MSIMA by acting as an effective barrier to superdislocation motion 28. The extraordinary ductility stems from the synergistic effects of reinforced GB and high work-hardening capacity (Fig. 2b), which arises from the enhanced dislocation activities facilitated by the DINL3. Combining these together, superb strength-ductility synergy can be successfully achieved in our present MSIMAs.

Furthermore, their fatigue properties are systematically evaluated, as shown in Fig. 2c. Obviously, conventional intermetallic alloys, such as the (Ni Al) B 7, Nb Al Ti Mo 16, and (Co Ti ) B alloys, 3 1-x x 42.5 12 44 1.5 76 24 99 1 display intrinsically poor fatigue resistance. Their fatigue limits (σ, defined as the maximum stress at a f cycling lifetime ≥107) fall below 400 MPa, well beneath their yield strengths. In contrast, the CG- and UFG-MSIMAs demonstrate increased fatigue limits of 800 and 1,100 MPa, respectively, reaching 2.7 and 3.7 times that of the (Co Ti ) B alloy (300 MPa). More intriguingly, the fatigue limits of both the CG- 76 24 99 1 and UFG-MSIMAs exceed their yield strengths, exhibiting an anomalously high fatigue efficiency (defined as the fatigue limit to yield strength ratio σ/σ 29) of 1.1. Figure 2d presents the fatigue efficiency of the f y MSIMAs as a function of yield strength, alongside the data from many other advanced alloys, including intermetallic alloys 7,16,30, twinning/transformation-induced plasticity (TWIP/TRIP) alloys 31-43, superalloys 44-50, and body-centered cubic (BCC) alloys 51-60, as summarized in Suppl. Tables 1-4. In general, most high- strength alloys exhibit a fatigue efficiency below 1.0, making them go to fatigue failure below their yield strength 8. In contrast, for our present MSIMAs, a high fatigue efficiency of 1.1 can still be achieved even at an ultrahigh σ of 1 GAPa. This unique behavior clearly distinguishes them from traditional intermetallic y alloys that generally exhibit fatigue limits below 400 MPa with a fatigue efficiency below 0.9 6,14-16. Such superb fatigue resistance also surpasses most commercial alloys with a comparable yield strength.

Fatigue deformation mechanisms

To reveal the microstructural origins of these unusual fatigue performances, the fractography of MSIMAs and (Co Ti ) B alloys was carefully characterized, as shown in Fig. 3. In the MSIMAs, owing 76 24 99 1 to the enhanced GB cohesion by the DINL, the fracture mode transfers from the intergranular to transgranular cracking. The dispersed boride particles, rather than the GBs, act as the preferred sites for crack initiation (Fig. 3b, Suppl. Figs. 8 and 9b). The deformation incompatibility between the borides and matrix results in the localized stress concentration, leading to the cleavage of borides and subsequent crack initiation (Suppl. Fig. 8). Afterward, the crack propagates transgranularly, with the fracture surface sequentially displaying intersected facets, fatigue striations, and final rupture-induced dimples (Figs. 3c-e and Suppl. Fig. 9). More accurate characterization using a combination of focused ion beam (FIB) milling and transmitted Kikuchi diffraction (TKD) reveals that these facets originate from the crack propagation along the {111} planes (Fig. 3f). However, cracks do not propagate directly along a single (111) plane. Instead, they were deflected across multiple {111} planes, forming a zig-zag pattern (Fig. 3g). Such a crack propagation manner is a hallmark of highly planar-slip alloys, corresponding to the near-threshold crack growth with low propagation rates 61. The consecutive zig-zag deflections induce a rough fracture surface, which promotes the crack closure and significantly reduces the crack propagation rate 19,22,62-64. Subsequently, as the crack propagates and the corresponding stress intensity increases, fatigue striations start to appear (Figs. 3d and 3h-i). The FIB-transmission electron microscopy (TEM) analysis reveals the microstructure underneath the striations, showing dense developed deformation bands with spacing comparable to the striation intervals, further indicative of planar dislocation slips ahead of the crack tip. The fractography of the CG-MSIMA is similar to that of the UFG-MSIMA (Suppl. Fig. 9). However, grain refinement results in the reduced facet size and more frequent kinks along the crack paths, which thereby enhance the resistance to crack propagation (comparing Fig. 3c and Suppl. Fig. 9c). Likewise, the spacing of striations also decreases from 143 ± 43.9 nm in CG-MSIMA to 51 ± 16.6 nm in UFG-MSIMA (Suppl. Fig. 10), suggesting a slower fatigue crack propagation rate in the UFG-MSIMA.

In stark contrast to the transgranular cracking behavior observed in the MSIMAs, the (Co Ti ) B 76 24 99 1 alloy without the DINL displays a typical intergranular crack morphology across the entire fracture surface (Fig. 3j), spanning the crack initiation (Fig. 3k), propagation (Fig. 3l), and final rupture stages (Fig. 3m). The “rock candy” appearance, marked by distinct grain shapes, reveals the serious GB brittleness of it, as well as the poor fatigue resistance. Evidently, the formation of DINL near the GBs plays a crucial role in suppressing intergranular cracking and improving the fatigue resistance of ordered intermetallic alloys.

To further unveil the micro-mechanisms underpinning the exceptional fatigue performance of present MSIMAs, we conducted detailed TEM analyses on the UFG samples fatigued at 1,300 MPa for various cycles (Fig. 4). A sequence of cyclic deformation mechanisms—superdislocations, stacking faults (SFs), and nanotwins—was progressively activated in both the CG- and UFG-MSIMAs, governing their fatigue behavior (Fig. 4 and Suppl. Fig. 11). Dislocations, acting as the primary carriers of plastic strain, are first emitted from the DINLs and subsequently transform into superdislocations upon entering the ordered lattice (Fig. 4a and Suppl. Fig. 11a). As the cyclic loading progressed to 1,000 cycles, superdislocations accumulated and self-organized into parallel arrays within the OSG grains (Fig. 4b and Suppl. Fig. 11b). At the half-life stage (~60,000 cycles), multiple dislocation-slip systems were activated, causing the intersections of dislocation arrays that evolved into highly entangled nanoscale superdislocation networks (Fig. 4c and Suppl. Fig. 11c). Notably, persistent slip bands (PSBs) with ladder structures or dislocation cells, which typically induce strain localization during fatigue 65, are absent throughout the entire fatigue process in MSIMAs. This is because the highly planar-slip behavior of superdislocations effectively suppresses dislocation annihilation and recovery, thereby promoting the uniform distribution of strains. Such behavior is further supported by ex-situ EBSD analysis (Suppl. Figs. 11f–i), which shows that strains are initially localized at GBs at the onset of fatigue but become more uniformly distributed as the cyclic loading progresses.

Unexpectedly, in the later stages of fatigue deformation, numerous SFs and nanotwins, appearing as intersecting and narrow bands parallel to the {111} planes (Figs. 4d-f), were identified in both the UFG- and CG-MSIMAs (Fig. 4g-h and Suppl. Fig. 11e). Atomic-resolution HAADF-STEM images further reveal that these intersecting nanotwins have an ultrafine thickness of 0.8 to 2 nm. By combining the center-of-symmetry (COS) 66 and Burgers circuit analyses (Figs. 4i-j), it is revealed that the Shockley partials (1/6 <112>) and Frank partials (1/3 <111>) consistently appear as pairs along the SFs and twin boundaries. These results indicate that the SFs/nanotwins are induced by the slip of Shockley partial dislocations. These Shockley partials could originate either from the intersection of superlattice partial dislocations (1/2 <101> ), for example via the reaction ( 1/2 [101]+ 1/2 [011̅] = 1/3 [111]+ 1/6 [112̅]) or from the further SF/twin-dislocation interaction ( 1/2 [101] = 1/3 [111]+ 1/6 [12̅1]), as illustrated in Suppl. Fig. 12. The SFs and nanotwins can not only enhance the local strength by impeding superdislocations but also alleviate the localized stress concentration and sustain cyclic strain by deflecting dislocation trajectories and accommodating them slipping along the SFs/twin planes (Suppl. Fig. 12b).

It is noteworthy that the formation of SFs and nanotwins in MSIMAs is absent in other intermetallic alloys, such as the (Co Ti ) B alloy, at room temperature (Suppl. Fig. 13). The inherently high stacking 76 24 99 1 fault energy (SFE) and the long-range atomic ordering in intermetallic alloys impose strong constraints on the nucleation of SFs and deformation twins 67. In particular, the nucleation of deformation SFs and twins via 1/6<112> partial dislocations requires a local breakdown of the ordered lattice, leading to a pronounced increase in energy 4,67. For instance, the complex SFE (CSFE) of Co Ti alloy reaches 222 mJ/m2 68, far 3 exceeding SFEs of disordered FCC alloys, nearly twice that of pure Ni (125 mJ/m²) and Al (166 mJ/m²) 69-73. Such a high CSFE necessitates an extraordinarily high critical twinning stress, reaching 4.6 GPa for the (Co Ti ) B alloy (Suppl. Note 3), which far surpasses the threshold for GB cracking and renders twinning 76 24 99 1 unattainable 4.

Our present MSIMA also exhibits high resistance to SFs and twinning due to the high CSFE of 306 ± 33 mJ/m² (calculated by DFT, see Methods), even higher than that of Co Ti alloy 68. We attribute the 3 activation of faulting and twinning in these EMSIMAs during fatigue to the interfacial disordering. On the one hand, the DINL with a chemical composition of Co Ni Ti Al Ta V B (determined by APT) 55.8 23.8 3.1 8.8 2.5 4.5 1.5 exhibits a low SFE of approximately 29 ± 28.8 mJ/m2 (Suppl. Fig. 14). The FCC-type DINL locally disrupts the long-range atomic order and markedly reduces the SFE, thereby lowering the nucleation barrier for SFs and twins 67,74. On the other hand, the DINL enhances the cohesive strength of grain boundaries, acts as a dislocation source to trigger early-stage plastic deformation, and delocalizes strains as a buffer layer 3. It thereby enables the MSIMAs to withstand gigapascal-level cyclic stresses without intergranular cracking. The cyclic deformation at gigapascal-level stress drives the accumulation of high-density superdislocations in the OSG, which inevitably distorts the local atomic ordering and even triggers localized order-to-disorder transitions 75. In the UFG-MSIMA fatigued at 1,300 MPa, the integrated intensity ratio of the (100) to (200) diffraction spots decreased from 0.17 at the half-life to 0.03 at fracture, confirming the progressive disordering response during fatigue deformation (Suppl. Fig. 15). Such an order-to-disorder transition could also effectively reduce the SFE and disrupt the long-range ordering, further creating favorable conditions for SFs and twins. Building on the above, the gigapascal cyclic stress provides a sufficient shear driving force for the nucleation of twinning partials through the interactions of superdislocations, ultimately leading to the formation of stacking faults and deformation twins (Suppl. Fig. 12).

Generally, the poor fatigue performance of conventional intermetallic alloys stems from their intrinsic GB brittleness and limited capacity to release localized stress concentration, further compounded by their relatively low strength 4,6,7,16. These limitations are effectively tackled in our designed MSIMAs with the DINLs and enhanced yield strength. Firstly, the DINL modifies chemical bonding at GBs, shifting the bonding character at GBs from a mixture of ionic, covalent, and metallic to predominantly metallic, thereby effectively eliminating the GB brittleness 3. Secondly, interfacial disordering eliminates the strong APB barrier to dislocation nucleation, enabling the stress relaxation through multiple deformation modes, including intense superdislocations, stacking faults, and nanotwins 3. In particular, the generation of SF and nanotwins helps to release the localized stress concentration induced by dislocation pile-ups 76,77. Dense SF and nanotwin networks can not only impede the dislocation motion but also facilitate the dislocation slip along the SF/twin boundaries, thereby effectively accommodating the cyclic strain and alleviating the stress concentrations 20,78-81. Furthermore, once cracks nucleate, deformation twins ahead of the crack tip bridge the crack and generate high back-stress fields, hindering the crack propagation by reducing the crack-tip driving force 82,83. These mechanisms are reminiscent of those in TWIP steels with the FCC structure, where the deformation twinning enhances the resistance to crack initiation and propagation 84. Collectively, by serving as an interfacial “glue” that suppresses grain-boundary brittleness and as a ductile source that releases local stress concentrations, the DINL effectively eliminates intergranular cracking. Beyond the benefits of DINLs, the relatively high strength of MSIMAs hinders the generation of lattice defects and delays the early-stage damage, thereby extending their fatigue life 3,19. Benefiting from these synergistic effects, the fatigue tolerance of our MSIMA is improved.

By analyzing the S–N data using the Basquin relationship, we obtained the fatigue strength exponent, a key indicator of materials’ resistance to fatigue-damage accumulation 65,85 (Suppl. Table 5). The CG- MSIMA exhibits a fatigue strength exponent of -0.05, higher than those of other intermetallic alloys without DINLs (-0.08 to -0.06), confirming its enhanced resistance to fatigue damage. In addition, grain refinement further promotes a more uniform strain distribution and generates increasingly tortuous crack paths, thereby improving both crack-initiation and crack-propagation resistance, which elevates the fatigue strength exponent to -0.03 in the UFG-MSIMA. Consequently, the CG-MSIMA achieves a fatigue limit of 2.7 times that of the (Co Ti ) B alloy with a similar grain size. With optimized grain structure, these MSIMAs 76 24 99 1 achieve high fatigue limits at the gigapascal level—a benchmark for ordered intermetallic alloys.

More prominently, the fatigue resistance of our highly ordered MSIMAs even outperforms those of most disordered alloys (Fig. 2d and Suppl. Table 5). Such an unexpected yet remarkable discovery can be primarily attributed to the more uniform deformation in our MSIMAs. For example, in the disordered FCC alloys, dislocations can easily cross-slip and recover, leading to the formation of PSBs, dislocation walls, and dislocation cells. This will significantly localize the strain distribution and promote the accumulation of fatigue-induced vacancies, which in turn facilitates the crack nucleation 9,78. In contrast, in the ordered L1 -type superlattice structure, the complex configuration of superdislocations (i.e., closely spaced pairs of 2 unit dislocations bound together by an APB) restricts the cross-slip behavior and results in the highly planar slip of dislocations 86. As a result, the undesired strain localization in the form of PSBs and/or dislocation cells is suppressed. This behavior is also verified by significantly lower fatigue-induced surface roughness (40 nm for CG-MSIMA and 10 nm for UFG-MSIMA, Suppl. Figs. 16–17) compared to disordered alloys (ranging from tens of nanometers to several micrometers) 40,87,88.

To conclude, the long-standing bottleneck problems of intermetallic alloys—the intrinsic brittleness and poor fatigue resistance—were successfully addressed by introducing the DINL in a multi-element symbiotic L1 -type intermetallic alloy. This MSIMA has an ultrahigh fatigue limit and efficiency that 2 largely surpasses those of previously reported intermetallic alloys and even the advanced TWIP/TRIP alloys with the disordered solid solution structure. The DINL enhances the GB cohesion and effectively coordinates strain transfer between adjacent grains, thereby eliminating the brittle intergranular cracks— the biggest threat to intermetallic alloys. Furthermore, by leveraging the DINL, multiple cyclic deformation mechanisms, especially mechanical faulting and twinning, were evoked within the ordered grain interiors during the cyclic deformation. These unique mechanisms synergistically enhance their resistance to fatigue- crack initiation and effectively blunt propagating cracks. The interfacial disordering architecture provides a robust strategy for designing next-generation intermetallic alloys with superior strength, ductility, and fatigue resistance, offering significant promise for advanced structural applications, particularly in aerospace and nuclear energy systems.

Methods
Specimen preparation

The ingots, with a nominal composition of (Co Ni Ti Al Ta V ) B (at.%), were produced by arc- 47 30 4 10 4 5 99 1 melting a mixture of raw elements (~ 99.9 wt.% purity) in an argon atmosphere. Off-stoichiometry (i.e., the concentration of Co and Ni is above 75%) and B addition were carefully designed to promote the interfacial disordering 3,23. To ensure compositional homogeneity, the ingots were flipped and re-melted at least eight times before being cast into a water-cooled copper mold. The obtained ingots were then homogenized at 1,150 C for 24 hours under argon protection to remove dendrite structures. For the UFG-MSIMA sample, after homogenization, the specimens underwent cold rolling to 90%, followed by annealing at 950 C for 7 mins. For the CG-MSIMA sample, after homogenization, the material was repeatedly cold-rolled with a reduction ratio of 20% and annealed at 1,150 °C, until a total reduction ratio of 66% was achieved. For comparison, specimens with a nominal composition of the (Co Ti ) B were also prepared using a similar 76 24 99 1 process of casting, cold rolling, and annealing with the homogenization temperature set at 1,050 C. All samples, regardless of chemical composition, were air-cooled to room temperature after annealing.

Mechanical tests

Tensile specimens with a gauge section of 12.5 × 3 × 1 mm3 and high-cycle-fatigue specimens with a gauge section of 5 × 3 × 1 mm3 were fabricated by electro-discharge machining (Sodick, AQ360LX) from CG-MSIMA and conventional (Co Ti ) B samples. For UFG-MSIMA, the gauge section of tensile and 76 24 99 1 fatigue specimens was 12.5 × 3 × 0.7 mm3 and 5 × 3 × 0.7 mm3, respectively.

The quasi-static tensile tests were performed on a Material Testing System (MTS) tension machine at room temperature (~ 20 °C) with a constant strain rate of 1 × 10-3 s-1. Stress-controlled tension-tension fatigue tests, with R = 0.1, were conducted on an MTS Landmark Servohydraulic Test System at a frequency range of 30 Hz. Tests were stopped when the specimen failed or achieved a lifetime of 107 cycles. All tensile and high-cycle-fatigue samples were polished to obtain a mirror finish before testing to eliminate surface defects. At least three specimens were tested to ensure data reproducibility.

Microstructural characterizations

Microstructures of the specimens were characterized by a field-emission SEM (FEI Scios) equipped with EBSD and TKD (OXFORD acquisition system). EBSD and TKD maps were acquired at an accelerating voltage of 20 kV and a probe current of 13 nA, with scanning step sizes ranging from 0.05 to 1.0 μm depending on the magnification. Kikuchi patterns were collected using a 2×2 binning. A minimum of 8 bands was required for successful indexing. Data points with a confidence index below 0.1 were excluded from the analysis. The Aztec Crystal 6.2 software was used for EBSD data analysis.

TEM analyses were performed on an FEI Talos F200 equipped with a Super-X EDX detector, at an accelerating voltage of 200 kV. Atomic-resolution high-angle annular dark-field scanning transmission electron microscope (HAADF-STEM) images and EDX maps were obtained on an aberration-corrected FEI Titan G2 microscope at 300 kV fitted with super-X EPDX. TEM images and FFT were post-processed by Digital Micrograph 3.2.0 software. The atomic-resolution HAADF-STEM and EDX data were processed using the FEI Velox 3.9 software. To enhance the visualization of the atomic lattice, a Gaussian filter was applied to the atomic-resolution EDX maps. Samples for SEM and EBSD characterizations were prepared by electro-polishing in an HNO (~ 25 vol. %/ CH OH (~ 75 vol. %) solution under a direct voltage of 20 3 3 V at a temperature of -40 °C. For the TEM observations, specimens were first ground to a thickness of 50 μm by SiC grit papers and then punched into discs with a diameter of 3 mm. Subsequently, these discs were further dimpled to a thickness of 15 μm, followed by ion milling in a precision ion-polishing system (PIPS, Gatan 695).

The surface undulation of MSIMA samples after fatigue was characterized by a confocal laser scanning microscope (CLSM, Olympus OLS4000) and an atomic force microscope (AFM, Oxford Vero ES).

The three-dimensional (3D) atomic-scale elemental distribution was analyzed using a local electrode atom probe tomography (APT) system (CAMECA, LEAP 5000 XR). Needle-shaped samples for the 3D- APT analysis were annularly milled in a focused ion beam/scanning electron microscope (FIB/SEM, FEI Scios). To minimize Ga⁺ implantation, the final milling step was carried out at a low accelerating voltage (2 kV) and a beam current of 27 pA for 1 min. APT analysis was conducted at 70 K in voltage mode, with a pulse repetition rate of 200 kHz, a pulse fraction of 20%, and an evaporation detection rate of 0.3%. The acquired data were reconstructed and analyzed using Imago Visualization and Analysis Software (IVAS) version 3.8.2. To precisely quantify the compositional variations across the GB, a cylindrical region of interest (ROI) with a diameter of 10 nm was positioned perpendicular to the GB, and a one-dimensional (1D) compositional profile along the z-axis with a bin width of 0.5 nm was constructed.

First-principles calculations

The energy calculations are performed using the first-principles density functional theory (DFT) approach, implemented through the Vienna Ab initio Simulation Package (VASP, version 5.4.4) 89,90, incorporating spin polarization in all calculations. The exchange-correlation potentials were described using the generalized gradient approximation (GGA) functional in the Perdew-Burke-Ernzerhof (PBE) form 91. Projector augmented wave (PAW) potentials, as recommended by the VASP manual, were employed 92. An energy cutoff of 500 eV was chosen to ensure that energy fluctuations converged to less than 1 meV per atom. Partial occupancies of orbitals were handled using the Methfessel-Paxton scheme (order one) with a smearing width of 0.1 eV. The convergence criteria for electronic and ionic iterations were 10-6 eV and 0.05 eV/Å, respectively. During bulk structure relaxation, cell shape, volume, and atomic positions were fully optimized, whereas for planar fault optimization, the cell was fixed, and atomic positions were relaxed. The Brillouin zone was sampled with a 3×3×1 k-point mesh. The simulation cells for the L1 2 structure consisted of nine atomic planes along the [111] direction, eight atomic planes along the [100] direction, and a total of 144 atoms. The shear vectors for the antiphase boundary (APB) and complex stacking fault (CSF) in the L1 structure were defined as a/2[101] and a/6[121], respectively, where a denotes the lattice parameter 93. 2 The simulation cells for FCC structure in DINL consisted of nine atomic planes along the [111] direction and a total of 146 atoms (Supplementary Data 1-6). The shear vector for intrinsic stacking faults in FCC structure is defined as a/6[12̅1] (Suppl. Fig. 15). Stacking faults were modeled using the alias method, and their energies were calculated using the following formula 94:

𝛾 = (E_SFE - E_perfect) / A (1)

where E and E are the total energies of the structures with and without stacking faults, SFE perfect respectively, and A is the cross-sectional area of the stacking fault.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "In the manuscript titled “Record-high fatigue resistance in ordered intermetallic alloys with multi-element symbiosis” contributed by Qian Li, et.al., the authors designed a new type of intermetallic alloy achieving an exceptional fatigue resistance. Intermetallic alloys exhibited great potential in demanding applications due to their excellent properties arising from their strong chemical bonding and lower atomic mobility. However, despite the promising results, the widespread engineering application of intermetallic alloys remains limited by their inferior fatigue strength, a critical shortcoming for components subjected to cyclic loading in long-term or safety-critical applications. This bottleneck arises intricately from the easy occurrence of intergranular cracking owing to the intrinsic grain boundaries (GBs) brittleness. Therefore, one central challenge in this field is how to improve their fatigue properties.\nBy leveraging the multiple elements that modulate lattice chemistry, the disordered interfacial nanolayer (DINL) near GBs is formed, playing the pivotal roles in property enhancement. Compared to conventional intermetallic alloys with low fatigue tolerance, the DINL in newly designed alloy serve as mutiple functions, 1) “glue” GBs to enhance GB’s strengthen, 2) as a buffer layer to at the GB, delocalising the strain concentration and thus enhancing the defect accommodation, and 3) even as a dislocation sources at the early stage of deformation to initiate the plastic deformation. Each role can make significant contributions to the improved capability of sustained plastic flow, while collectively profoundly bolstering the alloy’s resistance to long-term cyclic loading. Integrating with systematic microstructural characterisation and theoretical modelling, the authors uncovered how such unique microstructures modulate the deformation mechanisms under cyclic loading and then how the fatigue properties can be tremendously enhanced. This work holds great potential to revolutionise the design of intermetallic alloys, and accelerate the broader industrial adoption of intermetallic alloys for critical structural components. This compelling contribution is of scientific and technological significance and deserves to be highly felicitated. I recommend the publication of this manuscript after minor revisions. Below are a few detailed comments and suggestions:\n1. The authors applied the integrated intensity ratio of the (100) to (200) diffraction spots to validate the progressive disordering response during fatigue deformation. Since there are many technical factors affecting this intensity ratio, are there any better methods to evaluate this phenomenon? I agree that progressive disordering would occur; however, a more robust approach is needed to verify this behaviour.\n2. Previous investigations indicated that there is a complex interplay between atomic ordering and stacking fault energy, both of which play crucial roles in dictating deformation mechanisms, particularly deformation twins (https://doi.org/10.1016/j.actamat.2021.117460). Given the significant work-hardening capacity and ultrahigh strength, this intricate interplay could also be leveraged for the interpretation of the occurrence of deformation twins.\n3. The Basquin relationship is a widely accepted model for characterizing fatigue performance. The authors are encouraged to extract the fatigue strength coefficient and exponent, offering quantitative insights into fatigue mechanisms and allowing for comparison with other alloys.\n4. Under the cyclic deformation where the mean stress is not zero, the maximum stress at a cycling life time is increasingly applied to represent the fatigue limit. However, the stress amplitude is the widely used parameter charactersing the endurance limit. The authors should make a clear indication of the fatigue limit in this investigation and make a proper comparison with other advanced alloy families.", "reviewer_2": "In this work, the authors successfully addresses the long-standing challenge of the low fatigue resistance in ordered intermetallic alloys. Impressively, by innovatively designing a multi-element symbiotic intermetallic alloy, the authors achieved a unltrahigh fatigue limit of 1,100 MPa, even exceeding its yield strength, which has never been reported before. By combining advanced microscopy, atom probe tomography, and first-principles calculations, the underlying origins have also been clearly revealed. Evidently, these achievements will be of great significance to advance the design and manufacturing of high-performance alloys for advanced structural applications. Therefore, I recommend the publication of this excellent work after several questions being clarified.\n1. The manuscript demonstrates the critical role of the disordered interfacial nanolayer (DINL) in enhancing fatigue resistance. What are the principles that guided its formation, e.g., the segregation of Co and the formation of a disordered nanolayer?\n2. Figure 2 shows that the fatigue efficiency of TWIP/TRIP alloys spans a wide range, from 0.5 to 1.2. Could the authors clarify the reason for such a large variation? Specifically, what factors lead to the low fatigue efficiency observed in some cases?\n3. What is the difference between deformation twins formed in a disordered face-centered cubic structure and those in an ordered L1₂-type superlattice structure? It will be better to clarify this point.\n4. The segregation of boron in the one-dimensional compositional profile in Fig. 2g is not clearly depicted. It would be better to enhance the clarity of this segregation.\n5. In Figure 2, pls double check and make sure that the symbols of TWIP/TRIP are used consistently.\n6. In Figure 4, pls change the “cycs” into the “cycles”.", "reviewer_3": "The manuscript reported a study on a multi-element intermetallic L12 alloy with high fatigue strength superior to other structural alloys currently in use. The reasons for such remarkable strength are attributed to complex sublattice occupations that raise the antiphase boundary (APB) energy and the formation of a disordered interfacial nanolayer (DINL) near grain boundaries, which collectively suppress fatigue-induced defects and alter the cracking mode from intergranular to transgranular. Additionally, the proposed mechanisms involving mechanical faulting and twinning-assisted plasticity in the ordered intermetallic matrix, which are facilitated by DINLs are claimed to be another reason for such fatigue resistance. These findings would be instructive for designing fatigue-resistant intermetallic alloys and advancing their practical applications. However, the following comments should be considered in the revised submission.\n1. The authors are requested to clarify the design principles for the alloy composition. What is the reason behind selecting the specific elemental combinations, and how does this differ fundamentally from the other alloys previously published by the authors? A clear explanation would help the readers understand the innovation of this work.\n2. The manuscript attributes the high work-hardening capacity to enhanced dislocation activities facilitated by the disordered interfacial nanolayer (DINL). However, it remains unclear what the actual volume fraction or spatial distribution of the DINL region is, and whether it is sufficient to account for such significant plasticity. Quantitative data or statistical analysis should be provided. Or did the author also observe the fatigued microstructure of chemically complex L12 phase without DINL layer by TEM.\n3. In Fig. 2, the stress-strain behavior of the MSIMA with DINL shows significantly higher values compared with the reference alloy. It is necessary to clarify whether other factors, such as grain size, phase fraction, or compositional variations, could also contribute to this enhancement. Please elaborate on whether this comparison alloy is representative.\n4. How is the stability of the DINL at elevated temperatures? Whether this alloy can be applied at elevated temperatures since the chemical composition appears to be tailored for elevated temperature applications with Ta addition.\n5. On Page 13, line 231, the statement “Superdislocations, emerging as the primary strain carriers, were first emitted from DINLs” appears inaccurate. The dislocations emitted from the DINLs are expected to be perfect dislocations in the disordered region, which then evolve into superdislocations when entering the ordered lattice. This distinction should be corrected for physical accuracy.\n6. On Page 14, lines 251–254, the sentence “These results indicate that the SFs/nanotwins are induced by the slip of Shockley partials, which originate either from the intersection of superdislocations or from the further SF/twin-dislocation interaction” is not clearly or correctly expressed. The proposed dislocation reactions are inconsistent with the described mechanisms. In particular, for the reaction 1/2[101]+1/2[01-1]=1/3[111]+1/6[11-2], the Burgers vectors do not balance on both sides. In addition, since superdislocations 1/2<101> were proposed as the sources for forming Shockley partials, APBs are expected to form during this process. Whether the authors also observed APBs along with SFs?"}, "author_response": "Response to the Reviewers\nWe sincerely thank the reviewers for their insightful and constructive comments. These\nsuggestions are appreciated and doubtlessly lead to a great improvement of the manuscript.\nRevisions were made accordingly here. Below are our point-to-point responses to the comments\nreceived. The amended manuscript is also enclosed; the changes that we made in the revised\nversion have been highlighted in yellow.\nResponse to Reviewer:\nReviewer #1\nIn the manuscript titled “Record-high fatigue resistance in ordered intermetallic alloys\nwith multi-element symbiosis” contributed by Qian Li, et.al., the authors designed a new type\nof intermetallic alloy achieving an exceptional fatigue resistance. Intermetallic alloys exhibited\ngreat potential in demanding applications due to their excellent properties arising from their\nstrong chemical bonding and lower atomic mobility. However, despite the promising results,\nthe widespread engineering application of intermetallic alloys remains limited by their inferior\nfatigue strength, a critical shortcoming for components subjected to cyclic loading in long-term\nor safety-critical applications. This bottleneck arises intricately from the easy occurrence of\nintergranular cracking owing to the intrinsic grain boundaries (GBs) brittleness. Therefore, one\ncentral challenge in this field is how to improve their fatigue properties.\nBy leveraging the multiple elements that modulate lattice chemistry, the disordered\ninterfacial nanolayer (DINL) near GBs is formed, playing the pivotal roles in property\nenhancement. Compared to conventional intermetallic alloys with low fatigue tolerance, the\nDINL in newly designed alloy serve as multiple functions, 1) “glue” GBs to enhance GB’s\nstrengthen, 2) as a buffer layer to at the GB, delocalizing the strain concentration and thus\nenhancing the defect accommodation, and 3) even as a dislocation sources at the early stage of\ndeformation to initiate the plastic deformation. Each role can make significant contributions to\nthe improved capability of sustained plastic flow, while collectively bolstering the alloy’s\n1\nresistance to long-term cyclic loading. Integrating with systematic microstructural\ncharacterisation and theoretical modelling, the authors uncovered how such unique\nmicrostructures modulate the deformation mechanisms under cyclic loading and then how the\nfatigue properties can be tremendously enhanced. This work holds great potential to\nrevolutionise the design of intermetallic alloys, and accelerate the broader industrial adoption\nof intermetallic alloys for critical structural components. This compelling contribution is of\nscientific and technological significance and deserves to be highly felicitated. I recommend the\npublication of this manuscript after minor revisions. Below are a few detailed comments and\nsuggestions:\nResponse: Thanks for your positive and constructive suggestions. Your comments will\ndoubtlessly help us improve the manuscript. We hope that our revisions can address your\nconcerns.\n1. The authors applied the integrated intensity ratio of the (100) to (200) diffraction spots\nto validate the progressive disordering response during fatigue deformation. Since there are\nmany technical factors affecting this intensity ratio, are there any better methods to evaluate\nthis phenomenon? I agree that progressive disordering would occur; however, a more robust\napproach is needed to verify this behavior.\nResponse: We sincerely appreciate your insightful comment. To the best of our knowledge,\nselected-area electron diffraction (SAED) has been regarded as a suitable and effective\ntechnique for evaluating the reduction in ordering degree. Compared with other methods such\nas X-ray diffraction (XRD), SAED is more sensitive to the subtle changes in structure. This\nmethod has been employed in several previous studies, which have demonstrated its reasonable\nreliability and sensitivity in capturing the disordering evolution in ordered superlattice\nstructures [1-2].\nReferences:\n[1] Yan, Y. Q. et al. Ductilization of 2.6-GPa alloys via short-range ordered interfaces and\n2\nsupranano precipitates. Science 387, 401-406 (2025).\n[2] Howard, H. C. et al. Chemically ordered dislocation defect phases as a new strengthening\npathway in Ni–Al alloys. Acta Materialia 289, 120887 (2025).\n2. Previous investigations indicated that there is a complex interplay between atomic\nordering and stacking fault energy, both of which play crucial roles in dictating deformation\nmechanisms, particularly deformation twins (https://doi.org/10.1016/j.actamat.2021.117460).\nGiven the significant work-hardening capacity and ultrahigh strength, this intricate interplay\ncould also be leveraged for the interpretation of the occurrence of deformation twins.\nResponse: We are grateful for your recommendation of this valuable reference. This work\nprovides important insights that further support our interpretation of the deformation\nmechanism of MSIMAs. We have included the suggested paper (Ref. 36) in the revised\nmanuscript and have provided the relevant discussions accordingly.\nModification\nOn page 15, we have added the following text and cited the recommended reference:\nThe inherently high stacking fault energy (SFE) and the long-range atomic ordering in\nintermetallic alloys impose strong constraints on the nucleation of SFs and deformation twins\n36.\nOn page 16, the discussion on the occurrence of deformation twins is also revised:\nOn the one hand, the DINL with a chemical composition of\nCo Ni Ti Al Ta V B (determined by APT) exhibits a low SFE of approximately 29\n55.8 23.8 3.1 8.8 2.5 4.5 1.5\n± 28.8 mJ/m2 (Suppl. Fig. 14). The FCC-type DINL locally disrupts the long-range atomic\nordering and markedly reduces the SFE, thereby lowering the nucleation barrier for SFs and\ntwins 36,43. On the other hand, the DINL enhances the cohesive strength of grain boundaries,\nacts as a dislocation source to trigger early-stage plastic deformation, and delocalizes strains as\na buffer layer 3. It thereby enables the MSIMAs to withstand gigapascal-level cyclic stresses\n3\nwithout intergranular cracking. The cyclic deformation at gigapascal-level stress drives the\naccumulation of high-density superdislocations in the OSG, which inevitably distorts the local\natomic ordering and even triggers localized order-to-disorder transitions 44. In the UFG-MSIMA\nfatigued at 1,300 MPa, the integrated intensity ratio of the (100) to (200) diffraction spots\ndecreased from 0.17 at the half-life to 0.03 at fracture, confirming the progressive disordering\nresponse during fatigue deformation (Suppl. Fig. 15). Such an order-to-disorder transition\ncould also effectively reduce the SFE and disrupt the long-range ordering, further creating\nfavorable conditions for SFs and twins. Building on the above, the gigapascal cyclic stress\nprovides a sufficient shear driving force for the nucleation of twinning partials through the\ninteractions of superdislocations, ultimately leading to the formation of stacking faults and\ndeformation twins (Suppl. Fig. 12).\n3. The Basquin relationship is a widely accepted model for characterizing fatigue\nperformance. The authors are encouraged to extract the fatigue strength coefficient and\nexponent, offering quantitative insights into fatigue mechanisms and allowing for comparison\nwith other alloys.\nResponse: We sincerely appreciate your constructive suggestion. Following your\nrecommendation, the stress–life (S–N) data were analyzed using the Basquin relationship：\n𝜎 = 𝜎′(2𝑁 )𝑏\n𝑎 𝑓 𝑓\nwhere 𝜎 isthe stress amplitude (0.45 of the maximum stress, 0.45𝜎 , for R=0.1); 𝜎′ is the\n𝑎 𝑚𝑎𝑥 𝑓\nfatigue strength coefficient; 𝑁 is the number of cycles to failure; 𝑏 is the fatigue strength\n𝑓\nexponent. Table R1 shows the 𝜎′ and 𝑏, which is included in the revised manuscript (Suppl.\n𝑓\nTable 5). The fatigue strength coefficient 𝜎′ is generally correlated with the ultimate tensile\n𝑓\nstrength, whereas the fatigue strength exponent b reflects the influence of microstructure and\nthe intrinsic fatigue-damage mechanisms. A larger b value typically signifies slower damage\naccumulation and improved fatigue resistance.\nThe present MSIMA exhibits higher fatigue strength exponents (-0.05 to -0.03) compared\n4\nto those of conventional intermetallic alloys without DINLs (-0.08 to -0.06), indicating\nsubstantially improved resistance to cyclic damage attributed to the enhanced grain-boundary\ncohesion and multiple deformation mechanisms. Moreover, the fatigue strength exponents of\nMSIMA are also higher than those of disordered TWIP/TRIP alloys (-0.10 to -0.17).\nThis\ndistinction that can be attributed to the superdislocation-dominated uniform deformation and\nthe stronger work-hardening capability provided by deformation twinning in the ordered\nsuperlattice structure.\nWe also noted that the UFG-MSIMA sample shows a higher fatigue strength exponent\nthan its CG counterpart. This enhancement is likely associated with grain refinement, which\npromotes more homogeneous deformation and suppresses strain localization. Additionally,\nfractography analyses reveal finer fatigue striations in the UFG sample, implying a reduced\nfatigue-crack propagation rate and further corroborating its superior fatigue performance.\nTable R1. The fatigue strength coefficient (𝝈′) and exponent (b) of the MSIMA samples\n𝒇\nand other materials.\nSamples 𝜎′ b Refs.\n𝑓\nUFG-MSIMA 858 -0.03 This work\nCG-MSIMA 779 -0.05 This work\n(Co Ti ) B 554 -0.08 This work\n76 24 99 1\n(Co Ni Ti Al Ta V ) B 678 -0.07 This work\n47 30 4 10 4 5 99.98 0.02\n(Ni Al)B 402 -0.07 [1]\n3 0.6\nNb Al Ti Mo 429 -0.06 [2]\n42.5 12 44 1.5\n316L stainless steel / -0.15 ~ -0.1 [3]\nCu-Al alloy / -0.17 ~ -0.14 [4]\nReferences:\n[1] Li, G., Guo, J. T., Wang, Z. G. & Shi, C. X. Fatigue behaviour of Ni3Al(B) alloys at room\ntemperature. Steel Research 65, 451-454 (1994).\n[2] Soboyejo, W. O. et al. An investigation of the fatigue and fracture behavior of a Nb-12Al-\n5\n44Ti-1.5Mo intermetallic alloy. Metallurgical and Materials Transactions A 30, 1025-1038\n(1999).\n[3] Li, Q., Yan, F. K. & Tao, N. R. Enhanced fatigue damage resistance of nanotwinned\naustenitic grains in a nanotwinned stainless steel. Scripta Materialia 136, 59-63 (2017).\n[4] An, X. H., Wu, S. D., Wang, Z. G. & Zhang, Z. F. Enhanced cyclic deformation responses\nof ultrafine-grained Cu and nanocrystalline Cu–Al alloys. Acta Materialia 74, 200-214 (2014).\nModification\nOn page 19, we have added the following text:\nBy analyzing the S–N data using the Basquin relationship, we obtained the fatigue strength\nexponent, a key indicator of materials’ resistance to fatigue-damage accumulation 34,54 (Suppl.\nTable 5). The CG-MSIMA exhibits a fatigue strength exponent of -0.05, higher than those of\nother intermetallic alloys without DINLs (-0.08 to -0.06), confirming its enhanced resistance to\nfatigue damage. In addition, grain refinement further promotes a more uniform strain\ndistribution and generates increasingly tortuous crack paths, thereby improving both crack-\ninitiation and crack-propagation resistance, which elevates the fatigue strength exponent to -\n0.03 in the UFG-MSIMA.\n4. Under the cyclic deformation where the mean stress is not zero, the maximum stress at a\ncycling life time is increasingly applied to represent the fatigue limit. However, the stress\namplitude is the widely used parameter charactersing the endurance limit. The authors should\nmake a clear indication of the fatigue limit in this investigation and make a proper comparison\nwith other advanced alloy families.\nResponse: Thanks for your insightful comment. We fully agree that the stress amplitude\nis conventionally used to represent the fatigue limit, particularly under the fully reversed\nloading condition. However, for the tension–tension fatigue test with a nonzero mean stress, a\ngrowing number of recent studies have adopted the maximum stress to characterize the fatigue\n6\nlimit, as it more accurately reflects the highest stress level that the material undertakes during\ncyclic loading [1-2]. Following your suggestion, the definition of the fatigue limit has been\nexplicitly clarified in the revised manuscript (Page 9). All the fatigue limits of other advanced\nalloys compared in this work were determined at a stress ratio of R = 0.1, or converted from R\n= -1 using the Goodman equation (Suppl. Note 2). For consistency, all fatigue limit values cited\nfrom the literature are also represented in terms of the maximum stress (see Supplementary\nTables 1-4).\nReferences:\n[1] Qu, Z. et al. High fatigue resistance in a titanium alloy via near-void-free 3D printing.\nNature 626, 999-1004 (2024).\n[2] Dan, C. Y. et al. Achieving ultrahigh fatigue resistance in AlSi10Mg alloy by additive\nmanufacturing. Nature Materials 22, 1182-1188 (2023).\nModification\nOn page 9, the fatigue limit was defined:\nTheir fatigue limits (σ, defined as the maximum stress at a cycling lifetime ≥107) fall\nf\nbelow 400 MPa, well beneath their yield strengths.\nThis study addresses the long-standing challenge of the low fatigue resistance in\nintermetallic alloys. By designing a multi-element symbiotic intermetallic alloy, the authors\nachieved a fatigue limit of 1100 MPa, exceeding its own yield strength. The paper combines\nadvanced microscopy, atom probe tomography, and first-principles calculations to reveal that\nthe disordered interfacial nanolayers near grain boundaries enhance cohesion, enable multiple\ncyclic deformation modes, and thus suppress intergranular cracking.\nThis obtained fatigue property is very impressive, which is a rare achievement for this\nclass of materials. The results are sound and important. The characterization and mechanistic\nanalysis are thorough and well executed. I recommend publication of this manuscript pending\nthe following revisions.\nResponse: We sincerely thank you for the positive and profound comments. They are\nvaluable in assisting us to enhance the clarity and depth of our study.\n1.The manuscript demonstrates the critical role of the disordered interfacial nanolayer (DINL)\nin enhancing fatigue resistance. What are the principles that guided its formation, e.g., the\nsegregation of Co and the formation of a disordered nanolayer?\nResponse: Thanks for your insightful comment.\nThe alloy design was based on a new compositional principle that couples multi-element\nsymbiosis with off-stoichiometric tuning, distinct from the previous design strategy of\nintermetallic alloys. Design of intermetallic alloys with long-range ordered superlattice\nstructures normally satisfies a critical stoichiometry. For instance, L1₂-structured alloys such as\nthe Ni₃Al and Co₃Ti require a 3:1 ratio between the A-sublattice (face-centered sites) and B-\nsublattice (corner sites) atoms to maintain their fully ordered structure. Our previous\nexperiments revealed that enriching the A-sublattice elements (e.g., Fe and Co), thereby\ndeviating from the ideal stoichiometry, induces the local segregation of A-type atoms and\ntriggers the formation of disordered FCC regions [1,2].\nsynergy, the (Co Ni Ti Al Ta V ) B alloy was developed with a (Co, Ni):(Ti, Al, Ta, V)\n47 30 4 10 4 5 99 1\nratio of 77:23. The concentration of each element was deliberately tailored to form the L1 grain\n2\ntogether with FCC DINL. In addition to the conventional B-site constituents of Al and Ti, Ta\nand V were introduced to elevate the antiphase boundary energy and thereby strengthen the\nstructural stability of the ordered phase [3,4].\nReferences:\n[1] Xiao, B. et al. Off-stoichiometry-guided design of high-strength chemically complex\nintermetallic-based alloys with outstanding ductility. Journal of Materials Science &\nTechnology 160, 28-33 (2023).\n[2] Yang, T. et al. Ultrahigh-strength and ductile superlattice alloys with nanoscale\ndisordered interfaces. Science 369, 427-432 (2020).\n[3] Gorbatov, O. I. et al. Effect of composition on antiphase boundary energy in Ni3Al-\nbased alloys: Ab initio calculations. Physical Review B 93, 224106 (2016).\n[4] Liang, Z.D. et al. High-Temperature CoNi-Based Superalloys Strengthened by γ′-\n(Ni,Co) (Cr,Al,Ti,X): The Effect of Refractory Elements. Metallurgical and Materials\nModification\nOn page 4, we have added the following text:\nTo suppress fatigue-induced intergranular cracking, we devised a design strategy that\ncouples chemical and structural tuning in multicomponent intermetallic alloys. We deviated\nfrom the stoichiometric composition 23 and exploited the cooperative interactions among\nvarious elements 3,24 to design a multi-element symbiotic intermetallic alloy (MSIMA) with a\nnominal composition of (Co Ni Ti Al Ta V ) B . Such a design drives the segregation of\n47 30 4 10 4 5 99 1\nface-centered cubic (FCC)-stabilizing elements to GBs, yielding a core-shell structure\ncomprising a L1 -type ordered superlattice grain (OSG) encapsulated by an FCC-type\n2\ndisordered interfacial nanolayer (DINL) (Fig. 1a).\n9\n2. The tensile curves of the studied alloys exhibit a yield plateau. Does it arise from the\nspecific role of the disordered interfacial nanolayer (DINL) in facilitating dislocation nucleation?\nA brief discussion on this point would improve the readers’ understanding of the deformation\nprocess.\nResponse: Thank you for the valuable comment. The present alloy indeed exhibits a yield\nplateau, which is a typical feature of yield point drop and Lüders deformation. Similar behaviors\nhave also been observed in other IMAs, ultrafine-grained metals or alloys, and some steels. In\ngeneral, the occurrence of Lüders deformation needs to satisfy two conditions: that is, (1) the\nlow initial mobile dislocation density and (2) the abrupt, rapid multiplication of mobile\ndislocations. These conditions are easily attained in the present fully recrystallized MSIMA.\n3. Figure 2 shows that the fatigue efficiency of TWIP/TRIP alloys spans a wide range, from\n0.5 to 1.2. Could the authors clarify the reason for such a large variation? Specifically, what\nfactors lead to the low fatigue efficiency observed in some cases?\nResponse: Thank you for the insightful comment. The TWIP/TRIP alloys considered in Fig.\n2 include FCC single-phase alloys (e.g., 316L stainless steel) and BCC/FCC dual-phase alloys\n(e.g., the FeMn C steel). FCC single-phase TWIP/TRIP alloys generally exhibit relatively\n17 0.8\nlow strength but high fatigue efficiency, owing to the progressive TWIP/TRIP effect during\ncyclic loading, which effectively suppresses the strain localization. In contrast, the BCC/FCC\ndual-phase alloys typically show higher strengths but reduced fatigue efficiency. In these dual-\nphase alloys, pre-existing phase boundaries often serve as preferential sites for fatigue crack\nnucleation due to the deformation incompatibility between the constituent phases. Therefore,\nthe observed large variation in fatigue efficiency can be attributed to differences in phase\nconstitution and microstructural features among the studied alloys. FCC single-phase alloys\nachieve high fatigue efficiency through homogeneous plastic deformation, while dual-phase\nalloys with significant phase boundary density are more prone to early fatigue crack initiation,\n10\nresulting in lower fatigue efficiency.\n4. Although the mechanical properties of the MSIMA are impressive, it would be preferable\nto remove emphatic terms such as “substantially” and “exceptionally” to maintain a more\nobjective and neutral tone.\nResponse: Thank you for the valuable suggestion. Accordingly, we have removed these\nemphatic terms such as “substantially” and “exceptionally” in the revised manuscript.\n5. What is the difference between deformation twins formed in a disordered face-centered\ncubic structure and those in an ordered L1₂-type structure?\nResponse: Thank you for the insightful question. The fundamental difference arises from\nthe atomic ordering in the superlattice. In disordered FCC structures, atoms occupy lattice sites\nrandomly, and partial dislocations can glide relatively freely along {111} planes, enabling the\nformation of deformation twins at relatively low stress.\nIn contrast, L1₂-ordered structures possess long-range atomic ordering, where the motion\nof partial dislocations not only creates stacking faults but also breaks the ordered lattice,\nresulting in a substantial increase in energy. Consequently, ordered intermetallic alloys exhibit\nhigh SFEs—often more than twice those of well-known high-SFE disordered metals such as Ni\nand Al. It has been well recognized that even Ni and Al rarely deform via twinning or stacking\nfaulting at room temperature. Therefore, reports of stacking faults or deformation twins in\nintermetallic alloys under similar conditions are scarce.\n6. The segregation of boron in the one-dimensional compositional profile in Fig. 2g is not\nclearly depicted. It is better to enhance the clarity of this segregation.\nResponse: Thank you for pointing this out. In the revised manuscript, the one-dimensional\ncompositional profile of boron has been added to Fig. 2g.\n11\n7. Change “cycs” to “cycles” in Figure 4a.\nResponse: Thank you for pointing this out. In the revised manuscript, all the “cycs” were\nchanged into “cycles”.\n8. Ensure that the symbols for TWIP/TRIP in Figure 2 are used consistently.\nResponse: Thank you for your comment. The symbols for the TWIP/TRIP alloys in Fig.\n2 have been corrected for consistency in the revised manuscript.\n1. The authors are requested to clarify the design principles for the alloy composition.\nWhat is the reason behind selecting the specific elemental combinations, and how does this\ndiffer fundamentally from the other alloys previously published by the authors? A clear\nexplanation would help the readers understand the innovation of this work.\nResponse: Thanks for your constructive suggestion.\nThe alloy design was based on a new compositional principle that couples multi-element\nsymbiosis with off-stoichiometric tuning, distinct from the previous design strategy of\nintermetallic alloys. Design of intermetallic alloys with long-range ordered superlattice\nstructures normally satisfies a critical stoichiometry. For instance, L1 -structured alloys such as\n2\nthe Ni Al and Co Ti require a 3:1 ratio between the A- and B-sublattice atoms to maintain their\n3 3\nfully ordered structure. Our previous experiments reveal that enriching the A-sublattice\nelements (e.g., Fe and Co), thereby deviating from the ideal stoichiometry, induces the local\nsegregation of A-type atoms and triggers the formation of disordered FCC regions [1,2].\nTherefore, building on the off-stoichiometric design concept and leveraging multi-element\nsynergy, the (Co Ni Ti Al Ta V ) B alloy was developed with a (Co, Ni):(Ti, Al, Ta, V)\n47 30 4 10 4 5 99 1\nratio of 77:23. The concentration of each element was deliberately tailored to form the L1 grain\n2\ntogether with FCC DINL. In addition to the conventional B-site constituents of Al and Ti, Ta\nand V were introduced to elevate the antiphase boundary energy and thereby strengthen the\nstructural stability of the ordered phase [3,4].\nReference:\n[1] Xiao, B. et al. Off-stoichiometry-guided design of high-strength chemically complex\nintermetallic-based alloys with outstanding ductility. Journal of Materials Science &\nTechnology 160, 28-33 (2023).\n[2] Yang, T. et al. Ultrahigh-strength and ductile superlattice alloys with nanoscale\ndisordered interfaces. Science 369, 427-432 (2020).\n[3] Gorbatov, O. I. et al. Effect of composition on antiphase boundary energy in Ni3Al-\n13\nbased alloys: Ab initio calculations. Physical Review B 93, 224106 (2016).\n[4] Liang, Z.D. et al. High-Temperature CoNi-Based Superalloys Strengthened by γ′-\n(Ni,Co) (Cr,Al,Ti,X): The Effect of Refractory Elements. Metallurgical and Materials\nModification\nOn page 4, we have added the following text:\nTo suppress fatigue-induced intergranular cracking, we devised a design strategy that\ncouples chemical and structural tuning in multicomponent intermetallic alloys. We deviated\nfrom the stoichiometric composition 23 and exploited the cooperative interactions among\nvarious elements 3,24 to design a multi-element symbiotic intermetallic alloy (MSIMA) with a\nnominal composition of (Co Ni Ti Al Ta V ) B . Such a design drives the segregation of\n47 30 4 10 4 5 99 1\nface-centered cubic (FCC)-stabilizing elements to GBs, yielding a core-shell structure\ncomprising a L1 -type ordered superlattice grain (OSG) encapsulated by an FCC-type\n2\ndisordered interfacial nanolayer (DINL) (Fig. 1a).\n2. The manuscript attributes the high work-hardening capacity to enhanced dislocation\nactivities facilitated by the disordered interfacial nanolayer (DINL). However, it remains\nunclear what the actual volume fraction or spatial distribution of the DINL region is, and\nwhether it is sufficient to account for such significant plasticity. Quantitative data or statistical\nanalysis should be provided. Or did the author also observe the fatigued microstructure of\nchemically complex L12 phase without DINL layer by TEM.\nResponse: Thank you for your constructive comments.\nAssuming each grain as a sphere, the volume fraction of the DINL region (Φ ) can be\nDINL\nestimated geometrically as\n𝑉 𝑅3\nDINL\nΦ = = 1−\nDINL 𝑉 (𝑅+𝑊)3\nwhere 𝑉 and 𝑉denote the volumes of the DINL and the entire grain, respectively; 𝑅 is the\nDINL\ngrain radius, and 𝑊is the thickness of the DINL. The average DINL thicknesses in both the\nof 10.6 ± 5.7 μm for CG-MSIMA and 0.8 ± 0.36 μm for UFG-MSIMA, the Φ in the CG-\nDINL\nand UFG-MSIMA yields 0.11 vol. % and 1.5 vol. %, respectively. In addition, the chemically\ncomplex L1 -type intermetallic alloy of (Co Ni Ti4Al Ta V ) B without well-\n2 47 30 10 4 5 99.98 0.02\ndecorated DINLs was fabricated following the same processing route as the CG-MSIMA alloy.\nIt shows very limited ductility (~10%) upon tensile deformation accompanied by serious\nintergranular cracking. This further demonstrates the critical role of DINL in such significant\nductility enhancement. Regarding the fatigued microstructure of the chemically complex L1\n2\nphase without the DINL layer, we have not characterized it by TEM. We will prepare the\nsamples and carefully analyze them in the future.\nModification\nOn page 21 of the revised Supplementary Information, a detailed calculation of the volume\nfraction of the DINL is added as Note 1:\nAssuming each grain as a sphere, the volume fraction of the DINL region (Φ ) can be\nDINL\nestimated geometrically as:\n𝑉 𝑅3\nDINL\nΦ = = 1−\nDINL 𝑉 (𝑅+𝑊)3\nwhere 𝑉 and 𝑉denote the volumes of the DINL and the entire grain, respectively, 𝑅 is the\nDINL\ngrain radius, and 𝑊is the thickness of the DINL. The average DINL thicknesses in both the\nCG- and UFG-MSIMA are approximately 2 nm. Based on the mean grain sizes (diameter, d)\nof 10.6 ± 5.7 μm for CG-MSIMA and 0.8 ± 0.36 μm for UFG-MSIMA, the Φ in the CG-\nDINL\nand UFG-MSIMA yields 0.11 vol. % and 1.5 vol. %, respectively.\nOn page 5 of the manuscript, we added this description:\nThe corresponding volume fraction of DINL is 0.11 % in CG-MSIMA and 1.5 % in UFG-\nMSIMA, respectively (Suppl. Note 1).\n3. In Fig. 2, the stress-strain behavior of the MSIMA with DINL shows significantly higher\n15\nvalues compared with the reference alloy. It is necessary to clarify whether other factors, such\nas grain size, phase fraction, or compositional variations, could also contribute to this\nenhancement. Please elaborate on whether this comparison alloy is representative.\nResponse: We appreciate the reviewer’s valuable comment regarding the rationale for\ncomparing the MSIMA and reference alloy. As shown in the Suppl. Figs. 2 and 7, both the\nCG-MSIMA and the reference (Co Ti ) B alloy possess comparable grain sizes and share a\n76 24 99 1\nsimilar L1 -ordered structure. Therefore, the pronounced difference in mechanical performance\n2\ncannot be attributed to different grain size or phase fraction. The markedly higher strength of\nthe MSIMA alloy primarily originates from the introduction of multiple alloying elements,\nwhich significantly increases the antiphase boundary (APB) energy and thus enhances the\ndegree of order strengthening. The improved ductility is mainly ascribed to the formation of\nDINL.\nThe (Co Ti ) B alloy, an off-stoichiometric variant of Co Ti intermetallic alloy, was\n76 24 99 1 3\nselected as the reference alloy because it is a prototypical L1₂-ordered system with a simple\nbinary composition and well-defined crystal structure. These chemical and structural characters\nensure a valid baseline for comparison with our MSIMA. The intrinsic brittleness and poor\nfatigue resistance of (Co Ti ) B typify the inherent limitations of conventional intermetallic\n76 24 99 1\nalloys, providing an ideal benchmark to demonstrate how multi-element symbiosis and\ninterfacial disordering fundamentally enhances the strength, ductility, and fatigue resistance.\n4. How is the stability of the DINL at elevated temperatures? Whether this alloy can be\napplied at elevated temperatures since the chemical composition appears to be tailored for\nelevated temperature applications with Ta addition.\nResponse: Thanks for your insightful comments. Here, we focus on addressing the poor\nfatigue resistance of intermetallic alloys at RT and have not evaluated their performance at\nelevated temperatures. Systematic studies on this interesting topic will be conducted soon and\nreported in a separate work.\n16\n5. On Page 13, line 231, the statement “Superdislocations, emerging as the primary strain\ncarriers, were first emitted from DINLs” appears inaccurate. The dislocations emitted from the\nDINLs are expected to be perfect dislocations in the disordered region, which then evolve into\nsuperdislocations when entering the ordered lattice. This distinction should be corrected for\nphysical accuracy.\nResponse: Thanks for your kind comments. This sentence has been revised to\n“Dislocations, acting as the primary carriers of plastic strain, are first emitted from the DINLs\nand subsequently transform into superdislocations upon entering the ordered lattice. (Page 14)”\n6. On Page 14, lines 251–254, the sentence “These results indicate that the SFs/nanotwins\nare induced by the slip of Shockley partials, which originate either from the intersection of\nsuperdislocations or from the further SF/twin-dislocation interaction” is not clearly or correctly\nexpressed. The proposed dislocation reactions are inconsistent with the described mechanisms.\nIn particular, for the reaction 1/2[101]+1/2[01-1]=1/3[111]+1/6[11-2], the Burgers vectors do\nnot balance on both sides. In addition, since superdislocations 1/2<101> were proposed as the\nsources for forming Shockley partials, APBs are expected to form during this process. Whether\nthe authors also observed APBs along with SFs?\nResponse: Thanks for your good comment.\nThe results indicate that the SFs and nanotwins are induced by the glide of Shockley\npartials. These partials may originate from (i) the interaction between superlattice partial\ndislocations ( 1 < 101> ), for example, through the reaction 1 [101]+ 1 [011̅]= 1 [111]+\n2 2 2 3\n1 [112̅] , or (ii) the interaction between a superlattice partial dislocation and a pre-existing\n6\nstacking fault or twin dislocation, such as 1 [101]= 1 [111]+ 1 [12̅1]. We have checked the\n2 3 6\ndislocation reaction 1 [101]+ 1 [011̅]= 1 [111]+ 1 [112̅], and it is balanced.\n2 2 3 6\nShockley partials, APBs inevitably form during this process because the superlattice partial is\nbounded by an APB, as illustrated in Suppl. Fig. 12. However, in practice, APBs in the MSIMA\nalloy after fatigue are extremely challenging to capture experimentally, which can be attributed\nto three factors. First, obtaining a clear atomic-resolution HAADF-STEM image of an APB\nrequires a pronounced Z-contrast between atoms on the A and B sublattices. In the MSIMA\nalloy, Co and Ni occupy the face-centered sites (A-sublattice), whereas Ti, Al, Ta, and V reside\non the corner sites (B-sublattice) of the L1₂ structure. The multicomponent occupancy on the\nB-sublattice—particularly the addition of V and Ta—reduces the atomic-size contrast between\nthe A- and B-sublattices, making the APB barely distinguishable in HAADF imaging. Second,\nthe subsequent motion of superlattice partial dislocations can remove pre-existing APBs during\ncontinued deformation. Third, nanotwins and stacking faults form at a later stage of fatigue,\nwhen the long-range order has been substantially reduced (Suppl. Fig. 15), making it more\ndifficult to capture the APBs.\nModification\nOn page 15 of the revised Supplementary Information, the sentence has been revised to :\nThese results indicate that the SFs/nanotwins are induced by the slip of Shockley partial\ndislocations. These Shockley partials could originate either from the intersection of superlattice\npartial dislocations ( 1 < 101>), for example via the reaction ( 1 [101]+ 1 [011̅]= 1 [111]+\n2 2 2 3\n1 [112̅]), or from the further SF/twin-dislocation interaction ( 1 [101]= 1 [111]+ 1 [12̅1]), as\n6 2 3 6\nillustrated in Suppl. Fig. 12.\n18"}, {"round": 2, "reviewer_comments": {"reviewer_1": "After careful reviewing authors' detailed responses, I believer the authors have addressed all comments properly. Then, I recommend the publication of this high-quality work.", "reviewer_2": "The authors have well responded the questions and my concerns. It is okay for me now.", "reviewer_3": "The authors have thoroughly addressed the raised concerns. I therefore recommend the revised manuscript for publication in Nature Communications."}, "author_response": "Response to the Reviewers' Report\nNCOMMS-25-60783B\nReviewer #1 (Remarks to the Author):\nAfter careful reviewing authors' detailed responses, I believer the authors have\naddressed all comments properly. Then, I recommend the publication of this high-\nquality work.\nResponse: We cordially thank the reviewer for their kind appreciation of the revision\nwork.\nReviewer #2 (Remarks to the Author):\nThe authors have well responded the questions and my concerns. It is okay for me now.\nResponse: We sincerely thank the reviewer for their kind appreciation of the revision\nwork.\nReviewer #3 (Remarks to the Author):\nThe authors have thoroughly addressed the raised concerns. I therefore recommend the\nrevised manuscript for publication in Nature Communications.\nResponse: We are very grateful to the reviewer for the supportive comments."}]

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