Hyper-Range Amorphization Unlocks Superior Damage Tolerance in Alloys

Domain: physics
Published: 2025-11-24
Editorial outcome: major_revision

Abstract
Shear bands dictate the failure mechanisms of alloys across various strain rates and limit the damage tolerance of the alloy. While short-range amorphization has the potential to mitigate shear effects, it has thus far been confined to the nanoscale. Here, we extend amorphization to the micrometer scale, fundamentally replacing shear-dominated failure in multi-principal element alloy micropillars. We implement continuous compression strain-training from low to high strain rates, generating a top-down high-density dislocation gradient that drives the formation of a topological disorder network, extending over one-third of the micropillar height, which we define as hyper-range amorphization. Within the amorphous bands, atoms exhibit dynamic disorder, and the lattice rearranges and recovers, dissipating shear stress. The alloy achieves an ultimate compressive strength of ceramic level (~6.5 GPa), while maintaining ~59.1% plasticity. This work reveals a strain engineering-based mechanical mechanism for extending amorphization, establishing it as a viable pathway to enhancing the structural stability and energy dissipation capacity of alloys.

Introduction
Damagetoleranceisafundamentalpropertythatdictatesthestructuralreliabilityofalloysincriticalengineeringapplications,includingshipbuilding,automotive,microelectronics,construction,andaerospace,contributingtoaglobalmarketexceeding€3trillionannually1–3.Enhancingdamagetolerancerequiresachievingadelicatebalancebetweenstrengthandductility,yetconventionalmetallicsystemsoftenexhibitatrade-offbetweenthetwo4,5.Along-standingchallengeinsolidmechanicsistheinherentlimitationimposedbyshearbands,whichgovernthefailuremechanismsofalloysacrossawiderangeofstrainrates4,5.

Shearbandsemergeashighlylocalizedregionsofintenseplasticdeformation,actingasprecursorstocatastrophicfailure6.Withinthesebands,severelatticedistortionleadstodislocationpile-upsoramorphization,andeventualmicrocrackformation7.Traditionally,amorphizationhasbeenviewedasdetrimentaltoplasticdeformationduetoitslackofslipsystems,potentiallyacceleratinglatticeinstability8.However,recentstudies9–11suggestthatunderspecificconditions,amorphizationcanfacilitateplasticloadtransfer,potentiallyimprovingdamageresistance.Conventionalamorphousorglassyalloysaregenerallyproducedviarapidquenchingoralloyingroutesthatkineticallysuppresscrystallization12–16.Incontrast,thepresentstrain-trainingstrategyinducesamorphizationdirectlywithincrystallinealloysthroughmechanicalloading,therebybypassingthermalprocessinglimits.Thisapproachisdistinctfromothernon-thermalamorphizationpathways,suchasionirradiation17,themeltquenchingprocess14,orsingle-rollermelt-spinning15.Yet,itsinherentlylocalizednature,usuallyrestrictedtothenanoscale,preventsamorphizationfromeffectivelycounteractingthestructuraldegradationassociatedwithshearbanding.

Here,wereportthesuccessfulreplacementofshearbandswithahyper-rangeamorphizationmechanisminmulti-principalelementalloy(MPEA)micropillars.Byapplyingcontinuouscompressionacrossmultiplestrainrates(strain-training),amorphizationextendsbeyondtheconventionalnanoscalelimit,reachinganunprecedented~1.34μm,equivalenttoone-thirdofthedeformedmicropillarheight.Doublesphericalaberrationcorrectedtransmissionelectronmicroscopy(DC-TEM)anddensityfunctionaltheory(DFT)calculationsrevealthatatop-downdislocationgradientmediatestheaccumulationoftopologicallatticedefects,whichinturntriggerlong-rangeatomicrearrangementandnanocrystalrecovery,dissipatingshearstressesupto~6GPa.Contrarytoconventionalperceptionsofdamagetoleranceinalloys,hyper-rangeamorphizationnearlydoublestheultimatecompressivestrengthandductilitycomparedtoconventionalhigh-strain-ratedeformation.

Results
Strain-traininginducedhyper-rangeamorphizationandsuperiorproperties

Webeginbyexploringafundamentalframeworkinthefieldofalloyplasticmechanics:thelimitationsofdislocationmigrationcapacityondeformability.Duringplasticdeformation,achievingabalancebetweenhighstrengthandductilitynecessitatessufficientspacefordislocationmigrationandproliferation(SupplementaryNote1).TheeutecticMPEAsystem,characterizedbysemi-coherentphaseboundaries,offersauniqueopportunitytomediateplasticitythroughstrongdeformationbonding.Asexpected,phaseboundaryalignmentparalleltothestressdirectionconstrainsdislocationmotion,facilitatinghighlydirectionalslip(Fig.1a).Thisarchitectureminimizesslipinterferencewhileensuringalow-energydislocationpathway,establishingtheframeworkforgradientdislocationpropagationandhyper-rangeamorphization.

Face-centeredcubic(FCC)andbody-centeredcubic(BCC)alternatingeutecticMPEAsweremanufacturedviaLPBF,followedbyfocusedionbeam(FIB)millingtoextract〈111〉-orientedFCCmicropillars(Fig.1b,c,andSupplementaryFigs.S1–S3).Toensurestructuralconsistency,multiplemicropillarsweresectionedfromasingleeutecticcluster,preservingtheverticalphaseboundaryalignment(Fig.1d).Pre-compressiontests(~5%)revealedthepotentialfor“top-down”dislocation-mediatedplasticity,asverifiedinmicropillarswithdistinctphaseboundaryconfigurations(Fig.1dandSupplementaryFigs.S4andS5).

In situ uniaxial compression tests (Fig. 2a–c, Supplementary Figs. S6–S11 and Supplementary Movies S1–S4) were performed at room temperature along the 〈111〉 direction of the FCC phase, spanning strain rates of ~8×10⁻⁴ s⁻¹ (SR), ~3×10⁻³ s⁻¹ (MR), ~2×10⁻² s⁻¹ (HR) and strain-training loading (ST, see Methods for details). The inherent heterogeneity of soft FCC and hard BCC phases prolonged the work-hardening stage even under intense shear, enabling all micropillars to achieve ultimate compressive strengths exceeding 2 GPa (Fig. 2d). Nevertheless, in SR, MR, and HR samples, shear band formation remained the dominant failure mode (Fig. 2a–c, Supplementary Figs. S6–S14). Increasing the strain rate, the strength and plasticity of the micropillars increased simultaneously. Remarkably, at the same strain rate hardening effect (strain rate limited to HR), the ST-loaded micropillars exhibited superior resistance to shear tilting, with minimal shear activity observed in the final deformation stages (Fig. 2c). This resulted in an extraordinary ultimate compressive strength of ~6.5 GPa (close to the theoretical value, Supplementary Note 2) and a ductility of ~59.1% (Fig. 2d).

Compared these results with reported FCC MPEAs18–22, noble metals (Au23, Ag24), traditional alloys (e.g., Ti alloys25,26,, Al alloys27–30, Mg alloys31, Cu alloys32,33, Ni alloys34–36, steels37–39), ceramics40 and BCC MPEAs41,42, the studied ST micropillar exhibited a superior strength-ductility balance (Fig. 2e). By fitting the logarithms of the micropillar strength (lnσ) and strain rate (lnε_), a high strain rate sensitivity (m≈0.13) was observed with stable linearity, excluding fluctuations in strain rate strengthening (Supplementary Fig. S12).

Post-deformation analysis revealed that amorphized regions (~1.34 μm and ~1.01 μm) formed in patterns reminiscent of shear bands, yet extended significantly beyond the nanoscale limits reported in previous studies9–11. This mechanism allows the alloy to overcome the strain rate strengthening effect in the range of maximum strain rates (2×10⁻² s⁻¹) and to withstand additional external loads of up to ~3 GPa, which corresponds to a near doubling of its load-carrying capacity. The slope of the engineering stress–strain curve increased markedly compared to the constant rate test samples, reflecting the sustained work-hardening ability that rejuvenated the “aging” alloy.

It has been reported that dislocation migration in the soft FCC phase is relatively easy. However, dislocations encounter significant interfacial energy at vertical phase boundaries, restricting the lateral expansion of plastic deformation zones. The hard BCC phase acts as a supportive scaffold, limiting excessive deformation in the soft FCC regions and enhancing the overall load-bearing capacity. This is the reason why no obvious surface protrusions were found in the four samples (Fig. 2a–c).

TEM further revealed that at lower strain rates, high-density dislocations and deformation-induced defects concentrated near the micropillar top surface (Supplementary Fig. S14). The limited activation energy at these strain rates hindered long-range dislocation migration, leading to local stress concentration and shear-band-dominated failure (Fig. 2b and Supplementary Fig. S14). At higher strain rates, deformation occurs within a shorter timescale, promoting early-stage shear band nucleation that intensifies shear tilting (Fig. 2a–c). Unlike single-crystal micropillars, where dislocations can escape through free surfaces, the combined constraints of phase boundaries and shear bands restrict dislocation mobility, leading to secondary shear band nucleation at the pillar base in later deformation stages (Fig. 2a–c).

In contrast, ST loading enables a progressive transfer of stress from the top to the base of the micropillar, forming a dislocation gradient distribution along the height (Fig. 2c, d). Near the loading surface, severe slip and fracture events occur due to concentrated stress and temperature rise (Fig. 2d). Mid-section deformation manifests as multidirectional slip and rotational distortion, whereas the pillar base experiences constrained stress propagation, resulting in lower stress levels but pronounced structural distortion (Fig. 2d). This multistage loading induces a continuously decreasing stress gradient, ultimately facilitating the gradient-mediated plastic deformation observed in hyper-range amorphization.

Evolution of hyper-range amorphization

The observed amorphization exhibits a hyper-range characteristic, extending over approximately one-third of the deformed micropillar height, with widths exceeding ~10 nm (Fig. 3a–f). This phenomenon, mediated by crystal defects, spans a significantly broader region than previously reported cases9,11,43–47 (Fig. 3c). This hyper-range amorphization initiates in the fractured microstructural regions near the loading end, where atomic instability under high stress induces cavity formation (Fig. 3a). These cavities disrupt the long-range atomic stacking in adjacent matrix regions, triggering a cascade of disorder across the matrix (Fig. 3a). Selected-area electron diffraction (SAED) confirms the amorphous nature of the alloy micropillar in these regions (Fig. 3b). At the atomic scale, precise control of large displacements or high strain rates during in situ multistage loading remains a challenge, complicating direct observation of the hyper-range amorphization process. However, the presence of shear cavities formed under high-speed compression provides indirect evidence that amorphization predominantly occurs during the final high-strain-rate deformation stage, with prior deformation stages laying the structural foundation for this transition.

Fig. 3d, e, and their corresponding FFT patterns (Fig. 3g–j) reveal a transitional gradient zone between the amorphous band and the matrix, where partial lattice order is retained, forming a semi-amorphous region between the crystalline and amorphous states. The amorphous band itself lacks long-range lattice periodicity, effectively absorbing or annihilating dislocations near the interface. As a result, the dislocation density in the vicinity of the amorphous band is significantly lower than in other matrix regions (Fig. 3e, i). Similar to crystalline-amorphous interfaces observed in metallic glass composites48–50, the amorphous band interface serves as a terminus for dislocation motion, reducing dislocation accumulation and enhancing structural stability.

Notably, the amorphization process exhibits a downward bending trajectory upon traversing phase boundaries, with the soft FCC phase acting as a stress dissipater, concentrating deformation within the FCC-dominant regions. Fig. 3f, j illustrate that at the boundary between the gradient zone and the matrix, distinct diffraction spots coexist with diffuse disordered diffraction rings, indicating that the amorphous region actively participates in stress redistribution and load dissipation.

Near the amorphous band tip, stacking faults and subgrain boundaries preferentially form (Fig. 3k–m), whereas such defects are absent in other amorphous regions. This suggests that stacking faults and subgrain boundaries may act as precursors for the transition from crystalline to amorphous states. Local stress concentration at the tip is mitigated through defect reorganization, facilitating the progressive expansion of the amorphous band (Fig. 3k, m). These findings establish a direct correlation between crystalline defects and amorphization onset, where stacking faults and subgrain boundaries provide both initial conditions for disordering and local slip pathways in the absence of conventional slip systems, promoting the propagation of amorphization. Energy-dispersive X-ray spectroscopy (EDS) analysis (Fig. 3n) further confirms that the amorphous phase originates from the matrix, with a pronounced density mismatch effect51.

To gain deeper insight into this phenomenon, we employed density functional theory (DFT) simulations to investigate the evolution of atomic structure, energy, and stress under ST loading. A topological lattice model was constructed based on the micropillar’s crystal structure (Fig. 4a, b). In the initial state, minor lattice mismatches at the semicoherent phase boundary induce a network of topological defects, primarily pentagonal and heptagonal rings, in contrast to the typical hexagonal arrangement in perfect metallic crystals (Fig. 4b).

The model was subjected to simulated ST-mode loading (Fig. 4c), and the corresponding structural, mechanical, and energy evolution is illustrated in Fig. 4c–h. As the applied stress approaches its peak (Fig. 4g), the accumulation of high-density dislocations surpasses conventional annihilation mechanisms, leading to extensive dislocation interactions and the nucleation of successive topological defects. This cascade of defect formation induces atomic-scale disorder (Fig. 4d, e). At this stage, atomic bonding energy fluctuations and geometric distortions become increasingly pronounced, highlighting the role of dislocation-mediated defect proliferation in facilitating amorphization. Furthermore, the loss of lattice periodicity broadens the atomic bond energy distribution, enabling effective stress redistribution and energy dissipation, as reflected by the rapid decline in system energy and stress levels (Fig. 4g).

Disordered atoms undergo repetitive oscillations and irregular spatial transitions to accommodate external stress, characterized by localized atomic expansion and dynamic rearrangement, which serve as effective energy dissipation pathways (Fig. 4f). Additionally, atoms in high-energy disordered states remain thermodynamically metastable, allowing for local structural recovery under subsequent stress, transforming into low-energy crystalline structures (Fig. 4h). While direct in situ observation of this dynamic process remains technically challenging, high-resolution imaging of the amorphous band tip (Fig. 4i–k) reveals residual lattice fringes and diffraction spots, indicating the presence of nanocrystalline domains embedded within fully amorphous regions. This finding directly validates the proposed mechanism of energy dissipation via atomic rearrangement and local lattice recovery. Furthermore, this dynamic stress relaxation process explains the absence of crack nucleation at the amorphous band-matrix interface, underscoring the superior structural stability and initial microcrack stability (Supplementary Note 3 and Supplementary Fig. S15) conferred by hyper-range amorphization.

Gradient dislocation-induced amorphization expansion for shear stress dissipation

These findings provide compelling evidence for the critical role of gradient dislocations in driving hyper-range amorphization, thereby regulating the mechanical behavior of the alloy. The synergistic coupling between eutectic phase boundaries and gradient dislocations plays a role in stress regulation and structural stability (Supplementary Notes 4, 8, and Supplementary Figs. S16–S24). To quantitatively assess this process, the applied flow stress (σ) is decomposed into normal stress (σ N) perpendicular to the shear band and shear stress (σ τ) along the shear direction (Fig. 5a)52. The evolution of σ N and σ τ is analyzed considering the shear band angle (Fig. 5b–e). The deformation of the micropillar is predominantly governed by shear stress σ τ, which increases significantly with strain rate.

For ST samples, shear stress accounts for more than ~90% of the total flow stress (Fig. 5c), highlighting the essential role of hyper-range amorphization and gradient deformation in dissipating shear energy. The resistance of ST-loaded micropillars to shear effects increases by ~389% and ~93% compared with SR and HR loading, respectively. The shear failure force sustained by ST samples surpasses that of most shear-band-dominated materials19,53–60 (Fig. 5d). Notably, shear failure occurs only at the final deformation stage, suggesting that amorphization initiates following cumulative defect accumulation during earlier stages and is ultimately triggered at high strain levels.

During each elevated strain rate of ST loading, new dislocations are inevitably generated within the alloy. The atoms in the crystal will undergo relative displacement, the original dislocations will move and proliferate, and the dislocation density inside the alloy will increase. As the strain increases by orders of magnitude, new dislocations will be generated on the basis of the dislocations generated in the first step, and so on, cumulatively, which leads to a significant increase in the dislocation density (Fig. 5e, f). With the increasing dislocation density during the strain-training process, complex interactions such as entanglement and intersection occur between dislocations, hindering further dislocation movement. When an external force tries to deform the crystal, it is necessary to overcome these obstacles formed by dislocation interactions, which leads to an increase in the alloy’s ability to resist deformation, thus increasing the alloy’s strength in the first few stages. Internal stress caused by the concentration of defects weakens the alloy’s ability to resist external forces and exacerbates the formation of shear voids.

Strain-training mitigates this issue by homogenizing dislocation distribution, promoting gradual internal stress release, and enabling dislocation migration rather than localized accumulation. This process is conceptually similar to multistep stamping in industrial forming, where each deformation step corrects and refines the microstructural evolution of the material61. Each strain increment effectively readjusts the alloy’s internal structure, facilitating progressive stress redistribution, expanding dislocation pathways, and enhancing its mechanical stability.

Through this top-down gradient dislocation architecture, stress is efficiently transferred from high-density dislocation regions to low-density zones, minimizing local stress concentrations (Fig. 5f–i). Zhao et al.11 established a quantitative relationship between critical dislocation density (ρ a) for amorphization and stacking fault energy (γ sf) (Eq. 1), where b is the Burgers vector:

ρ a=γ sf / (Gb3) (1)

For the FCC-based alloy (CoCrAlFeNi), with γ sf=~4.33 mJ/m²62 and b=~0.254 nm, the predicted ρ a threshold for amorphization is ~3.26×10¹⁵ m⁻². Experimental data show that the top and middle regions of ST-loaded micropillars exceed this critical value, thereby enabling amorphization (Fig. 5e). As shear stress propagates, the amorphous band extends laterally from the high-dislocation-density region into the FCC matrix, which is inherently more susceptible to defect formation. The gradient-driven amorphization terminates at ~1.35 μm from the micropillar base (Fig. 2d), where the dislocation density falls below the amorphization threshold (Fig. 5e–i). The contribution of amorphization strengthening to damage tolerance, calculated using the rule of mixtures (ROM) model63, reaches ~3.663 GPa (Supplementary Note 5, Fig. 5j, and Supplementary Table S1).

Studies by Huang et al.64 and Wang et al.9, in conjunction with our findings, confirm that amorphization facilitates local energy dissipation through dislocation emission and lattice recovery at crack tips, thereby reducing stored elastic energy and stabilizing the alloy structure. This process aligns with Griffith’s65 energy release rate (G c) and its relationship with the deformation band length (a) (Eq. 2):

G c= K² / E = τ² πa / E (2)

where the Young’s modulus E~50 GPa, a~1.34 μm, and τ~6 GPa. The calculated energy release window for hyper-range amorphization is G c~3.235 J/m², providing an effective mechanism for stress dissipation in alloy damage.

During the ST process, topological lattice defects accumulate at low and moderate strain rates, forming a defect-rich network that provides the foundation for strain-induced amorphization under high-rate compression (Figs. 2c and 5e). This final stage supplies the shear energy necessary to drive hyper-range amorphization, allowing the alloy to adapt to extreme strain conditions through amorphization rather than conventional shear banding.

Extending hyper-range amorphization to CoCrNi alloy systems

To verify the applicability of the ST strategy across different alloy systems, we employed a representative cast face-centered cubic (FCC) single-phase CoCrNi alloy and investigated its deformation mechanisms along the (110) and (111) crystallographic orientations—both of which correspond to preferential slip directions (Supplementary Fig. S19). Micropillars with a diameter of 2 μm, matching that of the eutectic samples, were fabricated via precise FIB milling, and crystallographic orientations were confirmed using EBSD. These samples were subjected to both ST loading to systematically assess their strengthening responses.

According to Zhang et al.19, CoCrNi micropillars oriented along (110) and (111) exhibit pronounced ductility under SR loading, accompanied by extensive dislocation escape toward free surfaces. Post-yield deformation often leads to localized strain accumulation, which rapidly propagates toward the surface, manifesting as pronounced shear slip bands. This surface-directed dislocation escape diminishes internal dislocation storage, thereby compromising load transfer and accelerating plastic localization.

Under the ST loading protocol, deformation in CoCrNi (110) and (111) micropillars is governed by a hybrid mechanism involving both crystal slip and amorphization (Figs. 6 and 7). These micropillars exhibit strength/ductility comparable to that of 500 nm nanopillars (Fig. 8), with the corresponding ST processes detailed in Supplementary Figs. S20 and S21. Notably, ST triggers hyper-range amorphization (Figs. 6 and 7) while also inducing stress accumulation at the pillar top. Together, these effects restrict dislocation escape and promote internal multiplication and slip, facilitating vertical stress redistribution. This process is accompanied by deformation twinning, low-angle grain boundary formation, and widespread Lomer–Cottrell (LC) lock development. However, due to the absence of phase boundaries, as present in eutectic systems, localized slip can still propagate from the non-amorphized adjacent regions.

Unlike eutectic micropillars, where amorphization is stabilized and enhanced by phase boundaries, the single-crystal CoCrNi alloy relies on the amorphization-induced barrier to dislocation escape for uniform stress redistribution. This promotes the cooperative evolution of dense dislocation networks, stacking faults, LC locks, and deformation twins. Despite distinct mechanistic pathways, both systems ultimately achieve enhanced strength–ductility synergy via the ST-induced coupling of shear shielding (or partial shield) and amorphization. These findings underscore that the strain engineering–dislocation regulation–stress reconstruction–deformation control strategy is not inherently dependent on eutectic structures, thus demonstrating the generality and cross-system applicability of the ST approach.

Methods
Materials
The eutectic MPEA with a chemical composition of Co17Cr16Fe16Ni34Al17 was manufactured using a BLK series L-PBF machine custom-made by Foshan Chenfeng Material Technology Co., Ltd., and Bright Laser Technologies. This machine is equipped with a Yb fiber laser with a maximum power of 500 W. The powders were prepared by Foshan Chenfeng Material Technology Co., Ltd. using gas atomization, with particle sizes ranging from 15 μm to 53 μm (average size of 30 μm), and were used to construct alloy plates with dimensions of 40 mm (length) × 10 mm (width) × 2.1 mm (thickness). During bidirectional printing, 4140 steel was used as the substrate, with a scanning speed of 1000 mm/s, a layer thickness of 40 μm, and a scan spacing of 80 μm.

The MPEA samples were cut into 2 mm (width) × 1 mm (thickness) samples by wire-cutting. The surfaces were polished with 2000#, 3000#, and 5000# grits and papers and then vibrationally polished to prepare EBSD samples. High-deformation-tendency (111) eutectic colonies were identified via EBSD analysis, followed by dual-beam focused ion beam/scanning electron microscope (FIB/SEM, FEI Helios NanoLab 600i) to prepare micropillar samples (Fig. 1 and Supplementary Fig. S1). A standard lift-out and annular milling procedure was employed. A rough milling step was performed at 30 kV with beam currents of 2.5–9 nA, followed by sequentially lower currents down to 0.28 nA to minimize tapering. Final surface cleaning was carried out at 5 kV and 48 pA to reduce Ga⁺ implantation and surface damage.

Efforts were made to ensure that the phase proportions and morphologies in the micropillars were as similar as possible, minimizing the potential errors caused by microstructural variability. The micropillars had a diameter of 2 µm and a height of 6 µm.

For FCC/BCC stacking structures, the phase boundary acts as an early obstacle in the dislocation slip process, making it difficult to traverse multiple phase boundaries, thus increasing the likelihood of stress concentration at the top and heterogeneous nucleation of shear bands occurring at the phase boundary. Selection of the target alloy and customization of the dislocation slip direction are the first steps in exploring the damage tolerance of an alloy.

We used a FIB instrument to expose the micropillar and confirmed that the phase boundary was parallel to the stress direction. Additionally, we observed the BCC and FCC lattices of the matrix, providing reference information that can be used as the basis for subsequent theoretical modeling. Unlike other phase boundaries, the BCC and FCC phases satisfy the Kurdjumov–Sachs orientation relationship. With parallel phase boundaries and strong bonding capabilities, dislocations preferentially migrate and diffuse along these low-energy channels over long distances. Under high-stress conditions, dislocations may act as dislocation sources, where dislocations can proliferate near the phase boundary through reactions or climb, avoiding the surface proliferation of dislocations. In our macroscopic tensile tests of eutectic MPEAs, dislocations still exhibited characteristics of proliferation along phase boundaries, providing a preliminary verification that interference from second-phase precipitates can be neglected in the design.

Microstructural characterization
Thin foils for TEM were extracted from deformed micropillars along the diameter direction using lift-out. Final thinning and polishing were conducted at 5 kV and 48 pA to achieve an electron-transparent thickness of <100 nm. TEM analysis was performed using Talos F200X and double spherical aberration-corrected Spectra 300 microscopes (Thermo Fisher Scientific) operated at 200 kV. The microstructure and defect evolution of both the substrate and micropillars were investigated. The MPEA samples exhibited nanoscale FCC and BCC phases with alternating lamellar structures and sharp, flat phase boundaries, with average sizes of ~140 nm (FCC) and ~70 nm (BCC). High-resolution TEM (HRTEM) images were acquired along the [111] zone axis to resolve lattice distortions and amorphous regions. Fast Fourier transform (FFT) and inverse FFT analyses were carried out using Digital Micrograph 3.40.2804.0 (Gatan) to determine crystallographic orientations and highlight lattice fringes. For filtered images, the applied reciprocal-space masks and filter parameters are provided in the corresponding figure captions. All images were recorded under identical camera length and exposure conditions to ensure reproducibility.

In situ mechanical testing
Micropillar compression tests were conducted using a high-resolution field emission scanning electron microscope, with the strain rates carefully controlled. Compression was performed at constant strain rates of 8×10⁻⁴ s⁻¹ (SR), 3×10⁻³ s⁻¹ (MR), and 2×10⁻² s⁻¹ (HR), as well as under varying strain rates (Increase from SR to HR, referred to as strain-training compression, ST). For each loading condition, tests were performed on three independent micropillars prepared from different regions of the bulk sample to ensure statistical independence.

Selection of SR. A strain rate of 8×10⁻⁴ s⁻¹ corresponds to the conventional quasi-static compression rate (~10⁻⁴/s level), facilitating the exploration of the material’s intrinsic compressive strength and plasticity.

Selection of MR. The strain rate of 3×10⁻³ s⁻¹ follows the “order-of-magnitude stepwise” principle. This controlled rate increase (<10×) prevents mismatches between video capture and load application systems while simultaneously introducing strain rate hardening effects.

Hardware constraints on HR. The 2×10⁻² s⁻¹ strain rate is chosen as the upper limit due to experimental system constraints. The video acquisition system requires a sufficiently high frame rate to capture the microcolumn deformation process accurately, while the motor control system must maintain load stability with nanometer precision. These two factors collectively determine the strain rate ceiling. Exceeding this threshold risks a mismatch between video capture and mechanical loading systems. In situ mechanical testing in small-scale environments may amplify the effects of strain rate hardening66. At this scale, a strain rate of 2×10⁻² s⁻¹ corresponds to a displacement speed of ~120 nm/s. In the work of Kim et al.67, a strain rate of 2×10⁻² s⁻¹ at the macroscale was referred to as a high strain rate.

Strain transition management of ST. Prior to the implementation of ST, we carefully designed the strain rate transition protocol based on comprehensive experiments with single-strain-rate samples. Three key criteria were established: (1) The strain rate transition point must reach the yield strain of the micropillar but remain lower than the strain threshold where shear localization occurs. (2) Sufficient dislocation density must be accumulated to reach the critical level required for triggering amorphization. (3) Reproducibility must be confirmed through repeated tests (three experiments, considering repeatability and experimental cost). The initial SR stage promotes uniform dislocation nucleation, the subsequent MR stage enhances dislocation gradient formation and interactions, and the final HR stage effectively triggers amorphization.

Micropillars loaded at slow strain rates are enriched with a high density of dislocations and deformation defects at the top. The low activation energy is not sufficient to support long-range migration of dislocations, leading to early localized stress concentrations and shear localization. To avoid premature shear or precursors in the SR stage of ST samples, compression was first loaded at a slow strain rate to the yield stage. Due to the small deviation of grain size and phase fraction within micropillars, the yield strain of the SR sample was ~5.5±1.5%, resulting in a strain transition point of ~6% for the ST sample. Immediately afterward, the strain rate was increased to an MR to enhance uniform dislocation proliferation. This stage is needed to ensure sufficient accumulation of dislocations to establish a topological defect foundation for subsequent amorphization. The selection process of the strain transition point and micropillars diameter is described in detail in Supplementary Notes 6–8 and Supplementary Figs. S22–S28. Consideration of the positive strain rate hardening effect will allow this stage to sustain longer compressive displacements. Furthermore, amorphization requires enhanced atomic migration to rapidly disrupt the crystalline order. Such processes must be accomplished within a short period. Consequently, once the strain approached the uniform deformation limit observed in the SR–MR loading stage, the strain rate was immediately elevated to the upper hardware limit (HR level), enabling rapid compression until pillar fracture or abnormal deformation of the substrate occurred.

System control was implemented using software-driven displacement coordinated with the indenter, ensuring smooth transitions between strain rates and a strain sensitivity of ~0.1%. Upon surpassing a predetermined threshold, the system automatically triggered a strain rate transition to the next stage. Due to the inherent inertia of the experimental apparatus and the acceleration protocol, a slight delay in the actual transition point may occur during the acceleration phase (affected by the magnitude of the strain rate jump).

Atomistic simulation
DFT calculations were carried out using the Vienna ab initio simulation package (VASP 5.4)68–71. The interaction between electrons and ions was treated using the projector augmented wave (PAW) method72. The exchange and correlation energies of electrons were described by the PBE functional73 within the generalized gradient approximation (GGA)74. The electronic wave functions were expanded with a plane-wave basis set to a kinetic cutoff energy of 350 eV, which is 1.3 times greater than the default value for all the elements in the alloy (see Supplementary Note 8 and Supplementary Figs. S29 and S30 for more detail). All atoms were fully relaxed until the total energy and the Hellmann‒Feynman forces on the atoms were less than 1×10⁻⁴ eV and 0.02 eV/Å, respectively. A 7×1×1 k-point grid was employed for Brillouin zone sampling of the supercell, which included 264 atoms. To construct a periodic supercell containing an interface between FCC and BCC high-entropy CrNiCoAlFe alloys, 144-atom FCC-like and 120-atom BCC-like special quasirandom structure (SQS) supercells were constructed with the alloy theoretic automated toolkit (ATAT)75–78. The relaxed lattice constants were a=2.54 Å, b=35.17 Å, and c=18.65 Å, and a=2.52 Å, b=34.28 Å, and c=15.72 Å for the FCC-like and BCC-like supercells, respectively. The interface was then generated by alternatively stacking the FCC-like and bcc-like supercells along the c direction, leading to a lattice mismatch of less than 3% along b.

The uniaxial compressive loading on the supercell is simulated here through a displacement-controlled method67,68, where the lattice vectors of the supercells are deformed incrementally by compression strains in the direction parallel to the phase boundaries with a compressive strain step of Δε=0.01. At each strain step, the applied compression strain ε is fixed to calculate the corresponding compression stress σ, and the other five independent elements of the strain tensors and the positions of all atoms inside the supercells are fully relaxed. The lattice constants of the compressed supercells, the atomic positions, and the compressive stress-strain-strain relation (σ–ε) are calculated at each step by this constrained relaxation. To simulate the strain-training transition process, the total target strain (20%) was divided into three incremental strain steps, thereby indirectly capturing the effect of strain-rate variation.

It is important to note that DFT-based simulations are typically performed on near-perfect crystal models and do not incorporate environmental temperature, sample heating, or realistic time-scale effects. In contrast, experimental conditions may induce plastic yielding at lower strains, or DFT simulations may exhibit premature phase transitions due to idealized assumptions. As such, the strain transition points in the simulations (e.g., 7% and 15%) differ from those in the experiments. These values were selected to trigger and observe comparable physical responses within the constraints of the DFT framework.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This study presents “strain training” loading strategy to enhance the compressive strength and ductility of eutectic high-entropy alloy micropillar. Through the strain training approach with changing strain rates, the authors achieved high strength and ductility based on amorphization and dissipation of shear stress, which alleviates shear band formation. Micromechanical investigation using micropillars inevitably carries a risk in terms of reproducibility, as it focuses on small scale mechanical properties at specific locations. It seems this work has a problem in this regard, as it presents contradicting results, which is specified in the comment below. I see it as a serious issue on which the reliability of the work depends, as the whole point of the work is about suggesting a methodology to replace shear-dominated failure mechanism. Therefore, I do not recommend publication of the manuscript unless the authors can adequately address this. Technical comments are as follows:\n1. According to Fig. 1g, under strain training loading till 20% strain at an 8*10^-4 s-1 strain rate does not induce any shear band. On the other hand, under slow rate loading which the authors are presenting for comparison to highlight the excellence of strain training loading, 12% strain at the same strain rate leads to a severe shear band formation, which soon induces failure. This is hard to understand, and a possible answer would be an error of reproducibility. Such a contradiction harms the credibility of the authors’ strain training strategy and all the supporting evidences to justify it – I doubt if the same results can be achieved from repeating the mechanical testing under the same condition.\n2. Fig. 1b: The EBSD map should be presented together to show the phase distribution within the micropillar. The colors for the phase maps in Fig. S1 are wrongly indexed.\n3. How is the theoretical strength limit of 3.5 GPa was determined? Is it possible that the flow stress exceeds the theoretical limit? It is usually determined as G/10 where G is shear modulus. I doubt that the shear modulus of the eutectic HEA would be 35 GPa [Scr. Mater. 186 (2020) 336].\n4. The Ashby map is missing some of the most important achievements in microscale mechanical properties [Nat. Commun. 13 (2022) 1102; Nat. Commun. 6 (2015) 7748], which actually outperform the current material.", "reviewer_2": "The authors successfully induced micrometer-scale amorphization in multi-principal element alloys through a novel strain-rate-controlled compression technique. They induced a high-density dislocation gradient that triggered a broad amorphous region. This process replaced traditional shear-dominated failure with a more stable, energy-dissipating mechanism. The resulting alloy demonstrated enhanced strength and plasticity. The results are interesting and have a scientific impact. Below are some questions and comments for this article.\n1. The authors stated the choice of 2x10^-2 s^-1 as HR is due to the limits of the video capture and motor control systems, is there any specific reason to choose the use of 8x10^-4 s^-1 and 3x10^-3 s^-1 as SR and MR? Also, in ST, how is the transition between strain rates (from 8x10^-4 s^-1 to 3x10^-3 s^-1 to 2x10^-2 s^-1) managed in real-time?\n2. What is the reason for choosing a kinetic cutoff energy of 350 eV that is 1.3 times greater than the default value, and how sensitive are the results to this parameter?\n3. For the ST-loaded micropillars, the amorphous band terminates at ~1.35 μm from the base. What determines this termination point? Is it linked primarily to the dislocation gradient itself or other microstructural constraints?\n4. Since lattice recovery is mentioned at crack tips, is there evidence that hyper-range amorphization can stabilize microcrack formation in earlier deformation stages?\n5. Could you elaborate on how the top-down dislocation gradient specifically interacts with phase boundaries to enhance stress redistribution?\nMinor comments:\n1. The caption of Figure 1: description for subfigure (c) is marked as (f). (g). Also, the description for (c, d) needs to be checked.\n2. Some abbreviations are not properly defined in the paper, such as stacking faults were only marked as SF in Figure 3, FIB was defined later in the paper after its first few appearances, etc.", "reviewer_3": "The manuscript titled “Hyper-Range Amorphization Unlocks Superior Damage Tolerance in Alloys” reported a novel strain engineering (strain training) strategy to achieve micrometer-scale amorphization in eutectic high-entropy alloy micropillars, leading to a significant enhancement of both strength and plasticity. This approach is interesting, and the experimental results provide important insights into strengthening and toughening mechanisms. However, several concerns require further considerations, especially regarding the universality of the proposed strategy. Specific comments are outlined below:\n1. This work proposed a strategy of strain-training, overcoming the shear effects and achieving a steady deformation. However, this study only focused on eutectic high entropy alloy micropillars, which is special. Does this strategy still work in a single-crystal system? The critical influence of the eutectic microstructures (e.g., phase boundaries, boundary orientations, boundary coherent, ...) has not been fully elucidated.\n2. The strain-training strategy is associated with sudden shift of strain rate at certain strains. Could you provide more clarifications or rationales for selecting the strain rate transition point? The effects of different transition timings on amorphization extent should be further discussed.\n3. Since it is a micropillar experiment, the influence of micropillar size should be considered.\n4. The strength calculation using the rule-of-mixture (ROM) require more clarifications, i.e. amorphous phase strength, fraction, etc..\n5. Some minor errors: (1) (Page 5) The strain rate sensitivity is not consistent with your supplementary Fig. S7. (2) (Page 12) ~3.26×10^15 lacks unit."}, "author_response": "Response to Referees\nNATURE COMMUNICATIONS\nRef. No.: NCOMMS-25-11276\nTitle: Hyper-Range Amorphization Unlocks Superior Damage Tolerance in Alloys\nDear Referees,\nWe sincerely thank you for taking the time and effort to review our manuscript and for\nproviding valuable feedback to improve its quality.\nIn response to your insightful comments, we have carefully conducted additional\nexperiments, simulations, and theoretical analyses to address all your concerns. Below,\nwe provide a point-by-point response detailing the revisions made in accordance with\nyour suggestions. All changes are highlighted in blue in the revised manuscript.\nThese revisions do not alter the overall content or structure of the paper. Should you\nhave any further questions or require additional clarification, please do not hesitate to\nlet us know. We will do our best to respond promptly and refine the manuscript to\nensure it is as complete and rigorous as possible.\nOnce again, we sincerely appreciate your constructive suggestions.\nThank you and best regards,\nYing Li\nProfessor\nBeijing Institute of Technology\nE-mail: bitliying@bit.edu.cn\n1 / 61\nTable of Contents\nResponse to Editor ....................................................... Error! Bookmark not defined.\nResponse to Referees ................................................................................................ 1\nResponds to the referee’s comments: ........................................................................ 3\nReferee #1. ........................................................................................................... 3\nComment 1 .................................................................................................... 3\nComment 2. ................................................................................................... 7\nComment 3. ................................................................................................... 9\nComment 4. ................................................................................................. 13\nReferee #2. ......................................................................................................... 15\nComment 1 .................................................................................................. 15\nComment 2 .................................................................................................. 23\nComment 3 .................................................................................................. 26\nComment 4 .................................................................................................. 27\nComment 5 .................................................................................................. 29\nMinor Comment 1 ........................................................................................ 34\nMinor Comment 2. ....................................................................................... 34\nReferee #3. ......................................................................................................... 36\nComment. 1. ................................................................................................ 36\nComment. 2. ................................................................................................ 46\nComment. 3. ................................................................................................ 53\nComment. 4.. ............................................................................................... 58\nMinor Comment 1. ....................................................................................... 61\n2 / 61\nResponds to the referee’s comments:\nReferee #1\nThis study presents “strain training” loading strategy to enhance the compressive\nstrength and ductility of eutectic high-entropy alloy micropillar. Through the strain\ntraining approach with changing strain rates, the authors achieved high strength\nand ductility based on amorphization and dissipation of shear stress, which\nalleviates shear band formation. Micromechanical investigation using micropillars\ninevitably carries a risk in terms of reproducibility, as it focuses on small scale\nmechanical properties at specific locations. It seems this work has a problem in\nthis regard, as it presents contradicting results, which is specified in the comment\nbelow. I see it as a serious issue on which the reliability of the work depends, as\nthe whole point of the work is about suggesting a methodology to replace shear-\ndominated failure mechanism. Therefore, I do not recommend publication of the\nmanuscript unless the authors can adequately address this.\nResponse: Thank you very much for taking the time and effort to review our paper\nand for your positive comments on our manuscript. The reproducibility concerns\nyou raised have been addressed through validation by the additional independent\nexperiments, confirming the accuracy of our results. Below, we provide a detailed\npoint-by-point response.\n1. According to Fig. 1g, under strain training loading till 20% strain at an 8*10^-4 s-1\nstrain rate does not induce any shear band. On the other hand, under slow rate\nloading which the authors are presenting for comparison to highlight the excellence\nof strain training loading, 12% strain at the same strain rate leads to a severe shear\nband formation, which soon induces failure. This is hard to understand, and a\npossible answer would be an error of reproducibility. Such a contradiction harms\nthe credibility of the authors’ strain training strategy and all the supporting\nevidences to justify it – I doubt if the same results can be achieved from repeating\nthe mechanical testing under the same condition.\nResponse: We sincerely appreciate your insightful analysis of this critical issue.\nWe acknowledge the annotation error in the original manuscript and deeply\napologize for any confusion it may have caused. Your identification of this\n3 / 61\napparent inconsistency prompted a thorough re-evaluation and clarification of the\nexperimental details in the revised version.\nThe strain-rate transition strategy in the ST experiments was developed based on\na systematic analysis of failure patterns observed in constant strain-rate tests, as\noriginally described in the Methods section. To prevent premature shear\nlocalization and its potential precursors, the first transition point was intentionally\nset near the yield strain. Immediately after, the strain rate was increased to the\nMR stage. The strain level of ~20% marked in the figure refers to the second\ntransition point, not the initial one (Figs. 1e–g). Additional details on the SR and\nST compression processes—including the initial nucleation of shear, the evolution\nof localization, and the exact locations of the strain-rate transitions—are provided\nin Supplementary Figs. S7–S11.\nWe have also included a lateral comparison at 20% strain (Figs. 1e–g): the SR\nsample begins to fail, the HR sample shows multiple shear band nucleation sites,\nwhereas the ST sample still maintains uniform plastic deformation.\nTo enhance the reproducibility of our findings, we performed two additional\nindependent strain training experiments (three in total, Supplementary Figs. S9–\nS11). The statistical analysis indicates that, with strict adherence to the sampling\ncriteria and compression protocol, the failure behavior of micropillars remains\nconsistent. Minor statistical variations are inherent to small-scale mechanical\ntesting and primarily stem from:\n Variations in surface roughness induced by FIB milling.\n Slight local deviations in crystallographic orientation.\n Stochastic distribution of dislocation sources.\n Subtle fluctuations in grain size and phase fraction.\nThese supplementary experiments collectively demonstrate that, despite the\nintrinsic scatter of micro-scale testing, the deformation mechanisms triggered by\nthe strain training (ST) strategy are highly reproducible. Once again, we apologize\nfor the oversight and thank you for encouraging us to present this critical\nconclusion with greater scientific rigor.\n4 / 61\nRevisions:\nWe have included this repeat experiment in the Methods section (Lines 431-432).\nFig. S7. Details of failure modes in micropillars under SR loading. (a) Early elastic\nand yielding stages. (b, c) Shear band nucleation and propagation under the SR loading\npath. Shear bands preferentially nucleate at a height of ~32 % from the top of the\nmicropillar and propagate toward the free surface at a characteristic angle of ~33° (the red\narrow shows the development of the shear band). When the substrate collapses, the\ncompression is automatically terminated by the contact stiffness monitoring system,\nindicating that inelastic deformation of the base has occurred.\nFig. S8. Details of failure modes in micropillars under ST loading. (a) Early elastic and\nyielding stages (SR stage). (b) Uniform deformation during the MR stage. (c) Final\ndeformation under the HR stage.\n5 / 61\nFirst replication of ST loading\na\n= 8 = 8 = 8\n11 μμmm SR→MR\nε= 0 ε= 2% ε= 4% ε= 6%\nb\n= 3 = 3 = 3\n1 μm MR→HR\n1 μm\nε= 9.5% ε= 13% ε= 16.5% ε= 20%\nc\n= 2 = 2 = 2\n1 μm\nε= 25% ε= 36% ε= 47% ε= 61.1%\nFig. S9. First-round reproducibility validation of microcolumn deformation under ST\nloading to mitigate shear action. (a) Early elastic and yielding stages (SR), highly\nconsistent with those in the main text. (b) Still uniform deformation after strain rate\ntransition, no shear slip seen. (c) HR loading stage.\nFig. S10. Second repeatability verification of suppressed shear localization under\nST loading. (a–c) The deformation behavior of the micropillar follows the same pattern as\nreported in the original manuscript and the first repeat test, showing no evident shear\ndeformation. Although image contrast is reduced due to surface potential fluctuations\ninduced by 20 kV electron beam charging, the overall deformation characteristics remain\nclearly identifiable. Statistical analysis across multiple samples confirms that the ST\nstrategy effectively mitigates shear-induced damage and demonstrates process reliability.\n6 / 61\nFig. S11. ST strategy exhibits outstanding mechanical stability, as evidenced by\nthree repeated tests showing minimal deviation in stress–strain responses.\nFigs. 1e-g. (e), HR (f) and ST (g) loading modes, where SR and HR samples fail\nprematurely due to shear localization, while ST samples exhibit full deformation.\n2. Fig. 1b: The EBSD map should be presented together to show the phase\ndistribution within the micropillar. The colors for the phase maps in Fig. S1 are\nwrongly indexed.\nResponse: Thank you very much for your insightful and helpful suggestions.\n7 / 61\nIn response to your valuable recommendation, we have added a statistically\nmeaningful EBSD phase distribution analysis to Fig. 1b (acquired at 20 kV with a\n100 nm step size and an indexing rate > 0.95). The results reveal that the\nmicropillars are composed of an FCC matrix (area fraction 67 ± 3%) and a BCC\nphase (33 ± 3%). It should be noted that for phases with characteristic dimensions\n<100 nm (i.e., the BCC phase), EBSD phase identification indeed has limitations,\nand local indexing failures may occur, as discussed in Ren et al., Nature 608, 62–\n68 (2022).\nTo help readers better understand the microstructure of the micropillars, we have\nincluded an enhanced contrast image in Supplementary Fig. S2, processed\nusing adaptive histogram equalization (Oxford AZtecHKL software, CLAHE\nparameters: clip limit = 2). In addition, Fig. 1d shows a TEM image illustrating the\ninternal nano-lamellar structure of the micropillar.\nWe acknowledge your important observation regarding the phase color indexed in\nSupplementary Fig. S1. We apologize for this oversight. The color coding has\nnow been recalibrated according to EBSD and the ICSD standard database. The\nupdated color assignments are as follows:\n FCC phase: Red (RGB 255, 59, 48)\n BCC phase: Green (RGB 48, 209, 88)\nThis color scheme follows the convention described by Goldstein et al. in Scanning\nElectron Microscopy and X-Ray Microanalysis (4th Edition).\nOnce again, we greatly appreciate your thorough and constructive feedback.\nRevisions:\n8 / 61\nFig. 1b. Laser powder bed fusion (LPBF) fabrication of MPEAs and corresponding inverse\npole figure (IPF) and phase distribution.\nFig. S1. Microstructure of the initial eutectic MPEA characterized using EBSD. (a)\nand (b) are the antipodal plots (IPF, i), image quality maps (IQ, ii), and phase maps (iii) of\nthe block samples relative to the overall and localized positions.\n3. How is the theoretical strength limit of 3.5 GPa was determined? Is it possible that\nthe flow stress exceeds the theoretical limit? It is usually determined as G/10 where\n9 / 61\nG is shear modulus. I doubt that the shear modulus of the eutectic HEA would be\nResponse: Thank you for asking this important question regarding the mechanical\ntheoretical limit.\nThe reference to the ~3.5 GPa “theoretical limit” in the original text specifically\npertains to the upper bound of engineering strength under HR loading conditions\n(strain rate ~ 2×10⁻² s⁻¹). We agree with your comment that the original wording\nmay lead to potential misunderstanding. In response, we have revised the term to\n“strain rate limit.” Following your insightful suggestion, and to eliminate any\nambiguity for readers, we have evaluated the theoretical strength limit σ of the\nLimit\ndual-phase CoCrAlFeNi eutectic MPEA based on the rule of mixtures.\nAccording to the rule of mixtures (Eqs. 1-3):\n𝜎 = 𝑓 𝜎 +𝑓 𝜎 (1)\n𝐿𝑖𝑚𝑖𝑡 𝐹𝐶𝐶 𝐹𝐶𝐶 𝐵𝐶𝐶 𝐵𝐶𝐶\nHere, the phase volume fractions f ≈ 67±3% and f ≈ 33±3% were determined\nFCC BCC\nfrom EBSD analysis. The theoretical strengths (σ and σ ) of the FCC and\nFCC BCC\nBCC phases were estimated based on their respective shear moduli derived from\n3D-printed eutectic CoCrAlFeNi alloys 1, 2:\n𝜎 = 𝐺 ≈ 8. 𝐺𝑃𝑎, (𝐺 ≈ 8 GPa) (2)\n𝐹𝐶𝐶 𝐹𝐶𝐶 𝐹𝐶𝐶\n𝜎 = G ≈ 7.5 GPa, (𝐺 ≈ 57 GPa) (3)\n𝐵𝐶𝐶 𝐵𝐶𝐶 𝐵𝐶𝐶\nAccordingly, the theoretical strength of the alloy is 𝜎 ≈ 7.3 8±\n𝐿𝑖𝑚𝑖𝑡\n.24 GPa (error transfer included).\nAlthough the strain rate enhancement effect has been fully exploited, it is difficult\nto break through the fundamental limitations of the dislocation motion pattern to\napproach the theoretical limit at the current strain rate threshold (2 × 10-² s-¹). This\nraises a critical question: Can the strength of the alloy be further enhanced without\nexceeding the imposed strain rate threshold?\nTo address this, we implemented a strain-training strategy that allows strength\n10 / 61\nenhancement without introducing additional strain rate hardening effects. We\nachieved a compressive strength of ~6.5 GPa—approaching the theoretical limit.\nRevisions:\nIn response to this comment, we have made the following revisions:\n The label “theoretical limit” in Fig. 1h and Page 4 (Lines 118 - 123) has been\nupdated to reflect the value σ ≈ 7.308 GPa.\nLimit\n The original term “theoretical limit” has been replaced with “strain rate limit”\nto more accurately describe the imposed mechanical constraints.\n The full derivation of the theoretical strength calculation has been included in\nSupplementary Note 2.\nFig. 1h Engineering stress-strain curves, demonstrating pronounced strain hardening and\nsuperior damage tolerance in ST-loaded micropillars. Post-deformation structural analysis\nreveals a gradient deformation profile: Top region—severe plasticity and dislocation-driven\nfragmentation, Middle region—stress concentration leading to FCC/BCC phase bending\nand partial fracture, and Bottom region—minimal deformation and primary stress support.\n11 / 61\nDislocation evolution in ST samples enables progressive hardening, validating the top-\ndown strain transmission design. The red dashed lines represent the strain rate limit and\nthe theoretical strength limit.\nSupplementary Note 2: “To evaluate the theoretical strength limit σ of the dual-\nLimit\nphase CoCrAlFeNi system eutectic MPEA, we perform a more rigorous evaluation\nbased on the rule of mixtures (Eqs. 1-3):\n𝜎 = 𝑓 𝜎 +𝑓 𝜎 (1)\n𝐿𝑖𝑚𝑖𝑡 𝐹𝐶𝐶 𝐹𝐶𝐶 𝐵𝐶𝐶 𝐵𝐶𝐶\nHere, the phase volume fractions f ≈ 67±3% and f ≈ 33±3% were determined\nFCC BCC\nfrom EBSD analysis. The theoretical strengths (σ and σ ) of the FCC and\nFCC BCC\nBCC phases were estimated based on their respective shear moduli derived from\n3D-printed eutectic CoCrAlFeNi alloys 1, 2:\n𝜎 = 𝐺 ≈ 8. 𝐺𝑃𝑎, (𝐺 ≈ 8 GPa) (2)\n𝐹𝐶𝐶 𝐹𝐶𝐶 𝐹𝐶𝐶\n𝜎 = G ≈ 7.5 GPa, (𝐺 ≈ 57 GPa) (3)\n𝐵𝐶𝐶 𝐵𝐶𝐶 𝐵𝐶𝐶\nAccordingly, the theoretical strength of the alloy is 𝜎 ≈ 7.3 8±\n𝐿𝑖𝑚𝑖𝑡\n.24 GPa (error transfer included).\nAlthough the strain rate enhancement effect has been fully exploited, it is difficult\nto break through the fundamental limitations of the dislocation motion pattern to\napproach the theoretical limit at the current strain rate threshold (2 × 10-² s-¹). This\nraises a critical question: Can the strength of the alloy be further enhanced without\nexceeding the imposed strain rate threshold?\nTo address this, we implemented a strain-training strategy that allows strength\nenhancement without introducing additional strain rate hardening effects. We\nachieved a compressive strength of ~6.5 GPa—approaching the theoretical limit.”\nReferences\n1. Ren, J. et al. Strong yet ductile nanolamellar high-entropy alloys by additive\nmanufacturing. Nature. 608, 62-68 (2022).\n2. Xiong, T. et al. High-strength and high-ductility AlCoCrFeNi eutectic high-\n2.1\nentropy alloy achieved via precipitation strengthening in a heterogeneous\nstructure. Scripta Mater. 186, 336-340 (2020).\n4. The Ashby map is missing some of the most important achievements in microscale\nmechanical properties [Nat. Commun. 13 (2022) 1102; Nat. Commun. 6 (2015)\n7748], which actually outperforms the current material.\nResponse: Thank you for suggesting a comparison with these two significant\nstudies, which have made our investigation more comprehensive.\nWe have carefully reviewed and compared the focus of these two works. All\nstudies share the core objective of overcoming the strength-ductility trade-off.\nLiu et al. 1 achieved strengthening in BCC alloy micropillars through extensive\ninterstitial solid solution hardening (grain boundary segregation of C/N and\ninterstitial O atoms). Zou et al. 2 demonstrated that strong texture in BCC alloy\nnanopillars inhibits grain boundary migration, leading to high strength and ductility.\nBoth studies proposed innovative approaches to enhance the strength of BCC\nalloys.\nIn contrast, our study explores strain engineering in BCC/FCC alloys to regulate\namorphization, replacing shear banding to achieve high strength and ductility.\nWe propose a mechanical chain (dislocation gradient → topological defects →\namorphization expansion) which shifts the perspective from “suppressing\namorphization” to “leveraging amorphization,” transforming it from a passive\ndefect into an active energy dissipation mechanism. While Liu et al. 1 and Zou et\nal. 2 focused on interstitial solid solution strengthening and nanoscale HEA\nmicropillars design, respectively, and successfully broke the performance limits,\nthey did not cover this area.\n1. Liu et al. Nat. Commun. 13 (2022) 1102. [Reference 33 in the main text]\n2. Zou et al. Nat. Commun. 6 (2015) 7748. [Reference 34 in the main text]\nRevisions:\nThe revised results have been incorporated into Fig. 1i on Page 3, as detailed\n13 / 61\nbelow.\nFig. 1i. Ashby plot of microscale mechanical performance, highlighting that ST-loaded\nmicropillars outperform most metals, alloys, and ceramics in strength-ductility synergy.\n1. Liu et al. Nat. Commun. 13 (2022) 1102. [Reference 33 in the main text]\n2. Zou et al. Nat. Commun. 6 (2015) 7748. [Reference 34 in the main text]\nWe sincerely thank you once again for your thoughtful and constructive\nfeedback, which helped us substantially improve the critical details of our\nstrain-training strategy. We hope that our revisions have fully addressed\nyour concerns and clarified any remaining ambiguities.\n14 / 61\nReferee #2\nThe authors successfully induced micrometer-scale amorphization in multi-\nprincipal element alloys through a novel strain-rate-controlled compression\ntechnique. They induced a high-density dislocation gradient that triggered a broad\namorphous region. This process replaced traditional shear-dominated failure with\na more stable, energy-dissipating mechanism. The resulting alloy demonstrated\nenhanced strength and plasticity. The results are interesting and have a scientific\nimpact. Below are some questions and comments for this article.\nResponse: We sincerely appreciate your time and effort in reviewing our\nmanuscript and for your positive evaluation. We have provided detailed point-by-\npoint responses to the issues you raised. Please feel free to contact us if you have\nany further comments.\n1. The authors stated the choice of 2x10^-2 s^-1 as HR is due to the limits of the\nvideo capture and motor control systems, is there any specific reason to choose\nthe use of 8x10^-4 s^-1 and 3x10^-3 s^-1 as SR and MR? Also, in ST, how is the\ntransition between strain rates (from 8x10^-4 s^-1 to 3x10^-3 s^-1 to 2x10^-2 s^-\n1) managed in real-time?\nResponse: Thank you for your insightful comments. In our previous version, the\ntransition of strain rates was not explicitly clarified. We have now supplemented\nthis section in Methods (In situ mechanical testing) and Supplementary Note\n6.\nRevisions:\nIn situ mechanical testing (Page. 15, Lines 404 - 463): Micropillar compression\ntests were conducted using a high-resolution field emission scanning electron\nmicroscope, with the strain rates carefully controlled. Compression was performed\nat constant strain rates of 8×10⁻⁴ s⁻¹ (SR), 3×10⁻³ s⁻¹ (MR), and 2×10⁻² s⁻¹ (HR),\nas well as under varying strain rates (Increase from SR to HR, referred to as strain-\ntraining compression, ST).\nSelection of SR\n15 / 61\nA strain rate of 8×10⁻⁴ s⁻¹ corresponds to the conventional quasi-static\ncompression rate (~10⁻⁴ /s level), facilitating the exploration of the material's\nintrinsic compressive strength and plasticity.\nSelection of MR\nThe strain rate of 3×10⁻³ s⁻¹ follows the \"order-of-magnitude stepwise\" principle.\nThis controlled rate increase (<10x) prevents mismatches between video capture\nand load application systems while simultaneously introducing strain rate\nhardening effects.\nHardware constraints on HR\nThe 2×10⁻² s⁻¹ strain rate is chosen as the upper limit due to experimental system\nconstraints. The video acquisition system requires a sufficiently high frame rate to\ncapture the micropillar deformation process accurately, while the motor control\nsystem must maintain load stability with nanometer precision. These two factors\ncollectively determine the strain rate ceiling. Exceeding this threshold risks a\nmismatch between video capture and mechanical loading systems. In situ\nmechanical testing in small-scale environments may amplify the effects of strain\nrate hardening 61. At this scale, a strain rate of 2 × 10⁻² s⁻¹ corresponds to a\ndisplacement speed of ~120 nm/s. In the work of Kim et al 62, a strain rate of 2 ×\n10⁻² s⁻¹ at the macroscale was referred to as a high strain rate.\nStrain transition management of ST\nPrior to the implementation of ST, we carefully designed the strain rate transition\nprotocol based on comprehensive experiments with single-strain-rate samples.\nThree key criteria were established: (1) The strain rate transition point must reach\nthe yield strain of the micropillar but remain lower than the strain threshold where\nshear localization occurs. (2) Sufficient dislocation density must be accumulated\nto reach the critical level required for triggering amorphization. (3) Reproducibility\nmust be confirmed through repeated tests (three experiments, considering\nrepeatability and experimental cost). The initial SR stage promotes uniform\ndislocation nucleation, the subsequent MR stage enhances dislocation gradient\nformation and interactions, and the final HR stage effectively triggers\n16 / 61\namorphization.\nMicropillars loaded at slow strain rates are enriched with a high density of\ndislocations and deformation defects at the top. The low activation energy is not\nsufficient to support long-range migration of dislocations, leading to early localized\nstress concentrations and shear localization. To avoid premature shear or\nprecursors in the SR stage of ST samples, compression was first loaded at a slow\nstrain rate to the yield stage. Due to the small deviation of grain size and phase\nfraction within micropillars, the yield strain of the SR sample was ~5.5±1.5%,\nresulting in a strain transition point of ~6% for the ST sample. Immediately\nafterward, the strain rate was increased to an MR to enhance uniform dislocation\nproliferation. This stage is needed to ensure sufficient accumulation of dislocations\nto establish a topological defect foundation for subsequent amorphization. The\nselection process of the strain transition point is described in detail in\nSupplementary Note 6. Consideration of the positive strain rate hardening effect\nwill allow this stage to sustain longer compressive displacements. Furthermore,\namorphization requires enhanced atomic migration to rapidly disrupt the crystalline\norder. Such processes must be accomplished within a short period. Consequently,\nonce the strain approaches the uniform deformation limit identified in the SR–MR\nexperiments (Supplementary Note 6), the system immediately elevates the strain\nrate to the hardware upper limit (HR level), enabling rapid compression until\nabnormal deformation occurs between the substrate and the micropillar.\nSystem control was implemented using software-driven displacement coordinated\nwith the indenter, ensuring smooth transitions between strain rates and a strain\nsensitivity of ~0.1%. Upon surpassing a predetermined threshold, the system\nautomatically triggered a strain rate transition to the next stage. Due to the inherent\ninertia of the experimental apparatus and the acceleration protocol, a slight delay\nin the actual transition point may occur during the acceleration phase (affected by\nthe magnitude of the strain rate jump).\nSupplementary Note 6. Strain rate transition design based on dislocation\ndynamics and the critical conditions for amorphization\n1. Dislocation dynamics under strain rate dependence\n17 / 61\nIn the initial stage of SR loading (4–5.5% strain), the micropillar undergoes yield\ndeformation, characterized by uniform activation of dislocations along a single slip\nsystem (Fig. 1d), without evident localization. Between 6–7% strain, solid-state\nflow dominates with further homogeneous dislocation proliferation. However, in\nthe 9–12% strain range, irreversible signs of shear band initiation begin to emerge\n(Supplementary Fig. S7), allowing us to identify the critical onset of localization.\nTo avoid invisible precursors of shear inside the material, we selected ~6% strain\n(close to yield) as the transition point to the MR stage. The control system can\ndetect strain with ~0.1% precision (Page. 15, Methods for system description).\nWhile the first transition point is defined by the need to avoid shear, the second\ntransition is critical for controlling the extent of amorphization. Based on Orowan\nequation 1, dislocation-mediated deformation follows:\n= 𝜌𝑏𝑣 (7)\nwhere ρ is the dislocation density, b is the Burgers vector, and v is the dislocation\nvelocity. Under low-to-intermediate temperatures, v follows thermally activated\nbehavior (Arrhenius-type 2), but to highlight the strain rate dependence, we adopt\na power-law model:\n𝑣 = 𝑣𝑒 𝑄− 𝑘 𝜏 𝑇 𝑏𝑙 , 𝑣 ∝ ( 𝜎 )𝑚 (8)\n𝐺\nwhere m is the strain rate sensitivity exponent and G is the shear modulus. The\nstrength contribution from dislocations 3 is given by:\n𝜎\n𝜎 = 𝜎 +𝛼𝐺𝑏√𝜌, ∝ 𝑏√𝜌 (9)\n𝑑 0\n𝐺\nCombining these gives:\n∝ 𝜌𝑏∙(𝑏√𝜌) 𝑚 = 𝑏1+𝑚𝜌1+ 𝑚 2, ∝ 𝜌1+𝑚 (10)\nThis indicates that dislocation density increases almost linearly with strain rate,\nconsistent with our experimental observation of strain rate hardening.\n2. Kinetic design of the second transition point\n18 / 61\nBased on dislocation behavior under constant-rate loading, we implemented the\nfirst strain rate jump near the yield point in SR loading. According to the Kocks–\nMecking model 4:\n𝑑𝜌 𝑘\n(𝜔)\n= 1 √𝜌−𝑘 𝜌 (11)\n𝑑𝜀 𝑏\n(𝜔)\nHere, ε is strain, 𝑘 is the dislocation multiplication coefficient (strain rate\n1\ndependent), and k is the dynamic recovery coefficient (typically considered\n2\nconstant). Let 𝜌( ) = 𝜌 be the initial condition at yield. According to Eq. 10, the\n𝑎 𝑎\ndislocation multiplication efficiency exhibits an approximately linear relationship\nwith the loading rate, indicating that an increased strain rate leads to a faster\naccumulation of dislocation density per unit strain in the micropillar. When the\n(𝜔)\nstrain rate jumps from (SR) to (MR), k increases to 𝑘 = 𝜆𝑘 , where λ is\n1 1 1 1\na strain rate sensitivity constant, and ω is the number of strain rate transitions.\nFor > , the new form becomes:\n𝑎\n(1)\n𝑑𝜌 𝑘\n= 1 √𝜌−𝑘 𝜌 𝑓𝑜𝑟 > (12)\n𝑎\n𝑑𝜀 𝑏\nThis equation is in a nonlinear separable form. After the first-rate jump, the\ndislocation density in the MR stage evolves as:\n𝜌 𝑏 ( ) = [√𝜌 𝑎 ∙𝑒 𝑘2(𝜀 2 −𝜀𝑎) + 𝑘 1 (1) ( −𝑒 𝑘2(𝜀 𝜀𝑎) )] (13)\n𝑏𝑘2\nTo determine the appropriate second transition point that ensures sufficient\nconditions for amorphization, we designed additional experiments\n(Supplementary Fig. S25) where loading is stopped after completing only the MR\nstage. As shown in the stress–strain curves (Supplementary Fig. S26),\nincreasing the strain rate improves the mechanical response of the micropillar.\nShear banding begins at ~26% strain, and after full crack development (~40%\nstrain), the substrate exhibits significant plastic deformation, evidenced by a\nmarked change in the stress–strain slope.\nConsidering the inertia of the loading system and the potential for shear precursors,\n19 / 61\nthe second transition to the HR stage is conservatively selected at a strain ε ≤\nb\n20%, which safely precedes catastrophic localization.\nSimilarly, the evolution of dislocation density ρ in the HR stage (strain rate )\nc\nafter the second strain rate jump can be analytically expressed as:\n𝜌 𝑐 ( ) = [√𝜌 𝑏 ∙𝑒 𝑘2(𝜀 2 −𝜀𝑏) + 𝑘 1 (2) ( −𝑒 𝑘2(𝜀 𝜀𝑏) )] 𝑓𝑜𝑟 > 𝑏 (14)\n𝑏𝑘2\nIn this expression, the exponential terms are monotonically increasing, indicating\nthat the dislocation density rises nonlinearly and consistently with both strain and\nstrain rate. The linear strain-rate hardening observed under conventional single-\nstage loading (Eq. 10) is restructured into an exponential-like hardening behavior\nby the ST strategy. Therefore, higher dislocation accumulation during each stage\nof ST promotes a greater final dislocation density.\nDislocation density is closely linked to the amorphization process 5, with the critical\ncondition defined as:\n𝛾\n𝜌 = (15)\n𝑛 𝐺𝑏3\nwhere ρ is the critical dislocation density for amorphization and γ\\gamma is the\nn\nstacking fault energy. When ρ ≥ ρ , amorphization is likely to initiate in the\nc n\nmicropillar. To promote a broader and more pronounced amorphous zone, it is\nnecessary to maximize the duration of dislocation accumulation during the uniform\ndeformation stage. Accordingly, we select the second strain rate transition point\nat ε = 20%. The subsequent HR stage then provides the localized stress and\nb\nlattice perturbation required to transition from unstable shear to long-range\namorphization—a “window zone” for this transformation.\nTo quantify the volume fraction of the amorphous phase f as a function of the\nam\ntransition strain ε , we define:\nb\n𝑓 ( ) = max ( ,\n𝜌𝑐 (𝜀𝑏 )\n) (16)\n𝑎𝑚 𝑏\n𝜌𝑛\nSince amorphization is a defect-induced nonlinear transformation, this relationship\n20 / 61\nis not linear 6. When ρ < ρ , the crystalline lattice retains recovery capability, and\nc n\nthe amorphous volume fraction is near zero. As ρ approaches ρ , local atomic\nc n\norder begins to collapse, and small amorphous nuclei emerge—this corresponds\nto the typical nanoscale amorphization. To trigger continuous amorphization,\nfurther dislocation accumulation (ρ ) is required, causing localized nuclei to\ns\ncoalesce into amorphous bands 6. This transition, characterized by a critical\nthreshold, smooth evolution, and eventual saturation, is well described by the\nSigmoid function 7, 8:\n1\n𝑓(𝑥) = (17)\n1+𝑒−𝑥\nThus, the amorphous volume fraction becomes a probabilistic measure of phase\ntransformation per unit region:\n1\n𝑓 ( ) = { −𝛽( 𝜌𝑐(𝜀𝑏) ), 𝜌𝑐≥𝜌𝑛 (18)\n𝑎𝑚 𝑏 1+𝑒 𝜌𝑛 𝜌𝑐<𝜌𝑛\n,\nHere, β is a steepness parameter in the Sigmoid function that reflects the rate and\nsensitivity of disorder propagation. In metallic systems, amorphization typically\ninitiates with local atomic rearrangement and grows progressively. When the\ndislocation density approaches the critical value, long-range atomic order is lost.\nA large β implies that small increases in ρ can rapidly trigger a cascade of\nc\namorphous nucleation events, whereas a smaller β indicates a more gradual\ntransformation that requires sustained dislocation buildup. Thus, β is a key\ndescriptor of the amorphous nucleation and propagation kinetics. This kinetic\nframework is similar to Li et al. 's theories 6.\nAdditionally, we performed in situ compression experiments in which the\nmicropillar was directly transitioned from SR to HR loading (skipping the MR stage)\nto increase the strain rate jump amplitude (Supplementary Figs. S27). Due to the\nlarge transition span, the control system required ~10 s to stabilize the strain rate\n(as shown by the proportion-integration-differentiation controller).\nHowever, the low dislocation density developed during the SR yield stage is\ninsufficient to form a dense defect network capable of driving rapid lattice\n21 / 61\nreconstruction in the HR phase. As a result, minor discrete shear steps were\nobserved at ~10% strain, followed by shear slip activation in both the middle (~40%\nheight) and base of the micropillar. According to the amorphization kinetics, the\ndislocation storage capacity of the direct jump loading path returns to the constant-\nrate HR loading mode. These results highlight the necessity of the intermediate\nMR stage in the ST strategy to ensure progressive regulation and effective defect\nevolution.\nRelated Figures:\nFig. S25. Deformation process of the micropillar under SR–MR limiting compression.\n(a) The first strain rate jump was applied at ~6% strain. (b) During the MR stage, shear\nslip steps began to appear at ~26% strain. (c, d) With continued loading, these slip steps\nprogressively evolved into mature shear cracks, ultimately compromising the structural\nintegrity of the micropillar.\nFig. S26. Stress-strain curve of the micropillar loaded by SR-MR.\n22 / 61\nFig. S27. Failure process of a micropillar subjected to a direct strain rate jump from\nSR to HR. (a) The micropillar underwent a direct transition from SR to HR loading without\nan MR stage. (b, c) The absence of progressive dislocation accumulation leads to\npremature shear activation and the formation of discrete slip steps, ultimately resulting in\nstructural failure.\nReferences\n1. Orowan, E. Problems of plastic gliding. Proc. Phys. Soc. 52, 8 (1940).\n2. Laidler, K. J. The development of the Arrhenius equation. J. Chem. Educ. 61,\n494 (1984).\n3. Courtney, T. H. Mechanical Behavior of Materials. (Waveland Press, 2005).\n4. Kocks, U. & Mecking, H. Physics and phenomenology of strain hardening: the\nFCC case. Prog. Mater. Sci. 48, 171–273 (2003).\n5. Zhao, S. et al. Amorphization in extreme deformation of the CrMnFeCoNi\nhigh-entropy alloy. Sci. Adv. 7, eabb3108 (2021).\n6. Li, B. Y., Li, A. C., Zhao, S. & Meyers, M. A. Amorphization by mechanical\ndeformation. Mater. Sci. Eng. R Rep. 149, 100673 (2022).\n7. Kyurkchiev, N. & Markov, S. Sigmoid functions: some approximation and\nmodelling aspects. LAP LAMBERT Acad. Publ. Saarbrucken 4, 34 (2015).\n8. Delogu, F., Schiffini, L. & Cocco, G. The invariant laws of the amorphization\nprocesses by mechanical alloying: I. Experimental findings. Philos. Mag. A 81,\n1917–1937 (2001).\n2. What is the reason for choosing a kinetic cutoff energy of 350 eV that is 1.3 times\ngreater than the default value, and how sensitive are the results to this parameter?\n23 / 61\nResponse: We thank the reviewer for their thoughtful question regarding the\nplane-wave cutoff energy selection in our density functional theory (DFT)\ncalculations. The cutoff energy critically balances computational efficiency and\naccuracy in plane-wave basis DFT, as it determines the resolution of electronic\nwavefunction expansion. While higher cutoffs improve the precision of total\nenergies, stresses, and lattice parameters, they also significantly increase\ncomputational cost.\nFor the Al₁₇Fe₁₆Co₁₆Cr₁₆Ni₃₄ system studied here, the pseudopotential-derived\ndefault cutoff energy is 269.532 eV (Ni). Following VASP’s documented guidelines\n(see VASP Wiki: Pulay stress), we adopted a cutoff of 350 eV (1.3 times the\ndefault value) to ensure robust stress-tensor convergence—a standard practice to\nmitigate Pulay-stress artifacts during strain-dependent simulations.\nTo explicitly validate our choice, we performed comparative calculations at both\n270 eV and 350 eV cutoffs. As detailed in Supplementary Note 9, the total\nenergies and stresses exhibit excellent convergence till the peak compressive\nstress, confirming that 350 eV achieves numerically stable results without undue\ncomputational overhead. This approach aligns with our prior work on high-entropy\nalloys and established literature benchmarks.\nRevisions:\nSupplementary Note 9. Selection of the plane wave cutoff energy in density\nfunctional theory calculations.\nTo investigate the structural and mechanical properties of Al₁₇Fe₁₆Co₁₆Cr₁₆Ni₃₄,\nwe performed density functional theory (DFT) calculations using the Vienna ab\ninitio Simulation Package (VASP). The exchange-correlation interactions were\ntreated within the generalized gradient approximation (GGA) using the Perdew-\nBurke-Ernzerhof (PBE) functional. The maximum default plane-wave cutoff\nenergy derived from the pseudopotential files was 269.532 eV (Ni). In line with\nVASP recommendations, we adopted an energy cutoff of 350 eV (1.3 times the\ndefault value) to ensure accurate stress tensor calculations, particularly for\nresolving interfacial electronic effects and local bonding evolution under strain.\nThis choice was validated in our previous work 1 on CoCrAlFeNi shear simulations\n24 / 61\nand is consistent with established high-entropy alloy studies.\nTo assess the sensitivity of energy, stress, and atomic configurations to the cutoff\nenergy, we conducted comparative DFT simulations at 270 eV and 350 eV during\nstrain training. As shown in Supplementary Figs. S35, S36, the stress–strain and\nenergy–strain responses remain highly consistent between the two cutoff values.\nThis agreement confirms that 270 eV provides physically meaningful results, while\n350 eV ensures well-converged accuracy for detailed mechanical and electronic\nstructure analysis.\nRelated Figures:\nFig. S35. Sensitivity of DFT mechanical response to the energy cutoff. (a) Results\nusing a cutoff energy of 270 eV, and (b) using 350 eV. Subplots: (i) stress–strain response,\n(ii) total energy–strain behavior of the alloy system.\nFig. S36. Effect of increasing the energy cutoff from the default 270 eV to 350 eV on\n25 / 61\nthe DFT mechanical response. Compared with the 270 eV results shown in (a–c),\nsimulations using a cutoff of 350 eV (d–f) exhibit a slightly higher density of pentagon and\nheptagon defects before 15% strain, though the overall distribution and locations remain\nsimilar. At 20% strain, the defect spreads and penetrates more deeply into the matrix, more\nclosely resembling experimental observations.\nReferences:\n1. Du, J. et al. Mitigating Strain Localization via Stabilized Phase Boundaries\nfor Strengthening Multi‐Principal Element Alloys. Adv. Sci. 2414783 (2025).\n3. For the ST-loaded micropillars, the amorphous band terminates at ~1.35 μm from\nthe base. What determines this termination point? Is it linked primarily to the\ndislocation gradient itself or other microstructural constraints?\nResponse: Thank you for your insightful comments on the termination of\namorphization. We demonstrated through DFT calculations that amorphization is\nclosely linked to crystalline defects (Fig. 3). In the previous version, we established\nthe critical threshold for amorphization based on the physical relationship between\ncritical dislocation density and stacking fault energy, as proposed by Zhao et al\n(Page. 12). When the dislocation density exceeds this threshold, amorphization\npropagates; otherwise, the matrix remains stable. This indicates that the\namorphization process is governed by the dislocation gradient, and the position of\nthe stop can be directly measured and estimated, which is consistent with the\nmeasurement results of the SEM image (Fig. 1h).\nBeyond the influence of external loading, this gradient is also enhanced by the\nstable phase boundaries in eutectic MPEAs (this is answered in detail in\nComment 5), which constrain dislocations, preventing cross-phase entanglement\nor surface escape within the micropillar. As a result, dislocations accumulate more\neffectively from top to bottom, facilitating the development of high-density\nmicrostructures.\nThe kinetics of ST-stimulated amorphization hyper-range diffusion are shown in\nSupplementary Note 6.\n26 / 61\nWe have incorporated this response into Pages 5, 10, and 12 of the manuscript,\nhighlighted in blue for clarity.\n4. Since lattice recovery is mentioned at crack tips, is there evidence that hyper-\nrange amorphization can stabilize microcrack formation in earlier deformation\nstages?\nResponse: We sincerely thank you for your insightful comments. The crack tip is\na region of extremely high stress concentration, where local stress levels\nsignificantly exceed those in the surrounding matrix 1. This extreme stress state\ninduces intense lattice deformation along specific slip systems, leading to a rapid\nincrease in dislocation density to the critical threshold for amorphization. As a\nresult, dislocation entanglement and pile-up hinder long-range dislocation glide.\nConcurrently, high shear stress promotes the formation of cavity-like structures\nnear free surfaces (i.e., initial cracks), facilitating the cooperative shear\ndisplacement of atomic clusters and local accumulation of free-volume atoms in\nthe lattice (Supplementary Fig. S15). The formation of amorphous nuclei within\nthe amorphization transition/affected zone further supports this phenomenon (Fig.\n2b , c ).\ni, ii i-iv\nOnce the free-volume concentration surpasses the critical threshold, the long-\nrange lattice order collapses, forcing the alloy to dissipate energy through\namorphization 2. Wang et al. 3 directly demonstrated that in multicomponent alloy\nsystems, the resistance of the lattice near the crack tip to dislocation slip serves\nas a driving force for amorphization. This amorphous zone disrupts the\ncrystallographic propagation path of cracks, effectively stabilizing initial cracks. For\ninstance, the formation of amorphous bridges has been shown to retard crack\npropagation and enhance material toughness. In our study, the observation that\ninitial shear cavities do not evolve into long-range cracks further substantiates the\nstabilizing effect of amorphization. Moreover, according to Rice’s criterion 4, the\nductility parameter (η) is positively correlated with the extent of amorphization,\nsuggesting that the amorphization capability is intrinsically linked to a material’s\ninherent ductility. Ductile materials (i.e., those with high η values) can leverage\namorphization to stabilize microcracks during the early stages of deformation.\n27 / 61\nThe role of amorphization in suppressing initial crack propagation has been\npreviously reported in bulk metallic glasses 5. For example, pre-existing\namorphous phases have been utilized to optimize fracture resistance, and\namorphous coatings have been deposited on conventional alloys to prevent crack\ngrowth 6. Using molecular dynamics simulations, Huang et al. 7 demonstrated that\namorphization at the crack tip is energetically more favorable than direct fracture.\nThis amorphization can be triggered at the early stages of crack propagation,\ntrapping the crack tip in a series of metastable states rather than allowing it to\nextend freely. Ji et al. 8 further showed that amorphous nuclei originate from multi-\ndirectional dislocation nodes. Our DFT calculations align with these findings (Figs.\n3b-h), confirming that extended-range amorphization retains the fundamental\ncharacteristics of localized amorphization. The ability to stabilize cracks essentially\ndepends on a material’s capacity to absorb and dissipate energy—an area where\namorphous structures have been widely validated 9-12. In our main text (Page 13),\nwe employ the Griffith criterion to demonstrate that hyper-range amorphization\nexhibits a remarkably high energy dissipation capacity. Therefore, tuning the\nextent of amorphization emerges as a promising strategy for enhancing the\nstrength and toughness of alloys.\nThe above content is supplemented in Supplementary Note 3.\nRelated Figures:\nFig. S15. Amorphization prevents the further development of shear cavities. (a)\nMorphology of shear cavities and the amorphous band. (b) Free volume distribution in the\namorphous band is caused by the loss of atomic order.\nReferences:\n28 / 61\n1. Stepanova, L. V. & Belova, O. N. Coefficients of the williams power\nexpansion of the near crack tip stress field in continuum linear elastic\nfracture mechanics at the nanoscale. Theor. Appl. Fract. Mech. 119, 103298\n(2022).\n2. Li, F. C. et al. Amorphous–nanocrystalline alloys: fabrication, properties, and\napplications. Mater. Today Adv. 4, 100027 (2019).\n3. Wang, H. et al. Deformation-induced crystalline-to-amorphous phase\ntransformation in a CrMnFeCoNi high-entropy alloy. Sci. Adv. 7, eabe3105\n(2021).\n4. Rice, J. R. Dislocation nucleation from a crack tip: An analysis based on the\nPeierls concept. J. Mech. Phys. Solids 40, 239–271 (1992).\n5. Pekin, T. C. et al. Direct measurement of nanostructural change during in\nsitu deformation of a bulk metallic glass. Nat. Commun. 10, 2445 (2019).\n6. Jiang, H. et al. Structures and Functional Properties of Amorphous Alloys.\nSmall Struct. 2, 2000057 (2021).\n7. Huang, S., Zhang, S., Belytschko, T., Terdalkar, S. S. & Zhu, T. Mechanics of\nnanocrack: Fracture, dislocation emission, and amorphization. J. Mech.\nPhys. Solids 57, 840–850 (2009).\n8. Ji, W. & Wu, M. S. Nanoscale origin of the crystalline-to-amorphous phase\ntransformation and damage tolerance of Cantor alloys at cryogenic\ntemperatures. Acta Mater. 226, 117639 (2022).\n9. Zhang, C. et al. Laminar Fe-based amorphous composite coatings with\nenhanced bonding strength and impact resistance. Acta Mater. 72, 239 (2014).\n10. Wang, W. et al. Toughening Fe-based Amorphous Coatings by\nReinforcement of Amorphous Carbon. Sci. Rep. 7, 4084 (2017).\n11. Wang, J. et al. Designing Advanced Amorphous/Nanocrystalline Alloys by\nControlling the Energy State. Adv. Mater. 36, 2311406 (2024).\n12. Zhao, S. et al. Amorphization in extreme deformation of the CrMnFeCoNi\nhigh-entropy alloy. Sci. Adv. 7, 3108 (2021).\n5. Could you elaborate on how the top-down dislocation gradient specifically\ninteracts with phase boundaries to enhance stress redistribution?\nResponse: Thank you for your insightful question. This issue pertains to a central\naspect of our work. As discussed in both the Results section and Supplementary\nNote 1, the uniquely oriented eutectic microstructure serves as the foundational\ndesign premise for both the gradient dislocation architecture and the subsequent\namorphization process. This orientation facilitates top-down stress redistribution\nand dissipation.\nWe conducted comparative in situ SEM compression and TEM analyses on\nmicropillars with different eutectic orientations, including vertical, inclined, and\nhorizontal configurations. The results demonstrate that only when the eutectic\nlamellae are aligned parallel to the loading axis does a gradient dislocation\nstructure form along the loading direction, subsequently triggering large-scale\namorphization. In contrast, pillars with inclined or stacked lamellar orientations\ntend to undergo interfacial shear sliding or premature shear distortion along the\nphase direction under lower applied stresses.\nTo comprehensively address your concerns, we have provided additional\ninvestigations and discussion (Supplementary Note 4).\nRevisions:\nSupplementary Note 4. Eutectic phase boundaries — dislocation gradients\nsynergistically regulate stress distribution.\n As a limiting boundary for dislocation migration.\nTypically, within nanostructured materials, the limited volume suppresses\ndislocation diffusion and interaction, thereby hindering the continuous nucleation\nof new dislocations and the sustained activation of deformation mechanisms.\nUnder external loading, deformation is often accommodated by grain boundary\nsliding or separation to relieve stress concentration between adjacent fine grains.\nDislocation-mediated intragranular plasticity becomes more difficult, which leads\nto localized stress accumulation at the micro- or nanoscale. This geometric\nconfinement sacrifices local damage resistance and mechanical stability at small\nscales, thus requiring the material to offer a \"relief outlet\" for dislocation\nproliferation.\n3D-printed CoCrAlFeNi system alloys, characterized by the combination of soft\nand hard phases, feature nearly straight phase boundaries that enable directional\n30 / 61\nmigration or proliferation of dislocations 1. In our previous work 2, we demonstrated\nthrough in-situ EBSD, quasi-in-situ TEM, and molecular dynamics that these\nphase boundaries mediate plasticity. Dislocations propagate within the matrix via\npartial atomic plane migration, while the phase boundaries, with a slight mismatch\n(lattice mismatch ∆a ≈ 2%), do not initiate slip preferentially 2.\nAs shown in Supplementary Fig. S16, the phase boundaries in ST micropillars\nrestrict lateral dislocation motion via elastic stress fields, forcing dislocations to\npreferentially glide along the boundary normal direction (i.e., loading direction).\nThis mechanism prevents local stress buildup caused by dislocation pile-up,\ncommonly observed in traditional nanocrystalline materials, and instead provides\na low-energy migration path that facilitates stress redistribution. When the phase\nboundaries in the micropillar are oriented perpendicular to the loading direction,\npremature stress concentration tends to occur, resulting in reduced effectiveness\nin stress regulation (Fig.1a and Supplementary Fig. S5).\nUnder the constraint of phase boundaries, low initial strain rates promote\nhomogeneous dislocation nucleation, intermediate rates enhance dislocation\nmobility and interaction, and high strain rates eventually trigger the critical\ndislocation density required for amorphization. This results in non-local plastic\nstrain being distributed over a larger volume, achieving exceptional ductility at high\nstrength.\nWhen the eutectic lamellae are aligned perpendicular to the loading direction\n(Supplementary Fig. S5), the space available for dislocation accumulation along\nthe loading axis is significantly restricted, and the vertical phase boundaries\nstrongly impede dislocation motion. As a result, dislocations tend to pile up near\nthe interface close to the loading surface, leading to localized stress concentration,\nwhich in turn promotes early nucleation of shear cracks and undermines the\noverall mechanical performance.\nIn contrast, when the eutectic structure is inclined with respect to the loading axis\n(Supplementary Fig. S17), although certain deformation pathways are\nsuppressed, slip systems within the FCC phase can still be activated along the\ninclined direction, resulting in off-axis shear. This non-cooperative deformation\n31 / 61\nbehavior leads to premature onset of shear, hindering the development of stress\ngradient transfer and energy dissipation. Consequently, it exacerbates strain\nlocalization and limits the improvement of mechanical stability.\n As heterointerfaces to synergistically regulate stress.\nThe FCC/BCC architecture forms a typical lamellar heterogeneous structure. The\nhigh-modulus BCC phase primarily bears the compressive load, while the softer\nFCC matrix dissipates shear stress through dislocation glide 1,2. Under external\nloading, sequential deformation of the soft FCC and hard BCC phases causes\nstrain incompatibility at the interface, which emits geometrically necessary\ndislocations into the interior of the FCC, strengthening the soft phase via long-\nrange internal stress 3. We observed dislocation proliferation at the FCC phase\nboundaries (Supplementary Fig. S18), where long-range stress fields combine\nwith a dislocation density gradient 3. Meanwhile, the dynamic rearrangement of\natoms relieves local strain energy, guiding amorphization and forming a \"stress\nfunnel.\" Simultaneously, the stiff BCC phase maintains structural integrity.\nRelated Figures:\nFig. S5. TEM morphology of the stacked structure micropillars after ~5%\ndeformation. (a) The overall morphology shows significant deformation and shear bands\nat the top, and the micropillars tilt as a whole under load. (b) A local enlargement of the\ntop.\n32 / 61\nFig. S16. Dislocation–phase boundary interactions in the ST micropillar. (a–c)\nBeyond the fractured regions, dislocations are sequentially confined by phase boundaries\nfrom top to bottom. (d–f) Geometric phase analysis reveals the constraint effect of phase\nboundaries on lattice distortion along the shear direction ε .\nxy\nFig. S17. Compressive failure behavior of inclined-oriented eutectic micropillars\nunder quasi-static SR loading. (a) Initial morphology of the micropillar. (b) At ~6% strain,\nevident shear slip is observed. (c) At ~15% strain, large-scale shear failure occurs along\nthe crystallographic orientation. (d) TEM characterization after failure reveals severe strain\nlocalization at the top region, while the middle and bottom segments exhibit limited\nparticipation in uniform deformation, indicating poor load transfer capability and\npronounced deformation localization.\n33 / 61\nFig. S18. TEM observations showing dislocation accumulation at FCC phase\nboundaries. (a) Bright-field and (b) dark-field.\nReferences\n1. Ren, J. et al. Strong yet ductile nanolamellar high-entropy alloys by additive\nmanufacturing. Nature. 608, 62-68 (2022).\n2. Du, J. et al. Mitigating Strain Localization via Stabilized Phase Boundaries for\nStrengthening Multi‐Principal Element Alloys. Adv. Sci. 2414783 (2025).\n3. Guo, Y. et al. Microstructural origins of enhanced work hardening and ductility\nin laser powder-bed fusion 3D-printed AlCoCrFeNi2.1 eutectic high-entropy\nalloys. Int. J. Plasticity. 179, 104050 (2024).\nMinor comments:\n1. The caption of Figure 1: description for subfigure (c) is marked as (f). (g). Also, the\ndescription for (c, d) needs to be checked.\nResponse: Thanks for the tip, we found the issue and made the following\ncorrections.\nRevision:\nFig.1 (c) Morphology of the focused ion beam fabricated MPEA micropillar. (d)\nTEM microstructure of the micropillar before and after 5% pre-strain.\n2. Some abbreviations are not properly defined in the paper, such as stacking faults\nwere only marked as SF in Figure 3, FIB was defined later in the paper after its\nfirst few appearances, etc.\n34 / 61\nResponse: Thank you for helping us identify the issue with abbreviations. We\nhave now clarified the notation for SF and corrected the abbreviation for FIB at its\nfirst occurrence in the main text.\nRevision:\n1. Fig. 3 caption (Page 9, Line 232): SF stands for stacking fault.\n2. Page 4 (Line 103): focused ion beam (FIB).\nWe sincerely thank you for devoting your time and effort to help us\nsignificantly improve the manuscript. Your feedback has not only\nstrengthened the quality of our work but also deepened our understanding\nof the subject. We hope that our responses have satisfactorily addressed\nyour concerns.\n35 / 61\nReferee #3\nThe manuscript titled “Hyper-Range Amorphization Unlocks Superior Damage\nTolerance in Alloys” reported a novel strain engineering (strain training) strategy\nto achieve micrometer-scale amorphization in eutectic high-entropy alloy\nmicropillars, leading to a significant enhancement of both strength and plasticity.\nThis approach is interesting, and the experimental results provide important\ninsights into strengthening and toughening mechanisms. However, several\nconcerns require further considerations, especially regarding the universality of\nthe proposed strategy.\nResponse: We extend our sincere gratitude for your positive evaluation and\nconstructive feedback on our study. Your professional insights have motivated us\nto further explore aspects such as the universality of our findings. Below, we\nprovide point-by-point responses to your comments.\n1. This work proposed a strategy of strain-training, overcoming the shear effects and\nachieving a steady deformation. However, this study only focused on eutectic high\nentropy alloy micropillars, which is special. Does this strategy still work in a single-\ncrystal system? The critical influence of the eutectic microstructures (e.g., phase\nboundaries, boundary orientations, boundary coherent, ...) has not been fully\nelucidated.\nResponse: We sincerely appreciate your thoughtful review and especially your\nvaluable suggestions regarding the critical role of the eutectic microstructure and\nthe applicability of the strain training (ST) strategy.\nTo address these points, we conducted comparative in situ SEM compression and\nTEM analyses on micropillars with different eutectic orientations, including vertical,\ninclined, horizontal, and single-phase (non-eutectic) configurations. The results\ndemonstrate that when the eutectic lamellae are aligned parallel to the loading\naxis does a gradient dislocation structure form along the loading direction,\nsubsequently triggering large-scale amorphization. In contrast, micropillars with\ninclined or stacked lamellar orientations are prone to premature interfacial shear\nand localized deformation under lower stress, thereby limiting the effectiveness of\namorphization-induced strengthening.\n36 / 61\nUnlike the eutectic micropillars where amorphous layers blunt cracks and delay\nfailure, single-crystal CoCrNi micropillars (with (111) and (110) orientations)\nwithout phase boundaries exhibit amorphization that inhibits dislocation escape to\nthe surface. This leads to the accumulation of homogeneously distributed\ndislocations and stacking faults, enabling simultaneous strength and ductility.\nAlthough the underlying microstructures differ, both systems benefit from ST-\ninduced load shielding and amorphization effects, confirming the broader\napplicability of the ST strategy for alloy strengthening.\nTo comprehensively address your concerns, we have provided additional\ninvestigations and discussion in two specific sections:\n(1) Eutectic phase boundaries — dislocation gradients synergistically regulate\nstress distribution (Supplementary Note 4).\n(2) ST strategy hardening of CoCrNi alloy (Supplementary Note 8).\nRevisions:\n1. Eutectic phase boundaries — dislocation gradients synergistically\nregulate stress distribution (Supplementary Note 4).\n As a limiting boundary for dislocation migration.\nTypically, within nanostructured materials, the limited volume suppresses\ndislocation diffusion and interaction, thereby hindering the continuous nucleation\nof new dislocations and the sustained activation of deformation mechanisms.\nUnder external loading, deformation is often accommodated by grain boundary\nsliding or separation to relieve stress concentration between adjacent fine grains.\nDislocation-mediated intragranular plasticity becomes more difficult, which leads\nto localized stress accumulation at the micro- or nanoscale. This geometric\nconfinement sacrifices local damage resistance and mechanical stability at small\nscales, thus requiring the material to offer a \"relief outlet\" for dislocation\nproliferation.\n3D-printed CoCrAlFeNi system alloys, characterized by the combination of soft\nand hard phases, feature nearly straight phase boundaries that enable directional\n37 / 61\nmigration or proliferation of dislocations 1. In our previous work 2, we demonstrated\nthrough in-situ EBSD, quasi-in-situ TEM, and molecular dynamics that these\nphase boundaries mediate plasticity. Dislocations propagate within the matrix via\npartial atomic plane migration, while the phase boundaries, with a slight mismatch\n(lattice mismatch ∆a ≈ 2%), do not initiate slip preferentially 2.\nAs shown in Supplementary Fig. S16, the phase boundaries in ST micropillars\nrestrict lateral dislocation motion via elastic stress fields, forcing dislocations to\npreferentially glide along the boundary normal direction (i.e., loading direction).\nThis mechanism prevents local stress buildup caused by dislocation pile-up,\ncommonly observed in traditional nanocrystalline materials, and instead provides\na low-energy migration path that facilitates stress redistribution. When the phase\nboundaries in the micropillar are oriented perpendicular to the loading direction,\npremature stress concentration tends to occur, resulting in reduced effectiveness\nin stress regulation (Fig.1a and Supplementary Fig. S5).\nUnder the constraint of phase boundaries, low initial strain rates promote\nhomogeneous dislocation nucleation, intermediate rates enhance dislocation\nmobility and interaction, and high strain rates eventually trigger the critical\ndislocation density required for amorphization. This results in non-local plastic\nstrain being distributed over a larger volume, achieving exceptional ductility at high\nstrength.\nWhen the eutectic lamellae are aligned perpendicular to the loading direction\n(Supplementary Fig. S5), the space available for dislocation accumulation along\nthe loading axis is significantly restricted, and the vertical phase boundaries\nstrongly impede dislocation motion. As a result, dislocations tend to pile up near\nthe interface close to the loading surface, leading to localized stress concentration,\nwhich in turn promotes early nucleation of shear cracks and undermines the\noverall mechanical performance.\nIn contrast, when the eutectic structure is inclined with respect to the loading axis\n(Supplementary Fig. S17), although certain deformation pathways are\nsuppressed, slip systems within the FCC phase can still be activated along the\ninclined direction, resulting in off-axis shear. This non-cooperative deformation\n38 / 61\nbehavior leads to premature onset of shear, hindering the development of stress\ngradient transfer and energy dissipation. Consequently, it exacerbates strain\nlocalization and limits the improvement of mechanical stability.\n As heterointerfaces to synergistically regulate stress.\nThe FCC/BCC architecture forms a typical lamellar heterogeneous structure. The\nhigh-modulus BCC phase primarily bears the compressive load, while the softer\nFCC matrix dissipates shear stress through dislocation glide 1,2. Under external\nloading, sequential deformation of the soft FCC and hard BCC phases causes\nstrain incompatibility at the interface, which emits geometrically necessary\ndislocations into the interior of the FCC, strengthening the soft phase via long-\nrange internal stress 3. We observed dislocation proliferation at the FCC phase\nboundaries (Supplementary Fig. S18), where long-range stress fields combine\nwith a dislocation density gradient 3. Meanwhile, the dynamic rearrangement of\natoms relieves local strain energy, guiding amorphization and forming a \"stress\nfunnel.\" Simultaneously, the stiff BCC phase maintains structural integrity.\n2. ST strategy hardening of CoCrNi alloy (Supplementary Note 8).\nTo verify the applicability of the ST strategy across different alloy systems, we\nemployed a representative cast face-centered cubic (FCC) single-phase CoCrNi\nalloy and investigated its deformation mechanisms along the (110) and (111)\ncrystallographic orientations—both of which correspond to preferential slip\ndirections (Supplementary Fig. S19). Micropillars with a diameter of 2 μm,\nmatching that of the eutectic samples, were fabricated via precise FIB milling and\ncrystallographic orientations were confirmed using EBSD. These samples were\nsubjected to both ST loading to systematically assess their strengthening\nresponses.\nAccording to Zhang et al. 4, CoCrNi micropillars oriented along (110) and (111)\nexhibit pronounced ductility under SR loading, accompanied by extensive\ndislocation escape toward free surfaces. Post-yield deformation often leads to\nlocalized strain accumulation, which rapidly propagates toward the surface,\nmanifesting as pronounced shear slip bands. This surface-directed dislocation\nescape diminishes internal dislocation storage, thereby compromising load\n39 / 61\ntransfer and accelerating plastic localization.\nUnder the ST loading protocol, deformation in CoCrNi (110) and (111) micropillars\nis governed by a hybrid mechanism involving both crystal slip and amorphization\n(Supplementary Figs. S20, S21). These micropillars exhibit strength/ductility\ncomparable to that of 500 nm nanopillars (Supplementary Figure S22), with the\ncorresponding ST processes detailed in Supplementary Figs. S23, S24. Notably,\nST triggers over-range amorphization (Supplementary Figs. S20, S21), while\nalso inducing stress accumulation at the pillar top. Together, these effects restrict\ndislocation escape and promote internal multiplication and slip, facilitating vertical\nstress redistribution. This process is accompanied by deformation twinning, low-\nangle grain boundary formation, and widespread Lomer–Cottrell (LC) lock\ndevelopment. However, due to the absence of phase boundaries, as present in\neutectic systems, localized slip can still propagate from the non-amorphized\nadjacent regions.\nUnlike eutectic micropillars, where amorphization is stabilized and enhanced by\nphase boundaries, the single-crystal CoCrNi alloy relies on the amorphization-\ninduced barrier to dislocation escape for uniform stress redistribution. This\npromotes the cooperative evolution of dense dislocation networks, stacking faults,\nLC locks, and deformation twins. Despite distinct mechanistic pathways, both\nsystems ultimately achieve enhanced strength–ductility synergy via the ST-\ninduced coupling of shear shielding (or partial shield) and amorphization. These\nfindings underscore that the “strain engineering–dislocation regulation–stress\nreconstruction–deformation control” strategy is not inherently dependent on\neutectic structures, thus demonstrating the generality and cross-system\napplicability of the ST approach.\nRelated Figures:\n40 / 61\nFig. S5. TEM morphology of the stacked structure micropillars after ~5%\ndeformation. (a) The overall morphology shows significant deformation and shear bands\nat the top, and the micropillars tilt as a whole under load. (b) A local enlargement of the\ntop.\nFig. S16. Dislocation–phase boundary interactions in the ST micropillar. (a–c)\nBeyond the fractured regions, dislocations are sequentially confined by phase boundaries\nfrom top to bottom. (d–f) Geometric phase analysis reveals the constraint effect of phase\nboundaries on lattice distortion along the shear direction ε .\nxy\n41 / 61\nFig. S17. Compressive failure behavior of inclined-oriented eutectic micropillars\nunder quasi-static SR loading. (a) Initial morphology of the micropillar. (b) At ~6% strain,\nevident shear slip is observed. (c) At ~15% strain, large-scale shear failure occurs along\nthe crystallographic orientation. (d) TEM characterization after failure reveals severe strain\nlocalization at the top region, while the middle and bottom segments exhibit limited\nparticipation in uniform deformation, indicating poor load transfer capability and\npronounced deformation localization.\nFig. S18. TEM observations showing dislocation accumulation at FCC phase\nboundaries. (a) Bright-field and (b) dark-field.\n42 / 61\nFig. S19. Fabrication of CoCrNi alloy micropillars. Equiatomic CoCrNi alloy was\nprepared by arc melting Co, Cr, and Ni five times to ensure chemical homogeneity. EBSD\nanalysis confirmed a single-phase FCC structure. Micropillars were fabricated by FIB\nmilling from grains oriented along the (a) (111) and (b) (110) direction (EBSD z-axis).\nFig. S20. Microstructural features of CoCrNi (111) single-crystal micropillars after\nstrain-training (ST) loading. (a, i–v) Overall morphology of the deformed pillar. Severe\nshape distortion is observed under transitional loading and extensive plastic flow (i).\nSurface deformation is pronounced near the top and gradually decreases along the loading\ndirection. Amorphization and shear cracks are simultaneously present (ii–iv). Intense\nplastic deformation induces deformation twins (iii, v). (b, i, ii) Amorphous regions hinder\ndislocation escape and tend to evolve into low-angle grain boundaries. Localized stacking\n43 / 61\nfaults and LC locks accumulate around these regions. (c, i, ii) The lattice strain distribution\nat the pillar base is inhomogeneous due to the presence of high-density dislocation defects.\n(d) Degree of structural distortion decreases from top to bottom. (e) High-resolution TEM\nimaging and FFT patterns confirm the presence of over-range amorphization. (f, i–iii) EDS\nanalysis at the amorphous front reveals reduced atomic density, predominantly enriched\nin Cr and Ni, suggesting possible element migration and segregation during the\namorphization process.\nFig. S21. Microstructural features of CoCrNi (110) micropillar after ST loading. (a)\nPost-deformation cross-sectional morphology showing severe distortion, similar to that\nobserved in the (111) micropillar. Both amorphous regions and shear cracks are present.\n(b) Microstructural gradient from top to bottom. Unlike the eutectic micropillar, the bottom\nregion also exhibits a high density of dislocations, while the top is severely distorted and\nthe middle region is enriched with LC locks. (c) Dislocation-mediated inward slip bands\nare emitted from the free surface near the crack tip, accompanied by the formation of\nstacking faults. (d, e) Localized amorphization at the crack tip. (f) High-density dislocation\nnetworks distributed throughout the lower portion of the micropillar matrix.\n44 / 61\nFig. S22. Mechanical response of CoCrNi micropillars under ST loading. The\ncomparative data in this figure are referenced from Ref. 4.\nFig. S23. Failure process of CoCrNi (110) micropillar under ST loading. (a–c) During\nthe strain-rate transition, typical slip and coarsened extension are observed. (c) Once the\nstrain exceeds ~30%, the coarsened deformation zone extends beyond the contact region\nof the indenter, preventing uniform plastic flow.\nFig. S24. Failure process of CoCrNi (111) micropillar under ST loading. (a–c) During\n45 / 61\nthe strain transition, multiple slip systems are activated, resulting in a complex multislip\ndeformation mode. The propagation of slip bands occurs at a relatively high rate.\n(c) When the applied strain exceeds ~39%, the deformation zone extends beyond the indenter\ncontact area, preventing uniform plastic flow.\nReferences\n1. Ren, J. et al. Strong yet ductile nanolamellar high-entropy alloys by additive\nmanufacturing. Nature. 608, 62-68 (2022).\n2. Du, J. et al. Mitigating Strain Localization via Stabilized Phase Boundaries for\nStrengthening Multi‐Principal Element Alloys. Adv. Sci. 2414783 (2025).\n3. Guo, Y. et al. Microstructural origins of enhanced work hardening and ductility\nin laser powder-bed fusion 3D-printed AlCoCrFeNi2.1 eutectic high-entropy\nalloys. Int. J. Plasticity. 179, 104050 (2024).\n4. Zhang, Q. et al. Size effects and plastic deformation mechanisms in single-\ncrystalline CoCrFeNi micro/nanopillars. J. Mech. Phys. Solids. 162, 104853\n(2022).\n2. The strain-training strategy is associated with sudden shift of strain rate at certain\nstrains. Could you provide more clarifications or rationales for selecting the strain\nrate transition point? The effects of different transition timings on amorphization\nextent should be further discussed.\nResponse: We sincerely thank you for your important comments regarding the\nstrain rate transition strategy.\nOur rationale for introducing strain-training (ST) in the loading experiments is to\nregulate the defect evolution trajectory of the alloy through a multistage kinetic\npathway, thereby optimizing its damage tolerance. In the previous version, we\ndiscussed the underlying principles of ST: the SR stage promotes uniform\ndislocation nucleation and the formation of low-density slip bands. The subsequent\nMR stage enhances directional dislocation proliferation and interactions. And the\nfinal HR stage triggers dislocation entanglement and instability, activating lattice\ntopological disorder.\nDensity functional theory (DFT) calculations reveal that this multistage regulation\n46 / 61\nscheme essentially manipulates the evolution of defect-state density via strain rate\npath design, allowing the system’s free energy to approach the amorphization\nbarrier along an energetically favorable transformation trajectory. Therefore, it is\ncrucial to extend the SR and MR stages as much as possible while avoiding any\npremature shear localization or its precursors.\nTo further clarify this, we conducted additional experiments and calculations to\nevaluate how the choice of transition points affects the amorphization dynamics.\nSpecifically, SR–MR tests (excluding the HR stage) help identify the limiting\nposition of the transition point, while SR–HR tests (direct transition without MR)\nare useful for evaluating the effect of strain rate span on the deformation behavior\nof micropillars.\nA detailed discussion of this analysis is provided below and has been added to\nSupplementary Note 6.\nRevisions:\nSupplementary Note 6. Strain rate transition design based on dislocation\ndynamics and the critical conditions for amorphization\n1. Dislocation dynamics under strain rate dependence\nIn the initial stage of SR loading (4–5.5% strain), the micropillar undergoes yield\ndeformation, characterized by uniform activation of dislocations along a single slip\nsystem (Fig. 1d), without evident localization. Between 6–7% strain, solid-state\nflow dominates with further homogeneous dislocation proliferation. However, in\nthe 9–12% strain range, irreversible signs of shear band initiation begin to emerge\n(Supplementary Fig. S7), allowing us to identify the critical onset of localization.\nTo avoid invisible precursors of shear inside the material, we selected ~6% strain\n(close to yield) as the transition point to the MR stage. The control system can\ndetect strain with ~0.1% precision (Page. 15, Methods for system description).\nWhile the first transition point is defined by the need to avoid shear, the second\ntransition is critical for controlling the extent of amorphization. Based on Orowan\nequation 1, dislocation-mediated deformation follows:\n47 / 61\n= 𝜌𝑏𝑣 (7)\nwhere ρ is the dislocation density, b is the Burgers vector, and v is the dislocation\nvelocity. Under low-to-intermediate temperatures, v follows thermally activated\nbehavior (Arrhenius-type 2), but to highlight the strain rate dependence, we adopt\na power-law model:\n𝑣 = 𝑣𝑒 𝑄− 𝑘 𝜏 𝑇 𝑏𝑙 , 𝑣 ∝ ( 𝜎 )𝑚 (8)\n𝐺\nwhere m is the strain rate sensitivity exponent and G is the shear modulus. The\nstrength contribution from dislocations 3 is given by:\n𝜎\n𝜎 = 𝜎 +𝛼𝐺𝑏√𝜌, ∝ 𝑏√𝜌 (9)\n𝑑 0\n𝐺\nCombining these gives:\n∝ 𝜌𝑏∙(𝑏√𝜌) 𝑚 = 𝑏1+𝑚𝜌1+ 𝑚 2, ∝ 𝜌1+𝑚 (10)\nThis indicates that dislocation density increases almost linearly with strain rate,\nconsistent with our experimental observation of strain rate hardening.\n3. Kinetic design of the second transition point\nBased on dislocation behavior under constant-rate loading, we implemented the\nfirst strain rate jump near the yield point in SR loading. According to the Kocks–\nMecking model 4:\n𝑑𝜌 𝑘\n(𝜔)\n= 1 √𝜌−𝑘 𝜌 (11)\n𝑑𝜀 𝑏\n(𝜔)\nHere, ε is strain, 𝑘 is the dislocation multiplication coefficient (strain rate\n1\ndependent), and k is the dynamic recovery coefficient (typically considered\n2\nconstant). Let 𝜌( ) = 𝜌 be the initial condition at yield. According to Eq. 10, the\n𝑎 𝑎\ndislocation multiplication efficiency exhibits an approximately linear relationship\nwith the loading rate, indicating that an increased strain rate leads to a faster\naccumulation of dislocation density per unit strain in the micropillar. When the\n48 / 61\n(𝜔)\nstrain rate jumps from (SR) to (MR), k increases to 𝑘 = 𝜆𝑘 , where λ is\n1 1 1 1\na strain rate sensitivity constant, and ω is the number of strain rate transitions.\nFor > , the new form becomes:\n𝑎\n(1)\n𝑑𝜌 𝑘\n= 1 √𝜌−𝑘 𝜌 𝑓𝑜𝑟 > (12)\n𝑎\n𝑑𝜀 𝑏\nThis equation is in a nonlinear separable form. After the first-rate jump, the\ndislocation density in the MR stage evolves as:\n𝜌 𝑏 ( ) = [√𝜌 𝑎 ∙𝑒 𝑘2(𝜀 2 −𝜀𝑎) + 𝑘 1 (1) ( −𝑒 𝑘2(𝜀 𝜀𝑎) )] (13)\n𝑏𝑘2\nTo determine the appropriate second transition point that ensures sufficient\nconditions for amorphization, we designed additional experiments\n(Supplementary Fig. S25) where loading is stopped after completing only the MR\nstage. As shown in the stress–strain curves (Supplementary Fig. S26),\nincreasing the strain rate improves the mechanical response of the micropillar.\nShear banding begins at ~26% strain, and after full crack development (~40%\nstrain), the substrate exhibits significant plastic deformation, evidenced by a\nmarked change in the stress–strain slope.\nConsidering the inertia of the loading system and the potential for shear precursors,\nthe second transition to the HR stage is conservatively selected at a strain ε ≤\nb\n20%, which safely precedes catastrophic localization.\nSimilarly, the evolution of dislocation density ρ in the HR stage (strain rate )\nc\nafter the second strain rate jump can be analytically expressed as:\n𝜌 𝑐 ( ) = [√𝜌 𝑏 ∙𝑒 𝑘2(𝜀 2 −𝜀𝑏) + 𝑘 1 (2) ( −𝑒 𝑘2(𝜀 𝜀𝑏) )] 𝑓𝑜𝑟 > 𝑏 (14)\n𝑏𝑘2\nIn this expression, the exponential terms are monotonically increasing, indicating\nthat the dislocation density rises nonlinearly and consistently with both strain and\nstrain rate. The linear strain-rate hardening observed under conventional single-\nstage loading (Eq. 10) is restructured into an exponential-like hardening behavior\n49 / 61\nby the ST strategy. Therefore, higher dislocation accumulation during each stage\nof ST promotes a greater final dislocation density.\nDislocation density is closely linked to the amorphization process 5, with the critical\ncondition defined as:\n𝛾\n𝜌 = (15)\n𝑛 𝐺𝑏3\nwhere ρ is the critical dislocation density for amorphization and γ\\gamma is the\nn\nstacking fault energy. When ρ ≥ ρ , amorphization is likely to initiate in the\nc n\nmicropillar. To promote a broader and more pronounced amorphous zone, it is\nnecessary to maximize the duration of dislocation accumulation during the uniform\ndeformation stage. Accordingly, we select the second strain rate transition point\nat ε = 20%. The subsequent HR stage then provides the localized stress and\nb\nlattice perturbation required to transition from unstable shear to long-range\namorphization—a “window zone” for this transformation.\nTo quantify the volume fraction of the amorphous phase f as a function of the\nam\ntransition strain ε , we define:\nb\n𝑓 ( ) = max ( ,\n𝜌𝑐 (𝜀𝑏 )\n) (16)\n𝑎𝑚 𝑏\n𝜌𝑛\nSince amorphization is a defect-induced nonlinear transformation, this relationship\n20 / 61\nis not linear 6. When ρ < ρ , the crystalline lattice retains recovery capability, and\nc n\nthe amorphous volume fraction is near zero. As ρ approaches ρ , local atomic\nc n\norder begins to collapse, and small amorphous nuclei emerge—this corresponds\nto the typical nanoscale amorphization. To trigger continuous amorphization,\nfurther dislocation accumulation (ρ ) is required, causing localized nuclei to\ns\ncoalesce into amorphous bands 6. This transition, characterized by a critical\nthreshold, smooth evolution, and eventual saturation, is well described by the\nSigmoid function 7, 8:\n1\n𝑓(𝑥) = (17)\n1+𝑒−𝑥\nThus, the amorphous volume fraction becomes a probabilistic measure of phase\ntransformation per unit region:\n1\n𝑓 ( ) = { −𝛽( 𝜌𝑐(𝜀𝑏) ), 𝜌𝑐≥𝜌𝑛 (18)\n𝑎𝑚 𝑏 1+𝑒 𝜌𝑛 𝜌𝑐<𝜌𝑛\n,\nHere, β is a steepness parameter in the Sigmoid function that reflects the rate and\nsensitivity of disorder propagation. In metallic systems, amorphization typically\ninitiates with local atomic rearrangement and grows progressively. When the\ndislocation density approaches the critical value, long-range atomic order is lost.\nA large β implies that small increases in ρ can rapidly trigger a cascade of\nc\namorphous nucleation events, whereas a smaller β indicates a more gradual\ntransformation that requires sustained dislocation buildup. Thus, β is a key\ndescriptor of the amorphous nucleation and propagation kinetics. This kinetic\nframework is similar to Li et al. 's theories 6.\nAdditionally, we performed in situ compression experiments in which the\nmicropillar was directly transitioned from SR to HR loading (skipping the MR stage)\nto increase the strain rate jump amplitude (Supplementary Figs. S27). Due to the\nlarge transition span, the control system required ~10 s to stabilize the strain rate\n(as shown by the proportion-integration-differentiation controller).\nHowever, the low dislocation density developed during the SR yield stage is\ninsufficient to form a dense defect network capable of driving rapid lattice\n21 / 61\nreconstruction in the HR phase. As a result, minor discrete shear steps were\nobserved at ~10% strain, followed by shear slip activation in both the middle (~40%\nheight) and base of the micropillar. According to the amorphization kinetics, the\ndislocation storage capacity of the direct jump loading path returns to the constant-\nrate HR loading mode. These results highlight the necessity of the intermediate\nMR stage in the ST strategy to ensure progressive regulation and effective defect\nevolution.\nRelated Figures:\n51 / 61\nFig. S25. Deformation process of the micropillar under SR–MR limiting compression.\n(a) The first strain rate jump was applied at ~6% strain. (b) During the MR stage, shear\nslip steps began to appear at ~26% strain. (c, d) With continued loading, these slip steps\nprogressively evolved into mature shear cracks, ultimately compromising the structural\nintegrity of the micropillar.\nFig. S26. Stress-strain curve of the micropillar loaded by SR-MR.\n52 / 61\nFig. S27. Failure process of a micropillar subjected to a direct strain rate jump from\nSR to HR. (a) The micropillar underwent a direct transition from SR to HR loading without\nan MR stage. (b, c) The absence of progressive dislocation accumulation leads to\npremature shear activation and the formation of discrete slip steps, ultimately resulting in\nstructural failure.\nReferences\n1. Orowan, E. Problems of plastic gliding. Proc. Phys. Soc. 52, 8 (1940).\n2. Laidler, K. J. The development of the Arrhenius equation. J. Chem. Educ. 61,\n494 (1984).\n3. Courtney, T. H. Mechanical Behavior of Materials. (Waveland Press, 2005).\n4. Kocks, U. & Mecking, H. Physics and phenomenology of strain hardening: the\nFCC case. Prog. Mater. Sci. 48, 171–273 (2003).\n5. Zhao, S. et al. Amorphization in extreme deformation of the CrMnFeCoNi\nhigh-entropy alloy. Sci. Adv. 7, eabb3108 (2021).\n6. Li, B. Y., Li, A. C., Zhao, S. & Meyers, M. A. Amorphization by mechanical\ndeformation. Mater. Sci. Eng. R Rep. 149, 100673 (2022).\n7. Kyurkchiev, N. & Markov, S. Sigmoid functions: some approximation and\nmodelling aspects. LAP LAMBERT Acad. Publ. Saarbrucken 4, 34 (2015).\n8. Delogu, F., Schiffini, L. & Cocco, G. The invariant laws of the amorphization\nprocesses by mechanical alloying: I. Experimental findings. Philos. Mag. A 81,\n1917–1937 (2001).\n3. Since it is a micropillar experiment, the influence of micropillar size should be\nconsidered.\nResponse: We appreciate your insightful comment regarding the critical role of\nmicropillar size in mechanical response. We systematically investigated the\nmechanical response of microcolumns with different diameters under strain\ntraining (ST) strategies. The results reveal a non-monotonic size dependence in\nboth strength and ductility under ST loading—a notable deviation from the\nclassical \"smaller-is-stronger\" trend. Furthermore, the critical threshold for stress-\ninduced amorphization is found to be independent of pillar diameter, governed\ninstead by local stress states and the density of dislocations reaching a critical\nlevel. Key supporting data and analysis are presented in Supplementary Note 7.\n53 / 61\nRevision:\nSupplementary Note 7. Effect of micropillar diameter on mechanical\nresponse under ST loading\n1. Experimental considerations\nThe selection of micropillar diameter must balance deformation resolution with\nloading precision. For diameters d < 1 μm, accurate strain rate modulation is\nhindered by SEM and sensor resolution limitations at the nanoscale. At d = 2.5\nμm, severe bulging under ST loading causes the top surface to exceed the\nindenter contact area, leading to non-uniform deformation (Supplementary Fig.\nS29). Nevertheless, it can still be concluded that no shear failure occurred during\nthe SR and MR stages, and that deformation capacity was enhanced at this size.\nFor d > 2.5 μm, the peak loads exceed the 50 mN limit of the indenter load cell,\nrisking instrumental damage and reduced data fidelity. Therefore, we selected 1\nμm, 1.5 μm, and 2 μm pillars to ensure both data accuracy and representativeness\nunder the constraints of instrument capability and material response.\n2. Size effect\nIn line with classical plasticity theory, small variations in sample dimensions can\ncause significant fluctuations in mechanical behavior due to dislocation source\nlimitations and reduced plastic flow volume at smaller scales. To evaluate how the\nST strengthening strategy responds to size changes, we conducted additional ST\ncompression experiments on 1 μm and 1.5 μm pillars alongside the 2 μm\nbenchmark.\nThe strength trend shows a non-monotonic dependence on micropillar diameter,\ninitially increasing then decreasing with size reduction (Supplementary Figs. S30,\nS31, and Main Text Figs. 1h, i). This indicates that flow stress under ST remains\nsensitive to the size effect. TEM and deformation analyses reveal that 1 μm pillars\nexhibit dense dislocation pile-up at the base, indicating restricted space for\ndislocation multiplication (Supplementary Figs. S32, S33). This limited capacity\nfor strained accommodation promotes local stress concentration, making shear\nstress more intense than in larger pillars. As a result, amorphization alone cannot\n54 / 61\nsufficiently dissipate stress, leading to localized shear band formation during the\nHR stage (Supplementary Fig. S33c). This confirms that at the near-nanometer\nscale, adequate pillar height must be retained to sustain dislocation activity.\nNevertheless, due to the extensive amorphization, the 1 μm pillar still exhibits a\nsubstantial suppression of shear-induced failure (Supplementary Figs. S33a, b).\nThe mechanical behavior of the 1.5 μm pillar closely resembles that of the 2 μm\npillar (Supplementary Fig. S34), with detrimental shear banding effectively\neliminated. This strengthening mechanism remains operative across all tested\nsizes, consistent with Eq. 1 in the main text, which indicates that the critical\ncondition for crystal-to-amorphous transition is not dictated by geometric\ndimensions, but rather by dislocation-induced stress accumulation.\nTaken together, the size-dependent mechanical response under the ST protocol\nis governed by a dynamic competition between limited dislocation multiplication\nspace and amorphization-mediated compensation.\n3. Limitations and future work\nThe current ST protocol was optimized for 2 μm micropillars. For other sizes,\nadjustments to the strain rate transition points may be required to maintain optimal\ncoordination between dislocation evolution and amorphization onset. Future\nstudies will explore the quantitative relationship between micropillar geometry and\nST strain rate transition to establish a scalable strain-engineering framework.\nWe sincerely appreciate your insightful input, which prompted this in-depth\nextension of our work.\nRelated Figures:\n55 / 61\nFig. S29. ST compression process of a 2.5 μm MPEA micropillar. (a, b) Under quasi-\nstatic SR and medium strain rate MR loading, the micropillar exhibits uniform deformation\nbehavior similar to that of the 2 μm micropillar. (c) At the high-strain-rate HR stage,\nhowever, pronounced lateral bulging occurs. The top surface area exceeds the indenter\nsize, and the edges lose load constraint, with only the middle part continuing to be loaded.\nThese results indicate that larger micropillar sizes cannot match the indenter.\nFig. S30. Compressive stress–strain response of the 1 μm diameter micropillar\nunder ST loading.\nFig. S31. Compressive stress–strain response of the 1.5 μm diameter micropillar\n56 / 61\nunder ST loading.\nFig. S32. Failure process of the 1 μm diameter micropillar under ST loading. (a)\nElastic deformation and early yielding stage under SR loading. (b) Homogeneous plastic\ndeformation stage during MR loading. Minor surface slip traces are observed near the\nstrain-rate transition point, indicating potential internal stress concentration. (c) Rapid\nshear slip is triggered under HR loading, leading to sudden failure.\nFig. S33. Post-deformation microstructure of 1 μm MPEA micropillar under ST. (a\ni–\n) Microstructural morphology after ST compression, showing the coexistence of\niii\namorphous bands and shear bands. A gradient deformation structure is observed, with the\ndegree of distortion progressively decreasing from top to bottom. (b) Enlarged stitched\n57 / 61\nimages reveal, in contrast to 2 μm micropillars, the presence of high-density dislocations\nat the bottom, indicating that the 1 μm micropillar lacks sufficient space for effective\ndislocation glide. To accommodate the deformation, both shear and amorphization\nmechanisms are activated. (c ) Nanoscale deformation twins are identified adjacent to\ni–iv\nthe amorphous regions.\nFig. S34. Failure process of the 1.5 μm diameter micropillar under ST loading. Note\nthat the oblique striations indicated by red arrows in (a) are not intrinsic crystalline defects\nof the material, but rather surface artifacts induced by focused ion beam (FIB) milling.\nThese linear marks typically originate from high-current rough cuts and angled scanning\n(~52°) during FIB processing, resulting in surface streaks with characteristic tilt angles of\n~38°. (b) During MR compression, these FIB-induced striations tend to converge but do\nnot correspond to shear bands. No shear band initiation is observed in the MR stage,\nconfirming that these artifacts have negligible influence. (c) The striations are present prior\nto loading and exert minimal influence on the HR stage, which exhibits rapid deformation\nnear the compressive limit. The overall mechanical behavior remains unaffected, with a\ndeformation pattern that closely resembles that observed in the 2 μm micropillar.\n4. The strength calculation using the rule-of-mixture (ROM) require more\nclarifications, i.e. amorphous phase strength, fraction, etc..\nResponse: Thanks to you for your insightful comments regarding the strength\nestimation methods used in our work. This concise summary of strength\nestimation based on the ROM model has now been included in Supplementary\nNote 5.\n58 / 61\nRevisions:\nSupplementary Note 5. Quantification process of strengthening the\nmechanism.\nThe rule-of-mixture (ROM) model is fundamentally a first-order approximation\nmethod, which approximates the overall strength of macroscopic bulk samples as\nthe sum of contributions from individual strengthening mechanisms. The\nmechanical behavior of the amorphous phase is indeed different from that of\ncrystalline phases. However, both our work and previous studies 1,2 have\ndemonstrated that, under specific constraint conditions, amorphous phases can\neffectively participate in load sharing and enhance the matrix, thus further\nsupporting the rationality of employing the ROM model to evaluate its contributions\nin this study.\nExperimentally obtained ultimate strength (σ) of the material was combined with\ntheoretical models to estimate the contributions from lattice friction (σ₀), dislocation\nstrengthening (σ ), phase boundary strengthening (σ ), and solid solution\nd p\nstrengthening (σ ). The residual difference was then attributed to the\nss\nenhancement effect due to amorphization (σ ). It should be noted that the\nh\nmechanical properties significantly differ from those of macroscopic bulk alloys\ndue to size effects, and thus, any estimation errors under quasi-static conditions\nare attributed to the scale effects (σ ). Specifically, we have Eq. 4:\ns\n𝝈 = 𝝈−𝝈 −𝝈 −𝝈 −𝝈 −𝝈 (4)\n𝒉 𝟎 𝒔 𝒔𝒔 𝒑 𝒅\nBased on the Hall-Petch relationship (Eq. 5) 3, the phase boundary strength (σ )\np\nfor lamellar structured alloys relates to the average lamella spacing (λ):\n𝝈 = 𝝈 +𝑲 (𝝀) 𝟏 𝟐 (5)\n𝒑 𝟎 𝟎\nHere, the lattice friction σ₀ ≈ 256 MPa 4, and the Hall-Petch constant K₀ for the\neutectic CoCrAlFeNi multi-principal element alloy (MPEA) composed of BCC and\nFCC phases 4 was determined to be ~226 MPa·μm1/2. According to TEM\ncharacterization, the average lamella spacing λ ≈ 210 nm. Moreover, referencing\nprevious calculations, the solid-solution strengthening (σ ) in the eutectic\nss\n59 / 61\nCoCrAlFeNi MPEA 5 was estimated as ~80 MPa.\nThe dislocation strengthening contribution to alloy strength (σ ) was evaluated by\nd\nthe Taylor relation (Eq. 6) 6:\n𝟏 𝟏\n𝝈\n𝒅\n= 𝑴𝜶𝑮𝒃𝝆𝟐 = 𝒇\n𝑭𝑪𝑪\n𝑴\n𝑭𝑪𝑪\n𝒂\n𝑭𝑪𝑪\n𝑮\n𝑭𝑪𝑪\n𝒃\n𝑭𝑪𝑪\n𝝆 𝑭𝑪𝑪𝟐+𝒇\n𝑩𝑪𝑪\n𝑴\n𝑩𝑪𝑪\n𝒂\n𝑩𝑪𝑪\n𝑮\n𝑩𝑪𝑪\n𝒃\n𝑩𝑪𝑪\n𝝆\n𝑩𝑪𝑪\n𝟏 𝟐 (6)\nWhere M is the Taylor factor (~3.09 for the fcc phase and ~2.71 for the bcc phase,\nmeasured from EBSD), α is a dimensionless prefactor (~0.20 for FCC 7, ~0.24 for\nBCC 8), G is the shear modulus (~81 GPa for FCC, ~57 GPa for BCC) 9, b is the\nmagnitude of the Burgers vector (~0.254 nm for FCC, ~0.248 nm for BCC,\nmeasured by XRD) 10, ρ is the dislocation density (obtained from Fig. 4e), and the\nvolume fractions f of the FCC and BCC phases (~67% and ~33%) were measured\nby EBSD and TEM (Fig. 1b, d). The quantified contributions of each strengthening\nmechanism are provided in Supplementary Table S1 and Fig. 4g.\nTable S1. Quantification results of the strengthening mechanism of ST micropillars.\nStrengthening Calculate (GPa)\nNo.\nmechanism SR MR HR ST\n1 σ 0.256 0.256 0.256 0.256\n0\n2 σ 1.131 1.131 1.131 1.131\ns\n3 σ 0.080 0.080 0.080 0.080\nss\n4 σ 0.340 0.407 0.566 0.877\nd\n5 σ 0.749 0.749 0.749 0.749\np\n6 σ - - - 3.663\nh\nReferences\n1. Qian, L. et al. Amorphous Thickness-Dependent Strengthening–Softening\nTransition in Crystalline–Amorphous Nanocomposites. Nano Lett. 23, 11288–\n11296 (2023).\n2. Jiang, K. et al. Microstructural origins of enhanced work hardening and\nductility in laser powder-bed fusion 3D-printed AlCoCrFeNi2.1 eutectic high-\nentropy alloys. Int. J. Plasticity. 159, 103463 (2022).\n3. Hansen, N. Hall–Petch relation and boundary strengthening. Scripta Mater.\n51, 801-806 (2024).\n60 / 61\n4. Son, S. et al. Hierarchically heterogeneous microstructure and mechanical\nbehavior of the multi-materials prepared by powder severe plastic deformation.\nMater. Res. Lett. 11, 915-924 (2023).\n5. Mao, Z. et al. Understanding the yield strength difference in dual-phase\neutectic high-entropy alloys. Mat. Sci. Eng. A-Struct. 867, 144725 (2023).\n6. Courtney, T. H. Mechanical Behavior of Materials (Waveland, 2005).\n7. He, J. Y. et al. A precipitation-hardened high-entropy alloy with outstanding\ntensile properties. Acta Mater. 102, 187–196 (2016).\n8. Zhang, X., Hansen, N., Godfrey, A. & Huang, X. Dislocation-based plasticity\nand strengthening mechanisms in sub-20 nm lamellar structures in pearlitic\nsteel wire. Acta Mater. 114, 176–183 (2016).\n9. Ren, J. et al. Strong yet ductile nanolamellar high-entropy alloys by additive\nmanufacturing. Nature. 608, 62–68 (2022).\n10. Du, J. et al. Mitigating Strain Localization via Stabilized Phase Boundaries for\nStrengthening Multi-Principal Element Alloys. Adv. Sci. 2414783 (2025).\n5. Some minor errors: (1) (Page 5) The strain rate sensitivity is not consistent with\nyour supplementary Fig. S7. (2) (Page 12) ~3.26×10^15 lacks unit.\nResponse: Thank you for kindly pointing out these errors. We have recalculated\nand corrected the strain rate sensitivity value to m ≈ 0.13, which has been updated\non Page 5. Additionally, we have added the missing unit ~3.26 × 10¹⁵ m⁻², which\nis now displayed on Page 12.\nWe once again extend our sincere gratitude for your thoughtful and\nconstructive evaluation. Your insights have helped us to refine the scope\nand strengthen the clarity and impact of our manuscript. We hope that the\nrevised version meets your expectations.\n61 / 61\nResponds to the referees’ comments:\nReferee #1\nMy concerns have been adequately addressed.\nResponse: We thank you for your positive feedback and recommendation. We\nare glad that our responses have fully addressed your concerns.\nReferee #2\nThe authors have clarified all of my questions and addressed the concerns raised.\nI therefore recommend the manuscript for publication.\nResponse: We sincerely appreciate your thoughtful comments and\nrecommendations for publication. We are pleased that all concerns have been\nadequately addressed.\nReferee #3\nThe response is thorough and diligent, containing substantial revisions and new\nexperimental work. I appreciate the considerable effort invested in addressing the\nmajority of the concerns raised, particularly the role of eutectic microstructures and\nstrategy applicability to single-phase case. These revisions have demonstrably\nstrengthened the paper's rigor and contribution. However, the newly added\ncontent naturally prompts additional questions for clarification and potential\nrefinement:\nResponse: We thank you for the encouraging feedback and recognition of our\nrevisions and added experimental work. We also appreciate the follow-up\nquestions raised in response, which we believe are helpful for further improving\nthe clarity of this work. Our detailed, point-by-point responses are provided below.\n1. (Fig. S16) Could you please elaborate the GPA analysis? High-resolution images\nand selected g vectors are necessary for such analysis.\nResponse: Thank you for your valuable suggestion.\nGeometric phase analysis (GPA) was performed using the GPA++ software to\n1 / 6\nmap the distribution of lattice strain (ε ) across the phase boundaries in the ST\nxy\nmicropillar. The high-resolution TEM image used for GPA, along with the selected\ng-vector, has now been added as Supplementary Figs. S16(d-i).\nAs mentioned in Supplementary Note 4, the phase boundaries act as elastic\nbarriers that suppress lateral dislocation motion and promote dislocation gliding\nalong the loading direction. This effectively alleviates local stress concentrations\ncaused by dislocation pile-up and enables more uniform stress redistribution\nthrough phase-boundary-assisted pathways.\nRevision\nThese clarifications have been included in Supplementary Figs. S16(d-i).\nFig. S16. (d–f) High-resolution images near the phase boundary. (g-i) Geometric\nphase analysis using GPA++ reveals the constraining effect of the phase boundary on\nlattice distortion along the shear direction (ε ). The selected g-vector is shown in the\nxy\nupper-right inset. The phase boundaries act as elastic barriers that suppress lateral\ndislocation motion and promote dislocation gliding along the loading direction. This\neffectively alleviates local stress concentrations caused by dislocation pile-up and\n2 / 6\nenables more uniform stress redistribution through phase-boundary-assisted\npathways.\n2. (Fig. S22) I noted an apparent error of labeling in ST sample, both marked as\n(111).\nResponse: Thank you for your careful observation. We have corrected the\nlabeling error in Fig. S22 for the ST sample. The revised figure has been updated\nin the Supplementary Information, with the correction clearly highlighted in blue for\nclarity.\nRevision\nFig. S22. Mechanical response of CoCrNi micropillars under ST loading. The\ncomparative data in this figure are referenced from Ref. J. Mech. Phys. Solids. 162,\n104853 (2022).\n3. It could be better to merge the stress-strain curves in Fig. S28, Fig. S29 and Fig.\n1 to show the size effects clearly. The figure numbers could be wrong in\nSupplementary Note 7, check it please.\nResponse: Thank you for your insightful suggestion. In response, we have\nreorganized the stress–strain data and merged the key curves from Fig. 1 and the\noriginal Supplementary Figs. S28 and S29 into a single composite figure, now\npresented as the revised Supplementary Fig. S28. We have also carefully\ncorrected the figure references in Supplementary Note 7.\n3 / 6\nAs stated in Supplementary Note 7, the strength trend shows a non-monotonic\ndependence on micropillar diameter (Supplementary Fig. S28). The initial yield\nstrength exhibits a generally monotonic size dependence, consistent with\nconventional expectations. However, the ultimate compressive strength displays\nan inverse trend: the 1.5 μm and 1 μm micropillars show lower peak strengths than\nthe 2 μm pillar. This indicates that the flow stress under the ST condition remains\nsensitive to sample size. TEM and deformation analyses reveal dense dislocation\npile-up at the base of the 1 μm pillar, suggesting that the space for dislocation\nmultiplication is severely limited (Supplementary Figs. S29 and S30). This\nrestricted plastic accommodation promotes local stress concentrations, resulting\nin higher shear stress than in larger pillars. Consequently, the amorphization\nprocess alone cannot sufficiently dissipate the stress, leading to localized shear\nband formation during the HR stage (Supplementary Fig. S30). This confirms\nthat at near-nanoscale dimensions, sufficient pillar height or volume is required to\nsustain dislocation activity.\nRevision\nFig. S28. Compressive stress–strain response of the 1 μm, 1.5 μm, and 2 μm\ndiameter micropillars under ST loading.\n4. Could you precisely clarify how the dislocation density in Fig. 4e is utilized for\nstrength calculation? I found the apparent discrepancy between Fig. 4g and Table\nS1.\nResponse: Thank you for pointing out this important issue.\nTo clarify, the dislocation density shown in Fig. 4e was extracted from post-\n4 / 6\ncompression TEM images and was primarily used to qualitatively correlate the\nmicrostructural evolution with the observed strain hardening. The alloy consists of\nboth BCC and FCC phases, and the simply summed dislocation densities for each\nphase are presented in Supplementary Table S2. For quantitative strength\nestimation, we adopted the Taylor hardening model combined with the rule of\nmixtures (ROM), as summarized in Supplementary Table S3. The detailed\ncalculation procedures for each strengthening mechanism are provided in\nSupplementary Note 5.\nThe apparent discrepancy arises because Fig. 4g presents the total strength\ncontribution from solid-solution strengthening, phase boundaries, dislocations,\nand amorphous phase, whereas Supplementary Table S1 summarizes the\ncontributions, including the dislocation-related component calculated as above,\nalongside other strengthening mechanisms. To enhance clarity, we have revised\nthe caption of Fig. 4g to explicitly distinguish this difference.\nRevision\nWe have revised the caption of Fig. 4g to explicitly distinguish this difference\n(Lines 281-282): All strengthening mechanisms and dislocation density\ndistributions are detailed in Supplementary Tables S1-3.\nTable S2. Information on the top-down dislocation distribution (×104/m2) of\nmicropillars.\nSR MR HR ST\nLocation\nFCC BCC FCC BCC FCC BCC FCC BCC\nTop 5 4 8 6 18 12 31 20\nMiddle 2.5 1.5 3 2 6 3.5 25 12\nBottom 2 1 2.5 1.5 2.5 2 10 5\nTotal 9.5 6.5 13.5 9.5 26.5 17.5 66 37\nTable S3. ROM-based dislocation strengthening (σ GPa) quantification details.\nd,\nType SR MR HR ST\nFCC 0.392 0.467 0.655 1.033\nBCC 0.234 0.283 0.385 0.559\nROM 0.340 0.407 0.566 0.877\n5 / 6\nWe once again sincerely thank you for your constructive feedback. Your\ninsightful comments have significantly improved the clarity and impact of\nour manuscript. We hope that the revised version now meets all your\nexpectations.\n6 / 6"}, {"round": 2, "reviewer_comments": {"reviewer_1": "My concerns have been adequately addressed.", "reviewer_2": "The authors have clarified all of my questions and addressed the concerns raised. I therefore recommend the manuscript for publication.", "reviewer_3": "The response is thorough and diligent, containing substantial revisions and new experimental work. I appreciate the considerable effort invested in addressing the majority of the concerns raised, particularly the role of eutectic microstructures and strategy applicability to single-phase case. These revisions have demonstrably strengthened the paper's rigor and contribution. However, the newly added content naturally prompts additional questions for clarification and potential refinement:\n1. (Fig. S16) Could you please elaborate the GPA analysis? High-resolution images and selected g vectors are necessary for such analysis.\n2. (Fig. S22) I noted an apparent error of labeling in ST sample, both marked as (111).\n3. It could be better to merge the stress-strain curves in Fig. S28, Fig. S29 and Fig. 1 to show the size effects clearly. The figure numbers could be wrong in Supplementary Note 7, check it please.\n4. Could you precisely clarify how the dislocation density in Fig. 4e is utilized for strength calculation? I found the apparent discrepancy between Fig. 4g and Table S1."}, "author_response": "Responds to the referees’ comments:\nReferee #1\nMy concerns have been adequately addressed.\nResponse: We thank you for your positive feedback and recommendation. We\nare glad that our responses have fully addressed your concerns.\nReferee #2\nThe authors have clarified all of my questions and addressed the concerns raised.\nI therefore recommend the manuscript for publication.\nResponse: We sincerely appreciate your thoughtful comments and\nrecommendations for publication. We are pleased that all concerns have been\nadequately addressed.\nReferee #3\nThe response is thorough and diligent, containing substantial revisions and new\nexperimental work. I appreciate the considerable effort invested in addressing the\nmajority of the concerns raised, particularly the role of eutectic microstructures and\nstrategy applicability to single-phase case. These revisions have demonstrably\nstrengthened the paper's rigor and contribution. However, the newly added\ncontent naturally prompts additional questions for clarification and potential\nrefinement:\nResponse: We thank you for the encouraging feedback and recognition of our\nrevisions and added experimental work. We also appreciate the follow-up\nquestions raised in response, which we believe are helpful for further improving\nthe clarity of this work. Our detailed, point-by-point responses are provided below.\n1. (Fig. S16) Could you please elaborate the GPA analysis? High-resolution images\nand selected g vectors are necessary for such analysis.\nResponse: Thank you for your valuable suggestion.\nGeometric phase analysis (GPA) was performed using the GPA++ software to\n1 / 6\nmap the distribution of lattice strain (ε ) across the phase boundaries in the ST\nxy\nmicropillar. The high-resolution TEM image used for GPA, along with the selected\ng-vector, has now been added as Supplementary Figs. S16(d-i).\nAs mentioned in Supplementary Note 4, the phase boundaries act as elastic\nbarriers that suppress lateral dislocation motion and promote dislocation gliding\nalong the loading direction. This effectively alleviates local stress concentrations\ncaused by dislocation pile-up and enables more uniform stress redistribution\nthrough phase-boundary-assisted pathways.\nRevision\nThese clarifications have been included in Supplementary Figs. S16(d-i).\nFig. S16. (d–f) High-resolution images near the phase boundary. (g-i) Geometric\nphase analysis using GPA++ reveals the constraining effect of the phase boundary on\nlattice distortion along the shear direction (ε ). The selected g-vector is shown in the\nxy\nupper-right inset. The phase boundaries act as elastic barriers that suppress lateral\ndislocation motion and promote dislocation gliding along the loading direction. This\neffectively alleviates local stress concentrations caused by dislocation pile-up and\n2 / 6\nenables more uniform stress redistribution through phase-boundary-assisted\npathways.\n2. (Fig. S22) I noted an apparent error of labeling in ST sample, both marked as\n(111).\nResponse: Thank you for your careful observation. We have corrected the\nlabeling error in Fig. S22 for the ST sample. The revised figure has been updated\nin the Supplementary Information, with the correction clearly highlighted in blue for\nclarity.\nRevision\nFig. S22. Mechanical response of CoCrNi micropillars under ST loading. The\ncomparative data in this figure are referenced from Ref. J. Mech. Phys. Solids. 162,\n104853 (2022).\n3. It could be better to merge the stress-strain curves in Fig. S28, Fig. S29 and Fig.\n1 to show the size effects clearly. The figure numbers could be wrong in\nSupplementary Note 7, check it please.\nResponse: Thank you for your insightful suggestion. In response, we have\nreorganized the stress–strain data and merged the key curves from Fig. 1 and the\noriginal Supplementary Figs. S28 and S29 into a single composite figure, now\npresented as the revised Supplementary Fig. S28. We have also carefully\ncorrected the figure references in Supplementary Note 7.\n3 / 6\nAs stated in Supplementary Note 7, the strength trend shows a non-monotonic\ndependence on micropillar diameter (Supplementary Fig. S28). The initial yield\nstrength exhibits a generally monotonic size dependence, consistent with\nconventional expectations. However, the ultimate compressive strength displays\nan inverse trend: the 1.5 μm and 1 μm micropillars show lower peak strengths than\nthe 2 μm pillar. This indicates that the flow stress under the ST condition remains\nsensitive to sample size. TEM and deformation analyses reveal dense dislocation\npile-up at the base of the 1 μm pillar, suggesting that the space for dislocation\nmultiplication is severely limited (Supplementary Figs. S29 and S30). This\nrestricted plastic accommodation promotes local stress concentrations, resulting\nin higher shear stress than in larger pillars. Consequently, the amorphization\nprocess alone cannot sufficiently dissipate the stress, leading to localized shear\nband formation during the HR stage (Supplementary Fig. S30). This confirms\nthat at near-nanoscale dimensions, sufficient pillar height or volume is required to\nsustain dislocation activity.\nRevision\nFig. S28. Compressive stress–strain response of the 1 μm, 1.5 μm, and 2 μm\ndiameter micropillars under ST loading.\n4. Could you precisely clarify how the dislocation density in Fig. 4e is utilized for\nstrength calculation? I found the apparent discrepancy between Fig. 4g and Table\nS1.\nResponse: Thank you for pointing out this important issue.\nTo clarify, the dislocation density shown in Fig. 4e was extracted from post-\n4 / 6\ncompression TEM images and was primarily used to qualitatively correlate the\nmicrostructural evolution with the observed strain hardening. The alloy consists of\nboth BCC and FCC phases, and the simply summed dislocation densities for each\nphase are presented in Supplementary Table S2. For quantitative strength\nestimation, we adopted the Taylor hardening model combined with the rule of\nmixtures (ROM), as summarized in Supplementary Table S3. The detailed\ncalculation procedures for each strengthening mechanism are provided in\nSupplementary Note 5.\nThe apparent discrepancy arises because Fig. 4g presents the total strength\ncontribution from solid-solution strengthening, phase boundaries, dislocations,\nand amorphous phase, whereas Supplementary Table S1 summarizes the\ncontributions, including the dislocation-related component calculated as above,\nalongside other strengthening mechanisms. To enhance clarity, we have revised\nthe caption of Fig. 4g to explicitly distinguish this difference.\nRevision\nWe have revised the caption of Fig. 4g to explicitly distinguish this difference\n(Lines 281-282): All strengthening mechanisms and dislocation density\ndistributions are detailed in Supplementary Tables S1-3.\nTable S2. Information on the top-down dislocation distribution (×104/m2) of\nmicropillars.\nSR MR HR ST\nLocation\nFCC BCC FCC BCC FCC BCC FCC BCC\nTop 5 4 8 6 18 12 31 20\nMiddle 2.5 1.5 3 2 6 3.5 25 12\nBottom 2 1 2.5 1.5 2.5 2 10 5\nTotal 9.5 6.5 13.5 9.5 26.5 17.5 66 37\nTable S3. ROM-based dislocation strengthening (σ GPa) quantification details.\nd,\nType SR MR HR ST\nFCC 0.392 0.467 0.655 1.033\nBCC 0.234 0.283 0.385 0.559\nROM 0.340 0.407 0.566 0.877\n5 / 6\nWe once again sincerely thank you for your constructive feedback. Your\ninsightful comments have significantly improved the clarity and impact of\nour manuscript. We hope that the revised version now meets all your\nexpectations.\n6 / 6"}, {"round": 3, "reviewer_comments": {"reviewer_3": "Authors have provided extensive clarifications in response to reviewer's comments."}, "author_response": ""}]

Decision letter