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[SMM-S-24-01516.pdf](https://mdr.nims.go.jp/filesets/77968dc0-2ef6-43fb-9f68-f8198c36f659/download)

## Creator

[Digvijay Singh](https://orcid.org/0000-0003-4357-957X), [Fumiyoshi Yoshinaka](https://orcid.org/0000-0003-0534-7815), [Susumu Takamori](https://orcid.org/0000-0003-3422-9391), [Satoshi Emura](https://orcid.org/0000-0001-5789-6408), [Takahiro Sawaguchi](https://orcid.org/0000-0002-9405-002X)

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[Revealing extraordinary plasticity in single crystal Fe-15Mn-10Cr-8Ni-4Si alloy](https://mdr.nims.go.jp/datasets/602dead6-b1e7-411e-bfac-65ed935fa25d)

## Fulltext

Graphical AbstractRevealing Extraordinary Plasticity in Single Crystal Fe-15Mn-10Cr-8Ni-4Si Alloy  Digvijay Singh1, Fumiyoshi Yoshinaka1, Susumu Takamori1, Satoshi Emura1, Takahiro Sawaguchi1*  1 National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan  Digvijay Singh singh.digvijay@nims.go.jp  Fumiyoshi Yoshinaka yoshinaka.fumiyoshi@nims.go.jp  Susumu Takamori takamori.susumu@nims.go.jp  Satoshi Emura emura.satoshi@nims.go.jp  Takahiro Sawaguchi* sawaguchi.takahiro@nims.go.jp  * Corresponding author       Manuscript (word or pdf w/ figures) Click here to view linked References 1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 mailto:singh.digvijay@nims.go.jpmailto:yoshinaka.fumiyoshi@nims.go.jpmailto:takamori.susumu@nims.go.jpmailto:emura.satoshi@nims.go.jpmailto:sawaguchi.takahiro@nims.go.jphttps://www2.cloud.editorialmanager.com/smm/viewRCResults.aspx?pdf=1&docID=80148&rev=0&fileID=1265565&msid=1b4fa710-b8d3-49c6-9af4-5671fea8a62ahttps://www2.cloud.editorialmanager.com/smm/viewRCResults.aspx?pdf=1&docID=80148&rev=0&fileID=1265565&msid=1b4fa710-b8d3-49c6-9af4-5671fea8a62aAbstract An extraordinary elongation of over 156 % was achieved in single crystal Fe-15Mn-10Cr-8Ni-4Si alloy, which shows γ → ε → α’ transformation. Tensile tests were conducted along the <414>, <111>, and <100> directions, and the deformed microstructure was characterized with electron-backscattering diffraction (EBSD). Elongation in <414> orientation (156%) was highest as ever recorded in metallic alloy—approximately twice that of conventional TWIP steels. Mechanical behavior in <414> is characterized with slow work hardening, with upward curvature in the stress–strain curve. EBSD analysis revealed 45% of ε-martensite and significant γ-twin microstructure components in <414>. In initial stage, deformation was marked by planar slip and a sluggish γ → ε transformation, activating a single ε-martensite variant with a Schmid factor of 0.5. Later stages witnessed crystal rotation towards <112>, generating multiple shear bands and distinct ε variants. Fractures predominantly occurred along the <011> direction, with the fracture surface exhibiting a ductile nature.  Keywords: Extraordinary plasticity, Damping steel, Single crystal, Martensitic transformation, Shear band        1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 In recent year, Fe–Mn–Si-based alloys are renowned for their shape memory effect, owing to their unique γ → ε martensitic transformation characteristics when it experiences specific stimuli, such as temperature or stress [1,2]. Recently, a novel Fe–15Mn–10Cr–8Ni–4Si seismic damping alloy is developed with excellent Low-cycle (LCF) fatigue properties [3,4]. The fundamental idea behind designing this alloy is to utilize the reversible γ ↔ ε transformation under cyclic loading, which was discovered in high Mn-steels and shape memory alloys [5–18]. Reversible γ ↔ ε transformation is considered beneficial because it can retardative crack initiation and propagation, leading to prolonged LCF life [17]. Recently we identified that the necessary thermodynamic condition for achieving reversible γ ↔ ε transformation is ΔGγ→ε ≤ 0 [4,18]. Fe-15Mn-10Cr-8Ni-4Si alloy represents a class of medium-Mn austenitic steels with low-stacking-fault-energy (𝜞SFE ~ 19 mJ/m2) and minimal difference between Gibbs free energy between γ- and ε-phase (ΔGγ→ε ~ 12.8 J/mol) [4]. With a low SFE, this alloy exhibits an impressive strength-ductility balance (51.4 GPa%), as evidenced by a total elongation of 77% [6]. This performance is comparable to that of twinning-induced plasticity (TWIP) steels, due to the diverse range of martensitic transformations (γ → ε, γ → ε → α', or γ ↔ ε) and deformation twinning mechanisms [19,20]. During the γ → ε martensitic transformation, organized movement of {111} <121>γ Shockley partials takes place within {111} close-packed planes.  It is widely acknowledged that the crystallographic orientation can greatly influence the tensile properties by directly affecting the dominant deformation mechanisms [21–26]. Wang et al. demonstrated orientation-dependent tensile behavior in 316L stainless steel single crystals made by laser powder bed fusion [21]. Dislocation slip  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 dominated <001> oriented samples, while twinning dominated <111> samples. UTS and elongation were 645 MPa and 36% for <001>, and 840 MPa and 58% for <111>. Following this, we have investigated the tensile properties of Fe-15Mn-10Cr-8Ni-4Si seismic damping alloy single crystal in different crystallographic directions. The deformation was made in <001>, <111>, and <414> directions. In single crystals, strain hardening behavior is closely associated with the activation of secondary slip systems. The influence of these secondary slips on strain hardening capacities varies significantly depending on the crystal orientation [24,27]. In this paper, we report the extraordinarily high ductility (156%) of novel Fe-Mn-Si damping alloy when deformed in the <414> crystallographic direction. In-depth EBSD microstructural analysis was performed to unveil the mechanisms behind this remarkable ductility. A single crystal Fe-15Mn-10Cr-8Ni-4Si (mass %) alloy was produced at 1823 K in an inert gas environment using the Bridgman method. Quasi-static tensile tests were conducted in the <001>, <111>, and <414> crystallographic directions at an initial strain rate of 1.4 × 10−3 s−1. The strain was measured by a non-contacting video extensometer (TRView X, Shimadzu Corp.). EBSD microstructure analysis was performed on the deformed specimen's normal planes after surface preparation, which involved electrolytic polishing using perchloric acid solution. A JEOL JSM-7900F field-emission scanning electron microscope was used for EBSD analysis at 20 kV acceleration voltage, with a 15 mm working distance and a step size of 50 nm to 1 μm. Figure 1 illustrates a detailed orientation-dependent mechanical property of the Fe-15Mn-10Cr-8Ni-4Si alloy. The engineering stress-strain curves for all three different orientations are depicted in Fig. 1(a). Notably, the <414> oriented single crystal exhibits  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 extraordinary ductility, approximately 156%, which is approximately twice as high as that of the <111> crystal (74%) and nearly three times larger (49.8%) than the crystal oriented along the <001> direction. The maximum ultimate tensile strength (UTS) of 520 MPa and the minimum UTS of 350 MPa are observed for the <111> and <414> oriented single crystals, respectively. Noteworthy, in the <414> orientation, a significant drop in the yield stress is observed, which markedly differs from the other directions. The strain hardening behavior of tensile-deformed single crystals is illustrated in Fig. 1(b, c). During the initial stage of deformation, the <001> orientation exhibits a maximum strain-hardening rate (SHR, dσ/deT), while the <414> direction demonstrates the minimum. In the <001> orientation, the SHR quickly peaks, leading to a shorter initial stage. Around true strain, eT ≈ 0.08, it transitions directly to stage 3 hardening, where the rate gradually decreases until necking at eT ≈ 0.37. The SHR vs. true strain (eT) curve for the <414> crystal is shown in Fig. 1(c). Initially, the curve indicates easy glide up to eT ≈ 0.13 (region-II). This is followed by a significant secondary hardening stage (region III), with the SHR gradually increasing to a maximum at eT ≈ 0.92. In the <111> orientation, strain hardening also shows three distinct regions similar to the <414> orientation. However, region III has higher strain hardening, extending to a shorter true strain of about 0.52. Furthermore, the Considère criterion for plastic instability (dσ/deT ≤ σ) is met only for crystals oriented along the <001> direction. Figure 1(d) provides a comprehensive comparison of the strength-ductility trade-offs observed across commonly used single and polycrystalline materials. Remarkably, the <414> oriented specimen exhibits an exceptional strength-ductility balance surpassing the 50 GPa% threshold.  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65  Figure 1. (a) Engineering stress-strain curves, and (b) strain-hardening rate/true stress vs. true strain curves of tensile deformed specimens in the [414], [111], and [001] directions. (c) Work hardening stages in [414] oriented tensile deformed single crystal. (d) Comparison of the strength-ductility trade-off reported in various steels and the current study. Figure 2 presents a detailed EBSD microstructural analysis of a tensile-fractured single crystal oriented in various directions. It includes a colored phase map for three different orientations, depicted in Fig. 2(a), 2(c), and 2(e), and their respective inverse pole figure (IPF) maps in Figure 2(b), 2(d), and 2(f). Following tensile deformation, martensitic transformation occurs in all crystallographic directions, albeit with a highly orientation-dependent nature. The orientation-dependent martensitic transformation can be predicted using Schmid’s criterion with Thompson’s tetrahedron, as shown in Fig. 2(g).  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 Figure 2(h) shows all potential shear systems for martensitic transformation within the [001]–[101]–[111] standard triangle, indicated by arrows pointing towards or away from the gravity center relative to the tensile loading direction.  Figure 2. (a, c &e) EBSD phase map and (b, d &f) corresponding inverse pole figure (IPF) map of [414], [111], and [001] oriented tensile fractured single crystals, respectively. (g) Illustration of regular Thompson’s tetrahedron. (h) Plausible operational shear systems (Shockley partials) in the standard [001]-[101]-[111] triangle under tensile deformation.  Table 1 displays the Schmid factors for the operative martensitic shears in each of the three crystal directions. In the <414> orientation, two primary partials demonstrate  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 notably high Schmid factors of 0.5 (Dα/a) and 0.39 (δA/𝑑̅). Similarly, the three shears in <111> and eight shears in <001> orientation display equal Schmid factors of 0.31 and 0.24, respectively.  Table 1. Plausible γ → ε shear systems (partial dislocations) in various crystallographic orientations and their corresponding Schmid factor values. Orientation Operational shear systems (Schmid Factor)  001 Cα/a: 0.24, Dα/a: 0.24, Dβ/b: 0.24, Cβ/b: 0.24, γA/c̅: 0.24, γB/c̅: 0.24, δA/d̅: 0.24, δB/d̅: 0.24 111 Dα/a: 0.31, Dβ/b: 0.31, Dγ/c: 0.31 414 Dα/a: 0.50, δA/𝑑̅: 0.39 After tensile fracture, approximately 46% single-variant ε-martensite is observed in the <414> oriented specimen (Fig. 2(a, b)). In the <111> tensile-deformed microstructure, a similar volume fraction of ε-martensite (47 ± 3%), consisting of two variants (εa, εc), is observed (Fig. 2(c, d)). Unlike <414>, Unlike <414>, the <111> direction also shows an additional α’-bcc martensite phase. It has been established that α-bcc martensite is observed at specific sites where two different ε-martensite variants intersect each other [20,28]. The IPF map in Fig. 2(d) reveals the formation of multiple Kurdjumov–Sachs (K–S) variants at the ε-ε martensite intersection: V22 ((1 1 1̅)γ // (0 1 1)α’, [1 0 1]γ // [1̅ 1̅ 1]α’) and V23 ((1 1 1̅)γ // (0 1 1)α’, [1 0 1]γ // [1̅ 1̅ 1]α’). The  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 determination of K–S variants at these intersection sites is discussed in detail in our previous paper [28]. In the <001> orientation, about 21% ε-martensite is observed in the deformed microstructure, with no α’-bcc phase present. To investigate the microstructural evolution and deformation mechanism of the <414> single crystal, tensile tests were halted at 20%, 100%, and 156% (fractured strain). Figure 3 presents the superimposed phase map with the image quality (IQ) map, IPF map, and kernel average misorientation (KAM) maps for these strain levels. A lamellar banded structure comprising a single ε-martensite variant (green) aligned in the primary operative shear (Dα/a) direction is observed along the parent γ-austenite phase (white). At 20% strain, only 3 ± 1% ε-martensite is detected in the microstructure (Fig. 3(a)). The IPF map (Fig. 3(b)) shows a single color with a slight gradient, revealing minimal angular rotation of the γ-phase. At 100% strain, ε-martensite reaches its maximum volume of about 45 ± 5%. The deformation twins in the microstructure increases with higher strain levels.  In addition to uniform deformation, significant localized shear band (SB) formation occurs in the <414> crystal during later stages of deformation. These SBs intensify with increased tensile straining, as depicted in Fig. 3 (d-i). SBs generate severe strain fields, resulting in heightened misorientation and crystal rotation, as demonstrated in the KAM map (Fig. 3(f, i)). High magnification phase and IPF maps of a highlighted region (Fig. 3(j-k)) show very fine Σ3-type γ-twins (pink) at the γ/ε interface, and {101̅2} type fine ε-twins growing in strained ε-martensite region. The IPF map (Fig. 3(k)) reveals two distinct colors (yellow and red) representing the 86˚ (+ve twin) and 94˚ (-ve twin) ORs with the original ε-martensite phase. Corresponding pole figures (Fig. 3(l-m))  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 establish the Shoji–Nishiyama crystallographic orientation relationship among the γ-ε phases: {1 1 1}γ // {0 0 0 1}ε, <1 0 1̅>γ // <2 1̅ 1̅ 0>ε.  Figure 3. a, d, g) IQ + phase maps, (b, e, h) IPF maps, and (c, f, i) kernel average misorientation (KAM) maps of [414] single crystal specimens with tensile strains of 20%, 100%, and fractured strain (156%), respectively. (j) High magnification IQ + phase map,  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 (k) highlighted IPF map, and (l-m) corresponding pole figure maps showing the orientation relationship between γ-austenite (γ), γ-twin, and ε-martensite variant c (𝜀𝑐). In addition to uniform deformation, significant localized shear band (SB) formation occurs in the <414> crystal during later stages of deformation. These SBs intensify with increased tensile straining, as depicted in Fig. 3 (d-i). SBs generate severe strain fields, resulting in heightened misorientation and crystal rotation, as demonstrated in the KAM map (Fig. 3(f, i)). High magnification phase and IPF maps of a highlighted region (Fig. 3(j-k)) show very fine Σ3-type γ-twins (pink) at the γ/ε interface, and {101̅2} type fine ε-twins growing in strained ε-martensite region. The IPF map (Fig. 3(k)) reveals two distinct colors (yellow and red) representing the 86˚ (+ve twin) and 94˚ (-ve twin) ORs with the original ε-martensite phase. Corresponding pole figures (Fig. 3(l-m)) establish the Shoji–Nishiyama crystallographic orientation relationship among the γ-ε phases: {1 1 1}γ // {0 0 0 1}ε, <1 0 1̅>γ // <2 1̅ 1̅ 0>ε. Figure 4 displays the optical and SEM images of the tensile fractured <414> single crystal specimen. Macroscopically, SBs and primary slip traces along the <112> direction are visible on the normal direction (ND) plane (Fig. 4(a)). However, the final crack propagation occurred on a different plane along the <011> direction (Fig. 4(b)). The fracture surface, aligned approximately 52˚ to the <414> tensile axis, exhibits ductile fracture with fine and shallow dimples. Fig. 4(c) reveals densely packed dimples, ranging from sub-micron to a few micrometers, and proliferating shear band zones in the central fracture surface.  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65  Figure 4. Optical and SEM micrographs of the tensile fractured <414> oriented single crystal. PS: primary slip band; SB: shear band The Fe-15Mn-10Cr-8Ni-4Si single crystal oriented in the <414> direction demonstrated remarkable ductility (156%), nearly doubling that of its polycrystalline alloy counterpart, while maintaining its tensile strength. The yield strength (σy) of a single crystal can be estimated using equation (1), which considers the critical resolved shear stress (τCRSS) and maximum Schmid factor (m) in a specific crystal direction among all operational shear systems [21].  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65                              𝜎𝑦 =𝜏𝐶𝑅𝑆𝑆𝑚                              (1) In the Fe-Mn-Si alloy, the Schmid factor is lowest (m = 0.31) in <111> and highest (m = 0.5) in the <414> direction for a given τCRSS. Therefore, the highest σy is expected for the <111>, while the lowest σy is expected at <414> crystal orientation. Multiple slip activation results in the highest strain hardening in the <001> orientation during the early stages of deformation. Conversely, the <414> direction experiences easy glide in the initial stage due to primary shears with a maximum Schmid factor (Dα/a = 0.5), resulting in minimal strain hardening.  The strain hardening response in the <414> orientation exhibits behavior similar to materials displaying TRIP/TWIP phenomena [22,29]. In the strain hardening curve (Fig. 1(d)), region-I corresponds to dynamic recovery caused by dislocation rearrangement, leading to a sudden drop in the hardening rate at the onset of plastic deformation. Region-II shows a sudden increase in the strain hardening rate, likely due to the activation of planar slip, the primary mechanism of initial stage deformation in the <414> orientation. The formation of ε-martensite and deformation twinning (γ-twin and ε-twin) primarily contributed to the large regime hardening observed in region-III in <414>. To validate the strain hardening mechanism of the <414> orientation, EBSD analysis was conducted at strain levels of 20% and 156%, as depicted in Fig. 5. At 20% strain, ε-martensite appears alongside the less rotated γ region, highlighted in green in the γ-IPF map (Fig. 5b). Additionally, to accommodate plastic strain, γ-twins form along the γ/ε interface, shown in red (Fig. 5b). The crystal rotation from <414> to <011>, depicted in the IPF triangle (Fig. 5e), suggests that martensitic transformation relaxes the rotation  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 through slip mechanisms. Despite the discrete presence of ε-martensite, continuous slip traces are evident in the (IQ + phase) map (Fig. 5a). Therefore, in the initial stages of deformation, plasticity is primarily driven by <011> (11̅1) planar slip.  Figure 5. Deformation axis rotation of [414] single crystal at the tensile strain of 20%, and fractured strain (156%). (f) Unique definition of the deformation axis in [001]-[101]-[111]. In the TF specimen, the <414> microstructure features a significant amount of ε-martensite and γ-twinning. As strain increases, the <414> tensile axis rotates towards <112> on the [001]–[111] boundary (Fig. 5e). Similar lattice rotation has been observed in several FCC materials under tensile loading [30–34]. Upon reaching the [001]–[111] boundary, <101> (1̅11) conjugate slip with a high Schmid factor, referred to as secondary slip, is activated along with the primary slip. Fig. 5(f) provides a detailed representation  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 of the multiple slip systems (primary and conjugates) and operative martensitic shears using the (001) pole figure. Suzuki et al. reported that the conjugate system with a lower Schmid factor can activate even if the crystal rotation keeps the tensile axis within the [001]–[101]–[111] standard triangle [30]. Due to the activation of secondary slip systems, localized SBs readily form in materials under high strain conditions [35]. These SBs result from a concentrated accumulation of dislocations, driven by non-uniform plastic deformation, which intensifies strain in specific regions. In addition to TRIP/TWIP phenomena, SBs causes lattice rotation towards the tensile axis, significantly impacting strain hardening [36,37]. In fcc single crystals, SBs drive the activation of conjugate slip systems during stage-II hardening [36]. Crystals with stable orientations are less prone to SB formation during deformation [37]. In <414> crystals, SBs enhanced strain hardening capacity by accommodating significant plastic strain. As dislocations move within SBs, they hinder dislocation motion, leading to strain hardening. Similarly, kink-bands were found to act as superior barriers to dislocation motion compared to grain boundaries, reducing localized failure at low strain levels and extending ductility in Mg-Zn alloys [38,39]. As strain increases, multiple shear bands interact, thereby amplifying strain hardening through the dynamic Hall-Petch effect. Additionally, the SBs in <414> enhanced ductility by providing channels for stress redistribution, thus delaying failure.  In summary, we discovered orientation-dependent exceptional ductility in a Fe-15Mn-10Cr-8Ni-4Si alloy single crystal, reaching approximately 156% in the <414> orientation. The strain hardening behavior in this orientation was characterized by planar slip and a sluggish γ → ε transformation during the initial stage, followed by the activation of multiple slip systems (secondary/conjugate slip) and shear bands due to crystal rotation in the later stage of deformation. The interaction of nanoscale twinning and shear bands  1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 resulted in extremely fine dimples, which contributed to slow crack propagation and delayed necking. Thus, the remarkable ductility demonstrated by the <414> orientation is attributed to the combined effects of planar slip, slow martensitic transformation kinetics, and the formation of ultra-fine hierarchical structures (twins and SBs). These factors synergize to prolong strain hardening and contribute to the exceptional ductility observed in the material. These findings open new possibilities for fabricating ultra-high ductility alloys.  Acknowledgments The authors acknowledge the support from the Japan Society for the Promotion of Science (JSPS-21H01659, JSPS-23K13227). The authors wish to express their gratitude to Irina V. Kireevad, and Yuriy I. Chumlyakov for providing us for single crystal specimens.            1  2  3  4  5  6  7  8  9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 References: [1] A. Sato, E. Chishima, K. Sonlap, T. 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