Introduction

Rolling is one of the most commonly employed industrial metal forming/deformation processes for continuous production of bulk sheets. Rolling at room/ambient temperature is referred as cold-rolling. Cold-rolling (CR) deformation results in generation of large quantity of defects across various length scales like stacking faults, dislocations, microbands and shear bands. Often, the one that stands-out among the observable microstructural defects is the shear band (SB). Shear bands are macroscopic defects that run across several grains or sometimes throughout the thickness of the material irrespective to the crystallographic plane alignment in the neighboring grain. SB are most commonly observed in face-centered cubic (FCC) materials having low stacking-fault energy (SFE)1. When the SFE is low, {111}<110 > dislocations tend to dissolve into two Shockley partial dislocations having mixed nature of screw and edge dislocations leading to a suppressed dislocation cross-slip. Suppression of cross-slip renders the dislocation to planar slip (restricted to a very specific crystallographic plane) leading to plastic instability in the flow behavior2. Plastic instability leading to flow localization has been attributed to the origin of shear bands1. Formation of SBs of this nature is often characterized as 18° ~ 45° inclined intensely strained band with respect to the rolling direction3. It has been well documented that the pure metals having moderate- to high- SFE are resistant to shear localization due to the pronounced dislocations cross slip, thus SB formation was rarely seen among these metals4,5.

On the other hand, body-centered cubic (BCC) materials are known to have poorly defined slip plane as screw dislocations core is spread on multiple planes along < 111 > slip direction. Due to this, screw dislocation undergoes significant cross-slip, for example, among {112} and {110} leading to a wave-like slip band6. Along with the possibility of dislocations cross-slip, high strain-rate sensitivity of BCC materials aids prolonged uniform deformation7. Thereby, SBs are least expected during cold rolling of BCC materials at least for the moderate strains unlike FCC materials. Senkov et al.8, has introduced a new class of high-entropy alloys (HEAs)9,10,11 based on refractory elements, often referred to refractory HEAs12,13,14,15,16. These alloys typically exhibit BCC structure, with the equiatomic HfNbTaTiZr alloy being a prominent example of a single-phase BCC HEA that has garnered significant attention17,18. Senkov and Semiatin19and Eleti et al.20, investigated cold rolling/formability of HfNbTaTiZr BCC-HEA and found the alloy can be heavily cold rolled (CR) until 86% without any major cracks, yet lacked the normal strain-hardening. Furthermore, their study has also revealed a large fraction of shear bands above CR65%. Texture softening is often considered one of the possible factors contributing to the steady-state hardening behavior observed in metallic materials during heavy rolling21. Eleti et al20. and Senkov and Semiatin19have systematically investigated the evolution of deformation texture in the HfNbTaTiZr BCC-HEA. Eleti et al20. reported that the alloy exhibits typical RD- and ND-fiber texture components during deformation, with no significant enhancement in texture intensity even under heavy rolling, CR90%. Similarly, Senkov and Semiatin19 observed the formation of typical RD-fiber components and found an estimated Taylor factors of M = 3.05 and 3.09 for 65% and 85% cold-rolled conditions, respectively. These findings collectively indicated lack of texture softening in the HfNbTaTiZr BCC-HEA during cold rolling. Having said that, however, the origins of shear band formation in the HfNbTaTiZr BCC-HEA has been not understood, so far. Also to note, formation of shear bands has an intricate connection to the dislocation (slip band) interaction with the grain boundaries (GB)22,23,24. By combining advanced characterization techniques such as focused-ion beam (FIB) and transmission electron microscope (TEM), a detailed site-specific analysis along the GB region may offer deeper insights into the microstructural evolution of shear bands in the BCC-HEAs. Therefore, in the present study, we try to investigate the origins of shear band formation in the HfNbTaTiZr BCC-HEA and address the intricacies of strain accommodation mechanism along the GB region.

Experimental procedure

An equiatomic HfNbTaTiZr alloy was vacuum arc melted using constituent pure metals having at least 99.95% (wt%) purity. The alloy ingot was vacuum sealed and homogenized at 1200 °C for 6 h, followed by water quench. Phase characterization of the homogenized alloy was analyzed using X-ray diffraction (XRD) technique with a Cu-Kα source. The homogenized alloy ingot having dimensions 30 mm (length) x 10 mm (width) x 5 mm (height) was cold rolled (CR) until various initial height reductions. The microstructural observation was carried out at the center of the thickness plane of the cold rolled samples using a field emission-scanning electron microscope (FE-SEM, ThermoFisher-Apreo) attached with an electron backscatter diffraction (EBSD) at a step size 0.2 μm until CR35% reduction and a step size of 0.1 μm for CR50% and CR65% reductions. EBSD analyses was also complemented using backscatter electron (BSE) and electron channel contrast imaging (ECCI) detecting systems. The deformed samples were subjected to colloidal silica suspension polishing using fully automated vibratory polishing machine for 15 h (Buehler – VibroMet 2). Site-specific transmission electron microscopy (TEM) analyses were carried out on the deformed HfNbTaTiZr BCC-HEA specimens to clarify the local deformation substructures formed near the slip bands and GB intersections. Two representative regions along the GB, exhibiting distinct slip band impingement behavior, were selected based on prior SEM and EBSD observations. These regions were marked in situ within the dual-beam focused ion beam (FIB-FEI, Scios) system for subsequent TEM lamella extraction. TEM specimens were prepared using a site-specific lift-out method to ensure that the regions of interest were precisely retained within the final electron-transparent area. A wedge-shaped lift-out geometry was adopted to minimize curtaining effects and to achieve a gradual thickness gradient suitable for high-resolution imaging. The rough milling was initially performed at 30 kV with beam currents ranging from 0.3 nA to 100 pA, followed by successive low-kV polishing steps. The final thinning was conducted at 30 kV and 49 pA to achieve an electron-transparent region (~70 nm), and a subsequent fine polishing at 5 kV and 49 pA was employed to remove amorphous surface layers and mitigate Ga⁺ implantation and beam-induced damage. High-resolution TEM and selected area electron diffraction (SAED) analyses were performed using a JEOL ARM-200 F microscope operated at 200 kV to characterize the dislocation substructures, within the slip band and GB intersection regions.

Results

Figure 1 shows microstructure and X-ray diffraction (XRD) profile of the homogenized HfNbTaTiZr alloy. Figure 1(a) shows the SEM-BSE image composed of near-equiaxed grains having an average grain size, d ~ 140 μm. Figure 1(b) shows the corresponding XRD profile suggesting a single-phase BCC structure. The homogenized HfNbTaTiZr alloy was subjected to cold rolling until various strains, and micro-hardness measurements and microstructure evolution were investigated at the center of the cross-sectional area of the rolled sheets, as shown in the schematic Fig. 1(c). The evolution of micro-hardness during cold rolling of HfNbTaTiZr is shown in Fig. 1(d). Prior to cold rolling, the HfNbTaTiZr alloy showed high hardness of 380 Hv. Hardness of the alloy increased with the increase of cold rolling reduction, until CR35% (464 Hv) followed by a flow softening-like hardness reduction. However, further reduction of CR50% or more showed a similar hardness 445±4 Hv values indicated a steady state. The unusual aspects of increased hardness followed by a flow softening during cold rolling is further investigated using detailed microstructural analyses.

Fig. 1
figure 1

Characterization of equiatomic HfNbTaTiZr high-entropy alloy (HEA). (a). SEM-BSE image, and (b) XRD profile of HfNbTaTiZr HEA after homogenization at 1200 °C for 6 h. Microstructural observation revealed a typical single-phase microstructure having an equi-axed grains with an average grain size, d~40 μm. Complimentarily, XRD analysis too confirmed a single-phase body-centered cubic (BCC, β) structure. (c) Schematic illustration of a rolled sheet subjected to further characterizations on the thickness plane. RD = rolling direction, ND = normal direction, TD = transverse direction. (d) Evolution of micro-hardness as a function of cold rolling reduction.

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Figure 2 shows EBSD and BSE images corresponding to microstructure evolution of HfNbTaTiZr alloy during cold rolling until various strain intervals. Figure 2(a-c) shows EBSD-image quality (IQ) maps revealing the deformed microstructures. Adding to that, BSE/ECCI imaging was used to understand further subtle details to investigate the microstructure evolution during cold rolling. Figure 2(a′-c′) shows BSE/ECCI images of HfNbTaTiZr alloy corresponding to the EBSD maps shown in Fig. 2(a-c), respectively. The microstructural investigations revealed unusually straight, coarse and widely spaced slip bands during early stages of plastic deformation (refer, Fig. 2a′). Generally speaking, slip bands in BCC crystals tend to show wave-like morphology due to the repeated cross-slip of screw dislocation between {112} and {110} slip planes. Therefore, the possibility of straight slip bands in the current BCC HfNbTaTiZr alloy is rather unusual. Further deformation has revealed an unusual deformed microstructure evolution during cold rolling of HfNbTaTiZr BCC-HEA. A detailed, and closer examination of the straight slip bands indicated slip band and GB intersection regions. Furthermore, it is also evident that the slip bands cut through the GB regions leading to translation of slip across the GB into the neighboring grain, refer Fig. 2(b, b′). Adding to that, formation of steps or ledges were observed to have formed at CR35% whence the slip band translates the GB and an increased propensity of ledges along the GB with an increased continuous slip across the GB regions, as shown in red arrows of Fig. 2(c,c′). Rather, the findings are quite striking for various reasons. Firstly, GBs are structural defects where the continuity of crystallographic planes is broken. Meaning, dislocations slip does not ensure a continuous slip across the GB due to the lack of slip plane continuity into the neighboring grain, thus leading to dislocations pile-up along the GB regions. Due to this, excessive stress concentrates along the GB region which is often released in the form of activating dislocation sources in the neighboring grains (also referred to strain transfer). Having said that, lack of dislocation sources in the neighboring grains promote crack initiation along the GB unto excessive stresses are released. Surprisingly however, we have observed straight slip bands cutting through the GB regions into the neighboring grains during deformation of HfNbTaTiZr BCC-HEA2.Although the present HfNbTaTiZr BCC-HEA showed less hardening with the increase of cold rolling reduction, flow localization/softening was observed above a moderate CR35% deformation. Incidentally, a detailed microstructural investigation has shown an excessive strain localized regions at CR35%, as shown in Fig. 2(c-c′). The micro-hardness measurements support the microstructural observations, indicating strain localization along shear bands, which results in a lack of hardening or a steady-state-like hardening behavior. This suggests that the material exhibits a propensity for localized deformation, limiting its ability to undergo uniform strain hardening. To confirm the GB character against the observed continuous slip observed in Fig. 2(b, b′), further EBSD analyses is shown in Fig. 3. Both the grains highlighted as G1 and G2 corresponding to the BSE image of Fig. 2b′ have a totally different orientation across the GB indicating the high-angle GB nature. A misorientation line profile analysis is shown in Fig. 3b indicated both the grains G1 and G2 are separated by a high-angle boundary having a misorientation of 52°. This clearly shows that the continuous slip across the GB is one of the most important stages of shear bands evolution during plastic deformation of polycrystalline materials. Later, further increase of plastic deformation has clearly indicated that the prior coarse slip bands have evolved into shear bands, as shown in Fig. 4. It is noteworthy that microstructures consisting of shear bands are mere characteristic features of FCC materials having low–SFE are observed with striking similarities in the present HfNbTaTiZr BCC-HEA, particularly at moderate deformation, like CR 35%. Although SBs are less commonly expected in BCC systems at moderate deformation levels, their occurrence cannot be entirely excluded under conditions of large rolling reductions, particularly when influenced by specific crystallographic orientations or enhanced frictional shear at the roll bite. For example, SB formation during cold rolling of BCC interstitial-free (IF) steel has been reported under high reductions25. IF steels have been shown to develop SBs through conventional dislocation glide and rigid-body rotation mechanisms, which produce characteristic “S”-shaped microbands that subsequently evolve into fully developed SBs25,26,27. Such microband-assisted shear localization processes, commonly associated with strain partitioning and texture softening during large rolling reductions, were not observed in the present HfNbTaTiZr BCC-HEA. This indicates that the mechanism of SB formation in the current alloy is distinct and requires further investigation.

Fig. 2
figure 2

Microstructural evolution of HfNbTaTiZr BCC-HEA during cold rolling for various strain intervals. (a) CR5%, (b) CR10% and (c) CR35%. (a-c) EBSD image quality (IQ) map showing low- and high-angle grain boundaries (L/H-AGB). LAGBs with misorientation (θ), 3° ≤ θ < 15° are shown in red, while HAGBs with θ ≥ 15° are shown in blue lines. (a′-c′) High magnification area specific BSE/ECCI images of the corresponding microstructures shown in (a-c).

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Fig. 3
figure 3

Microstructure and misorientation profile analysis of grain boundary region in a cold rolled HfNbTaTiZr BCC-HEA, cold rolled until CR10%. (a) EBSD-IPF map showing crystallographic orientation parallel to the rolling direction (RD). (b) Misorientation profile among two grains G1 and G2 across a grain boundary, as shown in the inset of (a).

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Fig. 4
figure 4

Microstructural evolution of shear band formation in HfNbTaTiZr BCC-HEA during cold rolling until various stain intervals. (a, b) EBSD-IQ maps showing intense strain localization and shear bands (SB) formation at (a) CR50% and, (b) CR60%.

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However, the intricate process like slip transfer across the GB and the associated mechanism to facilitate strain compatibility along the GB regions in a BCC material has been not fully understood. We have performed additional slip trace analysis to understand the observed coarse slip bands and their corresponding slip systems. Of all the above analyses, CR10% has proved to foster an important step in the early stages of shear band formation, refer Fig. 2(b,b′). Therefore, further analysis has been concentrated on this microstructural condition to investigate the mechanics of shear band formation in the present HfNbTaTiZr BCC-HEA. A detailed explanation for the forthcoming Schmid factor analyses using slip traces is summarized in Table 1 and elsewhere28. Figure 5 shows the slip traces for both the differently oriented grains G1 (RD//[13 10 16]) and G2 (RD//[11 10 25]) having a continuous slip band with a Schmid factor (SF) of 0.42 (refer Fig. 5). For G1, the highest SF = 0.42 was found for the (121)[1–11] slip system, whereas for G2 the highest SF = 0.44 was found for the (011)[1–11], while SF = 0.41 was found for the (121)[1–11] slip systems, respectively (refer Table 1). Meaning, we would like to highlight here that despite the grain G2 orientation is highly favored for (011)[1–11] slip having the highest SF = 0.44 value, the current observations of slip trace and slip system analyses indicated that the (121)[1–11] slip system having relatively less SF = 0.41 is active. The implications of present findings suggest the grains having less SF value is active compared to the one having higher SF value. According to the fundamentals of crystal plasticity, it is well understood that the active slip system in a crystal/grain favors the orientation satisfying highest SF value. The present observation suggested slip transfer across the GB regions of HfNbTaTiZr BCC-HEA may favor non-Schmid slip systems.

Fig. 5
figure 5

SEM-BSE image showing microstructure evolution and slip band – grain boundary interaction in a HfNbTaTiZr BCC-HEA during cold rolling until, CR10%.

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Having said that, however, the above analysis does not capture the intricacies of ledge formation or strain compatibilities along the slip band and GB intersection regions that lead to shear bands formation. To address this aspect, here in the present study, we have performed additional site-specific FIB assisted TEM investigations on the deformed samples of HfNbTaTiZr BCC-HEA using samples extracted from the slip band and GB intersection regions. We chose two areas along the GB region to perform the above analysis, as shown in Fig. 6. The resulting TEM investigations revealed two major findings. One, nanotwins formation is evident along the GB ledge region. Two, the referred coarse slip bands were composed of long, and straight dislocations. The present results indicate that the coalescence slip of several dislocations through the grains and slip transfer across the GBs can nucleate a shear band when the instantaneous hardening rate drops, like CR35% condition. Further analysis using the diffraction vector, g = 112 as shown in the inset of Fig. 6(d), has revealed that all the dislocations were < 1 1 1 > type having the burgers vector parallel to the dislocation line. That is, screw dislocations were dominantly observed in the present HfNbTaTiZr BCC-HEA. The present results are consistent with the previous reports29,30,31, however, contrast to the recent reports of edge dislocations dominance in a few BCC-HEAs32,33.

Fig. 6
figure 6

Microstructural characterization on the early stages of shear band formation in HfNbTaTiZr BCC-HEA. (a) EBSD-IPF map of HfNbTaTiZr BCC-HEA after CR10%. (b) BSE/ECCI image of the corresponding area in (a). (c, d) TEM images of area specific observations as indicated in (b).

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Table 1 Summary of slip systems and the corresponding Schmid factor, SF values for grains G1 and G2 having orientations [13 10 16] and [11 10 25] respectively. Here, SF value was determined as the product of Dot product of the applied stress direction and slip plane, and the Dot product of the applied stress direction and slip direction. Slip systems resulting a negative SF value are not shown in the table below.
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Discussion

The prominence of shear band formation in FCC materials having low-SFE is essentially due to the planar dislocation slip. Here, it is noteworthy that ordered intermetallic compounds having FCC-type L12 structure were also reported favoring SB formation34. As is known, plastic deformation of L12 structure occurs due to the planar slip of superlattice dislocations35. In either of these cases, as is well understood that the planar slip of dislocations necessarily lead to dislocation pile-up and high shear stress intensity along the grain boundary regions. The dislocations pile-up directs the shear stress to concentrate along the GB and is theoretically estimated using an empirical relation as, shear stress intensity, ks36,

$$\:{k}_{s}=\frac{\pi\:\mu\:{b}^{1/2}}{4\alpha\:}\:$$
(1)

where, µ is shear modulus (GPa), b is Burgers vector (nm), and \(\:\alpha\:=2(1-\nu\:)/(2-\nu\:)\) with ν is Poisson’s ratio. For a large number of dislocations in a pile-up, Eshelby, Frank and Nabarro37 has theoretically estimated that the ks decreases at a rate of “one over square root distance” away from the dislocation pile-up head into the neighboring grain. This is analogues to the experimental measurement of grain refinement strengthening coefficient, kHP using the Hall-Petch relation. To understand the unique microstructure evolution of the present HfNbTaTiZr BCC-HEA in the context of kHP and ks, the present alloy is compared with various metals/alloys. Figure 7 shows the experimental and theoretical measurements of shear stress intensity of various metals/alloys.

Fig. 7
figure 7

Experimentally obtained Hall-Petch coefficient, kHP values of various metals/alloys against the theoretical shear stress intensity, ks. Nb 36, Al 36, Ni 36, Cu 36, CoCrFeMnNi46, HfNbTaTiZr47,48, Ni3Fe49. kHP values are normalized by the Taylors factor, M for an equivalent measurement against FCC (M = 3.06) and BCC (M = 2.7) structures. A dotted line in the above plot is merely an imaginary line.

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It is worth reminding that moderate- to high- SFE materials like Al, Ni and Nb undergo grain sub-division and extensive grain elongation along the rolling direction, with no traces of macroscopic shear localization or SB formation4. Incidentally, all these metals exhibit low kHP values. Adding to that, it is found that Nb and HfNbTaTiZr BCC-HEA showed relatively similar kHP values. Furthermore, an equiatomic CoCrFeMnNi FCC-HEA having low-SFE and Ni3Fe intermetallic compound were reported having high kHP values, as well as severe shear localization during cold rolling34,38. Despite the above discussion and relatively high strain-rate sensitivity, m of BCC-HEA (m = 0.017) compared to pure Nb (m = 0.011) (refer, Supplementary materials Figure S1), the present BCC-HEA showed profound SB formation, contradictory to the conventional understanding of shear localization phenomena. The present analysis was limited to a relatively small number of grains, primarily due to the large initial grain size of the alloy. Nevertheless, it is important to note that similar slip band and grain boundary interaction features were independently observed in both the CR10% and CR35% samples, thereby mitigating concerns regarding statistical limitations. Therefore, the above evidence suggests that the SB formation in the HfNbTaTiZr BCC-HEA is unconventional, and the role of shear stress concentration along GB regions in promoting SB formation can be reasonably excluded. Generally speaking, strain transfer across the GB may occur either by means of dislocation activity on the conjugate or secondary slip planes (here referred as possible strain transfer) or by the translation of slip across the GB (continuous/irreducible strain transfer). The analyses of Fig. 5 and Fig. 6reveal that the flow localization on {112} slip planes, resulting from the coalescence of dislocation slip bands, promotes the SB formation along this slip system. The present systematic investigation revealed that the screw dislocations of HfNbTaTiZr BCC-HEA suffer the possibility of cross-slip from {112} onto {110} planes along the < 111 > directions. Consequently, the dislocations slip bands appear straight. As of energetics viewpoint, dislocation slip on {112} planes of Schmid and non-Schmid oriented grains suggest the screw dislocations core structure of HfNbTaTiZr BCC-HEA is non-degenerate. The present analogy, according to Bassani et al.39,, and Duesberry and Vitek40 supports the significant contribution of screw dislocations slip in HfNbTaTiZr BCC-HEA, as opposed to edge dislocations17,28,29,30. Figure 8 shows the schematic illustration of dislocation slip and the possibilities of strain transfer that lead to shear band formation in the HfNbTaTiZr BCC-HEA. As shown in Fig. 8(a, b), dislocation slip and subsequent pile-up occur along the initial grain boundary (GB). The transmission of such slip across the GB, in an unconstrained or translational manner, necessitates the continuity of strain between neighboring grains. Livingston and Chalmers41 proposed that the accumulated shearing stress of the incoming dislocations shall be relieved by the outgoing (strain transfer) slip system experiencing the maximum shear stress. On the contrary, based on the current findings, the strain transfer has been favored by the non-Schmid effects This can be achieved provided the incoming and outgoing dislocations maintain the continuity of slip direction across the GB as in simple shear deformation42 thus, leading to the formation of a ledge/step on the GB surface, as schematically illustrated in Fig. 8(b, c). Once the GB surface is sheared through by the outgoing dislocations into the neighboring grain, further slip would become localized along these prior slip bands which may eventually lead to the SB formation, as in Fig. 8d. Adding to that, Li et al44. systematically investigated the deformation twinning in pure Nb, an BCC structure, and found the presence of stepped twin boundaries (with a thickness of 3n layers) exacerbates the shear strain near the interface. Twins, as observed in the present study along the GB regions can function as local stress concentration points, facilitating the nucleation and propagation of shear bands assisted by the non-Schmid effects. The present analogy nullifies the possibility of excess defect generation along the GB surface, which also further favors the lack of additional strain-hardening as observed in Fig. 1(d).

Fig. 8
figure 8

Schematic illustration of shear band formation mechanism in a body-centered cubic HfNbTaTiZr high-entropy alloy during cold rolling.

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It is acknowledged that the evolution of critical resolved shear stress (CRSS) and its coupling with the influence of non-glide stresses on dislocation motion can further influence non-Schmid effects in BCC alloys; however, such quantitative analysis lies beyond the scope of the present experimental study. For instance, non-Schmid effects such as, tension–compression asymmetry and non-glide stress effects on dislocation motion are well established in BCC metals. Consequently, the system with the highest Schmid factor does not necessarily activate first44,45. Thus, the present observation of non-Schmid behavior is not unique to the HfNbTaTiZr BCC-HEA. Nevertheless, the possibility of nano-twin formation as a local strain-accommodation mechanism under non-Schmid stress states may represent a critical distinction influencing SB evolution. The nucleation of nano-twins can effectively relax local stress concentrations, promote strain transfer across grain boundaries, and stabilize dislocation activity within confined regions. This cooperative interaction between non-Schmid slip and nano-twin formation is likely to facilitate strain localization, ultimately leading to the observed SB formation. Therefore, the present study suggests that shear localization in the HfNbTaTiZr BCC-HEA results from the synergistic interplay between non-Schmid deformation behavior and nano-twin mediated strain accommodation.

Conclusions

In this study, the microstructural origins of shear band (SB) formation were systematically investigated at various strain intervals during cold rolling of an equiatomic HfNbTaTiZr body-centered cubic (BCC) high-entropy alloys (HEA). A detailed microstructural analysis revealed that intense macroscopic strain localization emerged at deformation levels above CR35%. With further deformation, well-developed macroscopic shear bands were evident even at moderate strain levels (CR35–50%). The mechanistic investigation of SB formation indicated the operation of continuous dislocation slip propagating from one grain into its neighboring grain. Moreover, non-Schmid effects were found to promote translational slip across grain boundaries (GBs). The shearing of screw dislocations through GB regions generated GB ledges and localized strain accumulation, which eventually evolved into shear bands. Subsequently, nano-twins were observed along the GB ledge regions, facilitating stress relaxation and accommodating plastic incompatibilities at slip band–GB intersection sites. Overall, the present study establishes that shear localization in the HfNbTaTiZr BCC-HEA arises from the synergistic interaction of non-Schmid deformation behavior and nano-twin nucleation along grain boundaries. These findings provide new insights into the fundamental deformation mechanisms governing shear band formation in BCC-HEAs.