Plasticity in brittle intermetallics enabled by framework of amorphous interfaces and preexisting dislocations

Abstract

Intermetallics are highly attractive for their exceptional strength and high melting points, offering significant potential as advanced structural materials. However, their inherent brittleness at room temperature severely limits practical applications. In this work, we introduce a structure of framework of amorphous interfaces (FAIs) and preexisting dislocations into nanocrystalline (NC) CoAl intermetallics to synergistically enhance both strength and plasticity. Micropillar compression tests reveal a high yield strength exceeding 6 gigapascals, a sustained work hardening to approximately 8.5 gigapascals, and a compressive plastic strain exceeding 15%. The FAIs accommodate the plastic deformation of NC CoAl grains, preventing intergranular fracture while promoting dislocation emission and propagation into CoAl through deformation-induced crystallization. Molecular dynamics (MD) simulations confirm that dislocations are emitted from crystalized regions (BCC-like local motifs) and reveal that preexisting dislocations impede dislocation motion via interactions and multiplication, promoting dislocation storage. Together, these mechanisms enable enhanced work hardening and large plasticity. This strategy offers an approach to achieving room temperature plasticity in brittle materials, which often show limited dislocation activity.

Amorphous interfaces enable plasticity in intermetallics at room temperature.

INTRODUCTION

Intermetallics play a critical role in a wide range of applications such as jet engines, gas turbines, and automotive systems (1–3). Their remarkable strength, high melting temperatures, and superior creep resistance make them indispensable for structural usage in extreme environments (1–7). The superior properties intrinsically arise from a unique bonding nature, a combination of metallic, covalent, and ionic bonds (1, 6). However, the same bonding characteristics result in intrinsic brittleness, manifested as low fracture toughness and limited plasticity, which severely restricts their applications (1). This challenge stems from the difficulty of activating plastic deformation via dislocation activities, including nucleation, glide/slip, and multiplication, which are manifested as high critical resolved shear stress required for dislocation nucleation, high frictional stress that resists dislocation glide/slip, and the lack of sufficient independent slip systems (8–11). As a result, dislocations are neither favorably formed nor multiplied, which not only limits plastic deformation but also suppresses work hardening, a key mechanism that typically delays fracture in ductile materials. Consequently, work hardening behavior is rarely observed in intermetallics (1). In addition to limited dislocation activities, intermetallics also struggle to accommodate plasticity through the formation of planar defects such as stacking faults and antiphase boundaries (APBs), because of the high energy barriers associated with their formation, which limits alternative deformation pathways in intermetallics (11–13). Moreover, grain boundaries (GBs) could also be another origin of fracture and brittleness in intermetallics. For example, brittle intergranular fracture is dominant in polycrystalline L1₂ Ni₃Al and B2 NiAl intermetallics mainly due to poor GB cohesion (14, 15). This issue is further compromised by moisture-induced hydrogen embrittlement, particularly in aluminide intermetallics, where hydrogen generated from environmental moisture can penetrate along GBs, and ultimately resulting in brittle fracture (16, 17).

To overcome these intrinsic and microstructural limitations, extensive studies have been devoted to achieving synergy between high strength and plasticity in intermetallics but with limited success. Some of the classical approaches are briefly summarized here. (i) GB chemistry modification: Microalloying is used to enhance GB cohesion and overcome GB embrittlement. For instance, boron (B) additions are proven to be effective in improving plasticity in various intermetallics like L1₂ Ni₃Al and B2 FeAl via suppressing intergranular fracture (14, 18, 19). However, this approach appears to be system specific. George et al. (15) reported that B-dopant is not effective in improving GB cohesion in B2 NiAl intermetallics. (ii) Composites: Metal/intermetallic composites have been explored across various systems, like cellular NiAl/Cr (Mo and Hf) (20–22), laminated NiAl/Ni₃Al/Ni (23), and laminated stainless steel/FeAl (24). These composites showed high strength and adequate plasticity. However, most of the deformation was sustained via the ductile metal, and the intermetallics remain largely undeformable at room temperature. (iii) Controlling planar defect energy: Recent studies on the B2 RM compounds (where R is a rare earth and M is a main group or transition metal) revealed some ductility in systems such as YAg, YCu, and DyCu, compared to other brittle B2 intermetallics (25). The 1/2 <111> mobile dislocations coupled with APBs are observed to accommodate plasticity. Atomistic simulations of APB energy further support this mechanism, suggesting the low energy barrier of forming APBs in B2 RM intermetallics (25). However, the high cost and low strength hinder their broader applicability. (iv) Transformation/twining induced plasticity: B2 to martensite B33 phase transformation is observed in Zr-Co-Ni alloys, showing plasticity enhancement compared to Zr-Co alloys only with B2 intermetallics (26–28). Mn is found to be beneficial to lowering the stacking fault energy in L1₀ TiAl alloys, thereby promoting deformation twinning in L1₀ TiAl alloys and accommodating plasticity (29, 30). However, similar behavior is rarely observed in other systems. (v) High entropy intermetallic (HEI) alloys: HEI is a recently proposed concept that is characterized by numerous principal elements, long-range chemical order, and elemental random mixing. A single-phase B2 ordered HEI with the chemistry of (CoNi)₅₀(TiZrHf)₅₀ was recently designed by Wang et al. (31), where the highly distorted lattice with complex chemical order promotes multiple slip systems and high flow stress. Although these attempts show some improvement in plasticity, uniform deformation and work hardening capability remain absent in these intermetallics.

Given the limited success of the above-mentioned strategies, the extreme refinement of grain size might also emerge as an alternative route to enhance the mechanical performance of intermetallics (32). The approach of nanocrystallization has been proven to be efficient in strengthening metallic materials by orders of magnitude (33). However, dislocation activities and accumulation are suppressed by GBs, resulting in deteriorated plasticity and work hardening capability of nanocrystalline (NC) metallic materials (34). Moreover, GBs in NC metals promote creep and GB decohesion, further diminishing plasticity (35, 36). As a result, achieving high strength and plasticity synergistically has been a long-standing pursuit for enhancing the mechanical properties of NC metallic materials. This challenge becomes even more pronounced in NC intermetallics, considering their brittleness and susceptibility to GB decohesion.

Our study demonstrates an approach to achieving simultaneous high strength and plasticity in NC intermetallics via introducing a framework of amorphous interfaces (FAIs) and preexisting dislocations into B2 CoAl intermetallics, one of the most brittle B2 aluminide intermetallics (12, 37). In situ micropillar compression tests reveal that NC CoAl with FAIs and preexisting dislocations has an ultrahigh yield strength of ~6 GPa and sustains prominent work hardening to a flow stress of ~8.5 GPa with a compressive plastic strain of more than 15%. This performance positions the NC B2 CoAl in our work among one of the most deformable B2 intermetallics ever reported. Transmission electron microscopy (TEM) studies along with molecular dynamics (MD) simulations further reveal that this rare synergy of high strength and room temperature deformability is driven by the cooperative effect of FAIs and preexisting dislocations. The amorphous interface effectively facilitates dislocation nucleation via deformation-induced crystallization (DIC) and subsequent propagation into CoAl to accommodate plastic deformation. In addition, the FAIs avoid susceptibility to intergranular fracture and promote uniform deformation. Simultaneously, the introduction of preexisting dislocations toughens CoAl intermetallics, rendering them largely deformable. These artificially introduced dislocations impede dislocation movement through trapping, interaction, and multiplication, thereby enhancing dislocation storage capability and promoting significant work hardening.

RESULTS

FAIs and preexisting dislocations

Figure 1 shows TEM micrographs and energy dispersive spectroscopy (EDS) analyses of NC CoAl/FAIs. Bright-field (BF) TEM image in Fig. 1A outlines the stacking of CoAl/ FAIs. The inserted selected area diffraction (SAD) pattern confirms the formation of polycrystalline B2 phases. High-resolution TEM (HRTEM) image along the interface between CoAl and FAIs is shown in Fig. 1B. The inserted fast Fourier transform (FFT) analyses within CoAl and FAIs show B2 [100] zone and amorphous ring separately. The zoom-in image within FAIs showing amorphous structure is further displayed in Fig. 1C. The high-angle annular dark-field (HAADF) scanning transmission electron microscopy (STEM) image and corresponding EDS map scan in Fig. 1 (D to F) and fig. S2 reveal the horizontal and vertical Al-rich FAIs. The high-resolution STEM (HRSTEM) image overlapped with EDS mapping around one node of FAIs in Fig. 1G and fig. S3 shows that the chemistry of FAIs is an Al-rich binary Al-Co alloy. The EDS line scan across FAIs and CoAl in Fig. 1H confirms that the chemistry of FAIs is Al–25 atomic % (at %) Co. Figure 2 shows the TEM and ASTAR analyses of preexisting dislocations for single-crystal (SC) CoAl and NC CoAl/FAIs. SC CoAl is mostly dislocation free as shown from the BF TEM image in Fig. 2A [along the [110] zone which can be seen from Inverse Pole Figure (IPF) in Fig. 2B]. The Kernel average misorientation (KAM) and corresponding geometrically necessary dislocation (GND) maps in Fig. 2 (C and D) illustrate the absence of GNDs. In contrast, high-density preexisting dislocation segments are identified in CoAl of NC CoAl/FAIs, from BF and dark-field (DF) TEM images in Fig. 2 (E to G). The KAM and extrapolated GND maps in Fig. 2 (H and I) further reveal the existence of high-density GNDs before deformation. The GND density measured from ASTAR and total dislocation density estimated via the line intercept method on DF images for NC CoAl/FAIs are summarized in Table 1. An HRTEM image along the [110] zone and corresponding inverse fast Fourier transform (IFFT) analyses in Fig. 3 (A to C) reveal the existence of high-density preexisting [100] and [111] type of dislocations in CoAl of CoAl/FAIs. An HRSTEM micrograph along the [100] zone axis and corresponding IFFT analyses in Fig. 3 (D and E) also confirm the existence of the [100] type of dislocations. Detailed explanation of the formation of FAIs and introduction of preexisting dislocations are provided in note S1.

Fig. 1. Microstructure of the as-fabricated NC CoAl/FAIs.

(A) BF TEM image of CoAl/FAIs and corresponding SAD pattern showing the B2 structure. (B) HRTEM image and corresponding FFTs in CoAl and FAIs. (C) Zoom-in HRTEM image of amorphous interface. (D to F) HAADF STEM image and EDS maps showing chemically abrupt FAIs. (G) Superimposed STEM image and EDS maps showing the Al-rich amorphous Al-Co alloy in FAIs. (H) EDS line scan across interface showing chemical segregation of FAIs.

Fig. 2. Dislocation-free SC CoAl and NC CoAl/FAIs with preexisting dislocations.

(A) BF TEM image showing dislocation-free SC CoAl. (B) IPF of (A) along the [110] zone. (C and D) KAM and corresponding GND maps showing no GNDs. (E) BF TEM image of NC CoAl/FAIs showing dislocation arrays. (F and G) Zoom-in DF TEM images showing preexisting dislocation array in CoAl of NC CoAl/FAIs. (H and I) KAM and corresponding GND maps showing high-density preexisting GNDs with a value of (1.05 ± 0.89) × 10¹⁶/m².

Table 1. GND density measured from ASTAR and total dislocation density measured from DF TEM images, before and after deformation.

	Dislocation density (/m2)
	Before deformation	After deformation
GND-ASTAR	(1.05 ± 0.89) × 1016	(1.16 ± 0.92) × 1016
DF-TEM	(4.28 ± 1.05) × 1016	(1.68 ± 0.95) × 1017

Fig. 3. Preexisting dislocations in CoAl of NC CoAl/FAIs.

(A to C) HRTEM image of CoAl along the [110] zone and IFFT images showing the [100] and [111] type of dislocations. (D and E) HRSTEM image of CoAl along the [100] zone and IFFT images showing the [100] type of dislocations.

High strength and plasticity synergy

The mechanical behaviors of NC CoAl/FAIs along with NC columnar CoAl, NC equiaxed CoAl, and SC CoAl were investigated via in situ scanning electron microscopy (SEM) micropillar compression tests. The comparison of true stress-strain curves in Fig. 4A shows that SC B2 CoAl, NiAl, and CuZn intermetallics exhibit intermittent load drops on stress-strain curves, indicating their brittle nature. Multiple shear bands were observed for SC CoAl from snapshots in Fig. 4 (C to F), implying unstable deformation. A major crack with sparse dislocation initiation was observed at the pillar bottom from postdeformation TEM analyses for SC CoAl in Fig. 5A. The entire pillar is inclined along the crack [which can be seen from SAD patterns of the deformed pillar and base in Fig. 5 (B and C)], coupled with the formation of multiple shear bands in the deformed pillar. These features further confirm the brittleness of SC CoAl. Compared to SC B2 intermetallics, both NC columnar CoAl and NC equiaxed CoAl show ultrahigh yield strength, more than 5.5 GPa in Fig. 4A. NC columnar CoAl shows a significant load drop after yielding, accompanied by plastic buckling and grain refinement (Figs. 4, G to J, and 5, D to G). The pre- and postdeformation microstructure and deformation mechanism for NC columnar CoAl can be seen in detail in fig. S4 and note S2. NC equiaxed CoAl also exhibits a major load drop after yielding but coupled with the formation of intergranular/intragranular fracture (Figs. 4, K to N, and 5, H and I). The pre- and postdeformation microstructure and deformation mechanism for NC equiaxed CoAl can be seen in detail in fig. S6 and note S3. In contrast, NC CoAl/FAIs exhibit a yield strength of ~6 GPa with a prominent and sustainable work hardening to ~8.5 GPa, which is rarely seen in intermetallics given their inherent brittleness. Instead of unstable deformation features like shear bands, buckling, and fracture, NC CoAl/FAIs show uniform deformation up to 15% with barreling on the pillar top (Fig. 4, O to R). Uniform layer thickness reduction of CoAl and grain refinement [which can be seen from SAD patterns of deformed and undeformed parts in Fig. 5 (K and L)] have been observed in the deformed pillar (Fig. 5, J to M), showing significant plasticity. See movies S1 to S4 for more details. The reproducibility of true stress-strain curves can be seen in figs. S5 and S7 to S9. Figure 4B shows the comparison of fracture stress versus uniform strain (${\mathrm{\epsilon }}_{\mathrm{u}}$, strain where unstable deformations like crack, buckling, and shear band occur) among different strategies for improving plasticity of B2 CoAl, NiAl, and FeAl intermetallics. NC CoAl/FAIs investigated in this work display extraordinarily high fracture stress and considerable plastic deformability compared to other studies, indicating the remarkable synergy of strength and plasticity.

Fig. 4. Mechanical behavior of NC CoAl/FAIs and comparison with various B2 type intermetallics.

(A) True stress-strain curves of NC CoAl/FAIs compared to NC columnar, NC equiaxed CoAl, and different SC B2 intermetallics, via in situ micropillar compression tests. NC CoAl/FAIs exhibits high yield strength and continuous work hardening. (B) A comparison of fracture stress and uniform strain among various B2 intermetallics and composites (74–83). (C to F) SEM snapshots of SC CoAl showing shear band. (G to J) SEM snapshots of NC columnar CoAl showing plastic buckling. (K to N) SEM snapshots of NC equiaxed CoAl showing fracture. (O to R) SEM snapshots of NC CoAl/FAIs showing uniform plastic deformation with barreling on the pillar top. See movies S1 to S4 for more details.

Fig. 5. Deformation mechanisms in SC CoAl, NC columnar CoAl, NC equiaxed CoAl, and NC CoAl/FAIs.

A comparison of deformation mechanisms in (A to C) SC CoAl, (D to G) NC columnar CoAl, (H and I) NC equiaxed CoAl, and (J to M) NC CoAl/FAIs. (A to C) BF TEM images and SAD patterns in inclined deformed pillar and base for SC CoAl. Crack between inclined deformed pillar and base and inclined deformed pillar with shear bands were observed. (D to G) BF TEM image of the deformed NC columnar CoAl pillar showing plastic buckling. The SAD patterns of the undeformed and deformed portions showing grain rotation. The IPF showing grain refinement in the deformed portion. (H and I) BF TEM images of the deformed NC equiaxed CoAl pillar showing intergranular and intragranular fracture. (J to M) BF TEM images of the deformed NC CoAl/FAIs pillar showing uniform deformation manifested by barreling and layer thickness reduction. The SAD patterns of the undeformed and deformed portions indicate grain refinement and grain rotation after deformation.

DISCUSSION

Preexisting dislocations mediated toughening

The crack initiation and extension are associated with dislocation activities. For ductile materials, the accumulation of dislocations around the crack tip enables the formation of a plastic zone, consequently blunting the propagation of the crack (38). This process is referred to as the shielding effect of dislocations or dislocation toughening (39). Usually, high-density dislocations around the crack tip are the prerequisite of the shielding effect, which can be achieved via dislocation multiplication in ductile materials (40). The shielding effect is rarely observed in brittle materials because of their limited capability of generating dislocations (41, 42). However, the introduction of high-density preexisting dislocations provides an opportunity to achieve dislocation toughening in brittle intermetallics. The contribution of individual dislocation on stress intensity factor ($K_{\mathrm{D}}$) can be estimated via the crack tip shielding model (41, 42). The sum of stress intensity factors ($\sum K_{\text{Di}}$) from individual dislocations i within the plastic zone of the crack tip can be described as (39)

$$\sum K_{\text{Di}}=\frac{\mathrm{\mu }b}{{\left(2\mathrm{\pi }{x}_i\right)}^{0.5}\times (1-\mathrm{\nu })}f({\mathrm{\phi }}_i,{\mathrm{\theta }}_i)$$ (1)

where μ, $b$, and $\mathrm{\nu }$ are the shear modulus, Burgers vector, and Poisson’s ratio, which can be taken as 126.4 GPa (12), 0.248 nm for the <111>/{110} glide system, and 0.16 (43) for CoAl, respectively. $f({\mathrm{\phi }}_i,{\mathrm{\theta }}_i)$ is a function of the angle between crack and slip planes (${\mathrm{\phi }}_i$) and the line joining crack tip to the dislocation (${\mathrm{\theta }}_i$), which can be taken as 1 if dislocations are ideally oriented (39). $x_i$ is the radius of the plastic zone for dislocation i and can be described as

$$x_i=\frac{1}{2\mathrm{\pi }}{\left(\frac{K_{\mathrm{I}0}}{\frac{{\mathrm{\sigma }}_{y0}}{1-2\mathrm{\nu }}}\right)}^2$$ (2)

where $K_{\mathrm{I}0}$ is the pristine toughness without any dislocations and ${\mathrm{\sigma }}_{y0}$ is the yield strength, which can be taken as 0.88 MPa$\sqrt{m}$ (12) and 1.8 GPa (from Fig. 4A and fig. S9) for CoAl respectively. Assuming that the crack tip area is $\mathrm{\pi }{x_i}^2$ and each dislocation i has the same configuration, the cumulative toughening effect ($\sum K_{\mathrm{D}}$) from preexisting dislocations with a density of 4.28 × 10¹⁶/m² is 4.66 MPa$\sqrt{m}$, almost five times that of $K_{\mathrm{I}0}$. This estimate indicates that the fracture toughness with preexisting dislocations is significantly increased compared to dislocation free CoAl, thereby reducing the susceptibility to fracture and enhancing plasticity.

Preexisting dislocations/FAIs mediated deformation

Compared to SC CoAl showing unstable deformation features like shear band and fracture, uniform deformation with barreling is observed in NC CoAl/FAIs, indicating that plasticity may be mediated by preexisting dislocations and FAIs. Systematic postdeformation TEM analyses were performed to understand the deformation mechanism in NC CoA/FAIs. As shown in IPF and GB maps before and after deformation in Fig. 6 (A to D), the CoAl layers within the barreled region exhibited a dominant reduction in vertical layer thickness accompanied by lateral expansion after deformation. The statistical analyses of grain sizes of columnar CoAl grains are shown in Fig. 6 (E to G). The average longitudinal grain sizes of CoAl (d_long) before and after deformation were 90.7 and 49.9 nm, respectively, illustrating the reduction of columnar grains along the vertical direction. The average lateral grain sizes of CoAl (d_lat) before and after deformation were 33.5 and 42.1 nm, illustrating the expansion of columnar grains along the horizontal direction. The aspect ratio of d_long/d_lat is thereby reduced from 2.8 to 1.3. Plastic strains (layer thickness reduction) of deformed CoAl layers are listed in Fig. 6H, where the highest strain level of an individual CoAl layer reached more than 40%. Statistical analyses of densities of low-angle grain boundaries (LAGBs) and high-angle grain boundaries (HAGBs) before and after deformation can be seen in the inserted table in Fig. 6. It is found that the density and the fraction of LAGBs decreased, whereas the density and the fraction of HAGBs increased, implying the geometric evolution of columnar grains during deformation. The GND density (KAM and the corresponding GND map are shown in fig. S10) and the total dislocation density (measured from DF TEM images using the line intercept method) before and after deformation in CoAl are shown in Table 1. The total dislocation density increased by around three times from (4.28 ± 1.05) × 10¹⁶/m² to (1.68 ± 0.95) × 10¹⁷/m², exhibiting the capability of dislocation multiplication in B2 CoAl, which is rarely observed in brittle intermetallics.

Fig. 6. IPF, GB maps, and grain size analyses in NC CoAl/FAIs before and after deformation.

(A and B) IPF and GB maps before deformation. (C and D) IPF and GB maps after deformation showing grain rotation and grain refinement. (E to G) Statistical analyses of longitudinal and lateral grain sizes and aspect ratio (d_long/d_lat) showing prominent reduction along longitudinal direction and expansion along lateral direction in columnar CoAl grains after deformation. (H) Statistical analyses of the plastic strains of deformed CoAl layers calculated from layer thickness reduction. The statistical analyses of the density of LAGBs and HAGBs before and after deformation can be seen in the inserted table.

The deformation microstructure for plastic strain between 15 and 30% (calculated from individual CoAl layer thickness reduction) is shown in Fig. 7. Two regions were selected within the barreling zone in Fig. 7A, defined as area (b) in the middle part of the barreling and area (c) in the lower part of the barreling. In area (b), columnar grains were uniformly deformed, showing reduction along the vertical direction and expansion along the horizontal direction (Fig. 7B). Preserved dislocations are identified under DF TEM after deformation (Fig. 7B), which demonstrates uniform deformation assisted by preexisting dislocations and corresponding dislocation evolution during deformation. In area (c), columnar grains were severely rotated along the periphery of the lower bulged area, to accommodate the geometric change due to barreling. High-density dislocation array was observed to be stored in the rotated columnar grains (Fig. 7C). HRSTEM and corresponding FFTs of one rotated columnar grain along the [111] zone were further shown in Fig. 7 (D to F). No fracture and shear band were observed, despite a rotation of ~8°, which is quite unexpected considering the brittle nature of CoAl intermetallics. These deformation features demonstrate that brittle CoAl intermetallics can be toughened by the presence of preexisting dislocations and can become plastically deformable without fracture. We wish to emphasize that the rotation of columnar grains serves to accommodate geometric evolution near the bulged portion of barreling. When deformation proceeds and interior columnar grains expand horizontally (Fig. 8A), these rotated columnar grains near the bulging will tilt back and show uniform deformation (vertical reduction and horizontal expansion) gradually (Fig. 7B). The rotation can be further assisted by migration of FAIs. As shown in Fig. 7 (G to J), when deformation continues, the horizontal amorphous interface is compressed, and the vertical amorphous interface is also protruded along GBs, which aids in GB sliding/rotation. Amorphous interface is also believed to have the capability of delaying crack nucleation and suppressing crack propagation via delocalizing the strain concentration within the amorphous structure, therefore toughening NC materials and preventing potential intergranular fracture (44–46). A similar yet different phenomenon occurring in this study is discussed in note S4.

Fig. 7. Deformation mechanisms in NC CoAl/FAIs at 15 to 30% strain (calculated from individual CoAl layer thickness reduction).

(A) BF TEM image showing barreling. (B and C) DF TEM images showing preserved dislocation arrays in columnar grains of the barreled pillar. (D) HAADF STEM image showing rotated columnar grains without any fracture. (E and F) Corresponding FFTs in (D) implying grain rotation in the columnar grain. (G) HRSTEM image around horizontal CoAl/FAIs showing reduced thickness in the amorphous interface. (H to J) HRSTEM image and corresponding EDS maps near the vertical CoAl/FAIs showing protruded amorphous interface. No intergranular fracture is observed.

Fig. 8. Deformation mechanisms in NC CoAl/FAIs at ~40% strain (calculated from individual CoAl layer thickness reduction).

(A) BF TEM image showing uniform layer thickness reduction without any fracture in the middle of deformed pillar within the barreling. (B) DF TEM image showing equiaxed grains after deformation. (C) HRSTEM image showing refined equiaxed grains with a grain size of ~10 nm. (D) IFFT dislocation analysis exhibits preserved dislocations in refined equiaxed grains after deformation. (E to H) HRSTEM image and corresponding EDS maps showing amorphous clusters around refined equiaxed grains after deformation. (I) HRTEM image and corresponding FFT in amorphous clusters showing formation of locally crystalized regions. No intergranular fracture is observed.

The deformed microstructure for plastic strain above 40% (in the interior part of the deformed pillar) is displayed in Fig. 8. It is observed that the CoAl layers undergo uniform layer reduction, accompanied by a transformation from columnar to refined equiaxed grains (Fig. 8, A and B). The sizes of the refined equiaxed grains could be lowered down to several nanometers as revealed by HRSTEM micrograph (Fig. 8C). It is worth noting that IFFT analyses in Fig. 8D and fig. S11 reveal sustained dislocation storage in deformed nanograins. This observation is rather unexpected, as it is commonly believed that storing dislocations in such small grains is difficult due to the limited space for dislocation slip and accumulation (34, 47). Moreover, GBs are generally ineffective for dislocation storage at this scale, because dislocation sources are absent within the grains, and the GBs act simultaneously as dislocation sources and sinks (34, 48). However, the presence of preexisting dislocations or defects within the grain interior can serve as barriers, impeding dislocation motion and promoting dislocation accumulation (34, 49–51). This blocking effect accounts for the increased dislocation density observed after deformation (Table 1) and contributes to the dislocation storage that underpins the work hardening behavior. More supporting evidence is shown in fig. S12, where immobile [100] dislocations are seen in the deformed CoAl layers, effectively trapping mobile [111] dislocations. To accommodate severe deformation of ultrafine equiaxed CoAl grains, the FAIs evolve into clusters that encircle individual grains (Fig. 8, E to H, and fig. S13), which also helps suppress crack formation along GBs. Notably, DIC is observed in the deformed FAIs (Fig. 8I and fig. S14). The formation mechanism of crystallized regions and the impact on dislocation multiplication are discussed in the following section. Together, these microstructural features suggest that CoAl, when engineered with preexisting dislocations and FAIs, can be toughened and sustain substantial plastic deformation.

Dislocation multiplication and storage

A critical aspect of this study is to understand the origin of increasing dislocation density in NC CoAl intermetallics and the mechanisms by which the high-density dislocations are retained to accommodate plastic deformation and enable work hardening. Besides the preexisting dislocations, the dislocation emission from FAIs and dislocation multiplication from preexisting dislocations are the other two factors contributing to dislocation storage in NC CoAl intermetallics. Therefore, MD simulations of nanopillar compression were performed to investigate the detailed dislocation activity in CoAl with preexisting dislocations and amorphous interface, which can contribute to dislocation nucleation or storage. Because of the spatiotemporal limitations of MD simulations, it is computationally prohibitive to mimic the compression deformation of polycrystalline CoAl grains with FAIs under experimental size and strain rates. Therefore, we focus on using MD simulations to identify deformation mechanisms and corresponding dislocation activity arising from preexisting dislocations in CoAl and the amorphous interface. Specifically, two CoAl nanopillars are considered, one with a pristine B2 CoAl structure and without any defects to serve as a comparison, denoted as “SC CoAl” (Fig. 9B), and the other one with amorphous interface layers and preexisting dislocations in CoAl, denoted as “CoAl/amorphous interface with dislocations” (Fig. 9, C and D).

Fig. 9. MD simulations of SC CoAl and CoAl/amorphous interface with dislocations.

(A) True stress-strain curves reveal that SC CoAl has a long elastic deformation region followed by a sharp load drop. In contrast, CoAl/amorphous interface with dislocations has more continuous plastic deformation with work hardening. (B to D) Initial configuration of SC CoAl pillar and CoAl/amorphous interface pillar with preexisting dislocations. (E to G) Atomic structure, dislocation, and shear strain analyses of SC CoAl pillar. The formation of slip bands is observed. (G) von Mises strain distribution and sparse dislocation emission along the slip band. (H to J) Atomic structure and dislocation analysis of CoAl/amorphous interface pillar with preexisting dislocations. (H) Onset of DIC along CoAl/amorphous interface. The crystallized regions are identified as BCC-like local motifs. (I) Dislocations nucleate from BCC-like local motifs and propagate into CoAl, followed by dislocation reaction and propagation. (J) Plasticity is accommodated by barreling and layer thickness reduction. High-density dislocation storage is achieved after deformation. (K and L) Comparison of total dislocation length of [100] and 1/2 [111] dislocations in CoAl in SC CoAl and CoAl/amorphous interface with dislocations at different strain levels. Limited dislocation length increase is coupled with formation of slip band in SC CoAl. Continuous dislocation length increase is observed in CoAl for CoAl/amorphous interface with dislocations. See movies S5 and S6 for more details.

The MD simulations show that the SC CoAl pillar has a higher yield strength (Fig. 9A), which is expected because of its perfect crystalline structure. Before yielding, no dislocation activity is observed (movie S5). After yielding, slip bands nucleate from the corners and propagate through the pillar, eventually leading to a substantial load drop (Fig. 9, A and E to G). Very few dislocations are nucleated along slip bands (Fig. 9, E to G), and the pillar undergoes minimal deformation. The formation of slip bands/shear bands in MD simulations is consistent with observations from in situ micropillar compression tests (Fig. 4, C to F, and movie S1). These findings indicate that SC CoAl is intrinsically brittle with limited ability to nucleate dislocations. As a result, the plasticity is primarily accommodated via shear band propagation.

In contrast, the dislocation activity in CoAl/amorphous interface pillar with dislocations is markedly different. First, significant dislocation activity occurred in CoAl with high-density preexisting dislocations, leading to dislocation accumulation and storage (movie S6). Immediately after yielding, dislocation reactions, multiplication, and propagation have been observed in the CoAl layer of the pillar (Fig. 10, M to P). These dislocations are further trapped and entangled by other dislocations and subsequently stored in CoAl (Fig. 10, N to P). The corresponding evolution of total dislocation density in CoAl layer is displayed in Fig. 10Q, which shows a gradual increase as deformation continues and confirms dislocation storage in CoAl. During the subsequent plastic deformation, the localized crystallization is induced in the amorphous interface (Figs. 9, H to J, and 10, A to H). The shear strain is believed to drive the formation of crystallized regions by rearranging atoms in amorphous phase in a more energetically favorable local configuration (52–55). As shown in Fig. 10I, the volume fraction of the crystallized regions [Body-Centered Cubic (BCC)–like local motifs] in the amorphous interface continuously increases as plastic deformation continues. Details about the evolution of the atomic packing in the amorphous interface before and after compression in MD simulations can be seen in fig. S15 and note S5. Meanwhile, dislocations continuously nucleate from these BCC-like local motifs along the CoAl/amorphous interface and subsequently propagate into CoAl (Fig. 10, J to L). The dislocations propagated into CoAl can further react, multiply, and entangle with each other, causing dislocation accumulation in CoAl (Fig. 9, I to L). Similarly, Wang et al. (56) found that dislocations are nucleated from amorphous metallic glass layers in ductile crystalline-amorphous nanolaminate. However, if there are no high-density preexisting dislocations in the nanocrystal layer, the dislocations will further glide across the nanocrystal layer, eventually being absorbed at opposite amorphous layers. Therefore, although plasticity can be accommodated, the work hardening is absent due to the sparse dislocation accumulation (56). In contrast, the dislocation reaction between propagated dislocations and preexisting dislocations in our study traps the dislocations and results in the dislocation storage, eventually promoting the work hardening capability. Dislocation emission from deformation-induced crystallized regions within the amorphous interface thus constitutes an additional source contributing to dislocation multiplication in CoAl. Correspondingly, instead of forming slip band as observed in SC CoAl (Fig. 9G), barreling at the center and uniform layer reduction become apparent in CoAl/amorphous interface pillar with dislocations (Fig. 9J), showing enhanced plasticity. The storage of high-density dislocations, from both dislocation multiplication of preexisting dislocations in CoAl and dislocation emission from deformation-induced crystallized regions within the amorphous interface, blocks further dislocation movement and results in work hardening (Fig. 9A). The dislocation analyses show that both 1/2 [111] and [100] dislocations are mostly preserved after deformation and display continuous increase of length in CoAl/amorphous interface with dislocations (Fig. 9, K and L). The [110] dislocation is not stable because of its nonplanar core structure and can be further dissociated (57). The dissociation of [110] dislocation leads to formation of 1/2 [111] partial dislocations (Fig. 9I), which can further glide to accommodate plasticity (Fig. 9J). One can also notice that most of the dislocations are 1/2 [111] and [100] types in the top CoAl layer after deformation (Fig. 9J), suggesting that 1/2 [111] and [100] dislocations are the dominating ones that accommodate plasticity. This finding is also consistent with experimental observations in B2 intermetallics (25, 57). When comparing the simulations of dislocation length evolution in SC CoAl and CoAl/amorphous interface with dislocations in Fig. 9 (K and L), the SC CoAl shows an abrupt increase in the total length for both types of dislocations, coupled with the formation of slip band (Fig. 9, E to G). In contrast, the CoAl/amorphous interface with dislocations shows much higher values of dislocation length and continuous increase in Fig. 9 (K and L), which also highlights the capability of dislocation multiplication and storage in CoAl/amorphous interface with dislocations.

Fig. 10. Dislocation emission from DIC and dislocation activities in CoAl for CoAl/amorphous interface with dislocations.

(A to I) Shear strain induced crystallization (BCC-like local motifs) in amorphous interface. (A to D) Angle view (both cross section and top surface) showing the evolution of strain distribution (von Mises strain) in amorphous interface. (E to H) Corresponding phase evolution map in amorphous interface, indicating formation of BCC-like local motifs with increase of shear strain level in amorphous interface. (I) The evolution of volume fraction of crystallized regions in amorphous interface versus true strain. (J to L) Dislocation nucleation from crystallized regions (BCC-like local motifs) in amorphous interface and subsequent propagation into CoAl. (M to Q) Dislocation activities and storage in top CoAl layer of CoAl/amorphous interface with dislocations. (M to P) Dislocation storage via reaction, multiplication, propagation, and trapping in top CoAl layer of CoAl/amorphous interface with dislocations. (Q) Dislocation length evolution in top CoAl layer of CoAl/amorphous interface with dislocations, showing dislocation accumulation. See movie S6 for more details.

High strength and work hardening

The origin of high yield strength in NC CoAl/FAIs can be considered from several strengthening factors. The first one is GB strengthening as informed from the Hall-Petch relationship (58). Strengthening via grain refinement is effective for grain sizes as small as tens of nanometers (47, 59). High-density GBs act as barrier for dislocation motion, making plastic deformation more difficult while increasing yield strength or hardness. The second one is resistance for dislocation transmission across interface between CoAl and FAIs, which has been frequently discussed in the context of strengthening in nanolaminates (60–62). The third one is strengthening due to high-density preexisting dislocations. Increased dislocation density makes it harder for further dislocation nucleation and movement via trapping and reaction, therefore increasing strength (63, 64). The yield strength can be estimated as

$${\mathrm{\sigma }}_y={\mathrm{\sigma }}_{y0}+{\mathrm{\sigma }}_{\mathrm{d}}+{\mathrm{\sigma }}_{\text{int}}+{\mathrm{\sigma }}_{\mathrm{\rho }}$$ (3)

where ${\mathrm{\sigma }}_{y0}$ is the yield strength of pristine CoAl (1.8 GPa for SC CoAl in Fig. 4A and fig. S9). ${\mathrm{\sigma }}_{\mathrm{d}}$ is the strengthening from grain refinement and can be estimated as (65)

$${\mathrm{\sigma }}_{\mathrm{d}}=\frac{k_y}{\sqrt{d_{\text{eff}}}}$$ (4)

where $k_y$ is the strengthening coefficient [125 Pa m^0.5 for B2 CoAl (66)] and $d_{\text{eff}}$ is the effective grain size of columnar CoAl intermetallics (equivalent equiaxed grain sizes with the same volume of columnar grains). ${\mathrm{\sigma }}_{\text{int}}$ is the interface barrier resistance for dislocation transmission between CoAl and FAIs. ${\mathrm{\sigma }}_{\mathrm{\rho }}$ is the strengthening from preexisting dislocations and can be described using Taylor equation (67, 68)

$${\mathrm{\sigma }}_{\mathrm{\rho }}=\mathrm{\alpha }\mathrm{\mu }b{\mathrm{\rho }}^{0.5}$$ (5)

where α is a constant (taken as 0.19), μ is the shear modulus [126.4 GPa (12)], $b$ is the Burgers vector (0.248 nm for <111>/{110} system of B2 CoAl), and ρ is the dislocation density. Taking $d_{\text{eff}}$ as 43.1 nm and ρ as 4.28 × 10¹⁶/m² for CoAl, ${\mathrm{\sigma }}_{\mathrm{d}}$ and ${\mathrm{\sigma }}_{\mathrm{\rho }}$ can be estimated as 0.60 and 1.24 GPa, respectively. ${\mathrm{\sigma }}_{\text{int}}$ can be estimated from MD simulations and taken as the applied stress when dislocation emission from amorphous interface is initiated (2.2 GPa at 2.3%, also the yielding point; see movie S6 for more details). The ${\mathrm{\sigma }}_y$ is thus estimated as 5.84 GPa, which is very close to the observed yield strength in experiments (Fig. 4A and fig. S8).

Work hardening in intermetallics at room temperature is rarely discussed due to their inherently limited dislocation activity. However, the successful multiplication and storage of high-density dislocations in the current study present a valuable opportunity to explore this phenomenon. As the NC CoAl/FAIs shows work hardening, the strengthening or increase of flow stress can be attributed to grain size evolution, change of ${\mathrm{\sigma }}_{\text{int}}$, and dislocation multiplication. The $d_{\text{eff}}$ can be taken as 41.1 nm from grain size analyses, showing little change. The ${\mathrm{\sigma }}_{\text{int}}$ is taken as 3.85 GPa from MD simulation which is the applied stress at 11.2% strain when dislocations start to emit from deformation-induced crystallized regions and propagate into CoAl (see movie S6 for more details). ρ is taken as 1.68 × 10¹⁶/m² for CoAl from postdeformation TEM analyses. Inserting these values, the flow stress can be estimated as 8.97 GPa, consistent with experimental observation.

In summary, we achieved a synergy between high strength and plasticity in intrinsically brittle B2 CoAl intermetallic via introducing FAIs and preexisting dislocations. The high strength originates from the Hall-Petch strengthening, interface barrier resistance between CoAl and FAIs, and high-density preexisting dislocations. Significant plasticity and work hardening capability are realized via dislocation accumulation, arising from dislocation emission associated with DIC in FAIs and dislocation multiplication from preexisting dislocations in CoAl. Dislocation interaction and accumulation in NC CoAl further retard dislocation motion, eventually promoting dislocation storage and enhanced work hardening capability. The FAIs also accommodate geometric evolution of NC CoAl grains and avoid the formation of intergranular fracture. The current study provides an effective approach to achieving plasticity at room temperature in a wide range of intrinsically brittle materials.

MATERIALS AND METHODS

Sample fabrication

Approximately 2-μm-thick CoAl 100-nm/Al 10-nm multilayers with Al 50-nm seed layer (NC CoAl/FAIs) and 2-μm-thick CoAl (NC columnar CoAl) were fabricated on hydrofluoric acid etched Si (110) substrates using magnetron sputtering (AJA ATC-2200-UHV system). A substrate bias of −128 V was applied onto Si (110) substrates for these two specimens. Two-micrometer-thick Co 12-nm/Al 18-nm multilayers with an overall atomic concentration of 50 at % Co and 50 at % Al were fabricated on SiO₂ followed by a postannealing of 675°C/5 hours in a vacuum furnace. The postannealed sample is NC equiaxed CoAl. The deposition rates for Co and Al were 0.09 and 0.15 nm/s, respectively, for the above-mentioned three specimens. The base pressure of the chamber was 8 × 10⁻⁸ torr. The Ar pressure was set to 3.5 mtorr during deposition. Bulk CoAl intermetallics were prepared via arc-melting high-purity Co (99.9 at %) and Al (99.9 at %) pellets under an Ar atmosphere, using an Edmund Buhler Arc Melter.

Microstructural and chemical characterization

The texture of as-fabricated samples was identified via x-ray diffraction studies on a Panalytical Empyrean X’pert PRO MRD diffractometer operated at 40 kV with Cu Kα1 radiation source. The cross-sectional TEM foils were prepared via focused ion beam (FIB) inside the Thermo Fisher Scientific Helios G4 UX dual beam SEM system equipped with an Omniprobe manipulator. Pt protection coating was deposited to avoid FIB damage. Two-kilovolt low-energy Ga⁺ ion beam polishing was conducted as a final step to minimize ion damage on TEM foils. Microstructure and chemistries were studied under TEM and STEM, on an FEI Talos 200× transmission electron microscope operated at 200 kV and an FEI Themis Z double aberration-corrected transmission electron microscope operated at 300 kV, both equipped with Fischione ultrahigh resolution HAADF detector and Super-X EDS detector. The NanoMegas ASTAR system, an automatic crystallographic indexing and orientation/phase mapping tool on FEI Talos 200×, was used to measure automated crystallographic orientation, GB misorientation, and KAM. Postanalysis including extrapolation of GND density from KAM was conducted on EDAS OIM Analysis software. The total dislocation density was estimated via the line intercept method, which is proven to be effective for dislocation densities less than 10¹⁸/m² (69).

Nanomechanical tests

Nanoindentation was performed to measure hardness on a Hysitron TI Premier nanoindenter using a Berkovich tip under displacement control mode. The indentation depth was limited to below 15% of film thickness. Mechanical behavior was also studied using in situ SEM micropillar compression tests on Hysitron PI 88xR PicoIndenter inside an FEI Quanta 3D FEG dual beam SEM system. The size of the fabricated pillar is 1 μm (diameter) × 2 μm (height), which was milled using FIB on the Thermo Fisher Scientific Helios G4 UX dual beam SEM system. The fabrication process of the pillars consisted of four steps. For the first step, an annular trench with an outer diameter of 50 μm and an inner diameter of 5 μm was milled into the sample for 2 μm, so that there is enough space for a diamond flat punch tip to achieve compression tests. The annular trench was milled with high ion current (9.1 nA under 30 kV) from outer to inner direction. For the second step, the annular trench is further expanded from outer to inner direction with lower current (0.26 nA under 30 kV) until the inner diameter of 2 μm is achieved. For the third step, a fine milling was conducted from outer to inner direction with low current (41 pA under 30 kV) until the diameter of the pillar reached 1.25 μm. For the final step, a finer milling was used from inner to outer direction with low current (41 pA under 30 kV) to further polish the diameter of the pillar into 1 μm meanwhile minimizing the taper effect. The dwell time of the final step was set to below 5 s to minimize FIB damage. The displacement control tests were conducted at a strain rate of 5 × 10⁻³ s⁻¹ using a 10-μm diameter diamond flat punch tip. The fabrication of deformed pillar TEM foils and postdeformation microstructural and chemical characterization were performed using the above-mentioned procedures.

MD simulations

MD simulations of compression were performed on four cylindrical specimens. The first two specimens consist of multilayered CoAl/amorphous interface pillars, one with preexisting dislocations (CoAl/amorphous interface with dislocations) and the other one without preexisting dislocations (CoAl/amorphous interface). These two samples contain three B2 CoAl layers separated by two amorphous interface layers, with the CoAl layers being 13.4-nm thick and the amorphous interface layers being 3.2-nm thick. The CoAl layers have a B2 crystal structure oriented along the [110] direction, aligned with the loading direction. The amorphous interface layers are created by randomly replacing 25 at % of Al atoms with Co atoms in amorphous Al, whose chemistry and amorphous structure closely replicate experimental findings. The other two samples are SC CoAl pillars with identical dimensions, one with preexisting dislocations (SC CoAl with dislocations) and the other one without preexisting dislocations (SC CoAl). They have the same orientation as the CoAl in CoAl/amorphous interface pillars. All pillars have a total height of 46.4 nm and a diameter of 23.4 nm.

Each digital specimen contains approximately 1.65 million atoms. The samples with preexisting dislocations were created using Atomsk (70), where the total dislocation length in SC CoAl with preexisting dislocations is 685.2 nm and the total dislocation length in CoAl/amorphous interface with preexisting dislocations is 698.3 nm, consistent with the dislocation density in our experiments. Multiple full and partial dislocation loops with varying lengths at random positions at a given layer were created and then merged. All pillars were firstly relaxed using conjugate gradient energy minimization and then equilibrated at 300 K for 50 ps followed by compression at a constant strain rate of 2 × 10⁸ s⁻¹. The two CoAl/amorphous interface samples were compressed for 1500 ps. The SC CoAl sample without preexisting dislocations was compressed for 900 ps. The SC CoAl sample with preexisting dislocations was compressed for 1100 ps. All simulations were conducted using Large-scale Atomic/Molecular Massively Parallel Simulator (71) and visualized with Open Visualization Tool (OVITO) (72). Atomic structures were color-coded using common-neighbor analysis. Shear strain was quantified using the atomic strain modifier in Open Visualization Tool (OVITO), which computes the local von Mises shear invariant based on atomic displacements. Atomic packing analysis, i.e., Voronoi polyhedron and its index <n3, n4, n5, n6>, was computed using Voronoi analysis in OVITO (72). Interatomic interactions between Al and Co atoms were modeled using the embedded-atom method potential developed by Purja Pun et al. (73).

Supplementary Materials

The PDF file includes:

Supplementary Text

Table S1

Figs. S1 to S23

Legends for movies S1 to S8

References

Other Supplementary Material for this manuscript includes the following:

Movies S1 to S8