Crystallographic Effect of TiAl Alloy Under High-Speed Shock Deformation

Abstract

In this paper, the molecular dynamics simulation method was adopted to systematically study the microstructure evolution behavior of TiAl alloys under impact compression under three typical crystal orientations ([001], [110], [111]). By analyzing the characteristics of structural phase transition, defect type evolution, dislocation expansion, and radial distribution function, the anisotropic response mechanism under the joint regulation of crystal orientation and impact velocity was revealed. The results show that the [111] crystal orientation is most prone to local amorphous transformation at high strain rates, and its structural collapse is due to the rapid accumulation and limited reconstruction of dislocations/faults. The [001] crystal orientation is prone to forming staggered stacking of layers and local HCP phase transformation, presenting as a medium-strength structural disorder. Under the strain regulation mechanism dominated by twinning, the [110] orientation exhibits superior structural stability and anti-disorder ability. With increases in the impact velocity, the defect type gradually changes from isolated dislocations to large-scale HCP regions and amorphous bands, and there are significant differences in the critical velocities of amorphous transformation corresponding to different crystal orientations. Further analysis indicates that the HCP structure and the formation of layering faults are important precursor states of amorphous transformation. The evolution of the g(r) function verifies the stepwise disintegration process of medium and long-range ordered structures under shock induction. It provides a new theoretical basis and microscopic perspective for the microstructure regulation, damage tolerance improvement, and impact resistance design of TiAl alloys under extreme stress conditions.

1. Introduction

In recent years, with the increasingly widespread application of high-performance structural materials in key fields such as aerospace, nuclear engineering, and advanced manufacturing [1], the mechanical challenges faced by materials in extreme service environments such as high pressure, low temperature, and high strain rate have been continuously intensifying [2]. Against this background, advanced alloy materials, with their excellent comprehensive performance, have become the core candidates for achieving the design of structural components with high reliability and extreme stability. However, under the action of high strain rate impact loads, materials often exhibit complex nonlinear responses and anisotropic microscopic behaviors [3], which poses higher requirements for their microstructure stability and damage tolerance [1,2,4]. Therefore, a thorough understanding of the microscopic evolution mechanism of materials during high-speed impact processes, especially the essential characteristics of crystal structure evolution and phase transition behavior, is of great significance for establishing the theoretical basis of new high-performance structural materials for extreme environments.

Shock waves are high-pressure and high-strain-rate waves generated under high-speed loading conditions. As a typical dynamic loading mode, the shock wave process can be triggered by natural or man-made events such as volcanic eruptions [5], earthquakes [6], meteorite impacts [7], and explosions, and it has the characteristics of extremely short duration, high intensity, and strong disequilibrium. The propagation process of shock waves can be divided into three main stages: generation, propagation, and attenuation, which correspond to the formation of the initial pressure wave, the transfer of energy into the interior of the material, and the attenuation of amplitude caused by time or distance, respectively. Among them, the length of the bulk, the velocity, and the material impedance determine the formation mode of the initial shock wave, while the propagation and evolution of the impact within the material profoundly reflect its intrinsic mechanical properties. During the impact loading process, the material undergoes multi-stage evolution such as elastic response, plastic yield, structural phase transformation, and even melting, accompanied by complex solid–solid and solid–liquid phase transformation behaviors [8,9,10]. Recent studies have shown that crystal orientation plays an important regulatory role in the microscopic deformation mechanism, dislocation excitation mode, and phase transition path of materials [11,12,13]. Therefore, deeply revealing the impact response behavior dependent on crystal orientation is conducive to constructing a design paradigm for the regulation of microstructure evolution and providing theoretical support for the directional functional development of high-performance structural materials.

TiAl alloys, as a type of lightweight high-temperature material with low density (ρ ≈ 3.9–4.2 g/cm³), high specific strength, excellent oxidation resistance, and good corrosion resistance [14,15,16], have shown broad application prospects in aerospace, automotive manufacturing, and energy equipment [17,18]. Despite its obvious performance advantages, the brittle nature of TiAl alloys at room temperature and the complex phase transformation behavior under impact loading severely restrict their stable service performance [19] in extreme stress environments [20,21]. At present, although a large number of studies have systematically studied the microscopic response laws of metals such as Cu [22], Fe [23], and Ni [24] under impact conditions, the atomic-scale understanding of TiAl alloys, especially the high-speed impact evolution mechanisms under different crystal levels, is still insufficient. There is still a lack of systematic understanding of its shearing behavior, structural transformation path, and defect evolution process. Due to the limitations of traditional experimental methods in terms of temporal resolution and spatial accuracy, the nanoscale structural evolution of materials under the action of shock waves is often difficult to capture in real time. Molecular dynamics (MD) simulation, as a powerful tool capable of characterizing the mechanical responses of materials at the atomic scale, has been widely applied in recent years to study the evolution behaviors of various materials under single-board impact [25,26,27], laser impact [28,29,30], and combined loading [30,31,32,33]. Some existing studies have revealed stress transfer, interface spalling, and cavity closure mechanisms of composite materials under impact conditions, emphasizing the influence of interface structure, material composition, and layer thickness on the impact damage behavior. For example, Tian et al. [34] found that the thickness difference of the flywheel would cause the Cu/Al multilayer board to fracture at different interface positions; Zhu et al. [35] pointed out that interfacial reflected waves can compact nano-voids, thereby inhibiting the accumulation of damage; Liao et al. [36] studied the spalling evolution process in Cu/Nb multilayer materials and found that due to the different spalling strengths between the two phases, the damage often starts from the soft phase Cu, and the Nb layer fails to reach the theoretical fracture stress. However, most of the above-mentioned studies have focused on traditional metals or metal multilayer systems, and few have involved TiAl-like alloys with amorphous or complex phase transformation behaviors, lacking systematic exploration of their microstructure evolution paths.

In response to the above research gap, based on large-scale molecular dynamics simulations, this paper systematically analyzes the high-speed impact response behaviors of TiAl alloys in three typical downgrain directions: [001], [110], and [111]. Through the dynamic tracking of the evolution of shear stress, phase transition trajectories, and atomic structure reconstruction mechanisms, the microscopic behavioral characteristics such as the formation mechanisms of different downward shear precursor platforms, dislocation excitation, and twining-induced phase transition paths were revealed for the first time. This research not only deepens the essential understanding of the microstructure mechanical behavior of TiAl alloys under high strain rate conditions but also provides an important theoretical basis and scientific guidance for the microstructure regulation, material optimization design, and high-performance component development of TiAl alloys in extreme service environments.

2. Methodology

2.1. Simulation Details

In this study, the large-scale parallel atomic/molecular simulation platform LAMMPS (Version 2023, USA) [37] was adopted to systematically simulate the microscopic deformation behavior of TiAl alloys with different crystal directions under impact loading in order to reveal their potential evolution mechanism. The interatomic interactions were modeled using the embedded-atom method (EAM) potential developed by Zope and Mishin [38], which has been extensively validated for its ability to capture the plastic deformation, fracture, and phase transformation behavior of TiAl alloys [39,40,41]. To further verify its applicability in this study, we computed the lattice constants and elastic moduli of TiAl at 300 K. The optimized lattice parameters (a₀ = b₀ = 4.001 Å, c₀ = 4.181 Å) closely match the experimental values (a₀ = b₀ = 4.0050 Å, c₀ = 4.0707 Å) [14], with a maximum deviation of 0.12 Å (~3%), confirming the reliability of the chosen potential. To investigate the influence of crystallographic orientation on shock response, TiAl single-crystal models with [001], [110], and [111] orientations aligned along the Z-axis were constructed. Owing to the orientation-dependent lattice arrangements, slight differences in geometric dimensions exist among the three models (see

Table 1. Detailed parameters involved in the simulated systems.

Z-Axis Orientation	Lx (Å)	Ly (Å)	Lz (Å)	Number of Atoms
[001]	141.75	141.75	1011.50	1,225,000
[110]	143.85	140.30	1010.92	1,228,440
[111]	143.85	140.33	1010.13	1,227,744

for details). The atomic ratio of Ti to Al was maintained at 1:1, with atomic positions randomly distributed within the lattice to avoid artificial ordering effects. Figure 1a shows the overall structure of the simulation model, in which a rigid piston with a thickness of 10 A is set to drive the main structure constructed of the same material to generate shock waves. The free surface is on the far right of the model. The impact process is achieved by pushing the piston upward at a constant speed, with the speed range set at 0.6–1.8 km/s, aiming to cover the possible elastoplastic responses and phase transformation behaviors of TiAl alloys. Furthermore, the three typical crystal orientations, [001], [110], and [111], were set as the impact directions to explore the influence of their anisotropy on the impact response mechanism.

To ensure that the simulation system has good initial stability, the conjugate gradient method is first adopted to minimize the energy of the configuration. Subsequently, the system was subjected to 50 ps equilibrium relaxation under the constant pressure and constant temperature (NPT) ensemble, allowing it to fully evolve under thermodynamic conditions of 300 K and nearly zero pressure, thereby effectively eliminating the residual stress in the configuration [42]. After reaching the equilibrium state, in order to accurately describe the non-equilibrium dynamic behavior of the material under high strain rate impact conditions, the micro-canonical (NVE) ensemble is instead adopted in the simulation stage to ensure the conservation of the total energy of the system. The application time of the impact load is 50 ps. In terms of the setting of boundary conditions, periodic boundaries are maintained in the X and Y directions, while aperioperiodic free boundaries are adopted in the Z axis (i.e., the direction of shock wave propagation) to truly reflect the wavefront propagation and the response behavior of the material interface. The entire simulation process adopts a time step of 0.001 ps to ensure the accuracy of the dynamic process under a high time resolution.

2.2. Visualization and Parameter Calculation

To systematically reveal the microscopic response mechanism of TiAl alloys under impact loading conditions, in this paper, the OVITO (3.8.0, Germany) [43] software was used to visualize and reconstruct the molecular dynamics simulation results to analyze the atomic configuration evolution characteristics at different time scales. Based on the wave propagation characteristics along the Z-axis direction, a one-dimensional spatial surface element division strategy is introduced to divide the sample into multiple statistical boxes. The thickness of each statistical unit is approximately twice the lattice constant to balance the spatial resolution and computational cost. In each statistical box, physical quantities such as local stress and temperature are calculated by volumetric weighted averaging of atomic-level data, thereby capturing the non-uniform response behavior induced by shock.

The atomic stress tensor is calculated based on the classical Virial formula [42], considering the conservation of momentum and the contribution of the interatomic interaction force. Its expression is as follows:

$${\sigma }_{\alpha \beta } = -1/{\Omega }_i ( {\sum }_i{\mathit{m}}_i{𝒱}_i^{\alpha }{𝒱}_j^{\beta }+{\sum }_i{\sum }_{i>j}{r}_{ij}^{\alpha }{f}_{ij}^{\beta } )$$

(1)

Here, α, β ∈ {x, y, z} represent the component directions of the stress tensor, and m_i, v_i, f_ij, Ω_i represent the atomic mass, velocity, force, and local volume, respectively. By statistically analyzing the stress tensors of all atoms in the region and conducting weighted averaging, the local average stress, shear stress [44], and equivalent stress distributions of von Mises [45] were further obtained.

$${\tau }_{shear} = 1/2 ( {\sigma }_{zz} - 1/2 ( {\sigma }_{xx} + {\sigma }_{yy} ))$$

(2)

$$P_m = \frac{1}{3} ({\sigma }_{zz}+{\sigma }_{yy}+{\sigma }_{xx})$$

(3)

$${\eta }^{Mises} = \sqrt{1/6[{({\eta }_{xx}-{\eta }_{yy})}^2+{({\eta }_{yy}-{\eta }_{zz})}^2+{({\eta }_{zz}-{\eta }_{xx})}^2+{\eta }_{xy}^2 +{\eta }_{yz}^2 + {\eta }_{zx}^2]}$$

(4)

Furthermore, in order to deeply analyze the structural reconstruction behavior during the impact process, an improved crystal structure identification method is adopted to classify the atomic local environment. Specifically, atoms with typical crystal structures (such as face-centered cubic FCC, body-centered cubic BCC, and hexagonal close-packed HCP) are classified as Grain Interior (GI), while the remaining atoms that cannot be identified as stable structures are classified as Grain Boundaries (GBs). This method can simultaneously identify planar defects, including Stacking Faults [4,46,47] (SFs) and Twin Boundaries (TBs), which is helpful for understanding the crystallographic evolution path induced by deformation.

Ultimately, by constructing a multi-scale mechanical analysis framework with consistent spatiotemporal resolution, the longitudinal pressure Pz, average pressure P_m, and shear response parameters in the direction of impact propagation were systematically obtained, providing solid data support and a physical explanation basis for revealing the yield behavior and stability mechanism of materials in a high-strain-rate environment.

3. Result and Discussion

3.1. The Impact Response Characteristics of Different Downward Grain Velocities, Pressures, and Shear Stresses

To deeply reveal the impact response mechanism of TiAl alloys under different crystal orientations, Figure 2 shows the evolution curves of particle velocity, pressure, and shear stress along the three crystal directions of [001], [110], and [111] at multiple time steps (6 ps, 8 ps, and 10 ps).

Based on the above results, the stress propagation characteristics and shear response paths exhibited by different crystal orientations under impact loading are significantly different, demonstrating a highly anisotropic dynamic behavior. Among them, the existence of the “early shear excitation” phenomenon in the [110] and [111] directions provides a microphysical basis for the subsequent proposal of orientation control strategies in dynamic forming and impact-resistant structure design. This research is expected to promote the development of the structure–performance co-design paradigm of “oriented programmable impact-resistant alloys”.

3.2. Analysis of Shock Wave Structure and Plastic Behavior

To deeply reveal the anisotropic response characteristics of TiAl alloys under different crystal orientations during the impact compression process, Figure 3 shows the molecular dynamics simulation results of the [001], [110], and [111] orientations under the condition of an impact velocity of 1.2 km/s. The shock wave propagates along the path from the contact surface of the left piston to the free surface on the right. The overall wave structure can be divided into the leading shock wave region and the subsequent plastic region. The former marks the transition of the material from the elastic state to the plastic state, while the latter is manifested as a platform region with approximately constant stress.

In this study, the width of the shock front is defined as the spatial range within which the impact stress decreases from 95% of its maximum value to 0. The plastic zone corresponds to the area where the impact stress weakens from the peak to within 95%. Furthermore, the plastic initiation point of TiAl alloys under impact loads can be identified by the minimum stress in the elastic wave platform. It should be pointed out that no elastic single-wave structure was observed at the current impact velocity, indicating that the material response has exceeded the elastic limit. For all crystal orientations, a typical elastic–plastic two-wave structure is presented. This phenomenon is attributed to the impact strength exceeding the elastic load-bearing capacity of the material, causing irreversible plastic deformation. It is notable that the width of the plastic wave region for the [001] orientation is significantly larger than that for the [110] and [111] orientations, while its elastic band is relatively narrower. This difference may stem from the fact that plastic waves almost catch up with or even surpass elastic waves, leading to the occurrence of the merging phenomenon of the two. This trend is also reflected in simulation work by [8], but this study further quantifies its correlation with the shock wave propagation speed and dislocation evolution. More importantly, the downward shock front of the [001] crystal lags significantly behind the other two orientations, reflecting a lower wave velocity. This behavior can be attributed to a more intense energy dissipation process, mainly caused by the triggering and expansion of large-scale twin boundaries (SF) and high-density dislocations. The subsequent defect structure and dislocation distribution maps further verified this conclusion. The comprehensive results show that the material loaded under [001] has a lower Heggenho limit (HEL) and is accompanied by more abundant plastic evolution behavior.

From the three-dimensional distribution characteristics of shear stress, sufficient shear stress was generated during the impact process to drive the occurrence of plastic flow. In the unimpacted area, the shear stress is approximately zero. When the shock wave front enters the interior of the material, the shear stress rises rapidly and reaches a relatively stable value in the elastic zone. However, when observed from right to left along the direction of impact propagation, the shear stress shows a significant attenuation trend from the elastic region to the plastic region, especially in the [110] direction. This attenuation phenomenon is mainly attributed to the local stress relaxation caused by the nucleation and slip process of dense dislocations. In contrast, the dislocation density under loading [110] is concentrated and has a smaller distribution range, resulting in the rapid release of shear stress over a short distance. In the [001] orientation, since dislocations are uniformly generated in the larger plastic regions, the process of shear stress release is relatively gentle, forming an approximate plateau shape, and the overall shear stress level is lower than that of the [110] and [111] orientations. Furthermore, in the front part of the shock wave, the shear stress platform of the [001] orientation is relatively narrow, while the remaining crystal orientations show a wider shock absorption transition zone. These observations indicate that shock wave structures (such as frontier width and wave velocity) and plastic responses are significantly sensitive to crystal orientation.

Although plastic deformation shows certain differences in different orientations, under the current impact velocity conditions, the dominant mechanism is the sliding behavior driven by dislocation activities rather than the deformation mechanism specific to the orientation. As previously mentioned, a small volume fraction of BCC phase transition was detected under [001] loading conditions at higher impact velocities. This transition is linked to the sliding behavior driven by dislocation activities. The physical mechanism behind this phenomenon is explored in the next section.

3.3. Shock-Induced Phase Transition and Dislocation Evolution Mechanism

Under extreme loading conditions such as impact compression, the interior of the alloy often undergoes intense non-equilibrium microscopic reconstruction processes, manifested as complex evolution behaviors with multi-scale and multi-mechanism coupling. Studies have shown that structural phase transition and dislocation dynamics are the key microscopic mechanisms regulating the macroscopic mechanical responses of materials [48,49,50]. Compared with static loading or quasi-static loading, the transient high-pressure and high-temperature environment induced by impact can significantly reduce the phase transition potential barrier and drive the lattice structure to undergo non-equilibrium reconstruction paths from face-centered cubic (FCC) to closely arranged hexagonal (HCP) and even amorphous states. Furthermore, the impact plasticity process at high strain rates is also accompanied by complex dislocation activities, including the excitation, slip, and annihilation of the Shockley partial dislocations. This type of dislocation behavior not only directly affects the local stress–strain distribution but also may trigger the dynamic regulation of stacking fault energy (SFE), thereby reshaping the structural evolution path and critical conditions. Figure 4 systematically compares the time evolution of substructure types (upper figure) and the changes in total dislocation length (lower figure) of three typical crystal orientations—[001], [110], and [111]—during the impact loading process, revealing significant crystal orientation correlations.

For the [001] crystal range (Figure 4a), the FCC structure rapidly decays within 0–15 ps (A0 stage), the proportion of HCP increases simultaneously, and the number of other types of atoms (i.e., atoms with non-ideal stacking structures) significantly increases, indicating that the stacking sequence of atomic layers undergoes severe instability driven by impact stress, inducing a local amorphous trend. During the same period, the Shockley type dislocation density rose rapidly and reached a peak at around 16 ps. Its timing sequence was highly consistent with the phase transition process from FCC to HCP. It reflects the typical behavior of Dislocation-Mediated Phase Transition (DMPT). This coupled evolution process indicates that in the initial stage of the impact, dislocations not only undertake the function of plastic deformation but also significantly regulate the phase transition path and rate at the atomic scale. At 20–30 ps (A1 stage), the proportion of HCP atoms gradually decreases, while the other structure continues to increase, indicating that the system evolves from the ordered parameter state to a highly amorphous state. At this stage, dislocation activity is reduced, and the structural relaxation is dominated by thermal perturbation and dislocation annihilation, which may point to supercooled melting or the formation of the pre-amorphous state. In the A2 stage (after 30 ps), the FCC residual structure tends to stably decay, and the types of other atoms continue to increase, reflecting that the system structure is gradually shifting to a highly amorphous stage, while the dislocation mechanism gradually gives way to the reconfiguration process driven by atomic topological disorder.

In contrast, the [110] orientation (Figure 4b) exhibits a stronger dislocation excitation ability. The Shockley dislocation maintains a longer life cycle after shock loading, and its total length peak is higher, suggesting that this crystal orientation has a lower activation threshold for the slip system. The continuous increase in the proportion of HCP atoms and the peak of the total dislocation length show a high degree of synchronization, further verifying the direct dynamic coupling mechanism between dislocation slip and the phase transition path. For this crystal, partial dislocation slip, stacking, and the induced layering structure constitute the main path of structural transformation. In contrast, the [111] orientation (Figure 4c) exhibits the strongest lattice stability. The attenuation process of the FCC structure is significantly delayed, and the increase in the proportion of HCP and other atoms is more gentle, indicating that the structural reconstruction path is more restricted and the degree of dislocation activation is relatively mild. Although the Shockley type dislocation remains the dominant defect type, its total length remains at a medium level, highlighting the orientation-dependent dislocation response of the shock response. By integrating the evolution characteristics of the three crystal directions, a coupled mechanism of phase transition and dislocation dynamics can be proposed: The differences in atomic stacking sequences for different crystals significantly affect the activation path, dislocation evolution behavior, and structural reconstruction rate of the slip system, thereby controlling the multi-stage transformation process from FCC to HCP and even amorphous structures.

3.4. Distribution and Evolution of Defect Types Under Impact

During the intense non-equilibrium deformation process induced by impact, the microstructure evolution of crystal materials exhibits a complex and highly anisotropic defect synergy mechanism. Specifically, TiAl alloys undergo a transition from face-centered cubic (FCC) to closely packed hexagonal (HCP) phases under high strain rate loading that is also commonly accompanied by Intrinsic Stacking Faults (ISFs), Extrinsic Stacking Faults (ESFs), and twin boundaries. The rapid formation and dynamic recombination of microstructure defects, such as TBs and others [51], do not evolve in isolation but constitute a multi-level defect system that is synergistically coupled with local structural disturbances, dislocation slip, and phase transition transformation.

In this system, ISF usually originates from the sliding of Shockley-type partial dislocations along the {111} close-packed plane. Its formation can be regarded as a transitional process from the original ABCABC… type stacking sequence to the local ABAB… type rearrangement, constituting the precursor state of HCP layer keratinization. This structural transition not only marks the initial driving mechanism of the FCC → HCP phase transition but also induces structural reconfiguration of surrounding atomic clusters by changing the local stacking energy, thereby activating a series of complex phase transition–slip co-evolution paths. In contrast, ESF manifests as stacking sequence anomalies caused by asymmetric atomic layer insertions, often occurring in regions with intense interactions of multiple slip systems or high strain gradients, reflecting the role of stacking defects in regulating the stability of local crystals. TB, as a kind of dislocation interface structure with orientation selectivity, usually connects the HCP and the parent phase region in a mirror-symmetric manner, indicating the geometric coupling relationship in the slippy-twin transformation process.

Figure 5 shows the schematic diagrams of the local atomic configurations of typical defects such as ISF, ESF, HCP, and TB, as well as their spatial distribution patterns in different crystal directions. Under the loading condition for the [001] direction, ISF shows a wide range of expansion, and the local lamella tilt and unfold along the {111} group slip plane, indicating that the FCC → HCP phase transition is mainly achieved through the layered rearrangement process induced by ISF. The corresponding HCP atomic configurations are concentrated in the ISF-rich region, verifying the dominant role of Shockley dislocation slipping-induced stack recombination in HCP lamellar nucleation. In contrast, the number of ESF and TB is limited and their distribution is scattered, indicating that the phase–phase-dominated local reconstruction in this direction is superior to the twin mechanism, which limits the expansion and penetration of TB.

In the [110] orientation loading, although ISF is still widely present, what is more characteristic is that TB is highly orderly arranged in a band-like through structure along the main sliding direction, showing a strong coevolution trend between the twinning structure and the HCP lamella. This defect organization pattern reveals that for the [110] orientation, through the expansion and reorganization of the local stacking sequence, it is easy to form a continuous twinning–slip zone structure, thereby enhancing the plastic energy dissipation capacity under impact loads. ESF is also active in this direction, mainly located near the twinning boundary and the parent phase interface, reflecting the complex stacking disturbance behavior under the interaction of the slip–twinning–phase transformation mechanisms. In contrast, the defect evolution mechanism exhibited by the [111] orientation is significantly different. Although ISF still occurs widely, its distribution is in a diffuse state and lacks a layered structure that is orderly and aggregated along the sliding surface. Meanwhile, the occurrence frequencies of the HCP and TB configurations decreased significantly, and the spatial distribution showed highly localized characteristics. This phenomenon indicates that in the [111] direction, due to the limited number of slip surfaces and the inhibitory effect of the bulk density of plate-highlands, the effective activation of local phase transitions and twinning expansion is restricted. ESF is only sporadically distributed in the high-stress area, further confirming that strain regulation in this direction is mainly carried out in the form of local stacking disturbance rather than structural reconstruction.

In summary, the defect evolution mechanism of TiAl alloys under impact compression conditions shows strong crystal orientation dependence and anisotropic response characteristics: In the [001] and [111] directions, the phase transformation mechanism dominated by ISF is more prominent, especially in the [001] direction, and a relatively stable HCP layer is formed; and in the [110] direction, the continuous strip structure at the twinning boundary is highly developed, demonstrating a typical twinning–slip cooperative path.

It can be seen from Figure 6 that in the low-speed impact stage (0.6–0.9 km/s), the material as a whole still mainly maintains the original face-centered cubic (FCC) structure, and the occurrence of defects is extremely limited. Only sporadic Intrinsic Stacking Fault (ISF) and Extrinsic Stacking Fault (ESF) characteristics could be observed in local areas, and no obvious phase transition behavior occurred. The above phenomena indicate that at this strain rate, plastic deformation is mainly achieved through local slip and small-scale layering errors, which belongs to a typical low-energy consumption micro-regulation mechanism. As the impact velocity increases to 1.2 km/s, different crystal downward deformation modes begin to exhibit obvious anisotropic response characteristics. Among them, in the [001] direction, the ISF laminates expand inclinedly along the {111} slip plane. In the local area, hexagonal closely arranged (HCP) structures and Twin Boundary (TB) bands can be seen coexisting, indicating that the FCC→HCP phase transition mechanism and twin behavior have been initially synergistically activated. The twinning boundary runs through the entire region along the main sliding path in the [110] direction, which is speculated to be closely related to the relatively low local stacking energy in this direction. In the [111] direction, although HCP and ISF defects also occur, their distribution is relatively diffuse and limited in scale, suggesting that the plateau particle packing density structure effectively suppresses the large-scale expansion and phase transition excitation of the defects.

After further increasing the impact velocity to 1.5–1.8 km/s, the material enters a violent defect reconstruction stage inside, presenting typical non-equilibrium structural evolution characteristics under high strain rates. In the [001] and [110] directions, multiple types of defects such as ISF, HCP, TB, and ESF show significant accumulation and spatially interlaced distribution, forming a complex multi-level coupling network. Even in the local regions, other lattice structures such as body-centered cubic (BCC) emerged, indicating that the nonlinear phase transition mechanism driven by lattice polarization and local displacement rearrangement gradually dominates the microstructure evolution process. Among them, the [001] orientation is more inclined to form bulk HCP structure clusters, while the [110] orientation is characterized by the continuous twinning system passing through. Both are conducive to the effective improvement of plastic energy dissipation and the realization of local stress relaxation. In contrast, although the [111] direction also shows obvious damage to the original FCC structure and an increase in the number of other structures, HCP and twinning structures are still relatively sparse, indicating that its deformation mechanism is more inclined to the energy release path dominated by local fractures and stacking faults.

In conclusion, the increasing impact velocity significantly accelerates the accumulation of defect density, induces multiphase transformation processes such as FCC→HCP/BCC, and amplifies the structural evolution differences among different crystal directions. This series of results reveals that under the loading condition of high strain rate, the competitive relationship between the multi-level defect-phase transformation co-evolution mechanism and the slip/twin/rearrangement path constitutes the core intrinsic dynamic mechanism for regulating the microstructure reconstruction behavior and macroscopic mechanical properties of TiAl alloys, providing a new micro-regulation strategy and theoretical support for constructing new high-performance crystal alloys.

Figure 7 shows the evolution characteristics of the radial distribution function g(r) corresponding to the three crystal directions under different impact velocity conditions, systematically revealing the behavior of the evolution of atomic arrangement order and the degradation of local structure stability under impact induction. Specifically, compared with the initial crystal state (black curve), the g(r) curve after impact loading generally shows changes such as peak broadening, amplitude weakening, and peak position shift, indicating that the atomic pair correlation and medium and long-range ordered structure gradually deconstruct under the loading action. This evolution trend suggests that during the impact process, the stress wave propagation within the system interacts with the local non-equilibrium state, which may drive the crystal structure to reconstruct towards quasi-disordered or amorphous states, showing obvious structural response anisotropy and microscopic dissipation characteristics.

In the [001] direction, as the impact velocity increases from 0.6 km/s to 1.8 km/s, the second and third nearest neighbor peaks decay rapidly, especially in the range of 1.5–1.8 km/s. Among them, the mid-range coordination peaks almost completely disappear, and the g(r) curve shows obvious shoulder peak-like local characteristics (see the red circle indication). This phenomenon reveals the significant perturbation of the local coordination configuration by the FCC → HCP phase transition and the Shockley displace-induced dislocation packing process, driving the atoms to undergo high rearrangement and weakening of the order parameter, further supporting the non-equilibrium dynamic reconstruction mechanism of “defect first–structural response lag”. In contrast, the response of the [110] direction is more complex. At a moderate impact velocity (1.2–1.5 km/s), although the second nearest neighbor peak shows a certain degree of attenuation, its clarity is still maintained, indicating that the local lattice order structure is retained to a certain extent. This is consistent with the “twinning dominant plasticity mechanism” proposed earlier: Twins, as the strain dissipation path, effectively disperse stress concentration and slow down the atomic disorderization process. Even under the impact of 1.8 km/s, this direction still shows a certain degree of mid-range orderliness, demonstrating superior structural stability and anti-collapse ability. In sharp contrast, the [111] orientation shows the most significant coordination structure disorder characteristics at high strain rates. When the impact velocity exceeds 1.5 km/s, except for the first nearest neighbor peak, the remaining mid-range coordination peaks almost completely degenerate. The g(r) curve approaches a typical quasi-amorphous state, showing highly atomic disordered behavior. This trend can be attributed to the fact that [111] is the direction of the closest packing, with limited sliding paths and weak dislocation expansion capabilities, resulting in rapid defect accumulation and hindered structural reconstruction, thereby inducing local strain concentration and highly discrete atomic arrangement.

Based on the coordination structure response characteristics of the above three crystal directions, the obvious crystal direction regulation effect can be summarized as follows: The [001] crystal orientation tends to be dominated by the local reconstruction of the formation layer dislocation stacking, the [110] crystal orientation shows a progressive structural transformation path regulated by twinned, while the [111] crystal orientation is more likely to trigger direct and rapid atomic disordered behavior, showing a strong amorphous tendency. This evolutionary heterogeneity reveals the intrinsic regulatory effect of crystal orientation on the stability and collapse resistance of the microstructure and also provides direct atomic-scale support for constructing the coupled model of “impact load–defect evolution–structural transformation”.

4. Conclusions

In this study, atomic-scale molecular dynamics (MDs) simulations were systematically performed to investigate the high-velocity shock response of TiAl alloys along three representative crystallographic orientations: [001], [110], and [111]. By dynamically tracking the evolution of defects, phase transitions, and atomic structural rearrangements, we elucidate the critical role of crystallographic orientation in governing the material’s response under extreme strain rate conditions.

The simulations reveal distinct orientation-dependent responses. The [111] orientation shows the highest propensity for amorphization due to its limited slip systems, whereas the [001] orientation promotes stacking fault formation and HCP phase transitions. In contrast, the [110] orientation maintains superior structural integrity under shock loading, primarily mediated by deformation twinning. With increasing impact velocity, the defect landscape evolves hierarchically—from isolated Shockley partial dislocations to more complex configurations comprising stacking faults, HCP clusters, and amorphous bands. The threshold velocities for the onset of amorphization vary significantly among orientations—approximately 1.5 km/s for [001] and exceeding 1.8 km/s for [110]. RDF analysis further corroborates the anisotropic nature of structural degradation, with [111] being the most prone to atomic disorder and [110] exhibiting the greatest resistance to collapse.

Based on these insights, we propose a unified microscopic framework—crystallographic regulation–defect synergy–phase transition drive—that captures the coupled evolution of defects and phase transformations under shock loading. This mechanistic understanding not only advances the fundamental comprehension of TiAl alloy behavior under extreme dynamic conditions but also offers atomic-scale design principles for orientation-tailored, high-performance structural materials.