A Study on the Mechanical Properties of Ni-Al Alloy Based on Molecular Dynamics Simulation

Abstract

With the wide application of Ni-Al high-temperature materials, the research on their performance has increasingly attracted attention. To further advance the development of Ni-Al high-temperature materials, it is necessary to conduct an in-depth analysis of the brittleness mechanism of Ni-Al intermetallic compounds and elucidate the fundamental nature of their brittleness. In this study, the tensile mechanical behavior and microscopic mechanism of single crystals NiAl (B2) and Ni₃Al (L1₂) at different temperatures were systematically studied by molecular dynamics simulations. It is revealed that although the mechanical properties of both NiAl and Ni₃Al degrade with increasing temperature, their deformation mechanisms exhibit fundamental differences. The high-temperature strength of NiAl is attributed to stable plastic flow dominated by 1/2 <111> screw dislocation. The early softening of Ni₃Al is associated with the formation of stacking fault formation, the phase transition to the HCP, and the slip of various incomplete dislocations (e.g., 1/6 <112> Shockley dislocation). Atomic strain analysis shows that regions of high strain exhibit a strong spatial correlation with the phase-transformed domains. This study reveals the distinct deformation mechanism of the two alloy phases at the atomic scale, providing a key theoretical basis for the rational selection of Ni-Al alloy in specific high-temperature applications.

1. Introduction

With the continuous increase in thrust-to-weight ratio of advanced aviation engines, the continuous increase in combustion chamber temperature, and the increasing complexity of engine structures, weight issues have become increasingly prominent [1,2,3]. Therefore, there is an urgent need to develop new high-temperature structural materials that offer superior thermal stability, lower density, higher strength, and improved processability to replace conventional traditional Ni-based superalloys in hot-section components of aviation engines [4,5,6]. Among the candidate materials, Ni-Al intermetallic compounds exhibit potential for high-temperature applications and primarily exist in five phase states: NiAl, Ni₃Al, NiAl₃, Ni₂Al₃, and Ni₅Al₃ [7,8]. The melting point of common Ni-Al intermetallic compounds can reach up to 1600 °C, approximately 300 °C higher than those of conventional Ni-based alloys. Among these, NiAl and Ni₃Al—exhibiting high melting points of 1638 °C and 1395 °C, respectively—have received widespread attention as a substitute for Ni-based superalloys in aviation engine blades or thermal barrier coatings [9,10]. Moreover, NiAl possesses not only high thermal conductivity but also excellent resistance to high-temperature oxidation [11,12].

Previous studies on Ni-Al alloys have primarily focused on alloy design, processing routes, and microstructural strengthening mechanisms. For example, Y. Zhu et al. [13] used a complex solid compound as a toughening phase to improve the ductility of Ni-Al alloy during annealing. M. Krasnowski et al. [14] found that nanocrystalline Ni-Al materials have higher hardness and stronger heat resistance than microcrystalline Ni-Al alloys. G.P.P. Pun et al. [15] conducted a study on the impacts of chemical composition and uniaxial mechanical stress on martensitic transformation in Ni-Al alloy with a high nickel content through molecular dynamics simulation. A.M. Roy [16] investigated the influence of interfacial stress on the morphological evolution of martensitic nano-structured Ni-Al alloy and determined the evolution of the microstructure and stress field. Y. Du et al. [17] examined the precipitation behavior of Cu-Ni-Al nanoparticles in high-strength low-alloy steel and its influence on mechanical properties through atomic probe tomography and tensile tests. Although there are many studies on Ni-Al alloy, most of them focused on material preparation, doping, and performance analysis of polycrystalline materials, while the research on single-crystal structure is relatively limited. Because single-crystalline structures are superior to polycrystalline structures in mechanical properties and high-temperature durability, it is of great significance to study the characteristics of the single-crystal nickel–aluminum alloy for manufacturing high-temperature blades, turbines, etc. In this study, molecular dynamics methods were adopted to investigate the biaxial tensile process of NiAl and Ni₃Al single-crystalline structures at the atomic scale, enabling a detailed examination of the evolution of their mechanical properties, crystal structures, and defects.

Despite these efforts, several critical issues remain insufficiently addressed. Most experimental and computational studies have focused on polycrystalline or nanocrystalline Ni-Al alloys, in which deformation behavior is predominantly governed by grain boundaries and phase interfaces, thereby obscuring the intrinsic mechanical response of the underlying crystal lattice. Furthermore, systematic investigations of temperature-dependent mechanical behavior in single-crystal NiAl and Ni₃Al under complex loading conditions are lacking, particularly at the atomic scale. Moreover, direct comparative analyses between NiAl and Ni₃Al single crystals remain limited. As a result, the fundamental mechanisms governing temperature-induced changes in elastic–plastic behavior, defect evolution, and crystal structure stability in these two materials remain incompletely understood.

Given that single-crystal materials generally exhibit superior high-temperature strength, creep resistance, and fatigue performance compared to their polycrystalline counterparts—particularly in demanding applications such as turbine blades and other hot-end components—systematic, temperature-dependent atomic-scale investigations of single-crystal NiAl and Ni₃Al are critically important for elucidating their intrinsic deformation mechanisms.

In this study, molecular dynamics simulations were employed to investigate the biaxial tensile deformation behavior of single-crystal NiAl and Ni₃Al across a broad temperature range. By analyzing stress–strain responses, crystal structure evolution, and defect formation mechanisms at the atomic scale, this study provides new insights into the intrinsic mechanical properties and thermal sensitivity of these intermetallic compounds. It should be noted that the findings pertain specifically to the NiAl (B2) and Ni₃Al (L1₂) phases investigated; direct extrapolation to Al-rich brittle phases is not warranted. The results aim to clarify the fundamental differences between NiAl and Ni₃Al single crystals under thermomechanical loading, thereby offering theoretical guidance for the design and application of Ni-Al-based single-crystal materials in high-temperature aerospace environments.

2. Modeling and Simulation Details

In this work, the tensile deformation of the Ni-Al alloy was simulated by using LAMMPS (LAMMPS (2023), Sandia National Laboratory, Albuquerque, NM, USA) [18], a large-scale atomic/molecular parallel simulator. LAMMPS was selected for its high efficiency in simulating atomic dynamics in large-scale metallic systems [19], its ability to accurately describe the interatomic bonding and lattice evolution, and its suitability for resolving defect evolution and atomic reconstruction in the tensile process. The simulation model of the Ni-Al alloy is displayed in Figure 1. To accurately capture the tensile deformation process, a two-end loading method was employed in this study. The NiAl system exhibits a body-centered cubic (BCC) crystal structure with a lattice constant of 2.86 Å, while Ni₃Al adopts a face-centered cubic (FCC) crystal structure with a lattice constant of 3.52 Å [20]—slightly below the typical values of 3.56–3.58 Å at room temperature [21,22,23]. In order to obtain suitable tensile models, supercells of 24 × 24 × 52 unit cells for NiAl (containing 59,904 atoms) and 19 × 19 × 42 unit cells for Ni₃Al (containing 60,648 atoms) were generated. Periodic boundary conditions were employed during the simulation process to avoid size effects. Prior to tensile loading, each system underwent a 60 ps relaxation process to relax internal stresses, eliminate artifacts from model construction, and ensure the system reached thermal equilibrium at the target temperature. The simulation adopted the canonical ensemble (NVT) to precisely control temperature and study its effect, a standard approach for high-strain-rate MD simulations with limited computational resources [24]. It is acknowledged, however, that the constant-volume constraint of NVT may influence stress fluctuations and dislocation kinetics, a limitation considered in the interpretation of the results [25,26,27]. The loading rate was 0.005 Å/ps (i.e., 0.5 m/s). The simulation time step was set to 1 fs.

The interatomic potential function serves as the cornerstone of MD simulations. The selection of potential functions directly influences the accuracy and reliability of the calculations. In this study, the embedded atom method (EAM) potential function developed by Zhou et al. [28] was utilized to describe Ni–Al interactions. The EAM potential is primarily used to depict the interactions among metals [29]. The formula for calculating the potential energy E is as follows:

$$E=F_{\alpha }\left(\sum _{j\ne i}{\rho }_{\beta }\left(r_{ij}\right)\right)+{\displaystyle \frac{1}{2}}\sum _{j\ne i}{\varphi }_{\alpha \beta }\left(r_{ij}\right)$$

(1)

where F represents the embedding energy, which is dependent on the electron density ρ of atoms, ϕ denotes a pair potential interaction, and α and β are the element types of atoms i and j, respectively, and r is the distance between atoms i and j. The multi-body characteristic of the EAM potential is determined by the embedded energy term. The two summation terms in the formula are performed for all neighboring atoms j within the range of atom i, which are within the cutoff distance.

To conduct a detailed analysis of the subtle changes at the atomic level, OVITO 3.0 software was employed for visualization processing [30]. OVITO was selected for its robust and well-validated built-in analysis modules. Specifically, the common neighbor analysis (CNA) was utilized to examine the changes in the crystalline structure [31]. The atomic strain was adopted to analyze the strain of atoms [32]. The dislocation analysis (DXA) was applied to identify the defect structures [32]. Additionally, coordination number analysis was used to track variations in local atomic coordination and interatomic bonding during tensile deformation [33].

3. Results and Discussion

3.1. Tensile Analysis of NiAl

The tensile curves of NiAl under different temperatures are displayed in Figure 2. Based on these results, the tension process could be divided into four stages (corresponding to regions I, II, III, and IV in the figure): (I) there is a linear increasing relationship between stress and strain; (II) the slope of the curve has decreased, yet the relationship between stress and strain still remains approximately linear; (III) as the strain increases, the stress gradually rises and attains its maximum value; and (IV) as the strain continues to increase, the stress declines and the calculation comes to an end. The stress–strain results indicate that the fracture strain of the NiAl structure gradually diminishes with the increase in temperature, which implies that the rise in temperature causes a weakening of the deformation resistance. Correspondingly, the elastic moduli of NiAl declines from 116.37 GPa at 500 K to 102.62 GPa at 1000 K and further to 86.34 GPa at 1500 K. This reveals that, as the temperature goes up, the elastic modulus of the NiAl alloy gradually decreases. This trend arises because higher temperatures amplify atomic thermal vibration, increase the average interatomic spacing, and weaken interatomic bonding—thereby diminishing the material’s capacity to resist elastic deformation [34,35]. Additionally, the illustration in the figure depicts the crystalline structure of NiAl under the corresponding strain conditions, indicating a significant phase transition process during the tensile process of NiAl. This is because as the temperature increases, the binding force between atoms weakens and the lattice energy decreases, which weakens the lattice stability and makes it easier for the system to break through the phase transition energy barrier to produce phase transition [36].

The variation in the crystalline structure of NiAl under corresponding strain conditions at different temperatures is shown in Figure 3. It can be observed that, as the temperature increases, the number of amorphous phase transition atoms increases at the same strain rate. During the elastic–plastic deformation stages (i.e., I–III stages), the B2-ordered BCC structure of NiAl progressively transformed into both FCC-like and amorphous phases, accompanied by a continuous decline in the BCC fraction. As temperature rose, the extent of transformation to the FCC phase diminished, while the amorphous content steadily grew. After the occurrence of cracks and voids in stage IV, the proportion of BCC structure gradually increased due to the rearrangement of atoms at both ends of the defect after fracture. Notably, the higher proportion of BCC in Figure 3c at 20% strain compared to Figure 3b was precisely due to the lattice reconstruction after defect formation. The primary reason is that, in the stress concentration area during the pre-fracture tensile process, local strain induces crystal structure rearrangement, and the mechanical load promotes the B2-ordered BCC phase to transform into metastable FCC-like and amorphous phases, thereby alleviating stress concentration and delaying crack propagation. Once a crack forms, the external load is partially released, and the system relaxes toward its thermodynamically favored state. Since the BCC lattice represents the lowest-energy configuration for NiAl [37], atoms near the fracture surfaces tend to revert to this stable phase during post-fracture relaxation.

When the tensile strain was 35%, the statistical results of dislocation defects in NiAl structures at different temperatures are presented in Figure 4. It could be found that the dislocation defects in the NiAl alloy during the tensile process were mainly screw dislocations with a Burgers vector of 1/2 <111> and gap dislocation loops with a Burgers vector of <100>. This is the same as the dislocation type in other BCC crystals [38,39]. The reason for this phenomenon is that, in BCC metals or their alloys, deformation is mainly achieved through double twisted nucleation and propagation on screw dislocations [40,41]. Notably, <100>-type dislocations significantly influence the mechanical properties of BCC metal, and they will transform into 1/2 <111> dislocations. As temperature increased, the number of dislocation defects in the system gradually decreased under the same tensile strain conditions, and the defect mesh in the system gradually became fragmented, indicating the formation of more small clusters or voids in the system.

The atomic strain distribution in NiAl at a tensile strain of 35% across different temperatures is shown in Figure 5. This result indicates that as the temperature increased, the strain intensity of atoms in the matrix gradually increased after tensile (i.e., the number of atoms with cyan color in the matrix gradually increases). Compared with the results in Figure 3, it was found that the position of the high-strain atoms in Figure 5 was highly consistent with the atomic phase transition region during the tensile process. This indicated that the change in interatomic bond length caused by the phase transition leads to an increase in atomic strain. Figure 6 shows the radial distribution function (RDF) of the NiAl structure after tensile. Specifically, Figure 6a depicts the changes at different temperatures, while Figure 6b illustrates the changes at different strains at 500 K. The first and second peak positions of the RDF correspond to the positions of the nearest and second-nearest neighbor shells around a central atom, respectively. As shown in Figure 6a, with increasing temperature, the intensity of the first RDF peak decreases, the number of the nearest-neighbor atoms reduces, the probability of bonding between adjacent atoms declines, and the short-range ordering degree of atoms gradually diminishes. Meanwhile, the other peaks gradually decline, the peak shape broadens, and the peak position shifts slightly to the right, suggesting that the structure of NiAl has changed, displaying characteristics of short-range order and long-range disorder. As shown in Figure 6b, during tensile deformation from 0 to 20% strain, the probability of atomic bonding gradually decreased and the interaction force gradually weakened. However, when the strain increases further from 20% to 30%, the first peak unexpectedly intensifies and shifts leftward (to shorter interatomic distances). This behavior coincides with the occurrence of crack defects (Stage IV), the external forces acting on the atoms in the system gradually decrease, leading to the transformation of the system stress during tension into interatomic forces. As a result, atoms rearrange into a more compact configuration, strengthening local bonding and leading to a gradual shortening of bond lengths [42].

3.2. Tensile Analysis of Ni₃Al

The stress–strain response of Ni₃Al under different temperatures are presented in Figure 7. Similarly to NiAl, the tensile deformation of Ni₃Al can be divided into four stages. With increasing temperature, both the deformation resistance and overall mechanical strength of Ni₃Al progressively decline. Under identical strain conditions, Ni₃Al exhibits lower stress and lower elastic modulus compared to NiAl. This is because NiAl is a B2-type ordered BCC structure, Ni and Al atoms occupy the body center and vertex positions, respectively, and the atoms in the lattice are relatively tightly packed, so the lattice has a strong intrinsic ability to resist deformation [43]. Ni₃Al possesses an ordered FCC structure of the L1₂ type, in which Al atoms occupy the cube corners and face centers, while Ni atoms fully occupy the octahedral interstitial sites. This arrangement results in a lower atomic packing density compared to the B2-ordered BCC structure of NiAl. More crucially, the L1₂ lattice contains numerous parallel {111} slip planes, and the interlayer bonding force between slip planes is weaker than that between lattices in the B2 structure, so it is easier to have relative slip between layers under the action of external force, which shows lower tensile stress and elastic modulus macroscopically [44]. Moreover, the strain of Ni₃Al during its elastic–plastic deformation stage was smaller. These results indicated that, at equivalent temperatures, Ni₃Al exhibits inferior mechanical performance relative to NiAl. The corresponding crystalline structure changes indicated that there was a significant phase transition in Ni₃Al before and after tensile fracture, and the FCC structure was restored to a certain extent after fracture.

The evolution of the crystalline structure in Ni₃Al under increasing tensile strain is illustrated in Figure 8. During the stretching process, the phase transition caused by atomic displacement first occurred in Ni₃Al alloy, followed by the progressive nucleation and growth of voids and cracks as strain accumulated. In contrast to NiAl, Ni₃Al exhibited a pronounced tendency to undergo a deformation-induced transition to an HCP phase, which is attributed to lattice slip during tensile process. As the temperature increased, the proportion of amorphous phase transformation under the same strain increased, and the proportion of HCP first increased and then decreased. In addition, as the temperature increased, the fracture strain of Ni₃Al gradually increased, the fracture toughness increased gradually, the fracture cavity in the system decreased gradually, and the strain should be higher when it was completely fractured.

The results of dislocation defects in the Ni₃Al system at different temperatures under a strain of 25% are presented in Figure 9. The dislocation types identified in the Ni₃Al alloy during stretching include Shockley dislocations with a Burgers vector of 1/6 <112>, Stair rod dislocations with a Burgers vector of 1/6 <110>, Hirth dislocations with a Burgers vector of 1/3 <100>, Frank dislocations with a Burgers vector of 1/3 <111>, Perfect dislocations with a Burgers vector of 1/2 <110>, and other dislocations. Among these, Shockley dislocations and Stair rod dislocations are the predominant ones. Additionally, as the temperature rises, the total dislocation length in the Ni₃Al system gradually shortens, the defect grid becomes fragmented, and the concentration of vacancy defects and small clusters progressively rises. Moreover, the dislocation density in Ni₃Al remains higher than that in NiAl at a temperature below 1500 K, and the defect mesh fragmentation is more severe at 1500 K. This is because NiAl possesses a B2-type ordered BCC structure, its effective slip system is mainly <111> {110} at room temperature to moderate temperature, the number of slip systems is small, and the dislocation is 1/2 <111> total dislocation, so the lattice resistance and antiphase domain boundary (APB) resistance to be overcome in slip are high [45,46]. High resistance will inhibit the initiation and proliferation of dislocations, so the dislocation density is low. On the contrary, Ni₃Al is an ordered FCC structure of L1₂ type, with effective slip systems of <110> {111} and <110> {112}, and the number of slip systems is far greater than that of NiAl’s B2 structure. At the same time, incomplete dislocation dominates the slip, and the critical cleavage stress (CRSS) required for slip is significantly lower than that of NiAl’s total dislocation [47,48]. The low resistance enables NiAl to quickly start a large number of slip systems under a small external stress, which leads to the continuous proliferation of dislocations and directly increases the dislocation density. This indicates that Ni₃Al is more likely to fail than NiAl under the same working conditions.

The atomic strain results in the Ni₃Al alloy were represented in Figure 10. Similarly to the behavior observed in NiAl, atoms exhibiting high local strain in Ni₃Al are predominantly localized in regions undergoing phase transformation. The highest strain concentrations occur at the crack tip, where the atomic-level strain reaches its maximum intensity. As the temperature increased, the number of high-strain atoms progressively rose, which was related to the enhanced extent of atomic-scale phase transitions within the system. Moreover, the range of atomic high-strain regions within the system gradually increased, indicating that the deformation resistance of the system gradually weakened as the temperature increased. Figure 11 shows the evolution of RDF for the Ni-Al atomic pairs in the Ni₃Al system. Figure 11a showed that as the temperature increased, the Ni–Al interatomic interaction gradually weakened, and the peak position gradually shifted to the right. This shift indicated thermal expansion and an increase in the average Ni–Al bond length. Figure 11b displays the RDF evolution at 500 K under different tensile strain. Within the strain range of 0%–10%, as the tensile progresses, the interatomic bond length in the system increases, leading to a gradual weakening of atomic interactions, a rightward shift in the peak position, and a reduction in peak intensity. However, when the strain reached 20%–25%, the structural changes induced by atomic stretching cease due to the rupture of chemical bonds at both ends of the crack. As a result, the system underwent contraction, the bond length decreased, the interatomic interaction strengthened, the peak position shifted to the left, and the peak intensity increased.

4. Conclusions

In this study, the mechanical properties of single-crystal NiAl and Ni₃Al under high-temperature tensile loading and their underlying microscopic mechanism were systematically investigated using molecular dynamics simulations. The results demonstrate that the superior high-temperature deformation resistance of NiAl originates from its stable plastic deformation mechanism, which is dominated by 1/2 <111> screw dislocation, whereas the early softening behavior observed in Ni₃Al is closely associated with complex stacking fault, HCP phase transformation, and slip systems governed by Shockley incomplete dislocation (1/6 <112>). Atomic strain analysis confirms that regions of high strain are highly consistent with phase transformation zones. The increase in temperature generally weakens atomic interactions and promotes the amorphous phase transition. This work clarifies the fundamentally distinct deformation mechanisms of Ni–Al-based intermetallic compounds at the atomic scale, providing a crucial theoretical foundation and design guidance for the rational selection and design of Ni-Al-based intermetallic compounds for specific high-temperature working conditions, such as high-stress components or impact-resistant coatings.