Role of chemical disorder and local ordering on defect evolution in high-entropy alloys

High-entropy alloys (HEAs) have stimulated great interest due to their remarkable mechanical and irradiation performance. Experiments suggest that delayed defect evolution in HEAs, compared to conventional metals and dilute alloys, is the main reason for their improved irradiation resistance. However, the mechanism responsible for the observation remains elusive. Here we show that the potential energy landscape of defects under the influence of random arrangement of different species is the reason for the delayed defect evolution. We arrive at the conclusion by investigating the diffusion of defects and defect clusters under three cases: the averaged-atom model, random model, and the model with local short-range ordering. Our results suggest that, compared to the average model, the chemical fluctuation inherent in HEAs can suppress interstitial motion more than vacancy motion. The effects are more pronounced when SRO develops. For defect clusters, the chemical disorder can reduce their jump frequencies significantly and enhance correlation effects, leading to suppressed defect motion. Notably, we find that with SRO, such defect motion can be entirely trapped in local regions. This work demonstrates that chemical fluctuations and SRO are the main reason responsible for the suppressed defect evolution in HEAs, which dictates a promising way to improve the irradiation performance of HEAs through manipulating its chemical disorder states, such as local ordering.

Figure 1

The SRO parameters for the random (a) and local ordering (b) model, respectively. The local ordering model is built through our MC procedure. Note the different scales of the $y$ axis in these two subplots, which demonstrate that SRO parameters in the random model are all around zero, while high values of SRO parameters are found in the local ordering model.

Figure 2

Calculated tracer diffusion coefficient (left column) and defect diffusion coefficient (right column) for a point defect (single vacancy [(a),(c)] and interstitial [(b),(d)]) in CuNiCoFe based on the three different models: average atom model, random model, and SRO model.

Figure 3

Correlation factors for defect jumps during diffusion of a vacancy and interstitial in the three considered models.

Figure 4

Fraction of different dumbbell compositions during interstitial diffusion at different temperatures in the random (a) and SRO (b) model.

Figure 5

Calculated activation energies for vacancy (left) and formation energies for interstitials (right) in the random and SRO model. The dotted line denote the values in the averaged atom model. The averaged values for each defect type are also shown.

Figure 6

Formation energies (eV/defect) of interstitial defect clusters in the considered HEA with random elemental arrangement. The dotted line represents the results obtained in the average atom model.

Figure 7

Trajectory of a ${\mathrm{I}}_{15}$ interstitial cluster simulated at 1000 K in the average atom model, random model, and the SRO model. The elemental distribution is shown on the left, with red balls representing Cu. The trajectories are given on the right.

Figure 8

Evolution of local atomic environment of a ${\mathrm{I}}_{15}$ cluster at 1000 K in the random model (left) and the SRO model (right). The number of different elements surrounding every atom in the cluster within a 5 Å radius is shown (the atoms in the clusters are excluded).

Figure 9

Jump frequency of interstitial clusters with different sizes in the average atom model (a), random HEA (b), and the HEA with local ordering (c).

Figure 10

Correlation factor of interstitial clusters with different sizes in the average atom model (a), random HEA (b), and the HEA with local ordering (c).

Figure 11

Jump frequency of vacancy clusters with different sizes in the average atom model (a), random HEA (b), and the HEA with local ordering (c).