Dislocation evolution and hardening of CoCrFeMnNi high entropy alloy under Fe ion irradiation at room temperature and 500 °C

As the increase in energy demand, fusion energy is attracting more and more attention. The material issue is currently a key limiting factor in the development of fusion energy, as service materials are required to withstand extreme conditions of irradiation, temperature, and stress, and to have long-term stable properties. High entropy alloy (HEA) is an emerging multi-major element solid solution alloy.¹ Compared with conventional alloys containing only one major element, HEA exhibits excellent mechanical properties^2,3 and high-temperature stability^4,5 due to its random arrangement of elements and complex local chemical environment characteristics. Therefore, HEA has become a research hotspot as a new candidate structural material in nuclear reactors.^6–8 

High-energy heavy ion irradiation is a widespread and effective tool for studying neutron-induced cascade damage. Cascade collision-induced interstitials are capable of forming dislocation defects, which can cause deterioration behaviors such as hardening and embrittlement of the material, thus seriously affecting the subsequent serviceability of the material. It is well known that dislocations are sensitive to irradiation temperature, and different irradiation temperatures significantly change their behaviors, such as nucleation, growth, migration, and dissociation. Due to its special crystal structure, the temperature effect of irradiation damage in HEAs is different from that in conventional materials. Zhang et al. found that small and dense dislocations were generated in the irradiated region of CoCrFeMnNi HEAs after Fe ion irradiation at room temperature (RT); moreover, the distribution law of dislocations in HEAs was similar to that in pure metals under RT irradiation conditions.⁹ With increasing irradiation temperature, the dislocation loops in HEA increased in size and decreased in density, and this trend was also consistent with that in pure metals.^10–12 However, at the same temperature, He et al. found that dislocation growth in HEAs was significantly inhibited compared to pure metals, due to their high chemical complexity reducing point defect survival and dislocation migration.¹³ Meanwhile, under Ni ion irradiation at 500 °C, Lu et al. observed that dislocations in CoCrFeNi HEAs and CoCrFeMnNi HEAs were generated only in the region from the surface to near the damage peak, but dislocations in the pure Ni migrated to the region beyond the range of injected ions, which they attributed to the dislocation migration in HEAs followed a three-dimensional (3D) short-range mode due to the rugged energy landscape rather than a one-dimensional (1D) long-range mode like pure metals.¹⁴ However, under Ni ion irradiation at 540 and 580 °C, Su et al.¹⁵ and Fan et al.¹⁰ observed dislocations migrating to regions beyond the range of injected ions similar to those of pure Ni in CoCrFeMnNi HEAs and CoCrFeNi HEAs, respectively, which they attributed simply to the promotion of dislocation migration by the high temperature. Obviously, at different irradiation temperatures, the dislocation behavior in HEAs is sometimes similar to and sometimes different from that of pure metals, and the physical law behind it is not yet clear. Therefore, it is necessary to study the temperature effect law of dislocation behavior in HEAs and elucidate the physical mechanism. In addition, irradiation-induced dislocations significantly influence the hardness of the material, which may threaten plasticity, and it is also of interest to study irradiation hardening.⁸ 

In this study, the CoCrFeMnNi HEA was chosen as the research object. It is one of the most widely studied HEAs in the field of irradiation damage¹⁶ and exhibits good irradiation resistance, such as inhibition of irradiation swelling¹⁷ and segregation,¹⁸ slowdown in defect accumulation¹⁹ and bubble growth,²⁰ resistance to thermal shock irradiation-induced cracking.²¹ The effect of irradiation temperature on the irradiation-induced defects and the surface nano-hardness of CoCrFeMnNi HEAs was investigated using 19.6 MeV Fe ion irradiation at RT and 500 °C. This study contributes to a deeper understanding of the influence mechanisms of temperature effects on irradiation-induced microstructural evolution in HEAs.

The CoCrFeMnNi HEA was prepared by vacuum arc melting with elemental purity higher than 99.95 wt. %. The CoCrFeMnNi HEA was remelted and flipped over five times to ensure a homogeneous distribution of constituent elements.^22,23 The material surface was polished to a mirror finish by using a sandpaper and a diamond paste and finally ultrasonically cleaned with alcohol before irradiation. The CoCrFeMnNi HEA was irradiated using 19.6 MeV Fe ions with the dose of 3 × 10¹⁶ ions/cm² at RT and 500 °C. Irradiation experiments were performed at the Low Energy Highly-Charged Ion Accelerator Facility (LEAF) at the Institute of Modern Physics, Chinese Academy of Sciences.

The distribution of displacement damage and ion concentration with depth in CoCrFeMnNi HEAs after Fe ion irradiation was simulated by Stopping and Range of Ions in Matter 2013 (SRIM-2013) code,²⁴ with the displacement energy of 40 eV for all alloying elements.²¹ Thin foils for transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM) observations were prepared by the focused ion beam (FIB, Helios G4 UX) technique and then observed with a JEM-2100F microscope operated at 200 kV. The dislocation loop distribution was observed using STEM bright-field (BF) images under the [110] zone axis, and the STEM mode has advantages in the observation of dislocation loops,²⁵ on one hand, it has the ability to exhibit all dislocation and dislocation loop structures in the thin foil, and on the other hand, it has improved signal-to-noise ratio, effectively reducing background contrast disturbances such as thickness fringes and bend contours which contained in conventional TEM BF images. The phase structure of CoCrFeMnNi HEAs after irradiation was observed using the selected area electron diffraction (SAED) method. The nano-hardness changes of CoCrFeMnNi HEAs were measured by a Nano-indenter (FISCHERSOPE HM500) equipped with a Berkovich triangular pyramid indenter. The maximum load was 300 mN and loading time was 20 s.

Figure 1 shows the distributions curves of displacement damage and ion concentration with depth of CoCrFeMnNi HEAs after 3 × 10¹⁶/cm² Fe ions irradiation calculated by SRIM. The displacement damage was calculated by both Quick Kinchin–Pease (K–P) and Full-Cascade (F-C) modes. Because of the difference in computational principles, DPA predicted by the F-C model are approximately 2.3 times higher than those predicted by the K-P model.²⁶ According to the report of Weber and Zhang,²⁷ for multi-element materials such as HEAs, the F-C model predicts a more accurate damage profile. Furthermore, considering that pure metals, conventional alloy, and early HEA irradiation studies mostly used the K-P model, the result of this model was also plotted in order to facilitate direct comparisons with other studies. According to SRIM calculations, the damage produced by Fe ions with 19.6 MeV can reach a maximum depth of 3.5 μm, with the peak dpa of 62.3 (F-C) or 28.3 (K-P), and the peak ion concentration of about 0.7 at. %.

Figure 2 shows the STEM BF images of defect distribution of CoCrFeMnNi HEAs after 3 × 10¹⁶/cm² Fe ion irradiation at RT and 500 °C. The defect content can be reflected in the contrast of STEM images, as the high strain field of defects leads to electron scattering and darkens the local contrast in the BF images. As can be seen from Figs. 2(b) and 2(c), RT irradiation-induced defects had significantly higher density compared to that of elevated temperature (ET) irradiation. For Fe ion irradiation at RT, irradiation-induced dislocations were produced from the surface to the end of ion range with 3.5 μm depth. When the irradiation temperature was increased to 500 °C, the irradiation-induced dislocations could only extend to about 3.3 μm depth. It is worth noting that the dislocation distribution was inhomogeneous along the depth in the irradiation-affected region of CoCrFeMnNi HEAs. Combined with the dpa curves in Fig. 2(a), the distribution of defects can be divided into three regions: the front region, the damage peak region, and the damage attenuation region. After RT irradiation, the contrast in front region, damage peak region, and damage attenuation region was shown as dark, bright, and darker, respectively. The distribution of defects in the damage attenuation region was in the form of black bands with the width of about 200 nm. After irradiation at 500 °C, the contrast in the front region was bright and the damage peak region became the darkest region. The morphology of the damage attenuation region was consistent with that of the region beyond the Fe ion range, indicating that there were no visible irradiation-induced dislocations in the damage attenuation region at 500 °C. Figures 2(d) and 2(e) show the SAED images of the damage peak region after RT and 500 °C irradiations, respectively. As can be seen from the SAED images, CoCrFeMnNi HEAs still maintained the pristine face-centered cubic single-phase structure after Fe ion irradiation at both RT and 500 °C in this study. The CoCrFeMnNi HEA has good phase structure stability under irradiation, and it can maintain the face-centered cubic single-phase structure even under the irradiation condition of 580 °C and 53 dpa (K-P).²⁸ 

Figure 3 shows the magnified images of the defect morphology in different regions of CoCrFeMnNi HEAs after 3 × 10¹⁶/cm² Fe ion irradiation at RT and 500 °C. When the irradiation temperature was RT, the front region consisted of dense small-sized dislocation loops, which appeared to be independent of each other and were not tangled together. The density of dislocation loops in the damage peak region was obviously reduced compared to that of the front region. Tangled small dislocation loops were observed in the damage attenuation region and were smaller and denser compared to that of the front region. These tangled and dense small dislocation loops generated a large localized strain field that caused strong electron scattering, and thus the damage attenuation region showed the black band-like morphology with a very dark contrast in STEM BF images. When the irradiation temperature was 500 °C, sparse large-sized dislocation loops and dislocation lines were observed in the front region. The dislocation defects in the damage peak region were denser than those in the front region, and the dislocation loops and lines in this region appeared to tend to form dislocation networks. No irradiation-induced dislocation defects were observed in the damage attenuation region.

It can be seen that there was “shadow contrast” within the elliptical dislocation loops in the STEM images [as shown by the blue arrows in Fig. 3], which were typical features of the 1/3 ⟨111⟩ faulted loops.²⁵ 1/2 ⟨110⟩ perfect loops and 1/3 ⟨111⟩ faulted loops are the two most common forms of dislocation loops in face-centered cubic materials. In the STEM observations, due to the interference between the electron beams diffracted by the perfect crystal planes and the faulted crystal planes, the “shadow contrast” is shown within the faulted loops.²⁵ Ni-based concentrated solid solution alloys tend to form 1/3 ⟨111⟩ faulted loops through interstitials after irradiation,^29,30 and with the increase of dose¹⁸ and temperature,¹¹ the faulted loops gradually transform into perfect loops. This transition is suppressed with increasing chemical complexity of the alloy,¹⁸ and for CoCrFeMnNi HEAs, the dislocation loops are still dominated by faulted loops even at 500 °C and 38 dpa (K-P) dose.¹⁸ Therefore, under the irradiation conditions of this study, most of the dislocation loops produced in CoCrFeMnNi HEAs were 1/3 ⟨111⟩ faulted loops. The contribution of faulted loops to irradiation hardening is greater than that of perfect loops, because the faulted loops are considered sessile, whereas the perfect loops are movable.²⁵  Figure 4 shows the rel-rod TEM dark-field images and the statistics of the dislocation loop size distribution of the damage peak region in CoCrFeMnNi HEAs after Fe ion irradiation at RT and 500 °C. The rel-rod TEM dark-field image is a unique method of characterizing 1/3 ⟨111⟩ loops in FCC structured materials, by which more accurate measurements of the diameter of 1/3 ⟨111⟩ loops can be made.²⁵ The insertion of the atoms of the 1/3 ⟨111⟩ loops between the pristine crystal planes causes additional diffraction, and the additional 1/2(311) diffraction point under the [110] zone axis [shown by the green arrows in the insets of Figs. 4(a) and 4(b)] can be selected to show the 1/3 ⟨111⟩ loops in the dark-field image. Figures 4(c) and 4(d) show the statistics of the size distribution of dislocation loops in the damage peak region under the two irradiation conditions, respectively, with more than 50 dislocation loops counted in each figure. As can be seen from the images, the size of dislocation loops increased significantly when the irradiation temperature increased. The average size of dislocation loops generated in the damage peak region of CoCrFeMnNi HEAs was 24.3 and 83.2 nm after irradiation at RT and 500 °C, respectively.

Besides dislocation defects, it is well known that void swelling under elevated temperature irradiation is also one of the key problems in the failure of nuclear materials. The production of void which is difficult at RT becomes easier at higher irradiation temperatures.⁸ According to the report of Yang et al.,²⁸ CoCrFeMnNi HEAs can form void swelling under 500 °C irradiation with 53 dpa (K-P), and the voids in HEA usually preferentially generated in the damage attenuation region rather than in the damage peak region.¹⁴ After irradiation at 500 °C in this study, the damage attenuation region of CoCrFeMnNi HEA was under-focus, focus and over-focus imaged to observe the swelling behavior, as shown in Fig. 5. No voids can be observed in CoCrFeMnNi HEA after 3 × 10¹⁶/cm² Fe ions irradiation at 500 °C. There were some randomly distributed nano-sized dark contrast features generated by the FIB in the figure, and they did not satisfy the void determination conditions of white in under-focus and black in over-focus. For conventional nuclear materials such as 316 stainless steel, voids with tens of nanometers in size generated after only 2 dpa irradiation at the same temperature, and the swelling reached more than 6% at 7.3 dpa (K-P).^31,32 CoCrFeMnNi HEA has excellent resistance to irradiation swelling.^14,17,28 On the one hand, the localized complexity of HEA leads to the reduction of the mean free path of electrons by several orders of magnitude, which prolonged the thermal spike time and promoted the recombination between interstitials and vacancies.³³ On the other hand, due to high chemical disorder and lattice distortion of HEAs, migration path of interstitials alters from a 1D long-range model to a 3D short-range model,¹⁴ and the migration energy of interstitials partially overlaps with that of vacancies,^34,35 which leads to more efficient recombination. Both ways are able to reduce the number of irradiation-generated vacancies in HEAs and inhibit the creation and growth of voids.

Figure 6 shows the curves of nano-hardness with the depth in CoCrFeMnNi HEAs before and after irradiation. Due to the large data scatter associated with surface roughness, hardness data within approximately 100 nm of the material surface were discarded (shaded in the figure). As can be seen in the figure, CoCrFeMnNi HEA hardened obviously after irradiation at both RT and 500 °C, and the degree of hardening was greater in the RT irradiation condition. Irradiation-induced defects are the main factor in the hardening of materials,⁸ including dislocations, voids, etc. These irradiation defects can act as barriers to dislocation movement and generate the pinning effect when the material undergoes plastic deformation, leading to significant hardening. Obviously, irradiation-induced dislocations were the main source of hardening for CoCrFeMnNi HEAs in this study. In general, the increment of hardness after irradiation is proportional to the product of the density and diameter of irradiation-induced defects.⁸ For CoCrFeMnNi HEAs after RT irradiation, although the size of the dislocations was smaller, the density of the dislocations was much greater than that of the dislocations produced by 500 °C irradiation, especially in the front region with the largest volume [Fig. 2], and thus the degree of hardening was greater in the RT irradiation conditions.

Using high-energy Fe ions to irradiate CoCrFeMnNi HEAs at RT and 500 °C, it was found that: the experimental depth of dislocation distribution by TEM (D_e) under RT irradiation was the same as the calculated depth of damage stopping by SRIM (D_c), whereas D_e under 500 °C irradiation was smaller than D_c due to the absence of irradiation-induced dislocations in the damage attenuation region. Moreover, the dislocation density of the damage peak region at RT was lower than that of the front region and the damage attenuation region, whereas the dislocation density of the damage peak region at 500 °C was the highest among the three regions.

Figure 7 shows the relationship between D_e and D_c for pure metals and HEAs statistically obtained from the references. The details and corresponding references in the figure are listed in Table S1 in the supplementary material. Figure 7(a) shows the results of room temperature (RT) irradiation for pure metals and HEAs. In most cases, D_e is the same as D_c after RT irradiation. For the two special points in the dashed black ellipse, the incident ion is Au ion, which can generate more high-energy cascade and significantly enhancing the migration of defects due to its very large mass, and this is not discussed in this study.^36–39, Figures 7(b) and 7(c) show the results of elevated temperature (ET) irradiation for pure metal and HEA, respectively. As can be seen in the case of ET irradiation, D_e is larger than D_c in pure metals, whereas D_e tends to less than D_c in HEAs. When the irradiation temperature is RT, the dislocation evolution law of HEAs is the same as that of pure metals, and the damage attenuation regions of both materials exhibit black bands with dense small dislocations.^9,40–42 In contrast, under ET irradiation conditions, pure metals and HEAs show different dislocation distribution laws. When the irradiation temperature is elevated, the dislocations of the damage attenuation region in pure metals gradually underwent long-range migration to the deeper region, and the black band became wider with dislocations became sparser, resulting in the formation of “defect diffusion zone.”^43–45 However, there is generally no long-range migration in HEAs, and the black band of the damage attenuation region disappear on ET irradiation, as a result shown in Fig. 2. Under ET irradiation, swelling of pure metals may also increase the depth of dislocations, but this effect is much smaller than that of thermal migration, which may reach regions more than twice of D_c even at less than 1% of swelling.^44,46 In addition, two reports^10,15 also found that dislocations in HEAs migrated to the region deeper than D_c after ET irradiation [points in the dashed magenta ellipse in Fig. 7(c)]. Compared to other reports, it was found that, in both experiments, the irradiation temperature was higher than the common 400–500 °C, and dislocations migrated to the region far beyond D_c were large-size dislocation lines. Considering that the 3D short-range migration mechanism in HEAs originates from the fact that high-energy barriers in rugged energy landscape can easily change the migration direction of small interstitial-type defects.^14,47–49 As for large-size dislocations, it is difficult to shift the migration direction of the whole dislocation because its different parts are located in different energy landscapes. It follows that when the migration of large-sized dislocations is further enhanced due to higher irradiation temperatures, the rugged energy landscape may have a limited role in hindering their long-range migration.

Irradiation temperature can significantly affect the migration and dissociation of irradiation-induced defects, and Fig. 8 shows the schematic of the temperature dependence of dislocation migration and dissociation in pure metals (such as Ni, W) and HEAs (such as CoCrFeNi systems). For pure metals, long-range migration and dissociation of defects are generally difficult to occur under RT irradiation, and D_e = D_c.^41,42 When the irradiation temperature is increased, the activation of long-range migration of small dislocations happens first in pure metals, since the migration of dislocations occurs much more easily than the dissociation for pure metals,⁶⁹ and small dislocations in the damage attenuation region migrate to the deeper region, resulting in D_e > D_c.^45,55 In addition, long-range migration of large-size dislocations may be able to occur under lower irradiation temperatures than the dissociation of small-size dislocations in pure metals, because that both small and large dislocations can be observed in the region exceeding D_c of pure metals (such as Ni, W, Cr, and Nb) after ET irradiation,^{43,44,46,55,70} suggesting that the small dislocations have not yet dissociated when the large dislocations begin to migrate over a long-range. When the temperature continues to increase, thermal dissociation of small dislocations is activated, in which case D_e remains larger than D_c, but only large-sized dislocations can be observed in the region exceeding D_c.^14,57 Finally, when the temperature is high enough, the thermal dissociation of large dislocations is also activated and all dislocation structures produced by ion irradiation will no longer exist.

In the case of HEAs (such as CoCrFeNi systems), the defects are also unable to migrate over a long range at RT, and D_e = D_c, as shown in this study for RT irradiated CoCrFeMnNi HEAs [Fig. 2(b)]. However, compared to pure metals, on one hand, the rugged energy landscape in HEAs inhibits long-range migration of dislocations;¹⁴ on the other hand, the lower dislocation binding energy in HEAs means that thermal dissociation becomes more easily.⁷¹ With the increase of irradiation temperature, the dissociation of small dislocations in HEAs may be activated before its long-range migration, followed by no small dislocations within the damage attenuation region and D_e < D_c, as shown in this study for CoCrFeMnNi HEAs after 500 °C irradiation [Fig. 2(c)]. When the temperature continues to increase, the hindrance of the rugged energy landscape to long-range migration may be weakened for large dislocations; and at this point, the large dislocations in the damage peak region can migrate to the region exceeding D_c, i.e., D_e > D_c. This is the case for the two points in the magenta ellipse of Fig. 7(c).^10,15 Finally, similar to pure metals, when the temperature is high enough that large dislocations can be thermally dissociated, the dislocation structure is absent in ion-irradiated HEAs.⁴⁶ It should be noted that in Fig. 8, a direct comparison between pure metals and HEA in different temperature ranges may not be appropriate due to the differences in the metal systems (e.g., pure W and CoCrFeMnNi HEAs at the same irradiation temperature), but the dislocation evolution law in the figure is of reference value between similar systems.

Not only interstitial defects, but also vacancies were affected by the irradiation temperature. Under room temperature irradiation, vacancies with high concentration in the damage peak region of CoCrFeMnNi HEAs could not long-range migrate to other regions, resulting in the accumulation of oversaturated and diffusely distributed vacancies in the dislocation-free zones between dislocations. Therefore, the new isolated interstitial atoms were difficult to stay stably in the damage peak region and became nucleation sites for new dislocations, and they were either recombined by these oversaturated vacancies or absorbed by pre-existing dislocations. The existing dislocations gradually merged with each other, eventually leading to a decrease in dislocation density [Fig. 3(b)]. For 500 °C irradiated CoCrFeMnNi HEAs, vacancies were capable of long-range migration so that vacancies with high concentrations in the damage peak region would reduce in concentration due to diffusion.¹⁴ This means that the new dislocation nucleation was still allowed in the dislocation-free zones of the damage peak region. Some of these newly created small dislocations disappeared due to thermal dissociation, but some remained by interacting with pre-existing dislocations and forming stable structures such as dislocation networks [Fig. 3(e)]. Therefore, the dislocation density in the damage peak region of CoCrFeMnNi HEAs was the lowest of the three regions after RT irradiation, whereas highest after 500 °C irradiation.

The CoCrFeMnNi HEA was irradiated by 19.6 MeV Fe ions with the dose of 3 × 10¹⁶/cm² at RT and 500 °C. The effect of irradiation temperature on the irradiation-induced defects and the surface nano-hardness of CoCrFeMnNi HEAs was investigated. The following conclusions were obtained:

1. CoCrFeMnNi HEAs exhibited excellent swelling resistance and phase structure stability, and no void swelling or phase transformation was observed after irradiation.

2. Since the increase in irradiation temperature significantly reduced the density of dislocations in CoCrFeMnNi HEAs, RT irradiation induced greater irradiation hardening compared to 500 °C irradiation.

3. The irradiation-induced dislocation distribution in CoCrFeMnNi HEAs can be sequentially divided into front region, damage peak region, and damage attenuation region along the depth. The dislocation density in the damage peak region at RT irradiation was the lowest among the three regions, and the damage attenuation region consisted of dense small dislocations. The dislocation density in the damage peak region was the highest at 500 °C irradiation, and no irradiation-induced dislocations were observed in the damage attenuation region.

4. As the irradiation temperature increases, there may be four stages of dislocation distribution in HEAs (such as CoCrFeNi systems): D_e= D_c, D_e< D_c, D_e> D_c and no dislocations after irradiation, where D_e and D_c are the experimentally observed depth and the theoretically calculated depth of dislocation distribution, respectively. However, D_e< D_c is not found in pure metals (such as Ni, W). The main reason for this phenomenon may be that the dislocation migration is inhibited and dissociation becomes easier in HEAs, allowing small dislocations in the damage attenuation region to be dissociated before they migrate to the deeper region, unlike in pure metals where dislocations are able to have long-range migration and dissociation has not yet started.

See the supplementary material for the details and corresponding references of points in Fig. 7.

This work was supported by the National Natural Science Foundation of China (Nos. U23B2099, 11427904, and 11975065). The authors appreciate support of the Low Energy Highly-Charged Ion Accelerator Facility (LEAF) at the Institute of Modern Physics, Chinese Academy of Sciences, for the irradiation experiments.

The authors have no conflicts to disclose.

Lisong Zhang: Conceptualization (lead); Data curation (lead); Formal analysis (equal); Investigation (lead); Methodology (lead); Software (equal); Writing – original draft (lead); Writing – review & editing (equal). Peng Zhang: Formal analysis (equal); Investigation (equal); Methodology (equal); Software (lead). Na Li: Data curation (equal); Writing – original draft (equal). Xiaonan Zhang: Formal analysis (equal); Investigation (equal). Xianxiu Mei: Conceptualization (equal); Funding acquisition (lead); Methodology (equal); Writing – review & editing (lead).

The data that support the findings of this study are available from the corresponding author upon reasonable request.