Irradiation-Induced Phase Stability in Ti- and Nb-Containing Nickel-Based High-Entropy Alloys at 500 °C

Abstract

This study investigates the irradiation response of two L1₂-strengthened HEAs, (Ni₂Co₂FeCr)₉₂Ti₄Al₄ (TiHEA) and (Ni₂Co₂FeCr)₉₂Nb₄Al₄ (NbHEA), subjected to 6.4 MeV Fe³⁺ irradiation at 500 °C up to 30 dpa. Transmission electron microscopy (TEM) and atom probe tomography (APT) consistently showed that the Ti-containing HEA maintains L1_2-ordered structure and compositional stability better than Nb-containing alloys under irradiation. This difference is attributed to the distinct solute–defect interactions. Ti imposes a weaker hindering effect on vacancy mobility, allowing vacancies to remain mobile and participate in thermal reordering processes that counteract ballistic mixing, whereas Nb acts as a strong vacancy trap, suppressing the diffusion required for structural recovery. Irradiation-induced dislocation loops in the two alloys further exhibited different characteristics. TiHEA showed larger loops at lower number density, and NbHEA exhibited a higher density of smaller loops, consistent with their respective stacking fault energies and loop mobility. Nanoindentation results indicated that TiHEA exhibited a slightly higher irradiation hardening rate (27%) than NbHEA (23%), likely associated with a stronger order-strengthening contribution, given the better preservation of precipitate order in TiHEA under irradiation. These findings show the critical role of solute addition in designing radiation-tolerant high-entropy alloys.

1. Introduction

High-entropy alloys (HEAs), also referred to as complex concentrated alloys (CCAs), have been reported to exhibit enhanced irradiation tolerance [1,2,3,4]. The presence of multiple principal elements is expected to increase lattice distortion and modify defect mobility, leading to irradiation resistance. Experimental studies have demonstrated reduced void swelling [5] and modified defect evolution [6] in various HEA systems under irradiation.

Although several single-phase high-entropy alloys (HEAs), particularly Ni-rich HEAs, have demonstrated excellent irradiation tolerance, their mechanical strength at elevated temperatures remains unsatisfactory, which hinders their applications in advanced nuclear reactors [7,8,9]. Beyond single-phase solid solution HEAs, precipitation-strengthened HEAs represent a particularly promising class of materials for nuclear applications. The introduction of ordered precipitates into an FCC matrix can significantly enhance mechanical strength at elevated temperatures while providing additional defect sinks under irradiation. Among these systems, Ni-based HEAs containing coherent L1₂-type γ′ precipitates have been proposed as structural materials for the high-temperature nuclear systems. A small lattice misfit between coherent γ′ precipitates and FCC matrix contributes to the high temperature strength and creep resistance [10].

However, the stability of γ′ precipitates under irradiation has become a major concern. Previous studies have shown that γ′ precipitates are susceptible to radiation damage [11,12]. The evolution of γ′ precipitates subjected to irradiation is governed by two fundamental processes: irradiation-induced disordering and irradiation-induced dissolution [13]. Previous studies have shown a temperature dependence in the coupling between disordering and dissolution [14,15]. Accordingly, theoretical models have emphasized the role of temperature in irradiation-induced precipitate evolution. Early theoretical models only considered recoil resolution and predicted gradual particle dissolution via size reduction but neglected compositional evolution and the coupling between disordering and dissolution [16,17]. More advanced theoretical frameworks, including kinetic models and Ginzburg–Landau-based approaches, incorporated irradiation-induced atomic mixing, irradiation-enhanced diffusion, and reordering models consistently predict a temperature dependent transition: at lower temperatures, ordered precipitates first undergo disordering before dissolution, whereas at higher temperatures, disordering and dissolution occur simultaneously. In addition, irradiation damage is not entirely irreversible, as irradiation-induced back diffusion can promote recovery of long-range order, leading to a dynamic competition between damage accumulation and self-healing processes [18,19].

In practical alloys, alloying elements such as Al, Ti, and Nb are commonly introduced to improve mechanical performance [20,21]. It is suggested that precipitate composition plays a critical role in irradiation tolerance [18,22,23]. Moreover, the high-density coherent γ/γ′ interfaces can act as efficient sinks for irradiation-induced point defects, and the sink strength of these interfaces is strongly dependent on the local chemical environment [24]. Therefore, studying the influence of alloying elements on irradiation effects is essential.

Ti and Nb are widely used as γ′-forming elements in Ni-based alloys, but they exhibit distinct characteristics. Ti forms ordered L1₂ precipitates together with Al and exhibits moderate solute–defect interactions, whereas Nb has larger atomic size mismatch with the Ni matrix and stronger solute–defect binding. However, their roles in irradiation-induced precipitates’ evolution are not yet clear. Direct experimental comparisons of Ti- and Nb-containing L1₂ precipitation-strengthened HEAs under identical irradiation conditions are needed.

In this work, we investigate the irradiation responses of (Ni₂Co₂FeCr)₉₂Ti₄Al₄ (TiHEA) and (Ni₂Co₂FeCr)₉₂Nb₄Al₄ (NbHEA) under Fe³⁺ ion irradiation at 500 °C up to 30 dpa. A non-equiatomic base alloy, Ni₂Co₂FeCr, with increased Ni and Co content was selected to ensure the formation of a high density of γ′ precipitates. Ni is the primary γ′-forming element, while Co is also largely incorporated into the γ′ phase. The increased Ni and Co content ensures that, after elemental partitioning into the γ′ precipitates, sufficient amounts of Ni and Co remain in the matrix. Transmission electron microscopy, electron diffraction, STEM-EDS, atom probe tomography, and nanoindentation were conducted to compare the irradiation-induced microstructural evolution and mechanical property changes. The direct comparison between Ti and Nb alloying provides new insights into the design of irradiation-tolerant precipitation-strengthened HEAs.

2. Materials and Methods

The alloys were designed using the Thermo-Cal software (2020a version) with the TTNI8 database. The phase field calculation confirmed the specific “compositional window,” where the matrix is FCC and the precipitates are L1₂. The alloys with nominal compositions of (Ni₂Co₂FeCr)₉₂Ti₄Al₄ and (Ni₂Co₂FeCr)₉₂Nb₄Al₄ were synthesized via arc melting and subsequently subjected to solution treatment at 1200 °C for 2 h, followed by water quenching. The compositions measured through SEM-EDS are displayed in

Table 1. The composition of TiHEA and NbHEA after water quenching measured through STEM–EDS. The values represent the average of measurements taken from three different areas, each with a size of 5 × 5 µm².

Alloy	Elements	Al	Cr	Fe	Co	Ni	Ti
TiHEA	Concentration (at.%)	3.9 ± 0.3	14.6 ± 1.9	15.4 ± 2.1	30.9 ± 3.3	31.5 ± 3.5	3.7 ± 0.7
	Elements	Al	Cr	Fe	Co	Ni	Nb
NbHEA	Concentration (at.%)	3.8 ± 0.3	14.6 ± 1.9	15.5 ± 2.1	31.3 ± 3.3	31.1 ± 3.5	3.7 ± 0.5

. After solution treatment, the alloys underwent cold rolling to an approximate thickness reduction of 70%, followed by aging at 800 °C for 24 h to produce Ti-containing and Nb-containing L1₂-ordered precipitates, respectively. Transmission electron microscopy confirmed the presence of coherent precipitates uniformly distributed within the FCC matrix in both alloys. Prior to irradiation, specimens were sectioned and mechanically ground using SiC papers with grit sizes ranging from 200 to 4000. Then, the samples were electrochemically polished in a mixed solution of C₂H₆O and HNO₃ (4:1 by volume) at 30 V and ambient temperature to eliminate surface residual stress.

Ion irradiation experiments were performed with 6.4 MeV Fe³⁺ ions at 500 °C, reaching a total fluence of 3.2 × 10²⁰ ions/m². The ion flux was maintained at approximately 0.8 × 10¹⁶ ions/m²s. The irradiation temperature was chosen to simulate operating conditions relevant to advanced nuclear energy systems. Damage profiles and implanted ion distributions were calculated using SRIM, as shown in Figure 1. The peak damage is approximately 30 dpa, located at a depth of ~3000 nm, while the region around ~1500 nm corresponds to an intermediate damage level.

Cross-sectional TEM samples and APT needles were prepared after irradiation through a standard lift-out technique using a dual-beam SEM/FIB (FEI Scios) [25]. Based on the damage distribution, the region at ~1500 nm was selected for detailed microstructural and compositional analyses, corresponding to an intermediate damage level of ~6 dpa. The error associated with depth determination is within 50 nm. The TEM specimens were additionally cleaned by low-energy Ar⁺ ion milling in a Gatan PIPS II Pro instrument to remove FIB-related damage before characterization. The thickness of the lamellae was estimated from CBED patterns obtained under the two-beam conditions [26]. Dislocation loop sizes were measured from projection TEM images acquired under two-beam conditions near g = (200). The average loop size, d, was calculated by averaging the measured values of the loop diameter. The error is the standard deviation of the values. The number density of the loops was obtained through n = N/V, where N is number of the loops and V is the estimated volume. At least three regions were used to calculate the number density, and the error is the standard deviation. Regions from the intermediate level irradiation (~6 dpa) in both alloys were selected for further comparison. This depth was selected to ensure a steady-state damaged region while minimizing surface effects.

The irradiation-induced microstructural evolution, including defect formation, γ′ precipitates, and dislocation structures, was examined using a JEOL JEM-2100F transmission electron microscope operated at 200 kV. The ordered L1₂ structure of γ′ precipitates was analyzed through selected-area electron diffraction (SAED) and central dark-field (CDF) imaging, while elemental distributions were characterized through STEM-EDS. Atom probe tomography analyses were performed using a CAMECA LEAP 5000XR instrument operated in voltage-pulsing mode with a pulse fraction of 20%. The specimen temperature and pulse repetition rate were maintained at 60 K and 200 kHz, respectively. Three-dimensional reconstructions were carried out using CAMECA IVAS 6 software. Ni, Al, and Ti/Nb were identified as the primary clustering species in the precipitates, and isoconcentration thresholds of 50 at.% for Ni + Ti or Ni + Nb were applied to distinguish precipitate regions. Detailed descriptions of APT methodology and data analysis can be found in Refs. [27,28]. The nanohardness values were measured using a Shimadzu DUH-211SKLA G200X nano-hardness tester with a Berkovich-type indenter. The indentation depth was set at 800 nm to ensure that the hardness values can represent the irradiated region, considering that the depth of the indentation affected zone is empirically 3–5 times of the indentation depth. A minimum of 8 indents, spaced approximately 50 μm apart, were made on each specimen to obtain an average hardness value and ensure statistical accuracy.

3. Results

3.1. Initial Microstructure

Before irradiation, the initial microstructures of γ′ precipitates in both Ti- and Nb-containing HEAs were examined. As shown in Figure 2, central dark-field TEM images reveal that γ′ precipitates are uniformly distributed within the grains of both alloys. The Ti-containing precipitates are spherical-shaped, and the Nb containing counterparts appear as a cuboidal shape. The precipitate sizes and number density of the two alloys are comparable, which enables a meaningful comparison of their irradiation responses. Both alloys exhibit an FCC matrix prior to irradiation, with clear elemental partitioning between the matrix and precipitates. In TiHEA, γ′ precipitates are enriched in Ni, Ti, and Al and depleted in Co, Fe, and Cr, whereas in NbHEA, the enrichment of Nb replaces that of Ti, forming Ni–Nb–Al-rich precipitates. The corresponding APT atom maps are shown in Figure 2c,d. Our previous studies indicated that the precipitates formed in both alloys correspond to Ni₃(AlTi)- or Ni₃(AlNb)-type γ′ phases, with Co, Fe, and Cr atoms partially occupying the Ni face-centered sublattice in the L1₂ structure [29].

3.2. The Evolution of Ordered L1₂ Structure Under Irradiation

The dose (depth)-dependent stability of the ordered L1₂ γ′ structure was evaluated for each alloy after irradiation. Selected-area electron diffraction (SAED) patterns obtained from TiHEA and NbHEA along the ⟨001⟩ zone axis at various depths are presented in Figure 3, with the corresponding dose levels estimated from SRIM simulations. The SAED patterns corresponding to 0 dpa were collected from depths greater than ~4000 nm where irradiation effects are negligible.

The weakening of superlattice diffraction spots was observed in both alloys. In TiHEA, the superlattice reflections are weaker than those observed in NbHEA prior to irradiation, and the reflection spots of L1₂-type chemical ordering remain clearly observable after irradiation to 30 dpa. At the intermediated irradiation (6 dpa) and peak damage region (30 dpa), the intensity of L1₂ reflection spots in both alloys becomes similar. Compared to TiHEA, the corresponding superlattice reflections in NbHEA are remarkably weakened, suggesting irradiation-induced disordering and degradation of long-range ordering. These diffraction results indicate that Nb-containing γ′ precipitates exhibit reduced irradiation stability compared with their Ti-containing counterparts.

Central dark-field (CDF) TEM was employed to characterize the disordering behavior of γ′ precipitates in TiHEA and NbHEA, as shown in Figure 4. By selecting L1₂ superlattice diffraction spots, the depth-resolved distribution of surviving ordered γ′ precipitates were revealed. In both alloys, the fraction of ordered L1₂ γ′ precipitates decrease significantly from the surface toward the peak damage region, while beyond this region the precipitates remain stable, with fractions comparable to those in the unirradiated condition.

Although selected-area electron diffraction indicates the precipitates in TiHEA are more stable under irradiation compared with NbHEA, under central dark-field TEM imaging TiHEA exhibits pronounced reduced precipitate contrast. NbHEA exhibits a relatively uniform precipitate distribution across the irradiated region, with limited apparent degradation even at greater depths. This inconsistency between selected-area electron diffraction and central dark-field TEM imaging can be explained by the different sensitivities between the two techniques. The more severely reduced superlattice reflections observed in the SAED patterns of NbHEA indicate a more pronounced reduction in the average long-range order parameter within the selected area. CDF contrast is largely controlled by local strain-induced diffraction effects rather than by the long-range order parameter alone. The stronger local contrast may arise from larger lattice mismatch between Nb-rich clusters and the Ni-based matrix, which promotes elastic strain fields and, consequently, stronger dynamical diffraction effects.

3.3. Irradiation-Induced Elemental Redistribution Through STEM-EDS

Both SAED and CDF-TEM observations indicate that the L1₂-ordered γ′ precipitates undergo disordering during irradiation, while disordering and precipitate dissolution do not occur in a strictly coupled manner. In order to characterize the compositional distribution of the precipitates, EDS elemental mappings for TiHEA and NbHEA were performed. Figure 5 and Figure 6 present STEM-EDS elemental maps acquired at ~1500 nm (6 dpa) for both alloys after irradiation at 500 °C. The EDS results show that, for both alloys, the precipitates largely maintained their compositional features, even though strong disordering was observed using TEM. This indicates that disordering occurs faster than dissolution for these alloys at 500 °C. Blurring of the surface between precipitates and the matrix can be observed for all the elements. The extent of image blurring varied between elements. In both alloys, blurring of Co is more obvious than that of other elements. Comparing the major clustering elements between the two alloys, the blurring of Nb is slightly more apparent than Ti, indicating a stronger dissolution of precipitates in NbHEA.

3.4. APT Characterization of γ′ Precipitates After Irradiation

In order to probe the detailed compositional variations after irradiation, atom probe needles were taken from ~1500 nm. Figure 7 and Figure 8 show the atom maps of individual elements in TiHEA and NbHEA. In TiHEA, Ni-Al-Ti enriched regions remain clearly identifiable after irradiation. The compositional proxigram was taken from iso surface at Ni + Ti 50 at.%, as shown in Figure 7. The composition is quite steady inside the precipitates, with a well-defined compositional gradient at the precipitate–matrix interface. In NbHEA, however, the Ni, Nb, and Al distribution becomes spatially diffuse. The interface can no longer be clearly identified in NbHEA. Figure 8b revealed a strong compositional gradient from the center of the precipitate towards the interface, demonstrating a strong dissolution effect. These results suggest that irradiation-induced damage promotes elemental mixing, leading to the incorporation of matrix elements into the precipitates. As a consequence, a more diffuse precipitate–matrix interface is formed, consistent with observations from both EDS and APT analyses. These atomic scale observations correspond to the STEM-EDS results and confirm the superior irradiation stability of Ti-containing precipitates compared to Nb-containing precipitates.

3.5. The Irradiation-Induced Dislocation Loops in TiHEA and NbHEA

Figure 9 shows bright-field TEM images of irradiation-induced dislocation loops in TiHEA and NbHEA after irradiation at 500 °C to a dose of 6 dpa. In both alloys, a high density of faulted dislocation loops is observed within the irradiated regions. The results show a clear difference in loop morphology and statistics between the two alloys. TiHEA is characterized by relatively larger dislocation loops (~38 nm in size) with a lower number density (1.7 × 10²¹/m³), whereas NbHEA has much higher density (2.5 × 10²¹/m³) of smaller loops (~30 nm in size). Stacking fault energy (SFE) estimated using Thermo-calc software (2020a version) in Figure 9d is found to be higher in NbHEA than in TiHEA both for the FCC matrix and bulk alloy composition.

3.6. The Hardness Change After Irradiation in TiHEA and NbHEA

Figure 10 presents the nanoindentation hardness and hardening rates of the Ti-containing and Nb-containing high-entropy alloys before and after irradiation. For both alloys, irradiation leads to pronounced increase in hardness, indicating significant irradiation-induced hardening. In TiHEA, the average hardness in the unirradiated condition is approximately 5.0 GPa, which increases to about 6.5 GPa, showing a hardening rate of approximately 27%. In NbHEA, the unirradiated hardness is slightly higher, at approximately 5.5 GPa, and increases to about 6.8–7.0 GPa after irradiation, corresponding to a hardening rate of approximately 23%.

4. Discussion

4.1. Role of Ti and Nb in the Irradiation Stability of γ′ Precipitates

The difference in the phase stability of TiHEA and NbHEA under 500 °C irradiation indicates the role of specific alloying elements in governing the phase stability under irradiation. As evidenced by the SAED and APT results, NbHEA suffered from more severe disordering and interfacial dissolution than TiHEA.

The stability of the ordered precipitates under irradiation is fundamentally determined by the dynamic competition between two opposing processes: ballistic mixing driven by displacement cascades, which destroys order, and thermal reordering mediated by point defect diffusion, which restores the thermodynamic equilibrium structure. The evolution rate of the order parameter (S) can be expressed as [30,31]:

$${\displaystyle \frac{dS}{dt}} = {({\displaystyle \frac{dS}{dt}})}_{irr} + {({\displaystyle \frac{dS}{dt}})}_{th},$$

(1)

where ${({\displaystyle \frac{dS}{dt}})}_{irr}$ represents the disordering rate caused by irradiation, governed by the damage rate and disordering efficiency, and ${({\displaystyle \frac{dS}{dt}})}_{th}$ stands for the thermally induced reordering. Since both alloys were subjected to identical irradiation conditions (damage rate and cascade density), the ballistic disordering term ${({\displaystyle \frac{dS}{dt}})}_{irr}$ is assumed to be comparable. Consequently, the superior stability of TiHEA should arise from a more efficient thermal reordering term ${({\displaystyle \frac{dS}{dt}})}_{th}$ compared to NbHEA. In the intermediate temperature regime, assuming that most point defects are absorbed by fixed sinks during irradiation, the diffusion coefficient can be estimated as [32]:

$$D \approx {\displaystyle \frac{2K}{k_{iv}}},$$

(2)

where K is the damage rate, k_iv is the sink strength for interstitials or vacancies. The solute atoms can interact differently with interstitials and vacancies, thereby influencing the point defect diffusion and microstructural evolution. First-principles calculations in Ni-based alloys and high-entropy alloys consistently show that both Ti and Nb exhibit attractive interactions with vacancies but with different characteristics. Nb forms a strong and highly localized vacancy trap, with a lower first nearest neighbor solute–vacancy binding energy compared to Ti [33], as listed in

Table 2. Solute–vacancy binding energies for the first n-n positions [33].

Solute in Ni Matrix	1 n-n (eV)
Ti	−0.06
Nb	−0.13

. Such strong binding immobilizes vacancies in the vicinity of Nb atoms, increasing the activation energy and suppressing long-range vacancy migration. As a result, vacancy-assisted reordering and interface recovery are kinetically hindered in NbHEA. In contrast, Ti influences vacancy behavior, primarily by increasing vacancy migration barriers rather than creating deep localized traps, and still allows vacancies to remain sufficiently mobile for thermally activated reordering [22,23].

Thus, the more severe irradiation susceptibility of γ′ precipitate in NbHEA is attributed to the strongly suppressed vacancy-mediated recovery. By contrast, TiHEA retains sharper precipitate–matrix interfaces and stronger chemical partitioning, reflecting a more favorable balance between irradiation-induced disordering and thermally activated recovery.

4.2. Irradiation Hardening Mechanisms in TiHEA and NbHEA

Both alloys exhibited irradiation-induced hardening; however, NbHEA showed a slightly lower hardening rate (23%) than TiHEA (27%) despite possessing a higher dislocation loop density.

The total change in hardness ($∆H$) can be approximated as a superposition of hardening due to irradiation defects (${∆H}_{defects}$) and the change in precipitation strengthening (${∆H}_{ppt}$):

$$∆H \approx {∆H}_{defects} + {∆H}_{ppt}.$$

(3)

According to the dispersed barrier hardening (DBH) model, the hardening contribution from both dislocation loops and precipitates can be estimated by $∆H=\Delta \sigma /3.03 = \alpha M\mu b \sqrt{Nd}/3.03$: (where M is the Taylor factor, taken as 3.06 [34]; μ is the shear modulus, taken as 80 GPa [29]; b is the magnitude of the Burgers vector, 0.25 nm; α is the obstacle strength, taken as 0.5 for precipitates and 0.4 for dislocation loops [35]; N is number density of loops; and d is diameter). According to the above equations, hardening contributions are listed in

Table 3. Estimated and actual hardening from dislocations and precipitates.

Alloy	${\mathbf{Estimated}∆\mathit{H}}_{\mathit{d}\mathit{e}\mathit{f}\mathit{e}\mathit{c}\mathit{t}\mathit{s}}$	${\mathbf{Estimated}∆\mathit{H}}_{\mathit{p}\mathit{p}\mathit{t}}$	$\mathbf{Estimated}∆\mathit{H}$	$\mathbf{Actual}∆\mathit{H}$
TiHEA	1.1	0.2	0.9	1.5
NbHEA	1.2	0.4	0.8	1.3

.

Defect Hardening (${∆H}_{defects}$): After irradiation, TiHEA is characterized by a lower number density (1.7 × 10²¹/m³) of larger dislocation loops (~38 nm in size), whereas NbHEA has much higher density (2.5 × 10²¹/m³) of smaller loops (~30 nm in size). This difference can be attributed to the influence of alloying elements on the stacking fault energy (SFE). In the lower-SFE TiHEA, faulted Frank loops are thermodynamically stabilized, and their unfaulting into perfect loops is suppressed. A limited number of sessile loops can preferentially grow through absorbing point defects. However, the higher SFE of NbHEA facilitates loop unfaulting and enhances loop mobility, sustaining a high population of smaller loops.

Precipitate Contribution (${∆H}_{ppt}$): In the unirradiated state, L1₂ precipitates strengthen the matrix primarily via order strengthening (the energy required for dislocations to create anti-phase boundaries, APB). In TiHEA, the L1₂ order is preserved better than that in NbHEA. Thus, the order-strengthening mechanism remains active, and the decrease in ${∆H}_{ppt}$ is smaller. In NbHEA, the severe disordering of the precipitates means the APB energy contribution is effectively lost. The precipitates transform from ordered barriers to disordered obstacles, which are easier for dislocations to cut through. Thus, a softening term (${∆H}_{ppt}$) relative to the initial hardness is estimated to be equal to the reduction rate of superlattice refection. In NbHEA, the hardening expected from the high density of dislocation loops is partially offset by the loss of order strengthening from L1₂-γ′ precipitates.

5. Conclusions

In this work, the microstructural evolution and mechanical response of Ti- and Nb-containing precipitation-strengthened high-entropy alloys were comparatively studied under Fe³⁺ irradiation at 500 °C. The Ti-alloyed HEA (TiHEA) exhibited superior precipitate stability under irradiation compared to the Nb-alloyed HEA (NbHEA). Nb-containing alloy (NbHEA) suffered from more severe disordering and dissolution of L1₂-γ′ phase and extensive interfacial dissolution than TiHEA. Ti solutes impose a much weaker hindering effect on vacancy mobility than Nb, allowing vacancies to remain sufficiently mobile to participate in thermally activated reordering (self-healing) that competes against ballistic mixing. Conversely, the strong vacancy-trapping effect of Nb atoms in NbHEA hinders diffusion-induced recovery, leading to the dominance of ballistic mixing. Despite the higher loop density in NbHEA, its irradiation hardening rate was slightly lower than that of TiHEA. This is attributed to the more effective loss of order strengthening associated with γ′ disordering. Overall, the results highlight the critical role of alloying element in controlling phase stability, defect evolution, and mechanical response under irradiation, providing valuable guidance for the design of radiation-tolerant precipitation-strengthened high-entropy alloys.