High Strength–Ductility Synergy of As-Cast B2-Containing AlNbTaTiZr Refractory High-Entropy Alloy Under Intermediate and Dynamic Strain Rates

Abstract

Understanding the mechanical behavior of materials under various strain-rate regimes is critical for many scientific and engineering applications. Accordingly, this study investigates the strain-rate-dependent compressive mechanical behavior of B2-containing (TiZrNb)_79.5(TaAl)_20.5 refractory high-entropy alloy (RHEA) at room temperature. The RHEA is prepared by vacuum arc melting and is tested over intermediate (1.0 × 10⁻¹ s⁻¹, 1.0 s⁻¹) and dynamic (1.0 × 10³ s⁻¹, 2.0 × 10³ s⁻¹, 2.8 × 10³ s⁻¹, 3.2 × 10³ s⁻¹, and 3.5 × 10³ s⁻¹) strain rates. The alloy characterized as hybrid body-centered-cubic (BCC)/B2 nanostructure reveals an exceptional yield strength (YS) of ~1437 MPa and a fracture strain exceeding 90% at an intermediate (1.0 s⁻¹) strain rate. The YS increases to ~1797 MPa under dynamic strain-rate (3.2 × 10³ s⁻¹) loadings, which is a ~25 % improvement in strength compared with the deformation at the intermediate strain rate. Microstructural analysis of the deformed specimens reveals the severity of dislocation activity with strain and strain rate that evolves from fine dislocation bands to the formation of localized adiabatic shear bands (ASBs) at the strain rate 3.5 × 10³ s⁻¹. Consequently, the RHEA fracture features mixed ductile–brittle morphology. Overall, the RHEA exhibits excellent strength–ductility synergy over a wide strain-rate domain. The study enhances understanding of the strain-rate-dependent mechanical behavior of B2-containing RHEA, which is significant for alloy processes and impact resistance applications.

1. Introduction

Numerous scientific applications require an understanding of the mechanical behavior of materials subjected to a diverse range of strain rates. These applications involve low- to high-strain-rate mechanical deformation causing structural damage, which may include impact loading, forging, machining, accidental vehicle impact, explosives, or earthquakes [1]. Before employing, it is vitally important to examine the mechanical response and evolving microstructure of materials subjected to a diverse range of strain rates [2].

Usually, the mechanical properties of metallic materials vary with strain rate, from creep deformation, which occurs at extremely low strain rates, to high-velocity (explosive) impact loads. In general, the strain-rate-dependent deformation behavior can be classified as either a quasi-static or dynamic process. Quasi-static deformation is considered as a sequence of states of equilibrium that can be treated by the well-known equations of the mechanics of materials. In contrast, dynamic deformation is characterized by a wave of stress that propagates through the material body [3]. The transition from the quasi-static to the dynamic deformation process reduces the influence of thermal activation on dislocations, and the viscous drag effect becomes more pronounced [4,5]. Both of these changes result in enhanced Peierls–Nabarro-type stresses and the consequent strengthening of the material. Here, the nucleation of high-density dislocations multiplies with the strain rate, further impeding dislocation glide. Consequently, there is a substantial buildup of dislocations that results in localized stress concentrations and crack propagation at the early stage of deformation [6]. Hence, the poor strength–ductility synergy with increasing strain rate remains a long-standing problem in the realm of metallic materials.

Inspired by the concept of HEAs, RHEAs were originally introduced in 2010, with a focus on high-temperature applications [7]. Subsequently, many RHEAs have been found to display better mechanical strength than superalloys, including Inconel 718 and Haynes 230 [8,9,10]. However, they still have quite high densities. Later, Al-containing RHEAs comprising a binary B2 as the secondary phase were developed, which exhibited high specific YS but lacked the ductility to be applied to structural applications [11,12]. It was found that the insufficient slip planes for dislocation motion in the binary B2 structure were responsible for this brittle behavior [13,14,15]. A recent study discovered a ductile Al₁₀Nb₁₅Ta₅Ti₃₀Zr₄₀ RHEA, which was characterized as a continuous BCC/B2 nanostructure. The RHEA exhibited a stable YS and ductility of around ~1050 MPa and 50% from 20 °C to 600 °C [16]. Another advancement in the development of low-density RHEAs revealed the existence of ordered B2 and Al₃Zr₅ nanoprecipitates in disordered BCC matrices, providing high strength and ductility [6].

Furthermore, the high compressive mechanical properties at non-equilibrium strain rates are imperative for penetration-resistance applications and alloy-forming processes. However, the influence of the B2 nanoprecipitates on dynamic mechanical behavior is largely unknown. Recently, a study revealed that B2 nanoprecipitates play a positive role in the dynamic mechanical behavior of TiZrNbVAl RHEA, with considerable improvements in strength and plasticity [17].

In our recent study, we prepared a low-density (TiZrNb)_79.5(TaAl)_20.5 RHEA in as-cast condition [18]. BSE-SEM showed a dendritic morphology consisting of fine dendrites and inter-dendrites rich in Ta-Nb and Al-Zr elements, respectively. XRD analysis revealed a single-phase BCC structure, while the SAED pattern demonstrated B2 atomic ordering. The findings suggest the coexistence of a coherent BCC/B2 two-phase nanostructure. This two-phase structure likely forms soon after melting; the high-entropy effect promotes the nucleation of the disordered BCC phase. Later, during solidification, some of this phase transforms into an Al-Zr-rich B2 structure in the temperature range from 600 °C to 1200 °C, producing coherent BCC/B2 nanostructures [18]. As Al-Zr have negative enthalpy of mixing, this causes decomposition of the TiZrNbVAl RHEAs into Al-Zr and Nb-Ta-Ti phases above 600 °C to 1200 °C [19]. The interlaced B2/BCC at nanoscale contributes to the strength and ductility of the RHEA. As a result, the (TiZrNb)_79.5(TaAl)_20.5 RHEA has shown better YS (~1301 MPa) and ductility (>69%) under quasi-static conditions (1.0 $\times$ 10⁻³ s⁻¹) than the TiZrNbVAl RHEA, which has shown a YS of 744 MPa [17].

In this study, the mechanical behavior of a RHEA is evaluated at intermediate (1.0 × 10⁻¹ s⁻¹, 1.0 s⁻¹, 1.0 × 10³ s⁻¹) and dynamic (1.0 × 10³ s⁻¹, 2.0 × 10³ s⁻¹, 2.8 × 10³ s⁻¹, 3.2 × 10³ s⁻¹, and 3.5 × 10³ s⁻¹) strain rates. In summary, this study enhances understanding of the mechanical performance and deformation mechanism of B2 containing as-cast RHEA across non-equilibrium strain rates.

2. Materials Details

2.1. Fabrication Process

Sample ingot was prepared from designed composition with proportion of elements in atomic percent given in

Table 1. (TiZrNb)_79.5(TaAl)_20.5 RHEA design composition with proportion of elements in atomic %.

Elements	Al	Nb	Ta	Ti	Zr
Atomic %	23	25.5	8.5	27.5	26.5

. The alloy was synthesized by vacuum arc melting using high-purity (>99.9%) metal precursors in a high-purity argon atmosphere. The process began by creating a vacuum in the chamber with a pressure of 4.0 × 10⁻³ Pa, followed by backfilling it with argon at a purity level of 99.99%. To remove residual oxygen, a titanium getter was melted initially within the chamber. The ingot was then remelted 8–9 times and subjected to electromagnetic stirring to ensure chemical homogeneity, resulting in the as-cast RHEA specimen.

2.2. Quasi-Static and Dynamic-Compression Testing

Compression tests at intermediate strain rates were conducted using a 5985 universal testing machine (Instron Inc., Norwood, MA, USA). Here, cylindrical samples with dimensions ø (5 × 5) mm were subjected to 1.0 × 10⁻¹ s⁻¹ and 1.0 s⁻¹ of strain rate, which corresponded to ramp speeds of 0.50 mm/s and 5.0 mm/s. All tests were performed up to 60% of engineering strain. Dynamic compression tests were carried out on as-cast specimens at strain rates of 1.0 × 10³ s⁻¹, 2.0 × 10³ s⁻¹, 2.8 × 10³ s⁻¹, 3.2 × 10³ s⁻¹, and 3.5 × 10³ s⁻¹. These dynamic tests were performed using a conventional split-Hopkinson pressure bar (SHPB), the schematic architecture of which is shown in Figure 1. The bars utilized in the SHPB facility are made from 13 mm diameter maraging steel. Both the incident and transmission pressure bars are 1000 mm in length, whereas the striker bar measures 250 mm. The stress wave pulse generated upon impact was recorded by strain gauges positioned at the center of the incident and transmission bars. This pulse was then applied as in Equations (1)–(3) to determine the stress, strain, and strain rate of the specimen, which are based on the principle of one-dimensional elastic stress wave theory. The average strain rate during the plastic deformation stage was recorded as the dynamic loading strain rate for this study. Three specimens were tested for each condition, and all experiments were conducted at room temperature, ensuring consistency and reliability in the results. The mean compressive strength was calculated.

The nominal stress (σ), nominal strain (ε), and nominal strain rate ($\dot{\epsilon }$) can be calculated using the following equations [20]:

$$\mathsf{\sigma }={\displaystyle \frac{E{A}_0}{A}}{\mathsf{\epsilon }}_t$$

(1)

$$\epsilon ={\displaystyle \frac{2C_0}{L_0}}{{\displaystyle \int }}_0^t{\mathsf{\epsilon }}_{t}dt$$

(2)

$$\dot{\epsilon }={\displaystyle \frac{\mathrm{d}\mathsf{\epsilon }}{\mathrm{dt}}}=-{\displaystyle \frac{2C_0{\mathsf{\epsilon }}_{t }}{L_0}}$$

(3)

where $C_0$ stands for elastic wave velocity; $E$ is the modulus of elasticity; $L_0$ is the specimen length before an impact; $A_0$ is the cross-sectional areas of the pressure bar; $A$ represents the specimen cross-section area before an impact.

2.3. Materials Characterization After Intermediate and High Strain-Rate Compression

In our previous investigation, the (TiZrNb)_79.5(TaAl)_20.5 RHEA as-received microstructure was analyzed and discussed, and it revealed a hybrid BCC/complex B2 nanostructure [18]. In this study, the specimen deformed at strain rates of 1.0 s⁻¹, 1.0 × 10³ s⁻¹, and 3.5 × 10³ s⁻¹ was analyzed using a Quanta 400 scanning electron microscope in electron backscattered mode (SEM-BSE; FEI Inc., Hillsboro, OR, USA) equipped with an electron backscatter diffraction (EBSD) system (EDAX Velocity Super camera detector (Gatan Ametek, Berwyn, PA, USA). Additionally, surface examinations of specimens fractured after compression at a strain rate of 3.4 $\times$ 10³ s⁻¹ were conducted using the SEM in secondary electron (SE) mode. For SEM-EBSD analysis, the deformed specimen surfaces were prepared by grinding and polishing with 0.5 μm colloidal silica, followed by etching for 30 s in a solution containing 1.5 mL HNO₃, 3.0 mL HF, and 45 mL distilled water. For EBSD analysis, the specimens were further polished using ion beam milling (EM RES102, Leica Microsystems, Wetzlar, Germany). The images were scanned and captured from a specimen region of size 500 × 500 μm and a step size of 1.0 μm. The image was then processed for further analyses on EDAX APEX™ software (Gatan Ametek, Berwyn, PA, USA). High-angle grain boundaries (HAGBs) were defined as orientations higher than 15°, and low-angle grain boundaries (LAGBs) from 02° to 15°.

3. Results

3.1. Stress–Strain Curves at Intermediate Strain Rates

Figure 2a presents the true stress–strain response of the as-cast (TiZrNb)_79.5(TaAl)_20.5 RHEA, respectively, for strain rates of 1.0 × 10⁻¹ s⁻¹ and 1.0 s⁻¹. The mechanical behavior reveals three distinct findings. Firstly, the RHEA demonstrates remarkable YS and ductility, achieving a true strain of approximately 0.90. Secondly, the YS increases slightly with strain rate, from ~1351 MPa to ~1437 MPa in the as-cast condition, followed by large plastic strain without fracture. The deformation process is driven by the motion of dislocations, which rely on thermal activation to overcome energy barriers. As the strain rate increases, the time available for dislocations to bypass short-range obstacles decreases, reducing the effectiveness of thermal activation. Consequently, greater loads are required for dislocation motion, resulting in enhanced YS at higher strain rates [21]. The specimen compressed at a strain rate of 1.0 s⁻¹ exhibited a noticeable strain hardening effect, unlike the specimen deformed at 1.0 × 10⁻¹ s⁻¹ (see Figure 2a) and 1.0 × 10⁻³ s⁻¹ [18]. This signifies that the strain rate of 1.0 s⁻¹ marks the transition from the low to the intermediate strain-rate regime of the (TiZrNb)_79.5(TaAl)_20.5 RHEA. The stress remains stable during deformation (up to 90% strain), reflecting a great balance between strain hardening and softening. Such a result highlights the substantial formability of the alloy.

3.2. Stress–Strain Curve Under Dynamic Compression

Figure 2b presents the dynamic compressive behavior of as-cast (TiZrNb)_79.5(TaAl)_20.5 RHEA specimens. The results demonstrate high yield and ultimate compressive strength (UCS) of ~1750 MPa and ~1800 MPa at a 1.0 × 10³ s⁻¹ strain rate, showing a moderate increase of ~300 MPa from intermediate (1.0 s⁻¹) to dynamic (1 × 10³ s⁻¹) strain rate loading. It is also observed that the RHEA has shown a slight increase in yield strength under dynamic loading, from ~1759 MPa to ~1797 MPa at strain rates of 1.0 × 10³ s⁻¹ and 3.2 × 10³ s⁻¹, respectively. The maximum UCS of ~1849 MPa is also observed at a 3.2 × 10³ s⁻¹ strain rate. The stress–strain curves indicate an increase in strain hardening with strain rate up to 3.2 × 10³ s⁻¹. Later, the RHEA loses mechanical stability and the YS decreases to ~1201 MPa at a strain rate of 3.5 × 10³ s⁻¹. Overall, the RHEA demonstrates high strength manifested in the structural characteristics of as-cast (BCC/B2 nanohybrid) [18]. The RHEA’s low strain-rate sensitivity (SRS) further demonstrates a similar viscous-drag effect at room temperature, as observed in the ZrHfNbTaMo_x [22] and TaNbHfZrTi [20] RHEAs. In other words, the variation in the strain rate in the present dynamic regime is too small to strengthen the viscous-drag effect and further influence the flow stress. However, this effect may be enhanced with an increase in temperature that drives phonon interactions [20]. The achieved YS, UCS and peak strain in both intermediate and dynamic strain-rates, are given in

Table 2. Yield stress, UCS, and peak strain values for strain rates of 1.0 × 10⁻¹ s⁻¹, 1.0 s⁻¹, 1.0 × 10³ s⁻¹, 2.0 × 10³ s⁻¹, 2.8 × 10³ s⁻¹, 3.2 × 10³ s⁻¹, and 3.5 × 10³ s⁻¹ of as-cast (TiZrNb)_79.5(TaAl)_20.5 RHEA.

Sr#	Strain Rate (s−1)	Yield Strength (MPa)	UCS (MPa)	Peak Strain (mm/mm)
1	1.0 × 10−1	1351 ± 4.2	1552 ± 5.8	0.91 ± 0.001
2	1.0	1437 ± 5.4	1476 ± 6.7	0.91 ± 0.001
3	1.0 × 103	1759 ± 8.7	1788 ± 9.3	0.10 ± 0.001
4	2.0 × 103	1766 ± 11.1	1788 ± 10.8	0.17 ± 0.002
5	2.8 × 103	1778 ± 8.6	1796 ± 9.4	0.22 ± 0.001
6	3.2 × 103	1797 ± 10.7	1811 ± 12.1	0.26 ± 0.002
7	3.5 × 103	1201 ± 13.7	1348 ± 11.7	0.33 ± 0.004

.

4. Deformation Behavior

4.1. Effect of Intermediate Strain Rate Deformation on Microstructure

The EBSD analysis of specimens deformed at a strain rate of 1.0 s⁻¹ was conducted to examine the deformation characteristics, for which the results are shown in Figure 3. The grain orientation maps overlaid with high-angle grain boundaries (HAGBs) in Figure 3a reveal severely deformed grains elongated in the direction of plastic flow. Additionally, the overall grain size is reduced to ~12 μm compared with the grain size of ~19 μm observed in specimens deformed at a strain rate of 1.0 × 10⁻³ s⁻¹ [18]. Furthermore, the specimen deformed at the intermediate strain rate (1.0 s⁻¹) exhibits a non-uniform grain morphology with both large and small grains, indicating heterogeneous deformation. The inverse pole figure (IPF) in Figure 3d shows that the deformed grains are predominantly oriented along the (101) plane, which is a favorable slip plane in BCC metals. The kernel average misorientation (KAM) map, overlaid with HAGBs in Figure 3b, indicates a high degree of misorientation, highlighting significant intragranular strain. Figure 3c shows the frequency (%) of misorientation (°) for angles ranging from 2° to 62°. The high frequency of low-angle grain boundaries (LAGBs) represents dislocation-associated deformation, and the process is similar to the deformation reported in quasi-static condition.

4.2. Effect of High-Strain-Rate Deformation on Microstructure

Figure 4a–c illustrate the grain orientation, KAM, and LAGBs overlaid with HAGBs for the specimen dynamically deformed at a strain rate of 1.0 × 10³ s⁻¹. The RHEA subjected to 8% strain exhibits randomly oriented grains, as seen in Figure 4a. The KAM map shows a small misorientation caused by macrostrain, which is distributed randomly within the grains. Additionally, the misorientation and corresponding LAGBs indicate that strain occurs along the grain boundaries. LAGBs typically result from dislocation accumulation during plastic deformation, suggesting that dislocations form and accumulate at grain boundaries under a strain rate of 1.0 × 10³ s⁻¹. Figure 4d presents the grain orientation after deformation at a strain rate of 3.5 × 10³ s⁻¹. The black regions indicate unresolved planes, which are typically caused by severe plastic deformation [23]. Overall, the grains are highly deformed, as evidenced by the significant spread of misorientation in the KAM map and the corresponding LAGBs. The unresolved continuous line is likely indicative of an ASBs. ASBs are more pronounced in BCC RHEAs due to their composition of elements with low thermal conductivity and relatively simple deformation mechanisms [24]. It is significant to note the difference in magnitude of the deformed structures at strain rates of 1.0 s⁻¹ and 1.0 × 10³ s⁻¹. The specimen tested at the intermediate strain rate (1.0 s⁻¹) is deformed to 90% of true strain under equilibrium conditions. This means the structure is extremely deformed throughout the specimen, causing a very large number of misoriented grains with increased intensity of misorientation angle and dislocation density, whereas at a strain rate of 1.0 × 10³ s⁻¹, the specimen is dynamically deformed to < 10% of true strain only. This small strain represents a low-volume fraction of deformed grain structure with a lesser degree of deformation, reflected in the low misorientation angle and dislocation density.

Severe localized shear deformation at the bottom of the specimen that is characterized by bifurcated ASBs is clearly observed in the SEM-BSE images in Figure 5. The ASB is estimated to be approximately 10 μm thick and over 100 microns in length. The arrows (labeled 1, 2) in Figure 5a highlight branches of the bifurcated ASB. Similar shear effects were reported during the dynamic deformation of TiZrNbTa RHEAs at a strain rate of 6.5 × 10³ s⁻¹ [25]. It is important to note that the formation of ASBs that initiate local instabilities is a common failure mechanism in most RHEAs [26].

Figure 6 shows the fracture surface of the specimen after dynamic compression at 4.0 × 10³ s⁻¹. Two distinct regions of the fracture surface, labeled 1 and 2, are highlighted in the figure. Region 1 exhibits two separated shear traces on an otherwise inconspicuous surface. In contrast, region 2 displays river patterns, indicative of brittle fracture. Additionally, near-equiaxial and parabolic dimples were observed, as shown in Figure 6c. The presence of dimples generally indicates good plasticity in the material, and their shape is influenced by the state of stress. The formation of parabolic shallow dimples is attributed to a combination of tension, shear, and compression stresses [27]. In summary, the RHEA demonstrates a mixed ductile–brittle fracture mechanism under dynamic compressive loading.

Figure 7a,b compare the yield and UCS properties of the (TiZrNb)_79.5(TaAl)_20.5 RHEA with other RHEAs reported in the literature under dynamic loading experiments at different strain rates [1,6,22,28,29,30,31,32,33,34,35,36,37,38,39,40]. Different colors and shapes in the graphs represent unique structural characteristics of HEAs. For instance, the brown right triangles and blue diamond shapes correspond to the mechanical properties of HEAs with single-phase face-cantered-cubic (FCC) and BCC structures, respectively. The results demonstrate that in general, the (TiZrNb)_79.5(TaAl)_20.5 RHEA exhibits the best combination of yield and UCS across a range of dynamic strain rates. CoCrFeMnNi [28], Hf_0.75NbTaTiZr, and HfMo_0.5NbTaTiZr [22] RHEAs have shown competitive performance in mechanical properties but are disadvantaged in density (8.0 g.cm⁻³) compared with (TiZrNb)_79.5(TaAl)_20.5 RHEA, with density of ~6.9 g.cm⁻³.

Overall, the RHEA exhibits low SRS across intermediate and dynamic strain rates. The SRS is closely related to the disordered nature of the dislocation movement required to overcome short-range obstacles such as solute atoms. The presence of the B2 phase contributes to a reduction in SRS, as reported in previous studies [41]. The B2 acts as a long-range obstacle, with thermal activation amplitudes sufficiently large to diminish the impact of strain rate on dislocation motion, thereby reducing SRS [17]. The lower SRS results in a smaller increase in yield strength when transitioning from the quasi-static (1.0 × 10⁻³ s⁻¹) to the intermediate (1.0 s⁻¹) and dynamic (1.0 × 10³ s⁻¹) strain rates [42].

In general, the presently studied RHEA has demonstrated better strength properties compared with single-phase BCC RHEAs. The better performance may be due to the presence of a complex B2 structure, which effectively suppresses dislocation glide resistance during plastic deformation [17]. The RHEA grains were severely deformed after compression at 1.0 s⁻¹ strain rate up to 90% strain. No cracks were formed, indicating excellent forming ability. In dynamic deformation, the strain increases with the strain rate. This means the dynamic stress continues to deform the grain structure with increasing strain rate, from the widely dispersed dislocation cells (1.0 × 10³ s⁻¹) [19] along the grain boundaries to localized shear bands (3.5 × 10³ s⁻¹) that are formed adiabatically and are given in the schematic diagram in Figure 8. RHEAs have relatively lower thermal conductivities [25]. Therefore, the strain hardening raises the stress and is lowered by the associated rise in temperature. At a 1.0 × 10³ s⁻¹ strain rate, the strain hardening overcame the thermal softening, and later, the RHEA became thermally instable after deformation at a higher strain rate, i.e., 3.5 × 10³ s⁻¹, which caused a deterioration in strength, realized in the formation of ASB with the presence of fine grains near the ASB, as confirmed in Figure 5.

The (TiZrNb)_79.5(TaAl)_20.5 RHEA demonstrates an advantage of low density and facile room temperature processing over Senkov (NbMoTaW [7] and TaNbHfZrTi [43]) RHEAs. It opens a path to understanding in depth the role of complex B2 in the dynamic deformation behavior of refractory complex B2/BCC nanostructures for optimized performance. This study also encourages further research on the high-temperature dynamic mechanical behavior of B2-containing RHEAs for advanced engineering applications.

5. Conclusions

• The room-temperature compressive mechanical behavior and deformation mechanism of as-cast (TiZrNb)_79.5(TaAl)_20.5 RHEA is studied over intermediate and dynamic strain rates. The RHEA exhibits remarkable mechanical performance under intermediate strain rates (1.0 × 10⁻¹, 1.0 s⁻¹) with exceptional YS (~1437 MPa) and fracture strain (>90%). The RHEA also maintained excellent balance between strength and ductility during the entire plastic deformation.
• Under dynamic strain rates (1.0 × 10³, 1.5 × 10³, 2.0 × 10³ s⁻¹, 2.8 × 10³, 3.2 × 10³, and 3.5 × 10³), the RHEA demonstrates an increase in YS of around ~350 MPa compared with compression at intermediate strain rates. Further, the RHEA demonstrates a mixed ductile–brittle fracture mechanism, evident from the presence of dimples, which indicate good plasticity, while the observed river patterns suggest brittle fracture.
• The analyses of deformed specimens over intermediate strain rates show the presence of multiple delocalized dislocation bands causing large plasticity, whereas under dynamic loading, the deformed microstructure changes from the formation of dislocation cells that coalesce into microbands and, finally, to adiabatic shear bands after experiencing thermal instability.
• The excellent strength–ductility synergy and low SRS highlight the alloy’s robust performance under diverse loading conditions. The RHEA shows great potential for room-temperature processing and anti-penetration applications.