# Advancing our understanding of the effect of Al/Mo substitution in the TiVNb alloy on the hydrogen storage properties

Nayely Pineda-Romero, Loïc Perrière, Erik Elkaim, Claudia Zlotea

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Nayely Pineda-Romero, Loïc Perrière, Erik Elkaim, Claudia Zlotea. Advancing our understanding of the effect of Al/Mo substitution in the TiVNb alloy on the hydrogen storage properties. Journal of Alloys and Compounds, 2024, 1005, pp.176255. 10.1016/j.jallcom.2024.176255. hal-04879544

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# Advancing our understanding of the effect of Al/Mo substitution in the TiVNb alloy on the hydrogen storage properties

Nayely Pineda Romero¹*, Loïc Perrière¹, Erik Elkaim² and Claudia Zlotea¹*

¹Univ Paris Est Créteil, CNRS, ICMPE, UMR 7182, 2 rue Henri Dunant, 94320 Thiais, France

²Synchrotron SOLEIL, L'Orme des Merisiers, Saint-Aubin, BP48, 91192 Gif sur Yvette, France

* Corresponding authors: nayely.pineda-romero@cnrs.fr, claudia.zlotea@cnrs.fr

## ABSTRACT

The refractory body-centered cubic (bcc) high entropy alloys (HEAs) have shown large hydrogen uptake around 2.0 H/M, but challenges remain, the most important being the high stability of the hydrides, which requires high desorption temperatures. In this work, we reported an efficient hydride destabilization with the addition Al/Mo as non-hydride forming elements in the ternary TiVNb alloy. The alloys in the Alₓ(TiVNb)₁-xMo_y (x = 0-0.20; y = 0-0.25) system were successfully synthesized with a single-phase bcc structure and decreasing lattice parameters as Al/Mo content increases. Pressure-composition isotherms curves reveals that the addition of Al/Mo destabilizes the hydride formation, and this could be attributed to multifactorial effects. The hydrogen uptake of the alloys is strongly associated with the amount of Al/Mo in the TiVNb, decreasing less drastically with Mo than with Al. The temperature of the phase transition during desorption strongly decreases with the Al/Mo content, in agreement with thermodynamic destabilization. Therefore, the present results demonstrate that the addition of Al/Mo to refractory HEAs can enhance the design of high-capacity alloys with tunable stability.

## KEY WORDS:

High entropy alloys; Hydrogen absorption/desorption; in situ synchrotron X-Ray diffraction; thermodynamics.
# 1. INTRODUCTION

Hydrogen has become a hot topic in recent years to address the urgency to decarbonize the energy sector and it is positioned as a promising candidate for the transition to net zero emissions [1–3]. However, one of the biggest challenges is hydrogen storage due to its low volumetric energy density (9.9 MJ/m³) at standard temperature and pressure (STP) conditions.

Metal hydrides (MHs) have been shown to be safer than compressed and liquified hydrogen, with high volumetric densities (>100 Kg/m³) under moderate temperature and pressure conditions, tunable thermodynamic properties and reversible hydrogen absorption/desorption. The main drawbacks that can be faced with MHs are low gravimetric capacities, slow kinetics, high temperatures required for hydrogen desorption and high cost of raw metals [4–6]. Consequently, to overcome the challenges associated with metal hydrides, the research and development of new materials are crucial.

The so-called High Entropy Alloys (HEAs) or Multi-Principal Element Alloys (MPEAs) are the result of a new strategy of the design of alloys. In contrast to traditional alloys that focus on the vertices, edges, or faces (boundaries), these materials focus on the center of the multicomponent phase diagrams. Therefore, HEAs allow for a wide compositional space, resulting in the possibility of having a variety of properties that could be useful in different applications [7]. HEAs are composed of five or more main elements (equimolar or near equimolar ratios) with concentrations between 5-35 at.% of each constituent element and a configurational entropy (ΔS$_{conf}$) equal to or greater than 1.5R (R = 8.314 J/mol·K). Moreover, HEAs mainly crystalize into single solid solutions with simple lattices as face-centered cubic (fcc), body-centered cubic (bcc) or hexagonal close-packed (hcp) [8–14]. The wider concept of MPEAs allows to include ternary/quaternary alloys, element concentrations higher than 35 at.%, and materials with single intermetallic phase (IM) or multiphases (solid solutions and IMs) [15].
Reports on *bcc* high entropy alloys have increased since 2014 due to their interesting hydrogen storage properties. The fine tuning of the composition and the preparation process have been identified as important parameters in the design HEAs/MPEAs [16,17]. The hydrogen storage properties of MPEAs/HEAs using transition and refractory metals (Ti, V, Cr, Mn, Fe, Ni, Nb, Zr, Hf, Co, Mo, Ta) have been reported on several studies [18] with capacities around 2.0 H/M or 3.5 wt.% and phase transition from *bcc* solid solution to a *fcc* or body-centered tetragonal (*bct*) dihydride phase during reaction with hydrogen [16,19,20].

However, hydrogenation properties still require modifications to obtain suitable values for different applications. One of the main challenges is to reduce the high stability of the hydrides, which implies high temperatures to recover the stored hydrogen. Nygård et al. reported that the equimolar *bcc* TiVCrNb has a reversible capacity of 1.96 wt.% and an enthalpy of hydride formation (ΔH<sub>abs</sub>) of -51.6 kJ/mol H<sub>2</sub> [21]. The thermodynamic properties of the *bcc* TiVZrNb, TiVZrNbHf and TiV<sub>0.5</sub>ZrNbHf compositions with hydrogen uptake about 2.0 H/M are around -60 to -70 kJ/mol H<sub>2</sub>, following a decreasing linear trend with increasing the unit cell volume of the as-cast alloys [22]. This suggests that steric effect might be used to tune thermodynamics of the reaction with hydrogen, as also supported by data-driven machine learning approach [23]. Based on a large experimental metal hydride database, the authors have highlighted a strong dependence of the equilibrium pressure on a volume-based descriptor that can be calculated only from the chemical composition of alloys/intermetallics [23,24].

Other approach to adjust the thermodynamics of metal hydrides is based on the use of non-forming hydride element such as, Mo [25,26]. The substituting of Nb by Mo was studied in Ti<sub>0.20</sub>Zr<sub>0.20</sub>Hf<sub>0.20</sub>Nb<sub>0.40</sub> from 10 to 40 at.% Mo [25]. All the compositions crystallize in a *bcc* lattice with decreasing lattice parameters as Mo content increases. The capacity increases with Mo content from 1.12 to 1.54 wt.% but declines at higher Mo concentration. Upon hydrogenation, an *fcc* hydride phase is formed up to 20 at.% Mo and at higher Mo content (≥
30 at.%) a bct hydride is found. The hydrides are reversible after desorption and the temperature of desorption decreases with increasing Mo content. The instability of the hydrides has been attributed to the smaller cell volume of the hydride, and theoretical calculations indicate that this correlates with weaker H-metal bonding at higher Mo contents [27]. The addition of 10 at. % Mo in the bcc Ti<sub>0.325</sub>V<sub>0.275</sub>Zr<sub>0.125</sub>Nb<sub>0.275</sub> alloy reveals interesting hydrogenation properties: the maximum capacity at room temperature increases from 2.47 wt.% [28] to 2.8 wt.% by Mo addition and this material desorbs hydrogen at lower temperature than the alloy without Mo [29].

The use of lightweight and non-forming hydride elements, such as Al, has been also proven to improve hydrogenation properties. In our previous work, we combined the use of Mo with fixed Al content (5 at.%) in the ternary TiVNb alloy [30] in an attempt to destabilize the hydride. The Al<sub>0.05</sub>(TiVNb)<sub>0.95-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) alloys crystallize in single-phase bcc structure with decreasing lattice parameters as Mo increases. The maximum gravimetric capacities are stable around 2.7-2.8 wt.% by Mo addition. A thermodynamic destabilization for the hydride phase is confirmed by Mo addition. The ΔH<sub>abs</sub> obtained for Al<sub>0.05</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.10</sub> was -37 kJ/mol H<sub>2</sub>, which is almost two times smaller than -67 kJ/mol H<sub>2</sub> for the pristine TiVNb [31]. Therefore, these findings showed that the addition of non-forming hydride elements, such as Al and Mo, allows tailoring hydrogenation properties and obtaining promising materials for practical functioning at room temperature.

The aim of the present work is to clarify and complete our understanding on the effect of combined Mo/Al additions in the pristine TiVNb alloy up to the equimolar composition AlTiVNbMo. The Al<sub>x</sub>(TiVNb)<sub>1-x</sub>Mo<sub>y</sub> (x = 0, 0.10, 0.20; y = 0.05, 0.10, 0.175, 0.20, 0.225, 0.25) alloys have been synthesized and thoroughly characterized. For the sake of comparison, we are referring to our previous results in both Al<sub>x</sub>(TiVNb)<sub>1-x</sub> (x = 0 [31], 0.05 [30], 0.10 [31], 0.175 and 0.25 [30]) and Al<sub>0.05</sub>(TiVNb)<sub>0.95-y</sub>Mo<sub>y</sub> systems (y = 0.05, 0.10 [32]). Typical HEAs
empirical parameters (atomic size mismatch, valence electron concentration-VEC) and
physicochemical ones (density, structure type, lattice parameters), hydrogen uptake at room
temperature will be described as a function of the metal content (Al, Al/Mo, Mo). The
thermodynamic properties of the hydride formation and desorption will be addressed and the
relationship with various empirical parameters (volume per atom - Vbc/Z, Pauling
electronegativity mismatch - Δχ and enthalpy of mixing - ΔHmix) will be discussed. Furthermore,
the desorption properties will be addressed with the help of thermal desorption spectroscopy
and in situ synchrotron X-ray powder diffraction.
## 2. MATERIALS AND EXPERIMENTAL PROCEDURE

The Alₓ(TiVNb)₁-x-Moᵧ (x = 0, 0.10, 0.20; y = 0.05, 0.10, 0.175, 0.20, 0.225, 0.25) alloys have been synthetized by the arc melting method using lumps of Ti (99.99%, Alfa Aesar), V (99.9%, NEYCO), Nb (99.95%, Alfa Aesar), Al (99%, STREM Chem) and Mo (99.9%, ChemPur). A two-step strategy was employed for the alloy's preparation. The high-melting temperature elements Ti, V, Nb and Mo (Tm<sup>Mo</sup> = 2622 °C, Tm<sup>Nb</sup> = 2477 °C, Tm<sup>V</sup> = 1910 °C, Tm<sup>Ti</sup> = 1670 °C [33]) were first melted in an Ar atmosphere (~300 mbar) and by flipping the ingot between each melt to ensure good homogeneity. Afterwards, Al (Tm<sup>Al</sup> = 660 °C) was added to the corresponding pre-alloy and melted under the same conditions, but with fewer remelts.

The crystalline structure analysis of the alloys was carried out by laboratory X-ray diffraction (XRD) using a D8 Advance Bruker diffractometer (Cu Kα radiation λ=1.5406Å, Bragg-Brentano Geometry). The as-cast alloys were analyzed as cut buttons, while the hydride and desorbed samples were used as powders.

The microstructure of the as-cast alloys and hydrides were analyzed by a Zeiss Merlin high-resolution scanning electron microscope (SEM) equipped with an EDX (Energy-Dispersive X-ray Spectroscopy) detector from Oxford Instruments. Zeiss SmartSEM was used as the operating system for SEM and Aztec software for EDX analysis. For sample preparation, powder samples were embedded in epoxy resin and bulk samples were fixed with MCP 70 alloy. Finally, the samples were finely polished and coated with 1.5 nm Pd layer. An accelerated electron voltage of 10 keV was used and the element quantification was performed using Al(K), Ti(K), V(K), Nb(L) and Mo (L) signals.

Hydrogen sorption properties were determined using two types of measurements: absorption kinetics at 25 °C and pressure-composition isotherms (PCIs) at different temperatures (in the range of -10 to 200 °C), using a homemade volumetric apparatus equipped
with pressure gauge and thermostatically calibrated volumes (Sievert's method). Small sample pieces with carefully measured mass (400-500 mg) were placed into the stainless-steel sample holder with metal gaskets to prevent gas leakage and activated at 410 °C under dynamic vacuum (~1x10⁻⁵ mbar) overnight prior to starting the hydrogenation process. For the absorption kinetics, the sample holder was placed in a thermostated water bath (25 °C) and the sample is exposed to a single aliquot of hydrogen reaching a final pressure around 38±4 bar H₂. For PCIs acquisitions the sample holder was placed in a resistive furnace for high temperatures (above 25 °C), in a refrigerated circulating bath for temperatures below 25 °C, or in a mixture of liquid water with ice for 0 °C. The samples were exposed to high-purity hydrogen in small doses, step by step, until a maximum final pressure was reached for the last point. Then, PCIs of desorption were carried out by decreasing the hydrogen pressure stepwise until a minimum pressure is reached for the last point. Each PCI was followed by a complete desorption at 410 °C under dynamic secondary vacuum for 12 h. For some compositions, each PCI was performed with fresh sample portion. The storage capacity was calculated using the real gas state equation for H₂ obtained from a homemade macro-Excel file with GASPAK Version 3.32 software (Cryodata Inc., Horizon Technologies). The structural characterization of the full hydrides was analyzed by XRD.

Finally, to determine the hydrogen desorption properties, all hydrides were analyzed by thermal desorption spectroscopy (TDS). TDS was carried out using a homemade instrument with a quadrupole mass spectrometer (QMS). The procedure consisted of loading about 10 mg of the hydride into the sample holder and then introducing it into a silica tube, which is connected to the QMS operating under secondary vacuum (~1x10⁻⁶ mbar). The desorption profile was recorded by heating the sample at a constant rate of 5 °C/min up to 450 °C, the temperature is monitored by a thermocouple in close contact with the sample. Additionally, *in situ* SR-XRD during hydrogen desorption of Al₀.₁₀(TiVNb)₀.₈₅Mo₀.₀₅ was performed on the
CRISTAL beamline of the SOLEIL facility (λ= 0.67156Å) using 1.2 mm diameter quartz capillaries by applying a constant heating ramp of 2 °C/min from 25 °C up to 450 °C under a dynamic secondary vacuum. To reduce the X-ray absorbance, the powdered samples were mixed with fumed silica in a 1:2 mass ratio, respectively. The patterns were recorded every 4 °C in the scanning range from 10 to 60 2θ (°). The structural analysis of different phases was performed by Rietveld analysis [34] implemented in Fullprof Suite refinement software with the following features: the background is described as linear interpolation between a series of refinable experimental points and the peak shape is Thompson-Cox-Hastings pseudo-Voigt function. The fitted parameters were the following: scale factor, background points, zero offset, lattice parameters, full width at half maximum (FWHM) parameters (U, V, W from Caglioti formula [35]), shape parameters (X) and overall B-factor.

For comparison purposes, previous results for the compositions Al<sub>x</sub>(TiVNb)<sub>1-x</sub> (x = 0 [31], 0.05 [30], 0.10 [31], 0.175 and 0.25 [30]) and Al<sub>0.05</sub>(TiVNb)<sub>0.95-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) [32] will be included in the discussion.
## 3. RESULTS AND DISCUSSIONS

### 3.1 Structural and microstructural characterization of the pristine alloys

The compositions $Al_x(TiVNb)_{1-x}Mo_y$ with $x = 0, 0.10, 0.20$; $y = 0.05, 0.10, 0.175, 0.20$, 0.225, 0.25 were synthesized by arc melting method without any annealing treatment. The physicochemical (molar mass and density) and empirical parameters (as defined in reference [36]), such as the lattice distortion ($\delta$) and the valence electron concentration (VEC), are listed in Table 1 for the complete $Al_x(TiVNb)_{1-x}Mo_y$ ($x, y = 0-0.25$) series.

<table><caption>**Table 1.** Molar mass, bulk density, lattice distortion, VEC, structure and lattice parameters of the as-cast $Al_x(TiVNb)_{1-x}Mo_y$ ($x, y = 0-0.25$) alloys. To complete this series previous results have been cited from references [30–32].</caption><thead><tr><th>Composition</th><th>M (g/mol)</th><th>ρ (g/cm³)</th><th>δ (%)</th><th>VEC</th><th>Phase structure</th><th>a<sub>bcc</sub> (Å)</th><th>Reference</th></tr></thead><tbody><tr><td>TiVNb</td><td>63.9</td><td>6.4</td><td>4.4</td><td>4.7</td><td>Bcc</td><td>3.211</td><td></td></tr><tr><td>Al<sub>0.05</sub>(TiVNb)<sub>0.95</sub></td><td>62.1</td><td>6.3</td><td>4.3</td><td>4.6</td><td>Bcc</td><td>3.203</td><td></td></tr><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.90</sub></td><td>60.2</td><td>6.1</td><td>4.2</td><td>4.5</td><td>Bcc</td><td>3.197</td><td>[30,31]</td></tr><tr><td>Al<sub>0.175</sub>(TiVNb)<sub>0.825</sub></td><td>57.4</td><td>5.8</td><td>4.1</td><td>4.4</td><td>Bcc</td><td>3.194</td><td></td></tr><tr><td>Al<sub>0.25</sub>(TiVNb)<sub>0.75</sub></td><td>54.7</td><td>5.6</td><td>3.9</td><td>4.25</td><td>Bcc</td><td>3.189</td><td></td></tr><tr><td>Al<sub>0.20</sub>(TiVNb)<sub>0.60</sub>Mo<sub>0.20</sub></td><td>62.9</td><td>6.5</td><td>3.8</td><td>4.6</td><td>Bcc</td><td>3.181 (1)</td><td>Present work</td></tr><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.675</sub>Mo<sub>0.225</sub></td><td>67.4</td><td>6.9</td><td>3.9</td><td>4.8</td><td>Bcc</td><td>3.185 (1)</td><td>Present work</td></tr><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Mo<sub>0.10</sub></td><td>63.4</td><td>6.4</td><td>4.1</td><td>4.6</td><td>Bcc</td><td>3.197 (1)</td><td>Present work</td></tr><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub></td><td>61.8</td><td>6.2</td><td>4.2</td><td>4.6</td><td>Bcc</td><td>3.201 (1)</td><td>Present work</td></tr><tr><td>Al<sub>0.05</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.10</sub></td><td>65.2</td><td>6.6</td><td>4.2</td><td>4.7</td><td>Bcc</td><td>3.199 (1)</td><td></td></tr><tr><td>Al<sub>0.05</sub>(TiVNb)<sub>0.90</sub>Mo<sub>0.05</sub></td><td>63.7</td><td>6.4</td><td>4.3</td><td>4.6</td><td>Bcc</td><td>3.201 (1)</td><td></td></tr><tr><td>(TiVNb)<sub>0.95</sub>Mo<sub>0.05</sub></td><td>65.5</td><td>6.6</td><td>4.4</td><td>4.7</td><td>Bcc</td><td>3.210 (1)</td><td>Present work</td></tr><tr><td>(TiVNb)<sub>0.90</sub>Mo<sub>0.10</sub></td><td>67.1</td><td>6.8</td><td>4.3</td><td>4.8</td><td>Bcc</td><td>3.202 (1)</td><td>Present work</td></tr><tr><td>(TiVNb)<sub>0.825</sub>Mo<sub>0.175</sub></td><td>69.5</td><td>7.1</td><td>4.2</td><td>4.9</td><td>Bcc</td><td>3.193 (1)</td><td>Present work</td></tr><tr><td>(TiVNb)<sub>0.75</sub>Mo<sub>0.25</sub></td><td>71.9</td><td>7.4</td><td>4.1</td><td>5.0</td><td>Bcc</td><td>3.189 (1)</td><td>Present work</td></tr></tbody></table>

M = molar mass, ρ = bulk density, δ = lattice distortion, VEC = valence electron concentration, a<sub>bcc</sub> = lattice parameter of bcc phase

A simple comparison of the lattice distortion ($\delta$) values in Table 1 allows concluding that Al/Mo addition in TiVNb generally decreases this parameter, as also depicted in Figure SI1 (a) of the Supporting Information, SI. This occurs because the Al and Mo possess an intermediate atomic size among all constituent elements ($r_{Ti} = 1.46 Å > r_{Al} = 1.43 Å > r_{Nb} = 1.43 Å > r_{Mo} = 1.36 Å > r_{V} = 1.32 Å$ [36]), which reduces the atom size disparity in the alloys. Meanwhile, the
VEC parameter (Figure SI1 (b)) increases with Mo content and decreases with Al addition. However, despite these variations, all compositions are in the range of expected single-phase *bcc* solid solutions with $\delta < 6.6\%$ [37] and VEC < 6.87 [38], as experimentally corroborated by the XRD characterization (Figure 1).

<figure><img src="image_2.png" /><figcaption>**Figure 1.** XRD patterns of (a) Al<sub>0.10</sub>(TiVNb)<sub>0.90-y</sub>Mo<sub>y</sub> as-cast alloys with y = 0 [31], 0.05, 0.10 and 0.225 ; (b) (TiVNb)<sub>1-y</sub>Mo<sub>y</sub> as-cast alloys with y = 0 [31], 0.05, 0.10, 0.175 and 0.25 (λ = 1.5406 Å).</figcaption></figure>

All the HEA compositions crystallize in a single-phase *bcc* type structure (*Im*3*m* ) as shown by XRD patterns depicted in Figure 1 (a) for Al<sub>0.10</sub>(TiVNb)<sub>0.90-y</sub>Mo<sub>y</sub> (y = 0 [31], 0.05, 0.10, 0.225), in Figure 1 (b) for (TiVNb)<sub>1-y</sub>Mo<sub>y</sub> (y = 0 [31], 0.05, 0.10, 0.175, 0.25), and in Figure SI2 for the equimolar (AlTiVNb)<sub>1-y</sub>Mo<sub>y</sub> (y = 0 [30], 0.20). The series of Al<sub>0.05</sub>(TiVNb)<sub>0.95-y</sub>Mo<sub>y</sub> also form single-phase *bcc* lattices, irrespective of composition, as reported earlier [30,32]. This result is consistent with previous works on HEAs with the same elements such as Al<sub>0.05</sub>TiVNbMo [39], AlTiVNbMo [40] and Al<sub>0.10</sub>Ti<sub>0.40</sub>V<sub>0.20</sub>Nb<sub>0.20</sub>Mo<sub>0.10</sub> [41]. The a<sub>bcc</sub> listed in Table 1 were obtained from Rietveld analysis, showing a good fit for a single- phase solid solution. An almost linear decrease in the a<sub>bcc</sub> is observed with increasing Al/Mo content (Table 1 and Figure 2) in agreement with Vegard's law [42,43], which hints to the formation of substitutional solid solutions. The decrease in the average atomic radius (r<sub>Ti</sub> = 1.46 Å > r<sub>Al</sub> = 1.43 Å > r<sub>Nb</sub> = 1.43 Å > r<sub>Mo</sub> = 1.36 Å > r<sub>V</sub> = 1.32 Å [36]) explains this trend with Al/Mo addition. Similar volume contraction with Mo addition has been reported for the Ti<sub>0.30</sub>V<sub>0.25</sub>Zr<sub>0.10</sub>Nb<sub>0.25</sub>Mo<sub>0.10</sub> [29] and Ti<sub>0.20</sub>Zr<sub>0.20</sub>Hf<sub>0.20</sub>Mo<sub>0.40</sub> [27].
<figure><img src="image_3.png" /><figcaption>**Figure 2.** Lattice parameters of the *bcc* Alₓ(TiVNb)₁₋ₓMoᵧ (x, y = 0-0.25) phases as a function of Al/Mo content. The dotted line is a guide for the eyes.</figcaption></figure>

Additionally, the formation of the *bcc* structures in Alₓ(TiVNb)₁₋ₓMoᵧ (x = 0, 0.10, 0.20; y = 0.05, 0.10, 0.175, 0.20, 0.225, 0.25) systems can be explained by the stabilizing effect of Mo, Nb and V on the *bcc*-phase when alloyed with Ti (hcp at room temperature and *bcc* above 832 °C) as reported for several traditional *bcc* alloys [44,45] and HEAs [27]. Moreover, Al also acts as a *bcc* stabilizer as reported for similar compositions such as AlₓNbTiMoV [39].

The experimental densities (Table 1) were calculated using the experimental volume of the *bcc* unit cell, as determined by XRD, and the molar mass of each composition. The bulk density parameter (ρ) is slightly increasing with Mo content. Since their density is less than or equal to 7.1 g/cm³ (except for (TiVNb)₀.₇₅Mo₀.₂₅), these alloys can be considered “lightweight” refractory HEAs based on the criteria defined elsewhere [46]. Similar densities values have been reported, e.g., ρ = 6.4 g/cm³ for the equimolar AlTiVNbMo [40] and ρ = 6.1 g/cm³ for Ti₄₀Al₁₀V₂₀Nb₂₀Mo₁₀ [41].

SEM-EDX chemical mapping images and elemental analysis of Al₀.₁₀(TiVNb)₀.₉₀-Moᵧ (y = 0.05, 0.10), and (TiVNb)₁₋ᵧMoᵧ (y = 0.10, 0.25) are shown in Figure 3 (a-d) and Table 2, respectively.
<figure><img src="image_4.png" /><figcaption>**Figure 3.** SEM-EDX chemical mapping of the as-cast (a) Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub>, (b) Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Mo<sub>0.10</sub>, (c) (TiVNb)<sub>0.90</sub>Mo<sub>0.10</sub>, and (d) (TiVNb)<sub>0.75</sub>Mo<sub>0.25</sub>.</figcaption></figure>

<table><caption>**Table 2.** EDX elemental analysis of the Al<sub>0.10</sub>(TiVNb)<sub>0.90-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10), and (TiVNb)<sub>1-y</sub>Mo<sub>y</sub> (y = 0.10, 0.25) compositions.</caption><thead><tr><th>Composition</th><th>Element</th><th>Nominal (at. %)</th><th>Overall average (at. %)</th><th>Dendritic zone (at. %)</th><th>Interdendritic zone (at. %)</th><th>Δdendritic/ interdendritic</th></tr></thead><tbody><tr><td rowspan="5">Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Mo<sub>0.10</sub></td><td>Al (K)</td><td>10</td><td>9.3 (0.2)</td><td>9.2 (0.2)</td><td>9.5 (0.2)</td><td>0.3</td></tr><tr><td>Ti (K)</td><td>26.7</td><td>26.7 (0.5)</td><td>26.5 (0.4)</td><td>27.1 (0.4)</td><td>0.6</td></tr><tr><td>V (K)</td><td>26.6</td><td>26.7 (0.4)</td><td>26.7 (0.4)</td><td>27.1 (0.3)</td><td>0.4</td></tr><tr><td>Nb (L)</td><td>26.7</td><td>26.9 (0.4)</td><td>27.0 (0.3)</td><td>26.6 (0.5)</td><td>0.4</td></tr><tr><td>Mo (L)</td><td>10</td><td>10.4 (0.4)</td><td>10.6 (0.3)</td><td>9.7 (0.3)</td><td>0.9</td></tr><tr><td rowspan="5">Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub></td><td>Al (K)</td><td>10</td><td>9.7 (0.2)</td><td>9.6 (0.2)</td><td>9.7 (0.2)</td><td>0.1</td></tr><tr><td>Ti (K)</td><td>28.3</td><td>29.1 (0.4)</td><td>29.0 (0.5)</td><td>29.2 (0.3)</td><td>0.2</td></tr><tr><td>V (K)</td><td>28.3</td><td>28.4 (0.5)</td><td>28.0 (0.6)</td><td>28.6 (0.4)</td><td>0.6</td></tr><tr><td>Nb (L)</td><td>28.4</td><td>27.9 (0.5)</td><td>28.0 (0.5)</td><td>27.8 (0.4)</td><td>0.2</td></tr><tr><td>Mo (L)</td><td>5</td><td>4.9 (0.3)</td><td>5.4 (0.2)</td><td>4.7 (0.2)</td><td>0.7</td></tr><tr><td rowspan="5">(TiVNb)<sub>0.90</sub>Mo<sub>0.10</sub></td><td>Ti (K)</td><td>30</td><td>30.1 (1.0)</td><td>28.5 (0.5)</td><td>30.3 (0.8)</td><td>1.8</td></tr><tr><td>V (K)</td><td>30</td><td>30.2 (1.0)</td><td>28.7 (0.4)</td><td>30.3 (0.9)</td><td>1.6</td></tr><tr><td>Nb (L)</td><td>30</td><td>30.7 (1.2)</td><td>32.4 (0.7)</td><td>30.5 (1.1)</td><td>1.9</td></tr><tr><td>Mo (L)</td><td>10</td><td>9.1 (0.7)</td><td>10.4 (0.3)</td><td>8.9 (0.6)</td><td>1.5</td></tr><tr><td>Ti (K)</td><td>25</td><td>25.4 (0.8)</td><td>23.7 (0.7)</td><td>25.5 (0.6)</td><td>1.8</td></tr><tr><td rowspan="3">(TiVNb)<sub>0.75</sub>Mo<sub>0.25</sub></td><td>V (K)</td><td>25</td><td>25.5 (0.8)</td><td>24.3 (0.9)</td><td>25.6 (0.8)</td><td>1.3</td></tr><tr><td>Nb (L)</td><td>25</td><td>25.6 (0.5)</td><td>26 (0.7)</td><td>25.6 (0.5)</td><td>0.3</td></tr><tr><td>Mo (L)</td><td>25</td><td>23.5 (1.0)</td><td>26.0 (0.6)</td><td>23.3 (0.8)</td><td>2.7</td></tr></tbody></table>

A typical dendritic microstructure is observed for all as-cast alloys with the dendrites slightly enriched in Nb and Mo as compared to interdendritic zones. Despite this microstructure,
Table 2 shows that the average chemical composition agrees with the nominal one. These small chemical modulations are common in HEAs and have been attributed to the alloying element partitioning in the solid/liquid phase interface during the solidification process. Typical cooling rates of arc melting favor dendrite formation [39,47,48], which are enriched in elements with the highest melting points (2622 °C for Mo > 2477 °C for Nb > 1910 °C for V > 1670 °C for Ti > 660 °C for Al [33]).

The SEM-EDX analysis for Al~0.20~(TiVNb)~0.60~Mo~0.20~, Al~0.10~(TiVNb)~0.675~Mo~0.225~, and (TiVNb)~1-y~Mo_y (y = 0.05, 0.175) alloys showed similar results (Figure SI3 (a-d), respectively). Therefore, XRD and SEM-EDX confirm the formation of single-phase *bcc* solid solutions for all the Al_x(TiVNb)~1-x-y~Mo_y (x = 0, 0.10, 0.20; y = 0.05, 0.10, 0.175, 0.20, 0.225, 0.25) alloys with chemical compositions close to the nominal ones.

## 3.2 Hydrogen sorption at room temperature and hydride structure

Samples were activated prior to hydrogenation to remove surface contaminants [49]. The activated compositions were hydrogenated in a single aliquot of hydrogen gas at 25 °C under a final pressure of about 38.5±5 bar. The hydrogen absorption kinetics at room temperature are very fast (between 1-10 minutes) for all Al~0.10~(TiVNb)~0.90-y~Mo_y (y = 0, 0.05) alloys with a decreasing in the time required to reach 95% of the total capacity with Mo presence, as displayed in Figure SI xxx.

The capacities of all Al~x~(TiVNb)~1-x~Mo_y (x, y = 0-0.25) compositions obtained under these conditions are displayed in a ternary plot shown in Figure 4 and also given in Table 3. It is worth recalling that TiVNb has a maximum capacity of 2 H/M at 25 °C. Overall, the storage capacity in the Al~x~(TiVNb)~1-x-y~Mo_y system expressed as H/M decreases less drastically with the addition of Mo in the TiVNb as compared to Al. For the quaternary (TiVNb)~1-y~Mo_y (y = 0.05, 0.10, 0.175, 0.25), the capacity is almost the same for 5 and 10 at.% Mo. However, the capacity starts to slightly decrease for higher Mo content. When Al is added in the system, the capacity
strongly decreases and reaches values below 1 H/M with more than 15 at.% Al. Two main explanations can be suggested to understand this behavior: first, both Al and Mo additions decrease the *bcc* lattice parameters (Figure 2, Table 1) and consequently, the available interstitial volume for hydrogen atoms and second, both Al and Mo act as effective destabilizer of the hydride phase (see section 3.3) and subsequently, the pressure values to fully absorb hydrogen at room temperature become unachievable under our experimental conditions (too high pressure).

<figure><img src="image_5.png" /><figcaption>**Figure 4.** Hydrogen storage capacity of Alₓ(TiVNb)₁-xMoᵧ (x, y = 0-0.25) under 25 °C under a final pressure of around 38.5±5 bar as a function of Al/Mo content.</figcaption></figure>
<table><caption>**Table 3.** Hydrogenation properties (capacity, phase structure and lattice parameter of the hydride phase) for Al<sub>x</sub>(TiVNb)<sub>1-x-y</sub>Mo<sub>y</sub> (x, y = 0-0.25).</caption><thead><tr><th>Composition</th><th>Capacity (H/M)</th><th>Capacity (wt. %)</th><th>Phase structure</th><th>Phase fraction %</th><th>Space group</th><th>Lattice parameter (Å)</th><th>Reference</th></tr></thead><tbody><tr><td>TiVNb</td><td>2.0</td><td>3.2</td><td>fcc</td><td>100</td><td>Fm-3m</td><td>4.443</td><td></td></tr><tr><td>Al<sub>0.05</sub>(TiVNb)<sub>0.95</sub></td><td>1.84</td><td>2.96</td><td>fcc</td><td>100</td><td>Fm-3m</td><td>4.418</td><td></td></tr><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.90</sub></td><td>1.59</td><td>2.6</td><td>fcc</td><td>100</td><td>Fm-3m</td><td>4.375</td><td rowspan="3">[30,31]</td></tr><tr><td>Al<sub>0.175</sub>(TiVNb)<sub>0.825</sub></td><td>0.81</td><td>1.41</td><td>bcc</td><td>100</td><td>Im-3m</td><td>3.279</td></tr><tr><td>Al<sub>0.25</sub>(TiVNb)<sub>0.75</sub></td><td>0.69</td><td>1.25</td><td>bcc</td><td>100</td><td>Im-3m</td><td>3.225</td></tr><tr><td>Al<sub>0.20</sub>(TiVNb)<sub>0.60</sub>Mo<sub>0.20</sub></td><td>0.58</td><td>0.92</td><td>bcc</td><td>100</td><td>Im-3m</td><td>3.237 (2)</td><td>Present work</td></tr><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.675</sub>Mo<sub>0.225</sub></td><td>0.65</td><td>0.97</td><td>bcc</td><td>100</td><td>Im-3m</td><td>3.279 (1)</td><td>Present work</td></tr><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Mo<sub>0.10</sub></td><td>1.16</td><td>1.83</td><td>bcc</td><td>100</td><td>Im-3m</td><td>3.309 (1)</td><td>Present work</td></tr><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub></td><td>1.59</td><td>2.57</td><td>fcc +<br/>bcc</td><td>24<br/>76</td><td>Fm-3m<br/>Im-3m</td><td>4.340 (1)<br/>3.324 (1)</td><td>Present work</td></tr><tr><td>Al<sub>0.05</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.10</sub></td><td>1.79</td><td>2.75</td><td>fcc +<br/>bcc</td><td>85<br/>15</td><td>Fm-3m<br/>Im-3m</td><td>4.395<br/>3.335</td><td rowspan="2">[32]</td></tr><tr><td>Al<sub>0.05</sub>(TiVNb)<sub>0.90</sub>Mo<sub>0.05</sub></td><td>1.78</td><td>2.79</td><td>fcc</td><td>100</td><td>Im-3m</td><td>4.408</td></tr><tr><td>(TiVNb)<sub>0.95</sub>Mo<sub>0.05</sub></td><td>1.98</td><td>3.02</td><td>fcc</td><td>100</td><td>Fm-3m</td><td>4.434 (1)</td><td>Present work</td></tr><tr><td>(TiVNb)<sub>0.90</sub>Mo<sub>0.10</sub></td><td>1.98</td><td>2.96</td><td>fcc</td><td>100</td><td>Fm-3m</td><td>4.430 (1)</td><td>Present work</td></tr><tr><td>(TiVNb)<sub>0.825</sub>Mo<sub>0.175</sub></td><td>1.92</td><td>2.77</td><td>fcc +<br/>bcc</td><td>93<br/>7</td><td>Fm-3m<br/>Im-3m</td><td>4.417 (1)<br/>3.322 (1)</td><td>Present work</td></tr><tr><td>(TiVNb)<sub>0.75</sub>Mo<sub>0.25</sub></td><td>1.63</td><td>2.27</td><td>fcc +<br/>bcc</td><td>43<br/>57</td><td>Fm-3m<br/>Im-3m</td><td>4.405 (1)<br/>3.301 (1)</td><td>Present work</td></tr></tbody></table>

The gravimetric capacities of the system are also given in Table 3. It is worth noting that the values for Al<sub>0.05</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.10</sub> [32], (TiVNb)<sub>0.825</sub>Mo<sub>0.175</sub>, Al<sub>0.05</sub>(TiVNb)<sub>0.90</sub>Mo<sub>0.05</sub> [32], (TiVNb)<sub>0.90</sub>Mo<sub>0.10</sub>, and (TiVNb)<sub>0.95</sub>Mo<sub>0.05</sub> (2.75 wt.%, 2.77 wt.%, 2.79 wt.%, 2.96 wt.%, 3.02 wt.%), respectively, are among the highest capacities reported for lightweight HEAs/MPEAs materials such as, Al<sub>0.05</sub>(TiVNb)<sub>0.95</sub> (2.96 wt.%) [30], Ti<sub>0.30</sub>V<sub>0.25</sub>Zr<sub>0.10</sub>Nb<sub>0.25</sub>Al<sub>0.10</sub> (2.6 wt.%) [50], Ti<sub>0.30</sub>V<sub>0.25</sub>Zr<sub>0.10</sub>Nb<sub>0.25</sub>Mg<sub>0.10</sub> (2.7 wt.%) [51], and Mg<sub>35</sub>Al<sub>15</sub>Ti<sub>25</sub>V<sub>10</sub>Zn<sub>15</sub> (2.5 wt.%) [52].

Overall, the gravimetric capacity decreases with increasing Mo content, as expected. In other works, an increase in gravimetric capacity by Mo addition has been reported for Ti<sub>0.30</sub>V<sub>0.25</sub>Zr<sub>0.10</sub>Nb<sub>0.25</sub>Mo<sub>0.10</sub> (2.8 wt.%) [29] and Ti<sub>0.20</sub>Zr<sub>0.20</sub>Hf<sub>0.20</sub>Mo<sub>0.10</sub>Nb<sub>0.30</sub> (1.54 wt.%) [27]. However, higher Mo content typically decreases the gravimetric capacity, as observed for Ti<sub>0.20</sub>Zr<sub>0.20</sub>Hf<sub>0.20</sub>Mo<sub>0.40</sub> (0.92 wt.%) [27] and equimolar TiZrHfMoNb (1.18 wt.%) [25].
The XRD patterns of the obtained dihydrides for Al<sub>0.10</sub>(TiVNb)<sub>0.90-y</sub>Mo<sub>y</sub> (y = 0 [31], 0.05, 0.10 and 0.225) and (TiVNb)<sub>1-y</sub>Mo<sub>y</sub> (y = 0 [31], 0.05, 0.10, 0.175 and 0.25) are plotted in Figure 5 (a, b), respectively. The lattice parameters of the hydride phases are listed in Table 3, as obtained from structural refinements.

<figure><img src="image_6.png" /><figcaption>**Figure 5.** XRD patterns of hydride phases (λ = 1.5406 Å) obtained for (a) Al<sub>0.10</sub>(TiVNb)<sub>0.90-y</sub>Mo<sub>y</sub> (y = 0 [31], 0.05, 0.10 and 0.225), and (b) (TiVNb)<sub>1-y</sub>Mo<sub>y</sub> (y = 0 [31], 0.05, 0.10, 0.175 and 0.25) series.</figcaption></figure>

For the richest Al/Mo compositions, i.e., Al<sub>0.20</sub>(TiVNb)<sub>0.60</sub>Mo<sub>0.20</sub> and Al<sub>0.10</sub>(TiVNb)<sub>0.675</sub>Mo<sub>0.225</sub> (Figure SI4 a, b), hydrogen absorption induces the formation of *bcc* monohydrides (*Im*3*m*) with slightly increased lattice parameter (Table 3) compared to the as-cast materials (Table 1) due to hydrogen absorption and subsequent lattice expansion.

From Figure 5 (a) and Table 3, the hydrogenated Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub> composition is a mixture of 76% *bcc* monohydride with higher lattice parameter (3.324(1) Å) than the as-cast material (Table 1) and 24% *fcc* dihydride (*Fm*3*m*) with a lattice parameter of 4.340(1) Å. Increasing the Mo content to 10 at.%, i.e. Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Mo<sub>0.10</sub>, the hydride phase is almost a *bcc* monohydride with higher lattice parameter (3.309(1) Å) than the as-cast material (Table 1) and small traces of *fcc* dihydride phase.

Based on Figure 5 (b) and Table 3, the (TiVNb)<sub>1-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) hydrides are all *fcc* phases with close lattice parameters, i.e., 4.434(1) Å and 4.430(1) Å, respectively. The
(TiVNb)0.825Mo0.175 hydride is a mixture of a fcc dihydride with a lattice parameter of 4.417 (1) Å and a small amount of bcc phase (7% phase fraction) with higher lattice parameter (3.322(1) Å) than the as-cast material (Table 1). In addition, the (TiVNb)0.75Mo0.25 contains a mixture of 57% bcc monohydride with a lattice parameter of 3.301 (1) Å and 43% fcc dihydride with a lattice parameter of 4.405 (1) Å. Furthermore, SEM and EDX analyses for all the compositions after hydrogenation process revealed no phase segregation and chemical compositions close to the nominal ones (Figures SI5 and SI6; Tables SI1 and SI2).

### 3.3 Experimental PCIs and thermodynamic properties

The PCI experiments were performed for the most promising compositions that have shown the formation of a dihydride phase: Al0.10(TiVNb)0.80Mo0.10, Al0.10(TiVNb)0.85Mo0.05, (TiVNb)0.90Mo0.10, and (TiVNb)0.95Mo0.05. The hydrogen absorption/desorption PCI curves at different temperatures for the compositions Al0.10(TiVNb)1-yMoy (y = 0.05, 0.10), and (TiVNb)1-yMoy (y = 0.05, 0.10) are depicted in Figures 6 and 7, respectively.

All the compositions show a two-step reaction, as expected for bcc alloys [53]: the first plateau has a low equilibrium pressure ($P_{eq}$) below the detection limit (<10 mbar) and the first reaction forms a monohydride at about 0.8-0.9 H/M. While the second plateau has a higher $P_{eq}$, which increases with temperature and is associated with the formation of the dihydride phase. These results are consistent with the two-step transition, $bcc \leftrightarrow bcc$ monohydride $\leftrightarrow fcc$ dihydride, as shown for $Al_x(TiVNb)_{1-x}$ (x = 0 [31], 0.05 [30], 0.10 [31]) and $Al_{0.05}(TiVNb)_{0.95-y}Mo_y$ (y = 0.05, 0.10) [32]. The maximum capacities from the PCIs (25 °C) are 1.58 H/M (2.56 wt.%) for $Al_{0.10}(TiVNb)_{0.85}Mo_{0.05}$, 1.2 H/M (1.98 wt.%) for $Al_{0.10}(TiVNb)_{0.80}Mo_{0.10}$, 1.94 H/M (2.96 wt.%) for $(TiVNb)_{0.95}Mo_{0.05}$ and 1.99 H/M (2.97 wt.%) for $(TiVNb)_{0.90}Mo_{0.10}$. These results agree well with the single aliquot hydrogenation values obtained in section 3.2 (Table 3).
<figure><img src="image_7.png" /><figcaption>**Figure 6.** Pressure composition isotherms of dihydride formation for (a) Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub> and (b) Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Mo<sub>0.10</sub> in the temperature range of -10 to 50 °C.</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>**Figure 7.** Pressure composition isotherms of dihydride formation for (a) (TiVNb)<sub>0.95</sub>Mo<sub>0.05</sub> and (b) (TiVNb)<sub>0.90</sub>Mo<sub>0.10</sub> in the temperature range of 50-125 °C.</figcaption></figure>

From the PCI curves in figures 6 and 7 one can notice that the increase of Mo from 5 to 10 at.% into both Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub> and (TiVNb)<sub>0.95</sub>Mo<sub>0.05</sub> significantly increases the equilibrium pressure at the same temperature. Therefore, the hydride phase becomes more unstable with rising Mo amount and can sometimes spontaneously desorb hydrogen at room temperature. This can explain the presence of *bcc* monohydride phase in the XRD pattern of Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Mo<sub>0.10</sub>H<sub>1.16</sub> (Figure 5 (a)).

The ternary TiVNb alloy has sloped PCI curves, as shown in our previous work [31] and the addition of both Al/Mo element in this alloy does not significantly affect the slope of the PCI plateaus. On the contrary, the Mo addition in the ternary TiVNb alloy generates more sloped plateaus. Despite this, the application of the van't Hoff approach is feasible to determine the thermodynamic parameters for dihydride formation (ΔH<sub>abs</sub> and ΔS<sub>abs</sub>). These values
obtained for Al<sub>0.10</sub>(TiVNb)<sub>0.90-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) (Figure 8 (a)), (TiVNb)<sub>1-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) (Figure 8 (b)) are given in Table 4. The thermodynamic parameters for dihydride formation previously reported for Al<sub>x</sub>(TiVNb)<sub>1-x</sub> (x = 0 [31], 0.05 [30], 0.10 [31]) and Al<sub>0.05</sub>(TiVNb)<sub>0.95-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) [32] are also listed in Table 4 and included in the discussion.

<figure><img src="image_9.png" /><figcaption>**Figure 8.** Van't Hoff plots for dihydride formation in (a) Al<sub>0.10</sub>(TiVNb)<sub>0.90-y</sub>Mo<sub>y</sub> (y = 0 [31], 0.05, 0.10), (b) M<sub>x</sub>(TiVNb)<sub>1-x</sub> (M = Al [30,31], Mo; x = 0.05, 0.10).</figcaption></figure>

<table><caption>**Table 4.** Enthalpy/entropy of dihydride formation and empirical parameters (V<sub>bcc</sub>/Z, Δχ, ΔH<sub>mix</sub>) for the Al<sub>0.10</sub>(TiVNb)<sub>0.90-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) and (TiVNb)<sub>1-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) systems along with Al<sub>x</sub>(TiVNb)<sub>1-x</sub> (x = 0, 0.05, 0.10) and Al<sub>0.05</sub>(TiVNb)<sub>1-x-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) previously reported.</caption><thead><tr><th>Composition</th><th>ΔH<sub>abs</sub><br>(kJ/mol H<sub>2</sub>)</th><th>ΔS<sub>abs</sub><br>(J/K·mol H<sub>2</sub>)</th><th>V<sub>bcc</sub>/Z<br>(Å<sup>3</sup>)</th><th>Δχ<br>(Pauling)</th><th>ΔH<sub>mix</sub><br>(kJ/mol)</th><th>Reference</th></tr></thead><tbody><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Mo<sub>0.10</sub></td><td>-32 (±3)</td><td>-138 (±12)</td><td>16.34</td><td>0.105</td><td>-8.38</td><td>Present work</td></tr><tr><td>Al<sub>0.05</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.10</sub></td><td>-37</td><td>-125</td><td>16.37</td><td>0.106</td><td>-5.18</td><td>[32]</td></tr><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub></td><td>-42 (±4)</td><td>-156 (±13)</td><td>16.40</td><td>0.079</td><td>-8.24</td><td>Present work</td></tr><tr><td>Al<sub>0.05</sub>(TiVNb)<sub>0.90</sub>Mo<sub>0.05</sub></td><td>-45</td><td>-139</td><td>16.40</td><td>0.080</td><td>-4.85</td><td>[32]</td></tr><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.90</sub></td><td>-49</td><td>-154</td><td>16.34</td><td>0.023</td><td>-8.04</td><td>[31]</td></tr><tr><td>(TiVNb)<sub>0.95</sub>Mo<sub>0.05</sub></td><td>-53 (±1)</td><td>-142 (±1.5)</td><td>16.54</td><td>0.080</td><td>-1.03</td><td>Present work</td></tr><tr><td>(TiVNb)<sub>0.90</sub>Mo<sub>0.10</sub></td><td>-49 (±1)</td><td>-141 (±1)</td><td>16.41</td><td>0.106</td><td>-1.56</td><td>Present work</td></tr><tr><td>Al<sub>0.05</sub>(TiVNb)<sub>0.95</sub></td><td>-52</td><td>-141</td><td>16.43</td><td>0.023</td><td>-4.45</td><td>[30]</td></tr><tr><td>TiVNb</td><td>-67</td><td>-157</td><td>16.58</td><td>0.023</td><td>-0.42</td><td>[31]</td></tr></tbody></table>

V<sub>bcc</sub>/Z = initial bcc volume per atom, Δχ = Pauling electronegativity mismatch, ΔH<sub>mix</sub> = enthalpy of mixing of the solid solution.

To complete the discussion, we can compare our compositions with several TiVNb-based compositions, whose thermodynamic properties are reported in the literature (Table 5).
<table><caption>**Table 5.** Enthalpy and entropy of dihydride formation for several previous TiVNb-based compositions reported in the literature.</caption><thead><tr><th>Composition</th><th>ΔH<sub>abs</sub> (kJ/mol H<sub>2</sub>)</th><th>ΔS<sub>abs</sub> (J/K·mol H<sub>2</sub>)</th><th>Reference</th></tr></thead><tbody><tr><td>Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Cr<sub>0.10</sub></td><td>-34</td><td>-130</td><td>[54]</td></tr><tr><td>(TiVNb)<sub>0.65</sub>Cr<sub>0.35</sub></td><td>-39</td><td>-134</td><td>[55]</td></tr><tr><td>Al<sub>0.05</sub>(TiVNb)<sub>0.90</sub>Cr<sub>0.05</sub></td><td>-49</td><td>-139</td><td>[54]</td></tr><tr><td>TiVCrNb</td><td>-51.6</td><td>-129.3</td><td>[21]</td></tr><tr><td>TiV<sub>0.5</sub>ZrNbHf</td><td>-59.1</td><td>-87.4</td><td>[22]</td></tr><tr><td>TiVZrNbHf</td><td>-61.8</td><td>-88</td><td>[22]</td></tr><tr><td>(TiVNb)<sub>0.85</sub>Cr<sub>0.15</sub></td><td>-67</td><td>-172</td><td>[56]</td></tr><tr><td>TiVZrNb</td><td>-67.6</td><td>-90.3</td><td>[22]</td></tr><tr><td>Ti<sub>0.30</sub>V<sub>0.25</sub>Cr<sub>0.10</sub>Zr<sub>0.10</sub>Nb<sub>0.25</sub></td><td>-75</td><td>-161</td><td>[57]</td></tr></tbody></table>

The addition of 5-10 at. % Al in the TiVNb system destabilizes the hydride phase, as recently demonstrated [30,31], while a much larger amount of Cr has to be used (equal to or more than 25-at. % Cr [21,55]) to achieve the same destabilization level. Moreover, the addition of combined Al/Cr in TiVNb in the range 10-20 at. % also effectively destabilizes the dihydride phase [54]. The addition of Mo in Al<sub>0.05</sub>(TiVNb)<sub>0.95-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) also significantly destabilizes the dihydride formation, in agreement with machine learning predictions [32]. This work clearly demonstrates that it is possible to reduce the enthalpy of hydride formation when both Al and Mo are added into TiVNb, with Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Mo<sub>0.10</sub> (ΔH<sub>abs</sub>= -32 (±3) kJ/mol H<sub>2</sub>) having half of the enthalpy of dihydride formation of pristine alloy TiVNb (ΔH<sub>abs</sub>= -67 kJ/mol H<sub>2</sub>) [31]. Additionally, it is worth highlighting that the (TiVNb)<sub>1-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) system provides a strong destabilization while maintaining a high storage capacity. Finally, an overview of these results is given in the ternary plot displayed in Figure 9.
<figure><img src="image_10.png" /><figcaption>**Figure 9.** Enthalpy of dihydride formation, $\Delta H_{\text{abs}}$ (kJ/mol H$_2$), in Al$_x$(TiVNb)$_{1-x}$Mo$_y$ (x, y = 0-0.25) as a function of Al/Mo content.</figcaption></figure>

HEAs are complex materials whose properties are significantly influenced by a variety of parameters that are important to consider as they collectively determine the overall performance of HEAs. Understanding and optimizing the interplay between these factors is vital for the successful application of these complex materials. Various parameters might influence the enthalpy of dihydride formation such as, available interstitial's volume, electronegativity mismatch, electronic *d* band structure, thermodynamic stability of initial alloy etc. [23,58–60]. A multiple regression has been applied to 9 experimental data in this system, using 3 main empirical parameters considered as key parameters reported in the literature where interesting correlations with the enthalpy of hydride formation have been found: the initial *bcc* volume per atom - $V_{bcc}/Z$, Pauling electronegativity mismatch - $\Delta\chi$ and enthalpy of mixing of the solid solution - $\Delta H_{\text{mix}}$, the latter two parameters being defined previously [36]. The first parameter ($V_{bcc}/Z$) was chosen to describe steric effects, as suggested by data-driven machine learning approaches [23,24]; the second one ($\Delta\chi$) was selected to highlight electronic correlations, as proposed recently by density functional theory calculation for closely related *bcc* HEAs [61]; the third descriptor is based on Miedema's inverse stability rule [59,60], which predicts that more stable alloys form less stable hydrides and *vice versa*. The results of the
multiple regression are shown in Figure 10 and the correlations with different parameters can
be ranked using the statistical Prob>|t| values for a significance level of 0.05 [62]. These values
are 0.72, 0.02 and 0.002 for V<sub>bcc</sub>/Z, ΔH<sub>mix</sub> and Δχ, respectively. Thus, the strongest correlation
with ΔH<sub>abs</sub> is shown with Δχ followed by ΔH<sub>mix</sub> and V<sub>bcc</sub>/Z and a detailed discussion will be
addressed below.

The electronegativity of Mo (2.16 on Pauling scale) is higher in comparison with the
average value of the ternary TiVNb alloy (1.59 on Pauling scale) which signifies that the
interaction of H (2.2 on Pauling scale) with Mo containing alloys is expected to be less ionic
than the interaction of H with TiVNb. On the contrary, Al has an electronegativity of 1.61 on
Pauling scale and very close to the ternary TiVNb (1.59) which makes considerations based on
this parameter less straightforward.

Regarding the ΔH<sub>mix</sub>, both Al and Mo stabilize the <i>bcc</i> phase of the initial TiVNb
composition (Table 4), but Al has a stronger influence with more negative values of the
calculated enthalpy of mixing. In another work, it was also found that the calculated enthalpy
of a binary solid solution decreases with increasing Mo content in Ti-Mo alloys, favoring the
formation of <i>bcc</i> phase [63]. Consequently, the formation of less stable hydrides by addition of
Al and Mo in TiVNb might be also reasonable explained by Miedema's inverse stability rule.

The *bcc* volume per atom decreases with Al and Mo additions in TiVNb (Table 1 and Figure 2) which reduces the available interstitial volumes and consequently, the required H₂ pressure to form a hydride phase increases. Thus, a decrease of the enthalpy of hydride formation is expected, in agreement with previous observations in metal hydrides AB, AB₂ and AB₅ [64–66].
<figure><img src="image_11.png" /><figcaption>**Figure 10.** The multiple regression results vs. experimental values of enthalpy of dihydride formation, $\Delta H_{\text{abs}}$ in $Al_x(TiVNb)_{1-x}Mo_y$ alloys.</figcaption></figure>

This analysis provides a first working hypothesis to explain the variation of the enthalpy of hydride formation in HEAs: it seems that a multifactorial analysis might be used to better describe the variation of thermodynamics in multicomponent metal hydrides.

## 3.4 Hydrogen desorption properties

The desorption PCIs of $Al_{0.10}(TiVNb)_{1-y}Mo_y$ ($y = 0.05, 0.10$) and $(TiVNb)_{1-y}Mo_y$ ($y = 0.05, 0.10$) (Figure SI7) were used to determine the desorption enthalpies from the hydride phase by applying the van't Hoff approach (Figure SI8) and are listed in Table SI3. The thermodesorption properties of the hydrides were determined by TDS with a heating rate of 5 °C/min (Figure SI9-SI11). A direct comparison of the onset temperatures is elusive since some of the hydride phases are unstable at room temperature and spontaneously desorb hydrogen (see Section III in supplementary material). However, for the $(TiVNb)_{1-y}Mo_y$ series, where most of hydrides are stable at room temperature, the $T_{\text{onset}}$ was found to linearly decrease with increasing Mo content, in agreement with thermodynamic results (Figure SI11(b)).

To gain a better understanding of the effect of Mo on the desorption process and the evolution of the associated crystalline phases, *in situ* SR-XRD measurements were performed
during hydrogen desorption from Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub>H<sub>1.59</sub> at CRISTAL beamline (SOLEIL, France). The *in situ* SR-XRD patterns are plotted in Figure 11, along with the evolution of the lattice parameters (middle) and the phase fraction (right) as a function of temperature. The lattice parameters and phase fractions have been determined by Rietveld refinements and the results are listed in Table SI5.

<figure><img src="image_12.png" /><figcaption>**Figure 11.** *In situ* synchrotron X-ray diffraction (SR-XRD) thermo-diffractogram (λ = 0.67156 Å) (left) together with thermal evolution of *fcc/bcc* lattice parameters (middle) and thermal evolution of phase fractions (right) during hydrogen desorption for Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub>.</figcaption></figure>

The thermal evolution of the cell parameters of the Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub>H<sub>1.59</sub> hydride, obtained by *in situ* SR-XRD, showed a behavior similar to that of the Al<sub>0.05</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.10</sub> [32]. The initial dihydride *fcc* phase (marked by ♥) coexists with a minor *bcc* phase (marked by ♦), which is stable up to about 170 °C. From 200 °C to 250 °C, there is a pronounced decrease of *a*<sub>fcc</sub> and *a*<sub>bcc</sub> while the *bcc* weight fraction increases at the expense of the *fcc* phase (Table SI5). Above 280 °C, the *fcc* disappears and only the *bcc* phase is present with a continuous decline of the *a*<sub>bcc</sub> up to 450 °C. The *a*<sub>bcc</sub> at 450 °C (3.208(1) Å, Table SI5) is slightly larger than the as-cast (3.201(1) Å, Table 1), while the *a*<sub>bcc</sub> at 25 °C (3.198(1) Å, Table SI5) is close to that of the as-cast material. Therefore, a complete recovery of the *bcc* phase with lattice parameters close to the initial ones was achieved under *in situ* SR-XRD experiments. However, under these conditions only one phase transition *fcc* → *bcc* phase is visible, as also noticed for related hydrides in the TiVNb-based system [30,32]. It is worth to compare the temperature of
the phase transition as function of Al/Mo in TiVNb with our previous *in situ* SR-XRD experiments performed under similar conditions [30,32]. The phase transition occurs in the range 300-350 °C with a maximum at 330 °C for TiVNb whereas in the presence of 5 at.% Al this temperature domain changes to 160-330 °C with a maximum at 315 °C [30]. Mo has a more remarkable effect than Al since this phase transition occurs around 276 and 238 °C for Al<sub>0.05</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub> and Al<sub>0.05</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.10</sub>, respectively [32]. The latter value is very close to 233 °C observed presently for Al<sub>0.10</sub>(TiVNb)<sub>0.85</sub>Mo<sub>0.05</sub> (Figure 11). Therefore, these experiments clearly show that the addition of Al/Mo in exchangeable amounts has a similar effect on the desorption properties and results in a lower temperature required to desorb hydrogen as compared to pristine TiVNb. Moreover, these results are in good agreement with the experimental thermodynamic destabilization by Al/Mo addition.

4. CONCLUSIONS

In conclusion, new results are gathered here together with our previously reported findings into the TiVNb-based system to provide a complete view of the Al/Mo addition effects on the structure, microstructure, and hydrogen storage properties of the Al<sub>x</sub>(TiVNb)<sub>1-x-y</sub>Mo<sub>y</sub> (x, y = 0-0.25) alloys. All the compositions are a single-phase *bcc* structures with decreased lattice parameters as the Al/Mo content increases.

Overall, the alloys absorb hydrogen quickly at room temperature with storage capacities showing a strong dependence on Al/Mo content. The hydrogen storage capacity decreases with Al/Mo content but it reduced less drastically in the presence of Mo as compared to Al. Upon hydrogenation, the *fcc* dihydride phase was mainly obtained with low Al and Mo content (≤ 10 at.%). However, mixture of *fcc/bcc* phases or simple *bcc* monohydrides can be observed for Al or Mo > 10 at.% and Al/Mo ≥ 15 at.%. Therefore, Al<sub>0.10</sub>(TiVNb)<sub>0.90-y</sub>Mo<sub>y</sub> (y = 0.05, 0.10) compositions were not fully hydrogenated under the present conditions.
From PCIs measurements, it has been demonstrated that dihydride formation occurs *via*
a two-step reaction for Al/Mo ≤ 20 at.% and Al, Mo ≤ 10 at.%. The equilibrium plateau pressure
increases with the addition of Al and Mo, i.e., thermodynamic destabilization of the hydride
occurs in the presence of the Al and Mo with respect to the pristine TiVNb. It is worth noting
that the addition of 10 at.% Al along with 10 at.% Mo in TiVNb (Al<sub>0.10</sub>(TiVNb)<sub>0.80</sub>Mo<sub>0.10</sub>)
reduces the enthalpy of hydride formation by more than a half from -67 to -32 (±3) kJ/mol H<sub>2</sub>,
respectively. A multifactorial analysis of the enthalpy of dihydride formation was performed
on a total of 9 alloys in order to highlight the significance of several considered parameters. It
seems that the electronegativity mismatch is the most important parameter affecting ΔH<sub>abs</sub>
followed by the enthalpy of mixing in the *bcc* solid solution and the average volume per atom
in the *bcc* initial phase.

The phase transformation during desorption in this system *via in situ* experiments
revealed only one transition from fcc phase → desorbed bcc phase. After desorption, the bcc
phase is completely recovered with lattice parameters close to the as-cast materials for all the
compositions. The temperature of the phase transition strongly decreases with the Al/Mo
content, in agreement with thermodynamic destabilization. Therefore, the simultaneous
addition of non-forming hydrides elements, such as Al and Mo, to the refractory TiVNb alloy
can substantially destabilize the hydride phase and provides an original design strategy for
optimizing hydrogen storage performances of high entropy materials.

SUPPLEMENTARY MATERIALS

**Section I.** **Figure SI1.** Empirical parameters as a function of Al/Mo content, (a) lattice distortion (δ) and (b) valence electron concentration (VEC). **Figure SI2.** XRD patterns of equimolar (AlTiVNb)₁-yMo_y as cast alloys with y = 0 and 0.2 (λ = 1.5406 Å). **Figure SI3.** SEM-EDX chemical mapping of the as-cast (a) Al₀.₂₀(TiVNb)₀.₆₀Mo₀.₂₀, (b) Al₀.₁₀(TiVNb)₀.₆₇₅Mo₀.₂₂₅, (c) (TiVNb)₀.₉₅Mo₀.₀₅, and (d) (TiVNb)₀.₈₂₅Mo₀.₁₇₅.

**Section II.** **Figure SI4.** XRD pattern (λ = 1.5406Å) of hydrides phases obtained for (a) Al₀.₂₀(TiVNb)₀.₆₀Mo₀.₂₀ and (b) Al₀.₁₀(TiVNb)₀.₆₇₅Mo₀.₂₂₅, lattice parameters obtained by Topas profile fitting. **Figure SI5.** SEM-EDX chemical mapping of (a) Al₀.₂₀(TiVNb)₀.₆₀Mo₀.₂₀ hydride and (b-d)
Al₀.₁₀(TiVNb)₀.₉₀-γMoγ (y = 0.05, 0.10, 0.225) hydrides, respectively. **Figure SI6.** SEM-EDX chemical mapping of (TiVNb)₁-yMoγ hydrides with y = (a) 0.05, (b) 0.10, (c) 0.175, (d) 0.25. **Table SI1.** EDX chemical analysis for Al₀.₂₀(TiVNb)₀.₆₀Mo₀.₂₀ and Al₀.₁₀(TiVNb)₀.₉₀-γMoγ (y= 0.225, 0.10, 0.05) hydrides. **Table SI2.** EDX chemical analysis for (TiVNb)₁-yMoγ (y= 0.05, 0.10, 0.175, 0.25) hydrides.

**Section III.** **Figure SI7.** Pressure composition isotherms of hydride desorption for (a) Al₀.₁₀(TiVNb)₀.₈₅Mo₀.₀₅, (b) Al₀.₁₀(TiVNb)₀.₈₀Mo₀.₁₀, (c) (TiVNb)₀.₉₅Mo₀.₀₅, and (d) (TiVNb)₀.₉₀Mo₀.₁₀ in the temperature range of -10 to 125 °C. **Figure SI8.** Van't Hoff plots of desorption for (a) Al₀.₁₀(TiVNb)₀.₉₀-γMoγ (y = 0.05, 0.10), and (b) (TiVNb)₁-yMoγ (y = 0.05, 0.10). **Table SI3.** Thermodynamic properties of dihydride desorption for the Al₀.₁₀(TiVNb)₀.₉₀-γMoγ (y = 0.05, 0.10), and (TiVNb)₁-yMoγ (y = 0.05, 0.10). **Figure SI9.** Thermo-desorption spectroscopy (TDS) profiles for Al₀.₂₀(TiVNb)₀.₆₀Mo₀.₂₀H₀.₅₈. **Figure SI10.** Thermo-desorption spectroscopy (TDS) profiles for Al₀.₁₀(TiVNb)₀.₆₇₅Mo₀.₂₂₅H₀.₆₅ (top), Al₀.₁₀(TiVNb)₀.₈₀Mo₀.₁₀H₁.₁₆ (middle) and Al₀.₁₀(TiVNb)₀.₈₅Mo₀.₀₅H₁.₅₉ (bottom). **Figure SI11.** (a) Thermo-desorption spectroscopy (TDS) profiles for (TiVNb)₁-yMoγ (y = 0.05, 0.10, 0.175, 0.25 from bottom to top). (b) Variation of the desorption onset temperature as a function of Mo content (at. %). **Table SI4.** Lattice parameters of desorbed Alₓ(TiVNb)₁-xMoγ (x = 0, 0.10, 0.20 and y = 0.05, 0.10, 0.175, 0.20, 0.225, 0.25) alloys after TDS and as-cast for comparison. **Table SI5.** Lattice parameters of desorbed Lattice parameters of the Al₀.₁₀(TiVNb)₀.₈₅Mo₀.₀₅ alloy from in situ SR-XRD.

## AUTHOR CONTRIBUTIONS

Conceptualization: C.Z.; formal analysis: N.P.-R. and C.Z.; investigation, methodology: N.P.-R. and C.Z.; supervision, validation: C.Z.; writing—original draft: N.P.-R. and C.Z.; writing—review and editing: N.P.-R., L.P., E.E., and C.Z. All authors have read and agreed to the published version of the manuscript.

## FUNDING

This research was funded by ANR MASSHY project ANR-19-CE05-0029-01.

## ACKNOWLEDGMENTS

We acknowledge the SOLEIL Synchrotron for the time allocation on the CRISTAL beamline (experiment ID 20210367). Fabrice Couturas, Anis Bouzidi and Renato Belli Strozi from ICMPE are acknowledged for their help with the SR-XRD experiments and data treatment. We would like to thank Valérie Lalanne and Olivier Rouleau from ICMPE, for their constant technical help. We also thank Rémy Pires for his help with SEM-EDX.

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