Uncovering the Effect of Al Addition on the Hydrogen Storage Properties of the Ternary TiVNb Alloy

Abstract

The effect of Al addition on the structure, microstructure and hydrogen storage properties of the ternary TiVNb alloy was investigated from small amounts to equimolar composition. Al_x(TiVNb)_1−x (x = 0.05, 0.175 and 0.25) alloys are bcc single-phase materials with decreasing lattice parameters with increasing Al content. Al addition progressively decreases the hydrogen storage capacity but also destabilizes fcc dihydride formation for alloys with x ≤ 0.10. Among the different compositions, the most promising alloy was found to be that with x = 0.05 Al content that exhibited high initial storage capacity (2.96 wt.%), a less stable hydride (ΔH = −52 kJ/mol H₂ and ΔS = −141 J/K∙mol H₂), better desorption properties (desorption onset temperature around 100 °C) and enhanced reversible capacity during cycling (2.83 wt.%) compared to the ternary TiVNb. In situ and ex situ synchrotron X-ray powder diffraction, together with thermal desorption experiments, showed improved desorption properties with Al addition, together with a two-step reaction with hydrogen. These findings highlight the use of small quantities of lightweight Al in refractory multi-principal element alloys as a promising approach for enhancing the solid-state hydrogen storage performance of bcc-type alloys.

1. Introduction

In the context of climate change, hydrogen represents an opportunity to mitigate global warming as it is a zero-carbon fuel with three times more energy per mass than fossil fuels (142 MJ/kg). However, its energy per volume is very low and its storage necessitates close attention to ensure that it is a safe, affordable, and efficient technology. Available storage technologies include compressed gas and liquid hydrogen. In addition, metal hydrides used for solid-state hydrogen storage appear to represent a positive storage option that can meet required criteria. However, they face several challenges including sluggish absorption/desorption kinetics, unfavorable thermodynamics and fading reversible capacity [1,2]. Nevertheless, metal hydrides represent one of the safest options for stationary applications and show high volumetric hydrogen storage capacity in the range 60–100 Kg/m³, which is higher than that of compressed gas or liquid hydrogen. These properties make them good candidates for the investment of research effort to overcome their limitations and achieve efficient hydrogen storage systems which can operate in near ambient conditions [3].

A new concept in metallurgy was established in the work of Cantor et al. [4] and Yeh et al. [5] by the introduction of multi-principal element alloys (MPEAs). MPEAs have significant atom fractions of several elements in which the distinction between minor and major elements vanishes. Four or more elements at close to equimolar concentration are commonly used in MPEAs. Alloys with five or more elements are called high entropy alloys (HEAs). The main motivation for increasing the number of alloying elements is to maximize the configurational entropy to improve the stability of disordered solid solution phases, thus suppressing the formation of intermetallics [1,3]. MPEAs often form random solid solutions, adopting simple crystalline lattices, such as face-centered cubic (fcc), body-centered cubic (bcc) or hexagonal compact (hcp) structures. The large lattice distortion (δ) which may be created by sharing the same crystalline sites among many atomic constituents is a key factor that can create large interstitial sites to accommodate hydrogen. Among recent attempts to use MPEAs for hydrogen storage, bcc alloys have been proposed as good candidates to form metal hydrides [3]. The equimolar TiVZrNbHf alloy reported by Sahlberg et al. crystallizes in a bcc single phase and absorbs hydrogen within a single-step reaction forming a body-centered tetragonal (bct) hydride with a high capacity of 2.5 hydrogen atoms per metal atom (H/M) [6]. Later, Zlotea et al. described equimolar TiZrNbHfTa HEA, which also adopted a bcc lattice with a maximum capacity of 2.0 H/M [7]. The hydrogen absorption pathway is based on a two-step reaction: first, the alloy absorbs hydrogen at low pressure, transforming into a monohydride phase (capacity∼1.0 H/M), followed by a second transition at high pressure from the monohydride to a dihydride phase (capacity~2.0 H/M). Most reported HEAs follow this two-step reaction with variable limits for monohydride and dihydride capacities.

Non-equimolar compositions have also recently attracted attention, such as the bcc Ti_0.325V_0.275Zr_0.125Nb_0.275 alloy that showed a maximum capacity of 1.7 H/M [8]. Starting from this composition, the same authors demonstrated an increase in the total capacity up to 2.0 H/M by adding 10 atomic percent (at.%) of Ta, Cr or Mo (Ti_0.30V_0.25Zr_0.10Nb_0.25X_0.10 X = Ta [9], Cr [10] and Mo [11]), without affecting the bcc lattice of the initial solid solution phase. Despite these improvements, these alloys possess modest gravimetric capacities because of the use of quite heavy elements. Therefore, an important feature to consider in the field of HEAs for hydrogen storage is the increase in gravimetric capacity that can result from the introduction of lightweight elements [1]. However, only a limited number of reports have addressed the use of lightweight bcc HEAs for hydrogen storage purposes. The addition of low-density metals (Mg, Al, or Sc) to increase the hydrogen gravimetric capacity in HEAs involves complex interactions with the alloying elements, which has sometimes resulted in missing the target of enhancing capacity [1,3]. Strozi et al. investigated the bcc MgVAlCrNi alloy containing two lightweight metals (Mg and Al) but it exhibited a low hydrogen capacity (0.30 wt.%) that was attributed to the positive enthalpy of the hydrogen solution in most of the alloying elements, including Mg [12]. The MgAlTiFeNi alloy reported by Cardoso et al. crystallized in a bcc lattice and also exhibited limited H₂ capacity (0.94 wt.%) [13]. Ferraz et al. investigated the Mg₃₅Al₁₅Ti₂₅V₁₀Zn₁₅ alloy with the objective of increasing the gravimetric capacity by use of both Mg and Al. However, the material synthesized by reactive ball milling under an H₂ atmosphere was multiphase, composed of an fcc hydride together with MgH₂, with a total capacity of 2.75 weight percent (wt.%) [14]. Finally, the material behaved as a simple mixture of several hydrides. Montero et al. produced interesting results by the addition of Mg or Al to the Ti_0.325V_0.275Zr_0.125Nb_0.275 alloy. The Mg_0.10Ti_0.30V_0.25Zr_0.10Nb_0.25 and Al_0.10Ti_0.30V_0.25Zr_0.10Nb_0.25 alloys occurred as bcc single phases with a maximum capacity of 1.7 H/M (2.7 wt.%) [15] and 1.6 H/M (2.6 wt.%), respectively [16]. Furthermore, in our previous investigations, we demonstrated that the addition of 10 at.% of Al into the ternary bcc TiVNb alloy decreased the storage capacity from 2.0 H/M (3.2 wt.% for TiVNb) to 1.59 H/M (2.6 wt.% for Al_0.10(TiVNb)_0.90) [17]. Despite the observed decrease in capacity, other aspects of performance were positively affected by the presence of Al. Al addition (i) destabilized dihydride formation (from the monohydride to the dihydride phase), (ii) flattened the plateau pressure, which is a required condition for the practical use of metal hydrides, (iii) decreased the desorption onset temperature below 100 °C, and (iv) increased the stable capacity over absorption/desorption cycling.

Consequently, further research is needed to increase knowledge of the hydrogen storage properties of lightweight HEAs with a view to achieving high hydrogen capacity, fast hydrogen absorption/desorption kinetics close to standard temperature and pressure (STP) conditions, favorable thermodynamic properties, and high reversible hydrogen capacity during cycling [3]. Based on the above, we report on the hydrogen sorption properties of the lightweight series of alloys, Al_x(TiVNb)_1−x with x = 0.05, 0.175 and 0.25, synthetized by arc melting. Our objective was to vary the Al content to tune the microstructure, physicochemical, and hydrogen sorption properties. For this purpose, laboratory and large-scale experimental facilities were combined to clarify the effect of the amount of Al introduced into TiVNb. Well-known laboratory techniques, as well as ex situ and in situ synchrotron X-ray powder diffraction, were carried out and the results compared to our previous results for ternary TiVNb and quaternary Al_0.10(TiVNb)_0.90 alloys [17]. The thermodynamic experimental results obtained were compared with data-driven predictions facilitated by a previously published machine/statistical learning (ML) model [17]. We sought to perform a comprehensive analysis by varying the Al content introduced into TiVNb from low amounts to equimolar composition.

2. Materials and Methods

The Al_x(TiVNb)_1−x (x = 0.05, 0.175, 0.25) compositions were synthetized using an electric arc melting method from bulk pieces of Ti (99.99% pure, Alfa Aesar), V (99.9% pure, NEYCO Vacuum&Materials), Nb (99.95% pure, Alfa Aesar) and Al (99% pure, STREM Chem). Due to the high melting points of Ti, V and Nb (2468 °C for Nb > 1902 °C for V > 1660 °C for Ti) these three metals were pre-alloyed in an electric arc furnace under Ar atmosphere (around 300 mbar) and remelted 16 times by flipping the ingot between each melting to achieve good homogeneity. Due to its low melting point (660 °C) and to avoid loss of material, Al was added to the pre-alloyed TiVNb and melted under the same conditions as above with eight repetitions. The mass loss during the synthesis was less than 1% over the total mass of alloy (3 g).

The hydrogenation process was carried out using a Sievert apparatus, which consisted of a home-made manual manometric device with thermostatically calibrated volumes. A stainless-steel sample holder was filled with about 400–500 mg of sample cut into small pieces, sealed using metal gaskets to prevent gas leakage and connected to the Sievert apparatus. Before starting the hydrogenation process, an activation procedure was carried out by exposing the sample at 410 °C under dynamic vacuum (~1 × 10⁻⁵ mbar) overnight. Two types of measurements were carried out: absorption kinetics at 25 °C and pressure-composition isotherms (PCIs) for a range of temperatures. For the absorption kinetics, the sample holder was placed in a water bath under isothermal conditions (25 °C) and a final equilibrium pressure (P_eq) of around 38 ± 5 bar of H₂. For acquisition of the PCIs, small doses of hydrogen were used up to around 47 bar. A resistive furnace was used for the acquisition of PCIs at high temperatures (above 25 °C). The mass of the samples was measured carefully before hydrogenation to calculate the storage capacity using the real gas state equation for H₂ in GASPAK Version 3.32 (Horizon Technologies). For the absorption/desorption cycling experiments, absorption was performed at 25 °C and a final P_eq around 32 ± 4 bar H₂, while desorption was performed by heating at 410 °C under secondary vacuum (~1 × 10⁻⁵ mbar) for 12 h.

The crystalline structure of the samples was investigated by laboratory X-ray powder diffraction (XRD) using a D8 Advance Bruker diffractometer (Cu K_α radiation λ = 1.5406 Å, Bragg-Brentano Geometry). In addition, both ex situ and in situ structural characterizations were carried out by synchrotron X-ray powder diffraction (SR-XRD) at the Cristal beamline of the SOLEIL facility. The intermediate hydrides of TiVNb and Al_0.05(TiVNb)_0.95 compositions were placed into 0.2 mm diameter borosilicate capillaries and characterized by ex situ SR-XRD in the scanning range 10° to 60° (2θ) with the wavelength λ = 0.72907 Å. In situ SR-XRD measurements during hydrogen desorption were carried out for Al_0.05(TiVNb)_0.95 and TiVNb by applying a constant heating ramp with 2 °C/min from 25 °C up to 450 °C under dynamic secondary vacuum. Quartz capillaries with 1.2 and 1.0 mm diameter were used for Al_0.05(TiVNb)_0.95 and TiVNb, respectively. To reduce the X-ray absorbance, the powdered samples were mixed with fumed silica at a mass ratio of 1:2, respectively. The in situ SR-XRD patterns were recorded every 4 °C in the scanning range 10° to 60° (2θ) with the wavelength λ = 0.67156 Å. The phase distribution and related lattice parameters were determined from ex situ and in situ SR-XRD patterns using the Rietveld refinement method implemented in Fullprof software (Thompson-Cox-Hastings pseudo-Voigt function for the peak shape).

The microstructure was analyzed by scanning electron microscope (SEM) using a Zeiss Merlin instrument equipped with an EDX (energy-dispersive X-ray spectroscopy) detector from Oxford Instruments. The powder samples were embedded in an epoxy resin and bulk samples were fixed with MCP 70 alloy, then finely polished and finally coated with a 1.5–2.5 nm Pd layer. For elemental analysis, an accelerated electron voltage of 10 keV was used and the element quantification was performed using Al(K), Ti(K), V(K) and Nb(L) signals. The analyses were carried out for the as-cast and after cycling samples.

The hydrogen desorption properties were analyzed by thermal desorption spectroscopy (TDS) using a home-made instrument with a quadrupole mass spectrometer QMS, as described previously [18]. The procedure consisted of loading about 10 mg of the hydride sample into the sample holder and then connecting it to the quadrupole mass spectrometer working under secondary vacuum (~1 × 10⁻⁶ mbar). The desorption profile (partial pressures of evolved gases) was recorded by heating the sample at a constant rate of 5 °C/min up to 450 °C.

3. Results and Discussions

3.1. Synthesis and Microstructure of the Pristine Alloys

The physicochemical and empirical parameters, such as the lattice distortion (δ, as defined in reference [19]) and the valence electron concentration (VEC), for the Al_x(TiVNb)_1−x alloys with x = 0, 0.05, 0.10, 0.175 and 0.25, are listed in

Table 1. Physicochemical (molar mass and bulk density) and empirical parameters (lattice distortion and VEC) of the Al_x(TiVNb)_1−x (x = 0 [17], 0.05, 0.10 [17], 0.175, 0.25) alloys.

Composition	Al Content (at. fraction)	Molar Mass (g/mol)	Density ρ (g/cm3)	Lattice Distortion δ (%)	VEC
TiVNb	0	63.9	6.4	4.4	4.7
Al0.05(TiVNb)0.95	0.05	62.1	6.3	4.3	4.6
Al0.10(TiVNb)0.90	0.10	60.2	6.1	4.2	4.5
Al0.175(TiVNb)0.825	0.175	57.4	5.8	4.1	4.4
Al0.25(TiVNb)0.75	0.25	54.7	5.6	3.9	4.25

. For comparison, our previous data for the ternary TiVNb (x = 0) and the quaternary Al_0.10(TiVNb)_0.90 (x = 0.10) alloys from reference [17] are also included in the tables, figures and discussion throughout the whole document.

The lattice distortion (δ) decreased with increase in Al content (Table 1), which was assumed to be related to a decrease in the atom size disparity. The atomic radius of Al is very close to most of the elements (r_Ti = 1.46 Å > r_Nb = 1.43 Å > r_Al = 1.43 Å > r_V = 1.39 Å [19]), except for V which has the smallest radius. Therefore, increasing Al and decreasing V concentrations led to a decrease in the atom size disparity and, consequently, the lattice distortion reduced. Similar behavior was reported by Ma et al. for the (ZrTiVFe)_xAl_y series, where Al atomic size was very similar to that of the other alloying elements [20].

In the same way, the parameter VEC and the bulk density (ρ) were reduced with increasing Al content (Table 1), with the smallest density obtained for the equimolar composition, in agreement with Stepanov et al. [21]. The same lowering of the density by Al substitution to Cr was reported by Senkov et al. for CrMo_0.5NbTa_0.5TiZr and AlMo_0.5NbTa_0.5TiZr with densities of 8.23 g/cm³ and 7.40 g/cm³, respectively [22].

For all the compositions, δ and VEC were in the range where single-phase bcc solid solutions are expected: δ < 6.6% [23] and VEC < 6.87 [24].

All as-cast Al_x(TiVNb)_1−x alloys with x = 0 [17], 0.05, 0.10 [17], 0.175 and 0.25 crystallized in a single-phase bcc structure, as can be seen in the XRD patterns in Figure 1a. A preferential orientation along the (110) axis was observed with increasing Al content which decreased the relative intensities of the (200) and (211) peaks. The bcc lattice parameters, as determined by the Rietveld analysis, are listed in

Table 2. Lattice parameters of the Al_x(TiVNb)_1−x (x = 0 [17], 0.05, 0.10 [17], 0.175, 0.25) series in the form of as-cast alloys and hydrides.

Composition	Form	Phase Structure	Space Group	Lattice Parameter (Å)	Reference
TiVNb	as-cast	bcc	Im-3m	3.211(2)	[17]
(TiVNb)H2.0	dihydride	fcc	Fm-3m	4.443(1)
Al0.05(TiVNb)0.95	as-cast	bcc	Im-3m	3.203(1)	Present work
Al0.05(TiVNb)0.95H1.84	dihydride	fcc	Fm-3m	4.418(1)	Present work
Al0.10(TiVNb)0.90	as-cast	bcc	Im-3m	3.197(1)	[17]
Al0.10(TiVNb)0.90H1.59	dihydride	fcc	Fm-3m	4.376(1)
Al0.175(TiVNb)0.825	as-cast	bcc	Im-3m	3.194(2)	Present work
Al0.175(TiVNb)0.825H0.81	hydride	bcc	Im-3m	3.279(6)	Present work
Al0.25(TiVNb)0.75	as-cast	bcc	Im-3m	3.189(1)	Present work
Al0.25(TiVNb)0.75H0.69	hydride	bcc	Im-3m	3.225(10)	Present work

. Corresponding Rietveld refinements of the Al_0.05(TiVNb)_0.95, Al_0.175(TiVNb)_0.825 and Al_0.25(TiVNb)_0.75 compositions are given in Figure S1 of the Supplementary Information (SI).

In this series of alloys, only the equimolar composition AlTiVNb has previously been reported by Stepanov et al. [21] and Yurchenko et al. [25]. Both studies described the obtention of a single phase bcc lattice with unit cell parameters of 3.18 ± 0.03 Å [21] and 3.186 ± 0.001 Å [25]. These findings are in excellent agreement with the present result (3.189(1) Å for Al_0.25(TiVNb)_0.75 in Table 2).

It is of note that Al has been reported to favor the stabilization of bcc phases in HEAs/MPEAs containing refractory metals [26,27] despite its negative enthalpies of mixing with these elements [22]. It was been proposed that the similar atomic size of Al with the alloying refractory elements plays an important role in preventing the formation of intermetallic phases and favors the stabilization of random bcc solid solutions [22].

The bcc lattice parameters of the series Al_x(TiVNb)_1−x steadily decreased with increasing Al content, as shown in Figure 1b, although the trend was not linear, as also pointed out by Chen et al. in the series of Al_xNbTiMoV alloys [26]. The reduced lattice parameter with increasing Al content can be explained by decrease in the average atomic size with Al content since its atomic size radius ranges in the intermediate region: r_Ti = 1.46 Å > r_Nb = 1.43 Å > r_Al = 1.43 Å > r_V = 1.39 Å. Thus, the number of atoms with intermediate sizes increased at the expense of both the largest and the smallest atoms. This behavior has been reported in previous investigations on refractory alloys containing Al [16,17].

The microstructure and chemical distribution of the as-synthesized compositions with Al are detailed in Figure 2 and

Table 3. EDX chemical analysis of the Al_x(TiVNb)_1-x as-cast alloys with x = 0.05, 0.10, 0.175, 0.25.

Composition	Element	Nominal (at. %)	Overall Average (at. %)	Dendritic Zone (at. %)	Interdendritic Zone (at. %)
Al0.05(TiVNb)0.95	Al (K)	5	5.1 (0.4)	4.9 (0.2)	5.7 (0.2)
	Ti (K)	31.6	31.4 (1.9)	30.6 (1.3)	34.4 (0.5)
	V (K)	31.6	31.3 (2.0)	30.4 (1.2)	34.6 (0.4)
	Nb (L)	31.6	32.1 (4.2)	34 (2.6)	25.3 (1.1)
Al0.10(TiVNb)0.90	Al (K)	10	9.8 (0.4)	9.7 (0.4)	10 (0.1)
	Ti (K)	30	29.6 (0.8)	29.3 (0.7)	30.5 (0.3)
	V (K)	30	29.5 (1.1)	28.9 (0.6)	31.1 (0.1)
	Nb (L)	30	31.1 (2.1)	32.1 (1.5)	28.4 (0.5)
Al0.175(TiVNb)0.825	Al (K)	17.5	16.8 (0.2)	16.7 (0.1)	17 (0.3)
	Ti (K)	27.5	27.5 (0.5)	27.2 (0.5)	28 (0.3)
	V (K)	27.5	27.7 (0.5)	27.4 (0.3)	28.3 (0.1)
	Nb (L)	27.5	28 (1.1)	28.7 (0.5)	26.7 (0.6)
Al0.25(TiVNb)0.75	Al (K)	25	25.7 (0.8)	25.3 (0.6)	26.5 (0.7)
	Ti (K)	25	24.6 (0.8)	24.2 (0.5)	25.5 (0.8)
	V (K)	25	24.4 (0.3)	24.3 (0.2)	24.8 (0.3)
	Nb (L)	25	25.2 (1.9)	26.2 (1.1)	23.2 (1.6)

. The overall chemical compositions (Table 3) are in good agreement with the nominal compositions. Dendritic microstructures with a richer Nb composition than interdendritic zones, which are poorer in Nb compared to the nominal concentration, are encountered in all alloys. This microstructure is very common in refractory HEAs [21,26] and has been accounted for by the solidification process—the dendritic zones solidify first and are enriched in elements with the highest melting point, whereas interdendritic regions solidify afterward and contain less high melting temperature metals [28,29]. In our case, Nb had the highest melting point among all the elements (2468 °C_Nb > 1902 °C_V >1660 °C_Ti > 660 °C_Al), explaining why it was found in the dendritic microstructures. Interestingly, the difference in the Nb concentration between the dendritic and the interdendritric zones tended to lessen with increasing Al content, which suggests that the presence of Al helps to increase the overall chemical homogeneity of these alloys.

3.2. Hydrogen Absorption

First, absorption kinetics were recorded under 38 ± 5 bar H₂ pressure at 25 °C for the activated Al_x(TiVNb)_1−x alloys with x = 0.05, 0.175 and 0.25. The maximum capacities were: 0.69H/M (1.25 wt.%), 0.81H/M (1.41 wt.%) and 1.84H/M (2.96 wt.%) for Al_0.25(TiVNb)_0.75, Al_0.175(TiVNb)_0.825 and Al_0.05(TiVNb)_0.95, respectively. The capacities obtained under these conditions and the related XRD patterns for the whole series are plotted in Figure 3a,b. A decreasing trend in capacity was observed with increasing Al content, which might be related to the fact that Al does not have a high affinity for hydrogen compared to Ti/V/Nb which are hydride-forming elements [30]. Furthermore, the same fading in capacity with increasing Al content was observed by Ma et al. in (ZrTiVFe)_100−xAl_x [20], which was accounted for by the decrease in lattice distortion with Al addition.

In terms of gravimetric capacity, the Al_0.05(TiVNb)_0.95 alloy has one of the highest reported capacities (2.96 wt.%) among several Al containing alloys, such as, Mg₃₅Al₁₅Ti₂₅V₁₀Zn₁₅ (2.7 wt.%) [14], Al_0.10Ti_0.30V_0.25Zr_0.10Nb_0.25 (2.9 wt.%) [16], (ZrTiVFe)₉₀Al₁₀ (1.4 wt.%) and (ZrTiVFe)₈₀Al₂₀ (1.16 wt.%) [20]. The hydrogenated material obtained for the smallest Al amount (x = 0.05) was single-phase fcc dihydride (CaF₂-type, $Fm\overline{3}m$), in agreement with our earlier findings for the initial TiVNb and Al_0.10(TiVNb)_0.90 (x = 0.10) [17]. However, for higher Al content (x = 0.175 and 0.25), bcc hydrides were observed (Figure 3b). The lattice parameters for the hydride phases are given in Table 2. The Rietveld refinement for fcc Al_0.05(TiVNb)_0.95H_1.84 is given in Figure S2. For Al contents x ≥ 0.175, hydrogen absorption induced the formation of bcc hydride phases ($Im\overline{3}m$) with slightly increased lattice parameters compared to the as-cast materials due to hydrogen absorption and subsequent lattice expansion (Table 2 and Figure S3). Interestingly, the equimolar composition showed a strong preferential orientation along the (200) direction as indicated by the reinforcement of this diffraction peak relative to the other peak’s intensities (Figure S3b). This might be associated with the occurrence of a well-ordered microstructure following parallel planes induced by hydrogen absorption, as was also observed in an earlier study on TiZrNbTaHf [7].

3.3. Thermodynamics of Hydrogen Absorption

Hydrogen absorption PCI curves for ternary TiVNb and quaternary Al_0.10(TiVNb)_0.90 alloys have previously been reported [17]. Thus, only the PCI curves of Al_0.05(TiVNb)_0.95 recorded at several temperatures (25–175 °C) are provided in this section, whereas the discussion of thermodynamics will extend to previous results.

The PCI investigations for Al_0.05(TiVNb)_0.95 revealed the existence of two plateaus (Figure 4a): the first plateau occurred at low equilibrium pressure (below the detection limit of the pressure transducer) and ended around 0.9 H/M, irrespective of the temperature. The second plateau arose at high pressure, the value of which increased with temperature. This behavior was in good agreement with the two-step reaction (metal ↔ monohydride ↔ dihydride) already reported for both bcc traditional materials [31] and HEAs [7,10,17,32]. The maximum capacity at 25 °C reached 1.84 H/M (2.96 wt.%), in accordance with the full absorption kinetics at room temperature (Section 3.2).

The thermodynamic properties of the dihydride formation in Al_0.05(TiVNb)_0.95 have been determined from the Van’t Hoff plots in Figure 4b. For comparison, the Van’t Hoff plots of the ternary TiVNb and quaternary Al_0.10(TiVNb)_0.90 are also plotted, based on reference information in [17]. The enthalpies (∆H in kJ/mol H₂) and entropies (∆S in J/K·mol H₂) of dihydride formation for x = 0 [17], 0.05 and 0.10 [17] are listed in

Table 4. Thermodynamic properties of hydride formation in Al_0.05(TiVNb)_0.95 together with already reported experimental and ML predicted values for TiVNb and Al_0.10(TiVNb)_0.90 [17].

Composition	ΔHabs (kJ/mol H2)		ΔSabs (J/K·mol H2)
	Experimental	Machine Learning
TiVNb	−67 (±5) [17]	−58 [17]	−157 (±11) [17]
Al0.05(TiVNb)0.95	−52 (±1.5)	−52 [17]	−141 (±3)
Al0.10(TiVNb)0.90	−49 (±1) [17]	−51 [17]	−154 (±2) [17]

together with the earlier ML predictions [17] with very good agreement observed.

The increase in Al content in TiVNb led to a progressive decrease in the experimental enthalpy of dihydride formation from −67 to −52 and −49 kJ/mol H₂ for x = 0, 0.05 and 0.10, respectively. It is of note that a similar thermodynamic destabilization as observed for the addition of x = 0.05 Al content was achieved by the addition of x = 0.25 Cr in Cr_x(TiVNb)_1−x (equimolar composition) [33], suggesting that Al influences hydrogen thermodynamics more effectively than Cr.

Furthermore, the plateaus of the Al_0.05(TiVNb)_0.95 were less sloped than for the ternary pristine alloy, as reported previously [17]. Thus, the decrease in the experimental errors in the thermodynamic calculations with increasing Al content might reasonably be explained by the presence of more flat plateaus in the case of Al-containing alloys.

The two-step transformation induced by hydrogen absorption in Al_0.05(TiVNb)_0.95 was investigated by ex situ SR-XRD. The hydrogenation process was stopped at several hydrogen concentrations. The SR-XRD patterns of the intermediate hydrides with uptakes = 0.5, 0.9 and 1.3 H/M, the dihydride (1.8 H/M), as well as the fully desorbed alloy, are displayed in Figure 5a. Furthermore, for comparison, the ex situ SR-XRD results for the intermediate hydrides of the ternary TiVNb are also shown in Figure 5b. The lattice parameters were determined by Rietveld refinements for both compositions (Figures S4 and S5) and are listed in

Table 5. Lattice parameters of the Al_0.05(TiVNb)_0.95 and TiVNb alloys as dihydride (1.8 and 2.0 H/M), partially hydrogenated (0.5 and 1.3 H/M), monohydride (0.9 H/M), and desorbed phases.

Composition	Form	Phase	Phase Fraction (%)	Lattice Parameter (Å)
Al0.05(TiVNb)0.95H1.8	dihydride	fcc	100	4.412(1)
Al0.05(TiVNb)0.95H1.3	intermediate	bccmh ** fcc	35 65	3.349(1) 4.392(1)
Al0.05(TiVNb)0.95H0.9	monohydride	bccmh **	100	3.342(1)
Al0.05(TiVNb)0.95H0.5	intermediate	bccss * bccmh **	28 72	3.255(1) 3.312(1)
Al0.05(TiVNb)0.95	desorbed	bcc	100	3.206(1)
(TiVNb)H2.0	dihydride	fcc	100	4.439(1)
(TiVNb)H1.3	intermediate	bccmh ** fcc	65 35	3.353(1) 4.425(1)
(TiVNb)H0.9	monohydride	bccmh ** fcc	98 2	3.346(1) 4.423(1)
(TiVNb)H0.5	intermediate	bccss * bccmh **	50 50	3.262(1) 3.328(1)
TiVNb	desorbed	bcc	100	3.219(1)

* bcc_ss = bcc solid solution. ** bcc_mh = bcc monohydride.

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The partially hydrogenated Al_0.05(TiVNb)_0.95H_0.5 consisted of a mixture of two phases: a bcc solid solution with hydrogen with a slightly enlarged lattice parameter (3.255 Å, Table 5) than the initial one (3.203 Å, Table 2) and a bcc monohydride. The alloy completely transformed into a bcc monohydride at 0.9 H/M with a_bcc-mh = 3.342 Å which was around 4% larger than the initial value (3.203 Å, Table 2). At a hydrogen composition of 1.3 H/M, the material consisted of a mixture of two phases: a bcc monohydride (a_bcc = 3.349 Å) and an fcc dihydride (a_fcc = 4.392 Å). At full hydrogenation (1.8 H/M) the material was fully transformed into an fcc dihydride with a_fcc = 4.412 Å (Table 5). After desorption, the material recovered the bcc phase with a very close lattice parameter (3.206 Å, Table 5) compared to the initial alloy (3.203 Å, Table 2). These findings are in good agreement with the PCI curves and indicate that the alloy undergoes a first phase transition from bcc to bcc monohydride followed by a second transition from bcc monohydride to fcc dihydride, which correlates well with previous reports on classical bcc alloys and more recent bcc HEAs. Similar behavior was observed for the pristine TiVNb (Figure 5b and Table 5).

3.4. Hydrogen Desorption

The desorption properties of the hydrides were determined by thermo-desorption spectroscopy (TDS) with a heating rate of 5 °C/min. The TDS profile of Al_0.175(TiVNb)_0.825H_0.81 showed a main desorption event with a maximum desorption peak at 276 °C, whereas the Al_0.25(TiVNb)_0.75H_0.69 revealed two main desorption peaks with maxima at 305 °C and 375 °C (Figure 6a).

The TDS profile of Al_0.05(TiVNb)_0.95H_1.84 (Figure 6b) showed two main desorption events with a desorption onset temperature at around 106 °C. The first peak was rather broad with a maximum at around 150 °C, followed by a second sharp peak with a maximum at 249 °C. This two-step profile was in good agreement with our earlier results for Al_0.10(TiVNb)_0.90H_1.6 [17]_, though the range of desorption occurred at a lower temperature for the alloy with higher Al content. The desorption onset temperature for the fcc-forming hydride materials (x = 0, 0.05 and 0.10) was found to decrease with increasing Al content, as depicted in Figure 6c. Since VEC also depended on the Al content (it decreased with increasing Al content), this trend can be linked with VEC. Thus, the desorption onset temperature increased with the VEC value, as was also recently observed in the series of Ti_0.30V_0.25Zr_0.10Nb_0.25X_0.10 HEAs [34].

The crystalline structure of the desorbed samples was investigated to check their reversibility upon hydrogen absorption and desorption. The XRD patterns and the corresponding Rietveld refinements of the desorbed materials after the TDS measurements (Figure S6) demonstrated that all the desorbed compositions crystallized in a bcc lattice with a lattice parameter close to the as-cast material (Table S1). Interestingly, the composition Al_0.25(TiVNb)_0.75 after desorption showed a bcc structure with a very strong preferential orientation along the (200) direction (Figure S7), which was already introduced by hydrogen absorption (Figure S3b) and retained after desorption.

To better understand the effect of Al on the desorption process and the related crystalline phases’ evolution, in situ SR-XRD measurements for Al_0.05(TiVNb)_0.95H_1.84 and (TiVNb)H_2.0 are plotted in Figure 7a,b, respectively. Rietveld refinements were performed to acquire structural information about the variation in phase fraction and lattice parameters during desorption (Figure 7c–f). The full hydrides and desorbed phases of Al_0.05(TiVNb)_0.95H_1.84 and (TiVNb)H_2.0 are shown in Figure S8a–d, respectively.

The initial fcc dihydride Al_0.05(TiVNb)_0.95H_1.84 (marked with ● in Figure 7a) was stable from 25 °C up to around 150 °C, while the lattice parameter increased from 4.418(1) to 4.421(1) Å (Figure 7d and Table S2) because of the thermal expansion of the crystal lattice. Upon reaching around 160 °C, the fcc dihydride started to desorb hydrogen and a new bcc phase was formed (marked with * in Figure 7a). This temperature was close to the first desorption event observed in the TDS spectra (Figure 6b), though a shift in temperature could be noticed due to different heating ramps and vacuum environments. At above 160 °C, the two phases coexisted, and the lattice parameter of the newly formed bcc phase smoothly decreased, while the parameter for the fcc phase ceased to increase and even showed a decrease from 190 °C up to 300 °C (Table S2). This can be explained by the combination of the thermal expansion and the lattice shrinkage by hydrogen desorption occurring at the same time. This coexistence of two phases (bbc and fcc) has been observed previously by in situ neutron diffraction of Al_0.10(TiVNb)_0.90D_1.66 during deuterium desorption [17]. In the temperature range 250–350 °C the bcc phase strongly increased at the expense of the fcc dihydride (Figure 7c), while both the fcc and bcc lattice parameters decreased (Figure 7d and Table S2). This temperature range corresponded to the main sharp TDS event at high temperature (Figure 6b). Starting from 300 °C, the bcc lattice parameter began to steadily decrease up to the maximum temperature 450 °C. At 350 °C, the fcc phase completely disappeared and the material was transformed into a bcc phase. From the thermal evolution of the bcc lattice parameter above 300 °C, it seems that this material was still desorbing hydrogen at a maximum temperature of 450 °C. At this temperature, the bcc lattice parameter was 3.229(1) Å (Table S2), which was below the value of 3.342 Å obtained for the bcc monohydride at room temperature (Table 5). Moreover, the bcc lattice parameter value of 3.229(1) Å at 450 °C was larger than 3.206(1) Å obtained for the desorbed phase at room temperature (Table 5). However, this difference can be reasonably explained by the thermal expansion. This suggests that the material had completely desorbed hydrogen under these conditions, and that, at maximum temperature, the transformation from the fcc dihydride to the initial bcc phase was complete. Thus, from the in situ SR-XRD, only one transition was noticeable from fcc to a bcc phase without the bcc monohydride being discernible under these conditions.

The in situ SR-XRD profile from the desorption of the ternary (TiVNb)H₂ (Figure 7b) showed similar behavior, with the main difference being that the fcc to bcc transition was sharper and occurred in the short temperature range 300–350 °C, without any previous coexistence of the phases. The lattice parameters of bcc and fcc phases showed a similar trend to that for the Al_0.05(TiVNb)_0.95 alloy. Although initially the fcc lattice parameter increased with temperature, starting from 200 °C, it decreased, with this phase completely disappearing at around 350 °C. The bcc lattice parameter steeply decreased up to 350 °C and approached an almost stable value at higher temperature. Comparable to the Al_0.05(TiVNb)_0.95 alloy, the bcc lattice parameter at 450 °C was 3.234(1) Å (Table S2), which was smaller than 3.346(1) Å for the bcc monohydride at 25 °C (Table 5) but slightly larger than 3.219(1) Å for the desorbed material at room temperature (Table 5). This suggests that the desorption was almost complete for this alloy under the present experimental conditions. However, only one phase transition from fcc to bcc lattice was observed in the in situ SR-XRD, comparable to the Al_0.05(TiVNb)_0.95 alloy.

If both alloys undergo a reversible two-step reaction: bcc phase ⇄ bcc monohydride ⇄ fcc dihydride, as demonstrated by the PCI curves and the ex situ SR-XRD measurement, the question arises: why was the transition from the bcc monohydride to the desorbed bcc phase not visible from the in situ SR-XRD? Presently, the origin of this effect is not completely certain; however, it might be attributed to the fact that this transition is more difficult to detect in situ as it involves a very small displacement/rearrangement of atoms among the same bcc crystal structures, as was previously suggested for Ti_0.30V_0.25Cr_0.10Zr_0.10Nb_0.25 [10].

The most important difference between the in situ SR-XRD experiments on Al_0.05(TiVNb)_0.95 and the ternary TiVNb was the desorption that occurred at lower temperature for the Al-containing alloy compared to the pristine material. Moreover, both the fcc dihydride and the bcc desorbed phases coexisted within a large temperature range for Al_0.05(TiVNb)_0.95, as was also observed for Al_0.10(TiVNb)_0.90 during the in situ neutron diffraction reported in our previous work [17], whereas the ternary alloy showed a sharp transition from the dihydride to the desorbed phase at high temperature. This was in very good agreement with the TDS results and highlights the important role of Al in improving the hydrogen desorption process from these materials.

3.5. Hydrogen Absorption/Desorption Cycling and Kinetics

Lastly, Al_0.05(TiVNb)_0.95 was submitted to 10 cycles of hydrogen absorption/desorption to evaluate the reversible storage capacity, the crystalline structure stability, and the effect on the reaction kinetics. The cycling process consisted of complete hydrogen absorption with a final P_eq around 32 bar at 25 °C, followed by complete desorption at 410 °C under a secondary vacuum for 12 h. The kinetics were recorded during the absorption of hydrogen for each cycle (Figure 8a). The first cycle achieved maximum capacity in less than 40 min and, for subsequent cycles, the absorption kinetics improved by reaching maximum capacity in less than 10 min. This can be accounted for by the hydrogen decrepitation effect that causes the crumbling of alloys and decrease in the grain size. Thus, the kinetics were enhanced by two main processes: (i) new surface areas were created that allowed faster dissociation of hydrogen molecules, and (ii), the diffusion rate within the material increased due to shorter pathways within the grains.

The capacity slightly decreased from 1.84 H/M to 1.80 H/M after the first cycle and stabilized at 1.76 H/M at the 10th cycle. Interestingly, the stable capacity of Al_0.05(TiVNb)_0.95 in terms of wt.% (2.83 wt.%) was higher than both values for Al_0.10(TiVNb)_0.90 (2.55 wt.%) and TiVNb (2.8 wt.%) that we reported previously [17], as shown in Figure 8b. The Al_0.05(TiVNb)_0.95 alloy lost around 4.5% of the initial capacity until stabilization occurred (from 2.96 wt.% to 2.83 wt.%), which was higher than the 1.8% loss observed for Al_0.10(TiVNb)_0.90, but smaller than the 14 % loss for TiVNb. This suggests that Al addition in TiVNb positively affected the reversible capacity when limited Al amounts were used. An optimum capacity was found for x = 0.05 Al content in TiVNb with very good cycling stability and one of the best stable capacities reported so far. The improvement in the stability of the reversible capacity and the gravimetric capacity of hydrogen storage by Al addition as a lightweight element was in very good agreement with the behaviors reported for Ti-V-Zr-Nb with 10 at.% Al [16].

After the cycling process, the Al_0.05(TiVNb)_0.95H_1.76 hydride phase was analyzed by XRD (Figure S9b). This demonstrated that the fcc dihydride was present along with a monohydride bcc phase, suggesting that the absorption was not completed, explaining the observed decrease in capacity during cycling. The same behavior was previously observed for Al_0.10(TiVNb)_0.90 and TiVNb [17]. The Rietveld refinements for the cycled phases are given in Figure S9. In addition, the SEM-EDX chemical mapping of Al_0.05(TiVNb)_0.95H_1.76 after 10 cycles is shown in Figure S10. The same microstructure and overall chemical composition (Table S3) as the initial alloy were observed without any phase segregation, in good agreement with earlier results [17].

A simple comparison among these three compositions suggests that, not only the addition of non-absorbing lightweight elements in refractory MPEAs is important, but also that the composition must be carefully tuned to optimize both the structure/microstructure of the alloys and their hydrogen sorption performances.

4. Conclusions

The effect of Al addition on the structure, microstructure, and hydrogen sorption properties of a ternary TiVNb alloy was studied from low Al content up to equimolar composition. Al_x(TiVNb)_1−x alloys with x = 0–0.25 were prepared by arc melting as single-phase bcc materials with a decreasing lattice parameter by increasing the Al amount. They showed good chemical homogeneity with the typical Nb-rich dendritic microstructure of refractory HEAs.

The alloys absorbed hydrogen quickly at room temperature with the storage capacity showing a strong dependence on the Al content. Under the experimental conditions used, the richest Al alloys (x = 0.175 and 0.25) formed bcc solid solutions after hydrogen absorption with small hydrogen uptakes (0.81 H/M and 0.69 H/M, respectively), whereas the alloys with low Al content (x = 0.05 and 0.10) absorbed large amounts of hydrogen-forming fcc hydrides. Moreover, a strong thermodynamic destabilization of the hydride formation was observed with increasing Al content. The enthalpy of hydride formation reduced from −67 kJ/mol H₂ for the pristine TiVNb alloy to −52 and −49 kJ/mol H₂ for x = 0.05 and 0.10 Al quantities, respectively. Interestingly, the desorption onset temperatures from the fcc dihydride forming materials (x = 0, 0.05 and 0.10 Al) showed a strong decrease with increasing Al content, indicating thermal destabilization of the hydrides by Al addition.

The Al_0.05(TiVNb)_0.95 alloy underwent a reversible two-step reaction from the initial bcc to an intermediate bcc monohydride and, finally, to an fcc dihydride, as confirmed by the two plateaus observed in the PCIs and the ex situ synchrotron XRD. Moreover, a flattening of the plateaus by Al addition was observed which was an important factor in the reduction in errors in the calculation of the thermodynamic properties by the Van’t Hoff equation. Intriguingly, during hydrogen desorption, only one phase transition from the fcc dihydride to the bcc desorbed state was detectable by in situ SR-XRD for both Al_0.05(TiVNb)_0.95H_1.84 and TiVNb. The expected two-step reaction fcc dihydride → bcc monohydride → bcc desorbed phase was not noticeable, though complete desorption was confirmed by structural analysis. One plausible explanation might be that the bcc monohydride → bcc desorbed phase transition was more difficult to detect since it involved a very small rearrangement of atoms within the same bcc crystalline structure. However, the in situ SR-XRD clearly demonstrated the coexistence of an initial fcc dihydride with the bcc desorbed phase in a wide temperature range for Al_0.05(TiVNb)_0.95, in contrast with the thermal behavior of TiVNb.

Finally, the Al_0.05(TiVNb)_0.95 alloy possessed better cycling properties than both the TiVNb and Al_0.10(TiVNb)_0.90 alloys, with a stable reversible capacity of 1.76 H/M (2.83 wt.%), which is one of the best in the recent literature of MPEAs designed for solid-state hydrogen storage.