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    AxBy Intermetallics for hydrogen storage

    Subjects: Ergonomics
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    Submitted by: Julien O. Fadonougbo


    AxBy intermetallics show outstanding performances, notably for stationary hydrogen storage applications. Elemental substitution, whether on the A or B site of these alloys, allows the effective tailoring of key properties such as gravimetric density, equilibrium pressure, hysteresis and cyclic stability for instance.

    1. Introduction

    Humankind is on the verge of facing a worldwide energy crisis considering the soon-to-come fossil fuels shortage. The transition to environmentally friendly energy sources is a challenge that many countries are already tackling by reformatting their economy to implement alternative and sustainable solutions. As such, the hydrogen-based economy became one of the main candidates for the transition towards cleaner energy source, in the light of hydrogen’s positive impact on the environment and its intrinsic great potential as an abundant energy carrier: (i) high gravimetric energy density of 142 MJ kg−1 (against only 47 MJ kg−1 for petroleum) and (ii) high energy efficiency (fuel cells electrochemical processes show ~50–60% efficiency whereas that of combustion engines is as low as 25% for hydrogen-air mixtures, but still slightly higher than petrol-air) [1].

    The main requirements for a large scale application of metal hydrides for on-board applications are (i) low hydrogen release temperature in the typical working conditions of a PEM fuel cell, (ii) high hydrogen absorption and desorption rates, (iii) acceptable costs and most importantly (iv) high storage capacity of 8 wt% according to the recent European VII FP call [2], which sets the bar even higher than the 6 wt% targeted by the American Department of Energy (DOE) [3]. It is difficult to achieve the gravimetric capacity target, especially for intermetallic hydrides. Hence, the main application for the intermetallic hydrides would be stationary applications, which are essential parts of renewable energy systems.

    2. AB5-Type Alloys

    The AB5-type hydrides have been intensively studied during the last decades for their high potential for practical applications [4]. They have reversible and fast hydrogen absorption/desorption kinetics at near-ambient temperatures, simple activation process, and moderate pressure-temperature conditions of hydrogenation/dehydrogenation, which can easily be controlled. However, the maximum discharge capacity is limited to only around 1.5 wt% for the single CaCu5-type hexagonal structure [5][6].

    In this section, LaNi5 is taken as the reference material of the AB5 family, in the light of its remarkable properties and features in comparison with other AB5 compounds that were recently studied. Indeed, LaNi5-based hydrides show good hydrogen absorption/desorption characteristics under near-atmospheric conditions and excellent kinetics [7][8]. The amount of hydrogen desorbed from a typical LaNi5-type metal-hydride system ranges from much less than 1 wt% up to 1.2 wt% H2 between room temperature and 373 K, with a theoretical maximum reversible storage capacity of 1.5 wt% H2 (still below the DOE’s target) [9]. Despite attractive properties, LaNi5-based compounds have a high cost in comparison with other alloys and show a significant capacity loss (higher than 30% after 800 cycles under impure hydrogen gas containing 100 ppm of O2 [10]), therefore urging to develop other materials with higher discharge capacity, better cyclic stability and lower cost [11].

    The costly lanthanum in LaNi5 can thus be replaced by cheaper rare earth elements such as Ce [12], or by a cheaper rare earth mixture called mischmetal (Mm) consisting of La, Ce, Pr and Nd [13], which was investigated in many studies. MmNi5 possesses a hexagonal crystal structure similar to that of LaNi5 and tends to form stable hydrides. However, it shows a very high activation pressure (120 atm at 298 K), a high hydride formation pressure (30–60 atm at 298 K), large hysteresis between the absorption and desorption pressures and a maximum storage capacity of about 20% lower than that of LaNi5 [14][15]. Many groups have attempted to reduce the high hydride formation pressure in MmNi5 by partially substituting A and B components with various elements [16][17][18].

    To enhance the hydrogen storage capacity, Ca may partially replace Mm in MmNi5 because of its lightweight (at. wt. 40) in comparison to Mm (at. wt. 140, corresponding to the following composition: La 22%, Ce 52%, Nd 15% and Pr 11%). Hence, for H/M = 1.0 the storage capacity of MmNi5H6 is 1.38 wt%, while that of Mm0.66Ca0.34Ni5H6 corresponds to 1.5 wt% [19][20]. The studies on Mm1-xCaxNi5 were first reported by Sandrock [21] and Shinar et al. [22]. Sandrock’s results show that the hydride dissociation pressure decreased with increasing Ca content, while Shinar’s results indicate that the substitution of Ca for Mm or La caused an increase in hydride dissociation pressure. Such contradictory behaviour results from the variation in Ca content, as elucidated by Wang et al. [23]. Indeed, they reported that the dissociation pressure of the hydrides (at 298 K) increased when x < 0.3 but decreased when 0.3 < × < 0.9, which was attributed to the effect of geometrical and electronic factors. In addition, the first hydrogenation incubation time shortened and its absorption rate increased along with increasing x in Mm1 − xCaxNi5, and the hysteresis reduced.

    Different from Mm and Ca (A substitutes), substitutions for B element were reported to be effective in tailoring the plateau pressure. Among them, Al was used for reducing the plateau pressure, for instance from 50 atm for MmNi5 down to 0.5 atm for MmNi4.2Al0.8. However, the maximum storage capacity decreased from 1.44 to 1.3 wt% and the plateau slope increased [24]. Meanwhile, Fe is known to increase hydrogen storage capacity (1.5 wt% for MmNi4.6Fe0.4), and reduce sloping and hysteresis [25][26].

    3. AB2-Type Alloys

    AB2 Laves phase is another type of alloy with high potential for hydrogen storage. Usually, these alloys exist in three different crystal structures: cubic C15 (for instance MgCu2, ZrV2), hexagonal C14 (MgZn2, ZrMn2) and double hexagonal C36 (MgNi2). Laves phases with A = Zr show relatively high capacities (ZrV2H5.3, ZrMn2H3.6, ZrCr2H3.4), faster kinetics, longer lifetime and a relatively low cost in comparison to the LaNi5-based alloys. However, their hydrides are too stable at room temperature and more sensitive to contaminants [27]. This high stability of Zr-containing alloys is also seen in various type of materials, notably in amorphous structures in which hydrogen is irreversibly immobilized either in trapping sites [28][29][30], or by forming stable ZrH2 phase [31][32].

    In this section, we take Ti–Mn Laves phase alloys as the reference material of the AB2 family, because of their easy activation, good hydriding-dehydriding kinetics, high hydrogen storage capacity and relatively low cost. Besides, they display high plateau pressure at room temperature (over 20 atm) and a sloping plateau often accompanied with a large hysteresis that requires major improvements [33][34].

    In 2005, Toyota’s group demonstrated the use of Ti1.1MnCr alloys in a high-pressure metal hydride (MH) tank. This alloy has a maximum storage capacity of 1.9 wt%, but it has been reached only for a hydrogen pressure of around 350 atm at room temperature [35]. Kandavel et al. [36] substituted Zr in Ti1.1CrMn to provide favorable hydrogen sorption conditions and maximize the storage capacity. The increase in Zr content leads to a decrease in the equilibrium plateau pressure and faster absorption kinetics, together with an increase in the hydrogen storage capacity from 1.9 to 2.2 wt% for Ti1.1CrMn and (Ti0.9Zr0.1)1.1CrMn, respectively. Besides, Park et al. [37] conducted studies on Ti–Zr–Mn–Cr based metal hydrides and concluded that when Zr/Ti ratio increases, the lattice strain increases. This is partially responsible for a drastic increase of sloping, while the use of Cu was found very effective to mitigate the sloping.

    In 1995, Morii et al. [38] prepared and investigated (Ti, Zr)(Ni, Mn, X)2 alloys, where X is V or/and Fe. The results showed that V lowers both hysteresis and plateau pressure. On the other hand, Ni raises the plateau pressure and reduces the width of the plateau region, while Fe flattens and lengthens it.

    Improvements of the hydrogen storage properties of Laves phase AB2-type alloys at 303–308 K and 1–15 atm have been achieved by introducing non-stoichiometry at the A site of (Ti0.65Zr0.35)1 + xMnCr0.8Fe0.2 alloys. From pressure-composition-temperature (PCT) measurements, the maximum hydrogen storage capacity was found to be around 2.2 wt% at 35 atm and 305 K for (Ti0.65Zr0.35)1.1MnCr0.8Fe0.2, which is approximately 16% higher than that of the commercially available “Hydralloy C5” (Ti0.955Zr0.045Mn1.52V0.43Fe0.12Al0.03). These alloys show remarkable hydrogenation kinetics: the full capacity is reached within 10 min without any need for activation [39].

    Alloys without zirconium (such as Ti1.02Cr1.0Fe0.75Mn0.25) display 1.55 wt% of reversible hydrogen storage capacity when the temperature is as low as 233 K. However, without zirconium the effective hydrogen capacity is optimal only when the pressure is higher than 70 atm [40], proving the effectiveness of Zr in Laves phase alloys.

    Recent developments (<5 years) on AB2-type materials have highlighted their significant potential for high-pressure compressors, notably (Ti,Zr)(Mn,Cr)-based alloys. Indeed, Corgnale et al. [41] proposed a techno-economic analysis of metal hydride systems for efficient and novel high-pressure compressors. Among various materials, TiCr1.9, Ti1.1CrMn, TiCrMn0.4Fe0.4V0.2, and (Ti0.97Zr0.03)1.1Cr1.6Mn0.4, they suggested the last one as the best candidate for their novel two-stage hybrid electrochemical and metal hydride compression system, since pressures about 863 atm can be reached with a thermal power provided at approximately 423 K.

    Pickering et al. [42] further demonstrated the high capability of (Ti,Zr)(Mn,Cr)-based alloys for both hydrogen storage and high-pressure compression by producing industrial volumes (∼10 kg) of tailored AB2 intermetallics (A = Ti + Zr, B = Cr + Mn + Ni+Fe + V) by means of vacuum induction melting process. They successfully tuned the hydrogenation properties of the alloy, showing that at a fixed quite low Zr/(Ti + Zr) ratio the PCT properties of the materials can be adjusted in a wide range by the variation of V content which, in addition, results in the increase of the hydrogen storage capacity. Cheaper alternatives to pristine Ti and V nevertheless exist, notably by replacing those high purity raw materials by their low-cost and low-purity counterparts, namely Ti sponge and ferrovanadium (FeV), respectively. Such substitution in (Ti,Zr)(V,Fe,Cr,Mn) reduces the raw material cost by 83%, without altering the dissociation pressure (15 atm), nor the reversibility (1.4 and 1.5 wt% H2 after 1000 cycles, against an initial capacity of 2 and 1.7 wt% H2 for pristine and modified alloys, respectively) [43].

    The development of hybrid hydrogen storage system is equally appealing to the scientific community. For instance, rare earth elements (RE) such as La, Ce or Ho in Ti1.02Cr1.1Mn0.3Fe0.6RE0.03 have been shown in 2018 to yield better activation behaviour, larger storage capacity but lower desorption plateau pressure [44]. This study suggests Ti1.02Cr1.1Mn0.3Fe0.6La0.03 alloy as the best overall candidate since it can be fully activated at room temperature, and has a hydrogen storage capacity as high as ~1.7 wt%. Another example of hybrid system is reported by Puszkiel et al. [45], who demonstrated that mixing expanded natural graphite (ENS) into (Ti0.9Zr0.1)1.25Cr0.85Mn1.1Mo0.05 alloy not only improves the heat transfer properties, but also yields a reversible capacity of about 1.5 wt%, together with decent cycling stability and rapid reaction kinetics (25 to 70 s).

    Although all the above-mentioned (Ti,Zr)(Mn,Cr)-based Laves phase alloys are widely investigated in the light of their superior potential for high-pressure compressors (and hybrid hydrogen storage), Zr-based AB2 materials are nevertheless not to be discarded although they display significantly lower desorption plateau pressures. Wu et al. [46] thus elucidated the role of Ni addition on the hydrogen storage characteristics of Zr(V1 − xNix)2 (x = 0.02, 0.05, 0.1, 0.15, 0.25) intermetallic compounds. The hydrogen absorption capacity turns out to decrease, and the equilibrium pressure increases with increasing Ni content. The alloys exhibit fast absorption kinetics at room temperature and a remarkable cyclic stability even after 100 hydrogen absorption/desorption cycles.

    Owing to fast kinetics, high equilibrium pressure and impressive volumetric hydrogen storage density at ambient temperature, ZrFe2 based alloys are similarly good candidates for high pressure compressed hydrogen tanks. To bypass its rather large hysteresis, Mn, Ti, V and Cr addition [47][48] has been considered. On one hand, V addition is suggested to improve the hysteresis, while Ti helps to lower plateau sloping as well as to increase the plateau pressure. Zr1.05Fe1.6Mn0.4 shows a relatively high dehydriding pressure of 20.6 atm at 298 K, while (Zr0.5Ti0.5)1.05Fe0.95MnV0.05 delivers a maximum capacity of 1.64 wt% H2 and shows a dehydriding pressure of 6.8 atm at 298 K (calculated from Van’t Hoff plots) [49]. Additionally, the simultaneous Cr/V substitution for Fe decreases the equilibrium pressure (due to the enlarged unit cell), and Zr1.05Fe1.85Cr0.075V0.075 seems to exhibit decent overall hydrogen storage properties (1.54 wt%, and a desorption equilibrium pressure of 9.7 atm at 243 K) [50].

    4. AB-Type Alloys

    AB-type alloys are attractive materials for hydrogen storage because of their light molar mass and high weight capacities. TiFe alloys with cubic CsCl-type structures are the most known alloys of this class and stand among the best hydrogen storage materials up to this date [51].

    TiFe intermetallic compound is one of the most promising hydrogen storage alloys, due to its relatively high theoretical hydrogen storage capacity (1.9 wt%) at near-ambient conditions compared to other AxBy families. Besides, its economical merit based on the abundance and low cost of the constituting elements encourages extensive investigations on the TiFe system.

    The hydrogen sorption and desorption in TiFe was first described by Reilly and Wiswall in the year 1974 [52]. They reported two stable intermetallics of TiFe system (TiFe and TiFe2) and a third, Ti2Fe that forms only above 1273 K (dissociates to TiFe and Ti below that temperature). Only TiFe is known to make two ternary hydrides, TiFeH and TiFeH2.

    The hydrogen absorption in TiFe alloy depends on two factors: (i) the Fe/Ti ratio and (ii) the oxygen amount in the alloy. TiFe intermetallic exists over a narrow composition range of ~2.5 at% (from 49.5 to 52 at% Ti). Slightly less than 49.5 at% Ti results in a two-phase mixture of TiFe2 and TiFe, the first being of no use since it is a non‑hydride former. If Ti content is higher than 52 at%, the alloy consists of TiFe and (α or β) Ti solid solution [52]. Although Ti itself readily forms hydrides, they are highly stable and are non-reversible at the temperatures of interest (ambient).

    The lower plateau level and general shape of the curve is not significantly affected but the maximum hydrogen storage capacity substantially reduces with the increase in oxygen content [53]. Additionally, TiFe usually requires heating over 573 K for activation, which again suggests the low poisoning tolerance resulting in significant deterioration of hydrogen sorption even for trace amounts of gas species (oxygen and water vapor for instance) [54][55]. Most importantly, surface oxidation issues induce significant difficulties notably in the first hydrogenation. The problem with first activation can be resolved by partial replacement of the base element [56][57][58][59][60][61], mechanical alloying [62][63], surface modifications [64], groove rolling and high-pressure torsion [65]. Most of these studies did not lead to an improvement in hydrogen storage properties, and the result was usually a decreased maximum hydrogen absorption capacity and increased desorption temperature of the intermetallic hydrides.

    Very recently, in the year 2020, Yang et al. [66], documented the effect of Cr, Mn and Y substitution for Fe on the hydrogen storage properties. They concluded that Cr substituted alloys (TiFe0.9Cr0.1, TiFe0.9Cr0.1Y0.05) have lower equilibrium pressure and sloped plateaus, thus providing better hydrogenation kinetics as compared to Mn substituted alloys (TiFe0.9Mn0.1, TiFe0.9Mn0.1Y0.05), which have higher equilibrium pressure but flat plateaus and thus better dehydrogenation kinetics. Y substitution in Ti–Fe–Mn and Ti–Fe–Cr based alloys resulted in αY phase, which transforms to YH3 during hydrogenation.

    Ha et al. [67] investigated the contrast in the microstructure of as cast and heat treated TiFe-6 wt% ZrCr2 alloys. They reported that the as cast alloy has 65 wt% TiFe and 35 wt% TiFe2 (C14 Laves phase) while the heat-treated alloy has a portion of TiFe2 transformed to TiFe phase (84 wt%). The activation profile reveals that both the alloys can be activated at room temperature under 30.6 atm H2 but the as cast alloy displays enhanced absorption kinetics (activation starts without any delay while its heat-treated counterpart requires 40 h of incubation time). Both specimens show approximately equal maximum hydrogen storage capacity of 1.7 wt%. The first plateau for the annealed alloy is flatter in shape and the desorption isotherm shows less retained hydrogen as compared to the as cast alloy. In parallel, Jung et al. [68] conducted a study on tailoring the equilibrium plateau pressure of TiFe monohydride and dihydride via V substitution for both Ti and Fe, in order to achieve maximum reversible capacity under a narrow pressure range. When V substitutes for Ti, the monohydride plateau pressure rises whereas a pronounced opposite trend is seen if V substitutes for Fe. Interestingly, the plateau pressure for dihydride is lowered in both the cases.

    5. AB3-Type Alloys

    Research on AB3 alloys, whose structure consist of combined AB2 and AB5 (see equation below), was initially motivated by their strong potential for Ni-MH batteries [69]. Indeed, negative electrode materials based on AB3 can offer a higher hydrogen storage capacity than AB5-types alloys (already commercialized), but unfortunately suffer a severe degradation of their cyclic properties due to pulverization and oxidation/corrosion [70][71].

    AB5 + 2(AB2) = 3(AB3)


    Most frequently based on La2MgNi9, AB3 alloys however turn out to be promising for stationary hydrogen storage applications as well, considering their good activation and hydrogenation/dehydrogenation kinetics on one hand, and their relatively high storage capacity and low cost on the other (thus combining the best features of AB5 and AB2 respectively) [48]. Additionally, the phase composition of AB3 alloys (hence their properties) can be tuned by means of element substitution, heat treatment and different material processing methods, similarly to AB2 alloys [72].

    Pioneering work in the seventies [73][74] first reported hydrogen solubility and hydride forming ability of AB3 alloys based on rare earth elements (A side) and transition metals (B side). Later on, Kadir et al. further investigated such alloys, by providing exhaustive reports on the effect of rare earth elements on the hydrogenation properties of AB3 alloys (La, Ce, Pr, Nd, Sm, Gd) [75], as well as on the effect of La and Mg partial replacement by Ca and/or Y in La-Mg-Ni based alloys [76][77].

    The hydriding characteristics of LaNi3/CaNi3 and RT3 phases (R = Dy, Ho, Er, Tb, Gd; T = Fe or Co) showed that the hydrogen storage capacity of the AB3 phases exceeds that of the well-known hydrogen absorber LaNi5 [78]. Due to the special crystal structure of AB3 compounds, it is possible to combine Mg, Ca, and rare earth elements in the A side. Kadir et al. [79]synthesized (La0.65Ca0.35)(Mg1.32Ca0.68)Ni9 which absorbs ~1.87 wt% H2 at ~33 atm H2 and 283 K. Under identical pressure condition, Chen et al. [78] reached up to 1.8 wt% H2 at 293 K for LaCaMgNi9.

    In order to improve the performance of La–Mg–Ca–Ni AB3-type alloy, Lim et al. investigated the effects of partial substitution with Ce and Al on the hydrogenation properties of La0.65-xCexCa1.03Mg1.32Ni9-yAly alloys [80]. Their results indicated that the hydrogen storage capacity significantly decreased after Ce and Al substitution. Xin et al. [81] investigated the effects of Y partial substitution on overall hydrogen storage properties of (La0.65Ca0.35)(Mg1.32Ca0.68)Ni9. At 1 atm H2, the hydrogen desorption capacity of La0.60Y0.05Mg1.32Ca1.03Ni9 was approximately 1.624, 1.616, and 1.610 wt% at 298, 313, and 333 K, respectively. In addition, the equilibrium pressure could be tailored by altering the Y amount to range 1–10 atm.

    The effect of half replacement of Ca by R (R = Nd, Gd and Er) on the phase structure and hydrogen storage property of Ca2MgNi9 compound was investigated in 2019 by Zang et al. [82]. Results showed that alloys with Gd, Er or Nd instead of La have lower maximum storage capacity (1.4, 1.2, and 1.5 wt% H2, respectively, against 1.87 wt% H2 for La). Desorption behaviours of some remarkable AB3 alloys (plotted in Figure 7) show flatter plateau pressures than some AB2 and AB alloys while displaying comparable storage capacity (see detailed summary in Table 6).

    To summarize, partial substitution in the B site of Ni for elements with larger atomic radius increases the unit cell volume and results in a decrease of the absorption and desorption plateau pressures. As such, increasing Co concentration (in La0.7Mg0.3Ni3.4–xMn0.1Cox [83] or in La2Mg(Ni1–xCox)9 (x = 0.1–0.5) [84]), or Al and Mo (in La0.7Mg0.3Ni3.5–x(Al0.5Mo0.5)x with x = 0–0.8 [85]) decreased the desorption equilibrium pressure.

    6. Solid Solutions

    Metallurgically speaking, the term ”solid solution alloy” designates a primary element (solvent) into which one or more minor elements (solutes) are dissolved. Unlike the intermetallic compound, the solute does not need to be present at an integer or near-integer stoichiometric ratio and is present in a random (disordered) substitutional or interstitial distribution within the basic crystal structure. Several solid solution alloys form reversible hydrides, in particular those based on Pd, Ti, Zr, Nb and V solvents [86].

    Despite excellent properties such as fast absorption/desorption kinetics and large hydrogen gravimetric density of maximum 3.8 wt% at moderate temperatures, V-based alloys suffer major drawbacks preventing their rapid and widespread applications. These limitations are (i) the relatively difficult first activation, and (ii) the high thermal stability of its hydride phases yielding poor cyclic performance (reversible capacity down to ~2 wt% H2 at room temperature) [87].

    Upon hydrogenation, V forms a solid solution α followed by β phase (V2H with body-centered tetragonal structure) and then the γ phase (VH2 with CaF2 crystal structure), whose respective thermal stability drastically differs. Indeed, the β phase is so stable that its hydrogen desorption reaction never occurs under moderate conditions, its desorption pressure usually ranging 10−5–0.1 atm. On the other hand, the γ phase is not as stable as its β counterpart and its hydrogen absorption/desorption reaction occurs at moderate temperatures and pressures (over 1 atm at room temperature). Therefore, due to the stability of the β phase, only about half of the amount of hydrogen absorbed in vanadium metal can be used in the hydrogen absorption and desorption processes under practical conditions [87].

    Thermodynamic destabilization of the β phase of pristine V stands out as the main solution to tackle the issues mentioned above. Hence, similarly to any other AxBy alloy category, the use of alloying elements of diverse nature and simultaneous addition (binary, ternary and quaternary systems for instance) can destabilize the hydride phases, by altering the ionicity, electronic density of states and lattice parameters [88].

    Binary V-based systems cover a broad range of elements, with Ti being the most studied one in the light of its high solubility in V [89], the improved hydrogenation rates and increased terminal solid solubility (TSS) of hydrogen [88]. Although Ti is widely utilized, other elements such as Si, Al and Fe are also considered, but turn out to decrease the hydrogenation rates [90][91], while Mo addition increases hydrogenation-dehydrogenation pressure and decreases the hydrogen storage capacity for instance [92].

    To push further the enhancement brought by binary alloys, ternary systems have been developed, notably V–Ti–Cr which remains the most documented ternary alloy due to its excellent improvement of the cyclic stability (as compared to its former binary V–Ti counterpart) while maintaining high effective capacity at room temperature [93][94][95]. Storage capacity can be controlled and increased by tuning the compositional ratio of those three elements, for instance in a mixture of 60 at% V, 15 at% Ti and 25 at% Cr which reaches as high as 2.62 wt% [96] V–Ti–Cr alloys however show a steep slope of hydrogen absorption–desorption plateaus, requiring homogenization by heat treatment [97][98] and melt-quenching treatment [99][100]. Besides, the formation of an enriched Ti phase during heat treatment and the oxidation of Ti during melt-quenching both reduce the amount of stored hydrogen and complicate the activation process [101].

    In spite of the attractive storage capacity of V–Ti–Cr alloys, they remain expensive since the price of pure V is very high. Fe can thus be used as a replacement of V in ternary systems, and excellent storage capacity of 3.9 wt% with a reversible capacity of 2.4 wt% are reported for Ti43.5V49Fe7.5 (at 253 K) [102]. Fe also shows a great potential for tailoring plateau pressures, for instance in (V0.9Ti0.1)1 − xFe alloys (with x = 0–0.075) [103]. The reduction of costs by Fe addition has also been attempted for quaternary alloys, notably by Luo et al. [104] who synthesized V48Fe12Ti15Cr25. The maximum hydrogen storage capacity of this alloy reached 1.98 wt% at 315 K, which is lower than that of other V–Ti–Cr series alloys, due to smaller lattice constant and cell volume.

    The lattice constant of the alloys is closely related to the amount of hydrogen absorbed/desorbed [105][106]: V48Fe12Ti15Cr25 has smaller interstitial sites, which could lead to a lower hydrogen storage capacity, higher plateau pressure, and smaller hysteresis. Similar to Fe addition, the use of Ce is shown by Liu et al. [107] to improve the flatness of plateau of the Ti32Cr46V22 BCC alloy, as a result of the microstructural homogenization during heat-treatment (Ce also increases the hydrogen capacity by lowering the oxygen concentration). The heat-treated Ti32Cr46V22Ce0.4 alloy can release 2.00 and 2.52 wt% H2 at 343 and 298 K, respectively, under 1 atm.

    In general, quaternary alloys compile the advantages of the already optimized properties of ternary V–Ti–Cr alloys, and display an improved cyclic stability without noticeable change of the storage capacity after the addition of various atoms such as Fe [108], Nb [109] or even C [110]. However, even more complex systems exist, as shown by Yang et al. [111], who conducted partial substitution studies on V–Ti–Cr–Fe alloys using Co and Zr for improving the storage and cyclic properties. They found out that the hydrogen absorption-desorption capacities of the (VFe)60(TiCrCo)40-xZrx alloys decrease with increasing Zr content. The maximum desorption capacity reaches 2.10 wt% when x = 0, against 1.88 wt% when x = 2. This could be ascribed to the decrease of the volume fraction of the BCC phase while the other phases increase with the Zr content. At the same time, the rate of cyclic degradation decreases with higher Zr content, from 10.9% after 10 cycles (for x = 0) down to 4.5% (when x = 2). Moreover, as the Zr content increases, the hydriding incubation period shortens from 120 s for x = 0 down to 4 s for x = 2. Additionally, more than 90% of the maximum hydrogen absorption capacity is achieved in 400 s when x = 0, while only about 150 s when x = 2. Figure 8 shows the desorption behaviour of some representative solid solution alloys described in this section.

    The entry is from 10.3390/hydrogen1010004


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