Electrochemical hydrogen storage in metal hydrides (MH) is primarily applied in aqueous NiMH secondary batteries. The alternative MH electrode is aimed to replace the Cd-electrode which is widely used in Ni-Cd batteries and is highly toxic. Moreover, MH electrodes reveal significantly improved storage capacities and (dis)charge kinetics.

A schematic representation of a complete NiMH battery, containing a hydride-forming electrode and a nickel electrode, is shown in Figure 3. Both the electrically insulating separator and porous electrodes are impregnated with a strong alkaline solution (concentrated KOH) that provides the ionic conductivity between the two electrodes. The basic electrochemical reactions, occurring at both electrodes during charging and discharging can be represented by equations (11) and (12) respectively.

(11)

(12)

During charging, divalent Ni(II) is oxidized to the trivalent Ni(III) state and water is reduced to hydrogen atoms which are first adsorbed on the metal electrode, which are then subsequently absorbed by the hydride-forming compound. The reverse reactions take place during discharging. The net effect of this reaction sequence (Equations 11 and 12) is that electrical charge is chemically stored in the electrode in the form of hydrogen. The reaction takes place without any water consumption preventing drying out of the battery unlike Ni-Cd batteries. The hydrogen storage alloy currently used in NiMH batteries is a Mischmetal-based LaNi₅ compound. Researchers are still aiming to improve the properties of these materials mainly by focusing on the (dis)charge kinetics, storage capacity and corrosion resistance during over(dis)charging. Currently, the target sector of the NiMH battery technology is mainly mobile applications specifically Hybrid Electrical Vehicles (HEV) in the automotive industry [12].

Conceptually, ‘electrochromism’ is the reversible optical change of specific compounds in response to a change in the oxidation state of the involved electro-active species. The terminology was initially introduced by Platt in 1961 [13] for the color change of organic dyes under the influence of a strong electric field. The first extensive study was done by Deb in 1969 [14] in the field of inorganic metal oxide systems (WO₃). Numerous researchers both in industry and academically focused on different classes of electrochromic materials. The introduction of optically active metal-hydride thin films was reported by Huiberts et al. in 1996 and was based on yttrium- and lanthanum hydrides [15]. Metal-hydrides expanded the classes of electrochromic materials and were listed among advanced optically switching inorganic compounds. Unlike most of the electrochromic color changing materials, metal hydrides even showed all three optical states (reflecting, absorbing and transmitting states) during the reversible (de)hydrogenation processes without any color change [16]. In this section, the development of optically active, hydride-forming materials are highlighted.

In the extensive work of Griessen et al., metal-hydride switchable mirrors can be classified into three generations: rare earth (RE) metal hydrides, magnesium – rare earth (MgRE) metal hydrides and magnesium–transition metal (MgTM) hydrides [16]. The leading work of Huiberts et al. focuses on the hydrogenation of rare earth yttrium thin films via the gas phase and in situ optical characterization of the transmission of the sample. The pressure-composition isotherm of a 300 nm Y film capped with a top layer of 15 nm Pd, indicated by black squares, is shown in Figure 4 together with the intensity of transmitted light (red circles). The optical transmission has been measured at λ = 635 nm (ħω = 1,96 eV). Both the as-prepared metallic yttrium sample (α phase) and the di-hydrided YH₂ state (β phase) show almost no transmission whereas the tri-hydrided state at x = 3 (γ phase) has become optically transparent. Optical switching becomes saturated at composition YH_2,86 [16].

The first electrochemical hydrogenation experiments with combined optical analyses of yttrium film proved that optical switching can also be induced electrochemically in alkaline protective media, as is shown in Figure 6 [17]. Here a constant (de)hydrogenation current (1 mA) is applied to a 500 nm thick Y electrode that is capped with a 5 nm thick Pd-catalyst layer. Electrochemical (de)hydrogenation is performed in 6 M KOH electrolyte with Pt counter electrode and Hg/HgO reference electrode. Optical switching is reported to take place in the hydrogen concentration range of 2.7 < x < 2.8 for YH_x. Electrochemical switching could be accomplished in the order of seconds during charging and minutes during discharging, revealing the slower dehydrogenation kinetics.

The second generation Mg-RE switchable mirrors, Gd-Mg alloys with different Gd to Mg ratios were optically examined during gas phase hydrogenation [18]. The effect of alloy composition for Gd_1-xMg_x binary alloys (0 ≤ x ≤ 0.9) on the transmission is reported in Figure 6 for a wide wavelength spectrum in the range of 200 to 1000 nm.

All samples in Figure 6 were 200 nm films, capped with 10 nm Pd and hydrogenated under a constant pressure of 5 bar of hydrogen at room temperature. In this work, magnesium was intended to be used as alloying element for the rare earth element Gd. Moreover, the color neutrality and transparency for gadolinium-magnesium hydride were remarkable and of significant importance for optical applications. In Figure 7, all three possible optical states can be seen during the transition from metallic GdMg to semiconducting GdMgH₃.

At the beginning of the new century, magnesium became the material of renewed interest not only from an optical point of view but also from an energy storage aspect. Various transition metals were investigated in order to control the kinetics of the (de)hydrogenation reaction. Niessen et al. showed that meta-stable alloys can be prepared in thin film form by depositing Mg with Ti, V or Cr, though they are immiscible elements from a thermodynamic point of view, [20]. The hydrogen content of these materials also strongly affects the optical properties, which can be used in hydrogen sensors, smart solar collectors and switchable mirrors. The pioneering work of the third generation of switchable mirrors, Mg-transition metals, is highlighted in Ref. [17] where the reversible optical transition of Mg-rich Nb binary alloys was carried out by Richardson et al. in 2001 [21].

Figure 8 shows the electrochemical (de)hydrogenation of a 50 nm thick Mg-Ni film covered with a 10 nm Pd top layer in 8 M KOH. Nickel is an interesting alloying element, which is thermodynamically expected to destabilize Mg. Optical switching of this material has also been confirmed in the gas phase under a stream of 4% hydrogen in argon or helium. Visual coloration has been reported as pale yellow in the Mg-rich region and deep red in the Ni-rich region in the hydride, transparent, state.

In order to understand the main concepts behind the different colorations, the optical states and visual responses during the metal-to-metal hydride transition, the energy band gap between the valence band and conduction band of the individual metal-hydride must be considered.

For instance, MgH₂ is a colorless insulator with a band-gap of E_g = 5.16 eV. The optical band-gap of Mg-Ni alloy (Mg₂NiH₄), on the other hand, varies between 2.5 and 2.8 eV, depending on the presence of Mg rich areas which has been attributed to MgH₂ segregation. Since the band-gap energy of these materials seems to be an important and determining factor, an overview of the band gaps is shown in Figure 9 and was prepared by Karazhanov et al. [22].According to the band-gap (E_g) theory, all incident light is transmitted when E_g of the material is greater than the energy of the visible light which ranges from 390 nm violet (3.2 eV) to 760 nm red (1.6 eV). When E_g value is in the range of the energy of the visible light (1.6 < E_g < 3.2 eV), color change is expected, depending on the absorbed and emitted light wavelengths (gradient region). Finally, when E_g is less than the energy of the visible light, all the white light energy is absorbed and emitted, which results in a zero transparency (dark region). Figure 9 illustrates the basic idea about the expected optical response bands for magnesium containing metal-hydride compounds.

The specific fluorite-type of magnesium-based compounds alloyed with scandium has been studied electrochemically with compositions ranging from 50 to 100 at.%, both in thin film and bulk forms. The thin film electrodes were prepared with a thickness of 200 nm on quartz substrates capped with a 10 nm Pd layer whereas the bulk powders were prepared through arc melting in a molybdenum crucible Mg_xSc_1-xPd_0.24. The XRD results of the thin films confirm the formation of HCP solid solution irrespective of the composition. This is far from the equilibrium state reported in the phase diagram, where a Cs-Cl type solid solution is found between x = 50 − 65 at.% Mg and single phase α-Mg is only formed beyond 80 at.% of Mg. The XRD results of bulk powders, on the other hand, confirm the equilibrium phase diagram.

However, the electrochemical properties of Mg-Sc thin films and bulk materials are very similar. Both thin film and bulk materials show similar trends in storage capacity and the shape of the discharge curves are also very similar for the various compositions (Figure 12). The maximum storage capacity was reached at a composition of 80 at.% Mg and was 1740 mAh/g for thin films, corresponding to 6.5 wt.% of Hydrogen and 1495 mAh/g (5.6 wt.%) for bulk materials. The rate capability of the thin films is, however, much better than that of the bulk materials due to shorter diffusion lengths. [27,28].

It was found that the favorable discharge kinetics was retained up to 80 at.% magnesium. The improvement of the (de)hydriding kinetics compared to pure magnesium has been attributed to the more open fluorite structure of scandium-hydride instead of the more compact rutile structure of the conventional magnesium-hydrides. The plateau pressure is not a function of composition and almost equal to the equilibrium pressure of pure MgH₂. Diffraction data obtained from the bulk system elucidates that the fully charged hydride is a single-phase fcc-structured hydride and no phase segregation into MgH₂ and ScH₂ has taken place [28].

Recently, the crystal structure of a palladium-containing Mg₆₅Sc₃₅ compound has been studied upon hydrogenation by neutron diffraction [29]. The metallic phase adopts a pseudo-CsCl-type structure. When hydrogenated, the structure changes from pseudo-CsCl to fcc involving an elongation along the c-axis and shrinkage along the a-axis (Figure 13). Upon deuteration, a two-phase domain is subsequently observed between 0.85 and 1.55 [D/M] with two hydrides in equilibrium. Above this value, a pressure increase becomes apparent in the isotherm, leading to fully occupied tetrahedral sites and partially filled octahedral sites [30].

It has been argued that interstitial hydrogen is highly mobile, due to the large and partially occupied octahedral sites that that favorably influence the transportation properties. Using proton nuclear magnetic resonance (NMR), the crystallography has been proven to be responsible for the beneficial transportation properties of the fluorite-structured magnesium compounds. The frequency hopping rate for the fluorite-structured compounds at the relevant temperatures was found to be many orders of magnitude higher than that for the rutile-structured material. At any given temperature, the hopping rate of Mg_(1-y)Sc_yH_x is faster than in ScH₂ and much faster than in the ionic MgH₂, but still slower than in LaNi₅H_6.8 (Figure 14) [31].

First principle calculations were done on Mg_(1-y)Sc_y alloys by Pauw et al. [32]. Figure 15 shows that the decrease in the formation enthalpy with increasing Sc-content is more pronounced for the fluorite structure than for the rutile structure. The crossing point of the two curves is located around 20% Sc (y = 0.20 in Figure 16), which is the exact composition where the electrochemical results showed the maximum rate capability and hence electrochemical storage capacity. The first principle calculations are in good agreement with these experimental findings.

Since scandium is rather expensive, attempts to find cheaper substitutes were recently undertaken. Titanium dihydrides are also promising candidates because of having the same fluorite structure as that of scandium dihydride. Experiments performed with Mg_0.80Ti_0.20 thin film electrodes indeed showed a similar drastic improvement of the discharge (dehydrogenation) kinetics compared to pure magnesium hydride (MgH₂). A more elaborate study on the Mg–Ti system (Figure 16) revealed that the composition dependence of the hydrogen storage capacity was quite similar to that of the Mg-Sc system, also showing a maximum capacity around 80 at.% magnesium [33].

Pauw et al. [32] also investigated the formation enthalpies of metal hydrides in both the rutile and fluorite structure. The calculations showed that the addition of Ti induces a remarkable change in the crystal structure compared to pure MgH₂. The formation enthalpy of rutile Mg_0.50Ti_0.50H₂ is −0.421 eV/H₂, which is 22.7 kJ mol⁻¹ less negative than that of pure MgH₂. The rutile structure with composition Mg_0.25Ti_0.75H₂ is also less stable than MgH₂, which means that Ti destabilizes the rutile structure over almost the entire compositional range. The formation enthalpy of the fluorite structure shows a tendency towards more negative values over the entire compositional range, the same as for the Mg-Sc system. The fluorite structure becomes the most stable phase around y = 0.20, which is also very similar to Mg-Sc and these theoretical results are therefore in good agreement with the experimental results. The crystal structure of the resulting hydrides seems to play a crucial role in determining the (de)hydrogenation kinetics for this interesting new class of hydrogen storage materials [32].

In situ electrochemical XRD was carried out by Vermeulen et al. to investigate the crystallographic phase transformation during the electrochemical hydrogenation of the MgTi alloys [34]. While hydrogenating Mg_0.90Ti_0.10 thin films electrochemically, the hexagonal reflections decrease with increasing hydrogen content (Figure 17 a). Simultaneously, the bct reflections evolve, indicating the formation of a body centered tetragonal phase which is the same as for pure MgH₂. On the other hand, hydrogenation of Mg_0.70Ti_0.30 showed an increase in the fcc phase (Figure 17 b). Substantial electrochemical storage capacities of more than 1500 mAh g^-1 have been reported for the Mg-Ti system at room temperature. This corresponds to a reversible hydrogen content of about 6 wt.%. In addition, it has been shown that these meta-stable materials cannot only be produced by thin film evaporation and sputtering techniques but also by mechanical alloying.

Several attempts have been made by researchers to produce Mg-Ti alloys by mechanical alloying. Rousselot et al. investigated the phase transformation during electrochemical hydrogenation [35]. The meta-stable hcp Mg_0.50Ti_0.50 alloy was formed through mechanical alloying. XRD was performed on the electrochemically hydrogenated alloy. Upon cycling, the intensity of the hcp phase decreases and a new peak arises, which corresponds to hydrided fcc phase. The fcc phase is irreversibly formed and does not disappear after dehydrogenation, instead a slight decrease in the lattice constant was observed upon discharge. Kalisvaart et al. reported the formation of two fcc phases of the Mg-Ti alloy when the milling process was carried out with stearic acid as process control agent [36,37]. Among the two fcc phases only one was concluded to be electrochemically active. The storage capacity of this bulk material was reported to be 500 mAh/g.

Kyoi et al. synthesized Mg-Ti-H compound at 600 °C in a high-pressure anvil cell above 8 GPa by mixing MgH₂ and TiH_1.9 [38]. The metal atom structure was found to be close to the Ca₇Ge type fcc structure. This is in line with the findings of Vermeulen et al [34]. It could be concluded from in situ XRD studies that the hydrogen atoms occupy the tetrahedral sites.

The Mg-Ti alloy has a positive heat of mixing and must therefore be considered as an immiscible system [39]. The alloy can therefore, in principle, not be synthesized under equilibrium conditions and thus the only way to produce these meta-stable materials is by making use of non-equilibrium preparation methods. The stability of the meta-stable alloys is still questionable at high temperature and pressure conditions. Liang et al. reported an extended solubility of Ti in Mg, during mechanical alloying of Mg and Ti [40]. In the case of Mg with 20 at.% Ti, about 12.5 at.% Ti is dissolved in the Mg lattice when the mechanical alloying process reaches a stable state, while the rest remains as fine particles of size 50–150 nm in diameter. Dissolution of Ti in the Mg lattice causes a decrease in unit cell volume. The phase segregation begins at high temperature and the supersaturated Mg-Ti alloys start to decompose at 200 °C. Hydrogenation is an exothermic reaction and enhances the decomposition process.

Srinivasan et al. investigated the structure of bulk magnesium-titanium deuteride Mg_0.65Ti_0.35D_xprepared via mechanical alloying and deuterium gas-phase loading at 75 bars and 175 °C [41]. The structural analysis was made with XRD, NMR and neutron diffraction. Deuterium loading causes the formation of both interdispersed rutile MgD₂ – TiD_y nano-domains and a separate fluorite TiD_z phase. Exchange NMR technique indicates complete deuterium exchange between the MgD₂ and TiD_y phase within one second. Gas phase hydrogenation of the Mg-Ti alloys at higher temperatures increases the risk of phase segregation.

Even though the kinetics of the fluorite-structured compounds is faster compared to the MgH₂, the thermodynamic properties are not improved by the addition of these transition metals. For instance, the dehydrogenation enthalpy of Mg_xSc_(1-x)H₂ and Mg_xTi_(1-x)H₂ is higher compared to MgH₂. Thus, high temperature is required to dehydrogenate these compounds in case of gas phase loading. However, only Mg-Sc alloys are stable with respect to decomposition into its elemental metallic form. Since scandium is expensive, the next optional metal Ti is preferred considering the storage capacity compared to Cr and V. Due to the high stability of the Mg-Ti hydrides and low thermal stability of Mg-Ti alloy, the addition of a third element into the binary alloys is required to stabilize the host materials and, consequently, to destabilize the alloy hydrides. Careful selection of the third alloying element could assist in improving the thermodynamic properties while maintaining the storage capacities.

Increasing the plateau pressure of the metal-hydrogen system can be accomplished by destabilization of the metal hydride. Miedema proposed a mathematical model to calculate the effect on the thermodynamic stability of an additional element(s) to the metallic compound [42,43]. The Miedema model was used by Vermeulen et al. to screen elements that stabilize the metallic Mg-Ti system [44]. Al and Si were selected to destabilize the Mg-Ti alloy, due to their advantageous lightweight mass to maintain the high gravimetric capacity of the overall Mg-Ti alloy. The isothermal curves of the Mg_0.69Ti_0.21Al_0.10 alloy show two plateau regions, a plateau up to 4 wt% of Hydrogen at 10⁻⁶ bars and a second sloping plateau which lies at higher pressures (Figure 18).

Addition of higher Al content in the Mg-Ti lattice leads to a decrease in the reversible storage capacity. If only 10 at.% is incorporated in Mg-Ti alloy with a high Mg-content, it does not seem to significantly affect the reversible capacity. The declining capacity at high Al content could be due to kinetic limitations imposed by the material. Additionally, Si also clearly affects the isotherms compared to Mg-Ti alloys. The length of the first plateau decreases, when increases steeply and soon reaches a sloping plateau. Since Si is heavier than Al, the storage capacity is lower than that of the Al-containing ternary alloys. The occurrence of the second plateau was also observed in other systems, such as Mg₂CuAl_0.375, LaCo₆, Ti-Fe and NdCo₅. Interestingly, the Mg-Ti-Al-H system shows a plateau near atmospheric pressure and temperature.