Keywords: magnesium hydride, hydrogen storage, nanocomposites, metal hydrides. II III hydrure de magnésium, stockage de l'hydrogène, nanocomposites, hydrures métalliques

Mg/transition-metal nanomaterials for efficient hydrogen storage

Magnesium metal is a prominent element for solid-state hydrogen storage due to its large abundance in earth's crust and its high weight and volumetric hydrogen uptakes.

However, hydrogen sorption suffers from sluggish kinetics and the formed hydride is too stable for applications working under ambient conditions. The former issue can be solved by developing composites combining two hydrides, MgH2 and TiH2 at the nanoscale. These materials are synthesized by mechanical milling under reactive atmosphere. By this technique, the formation of nanocomposites and their hydrogenation can be obtained in a single-step.

Moreover, these materials can be produced at large scale for application purposes.

The work focused on three topics: i) the optimization of the TiH2 content in the (1y)MgH2+yTiH2 system. This was accomplished by optimizing the titanium content (0.0125≤y≤0.3 mole), while keeping good kinetics, hydrogen reversibility and cycle-life. The data show that y=0.025 is the best compromise to fulfill the most practical properties; ii) the extension to other transition metals for the system 0.95MgH2+0.05TMHx (TM: Sc, Y, Ti, Zr, V and Nb), evaluating the contribution of each additive to kinetics, hydrogen reversibility and cycle-life; iii) the conception of an automatic cycling device able to carry out hundreds of sorption cycles whit the aim of measuring the cycle-life of metal hydrides.

The work was done using manifold experimental methods. For synthesis, reactive ball milling under hydrogen atmosphere was primarily used. The crystal structure and the chemical composition of nanomaterials was determined from X-ray diffraction (XRD) analysis. Particle size and morphology were obtained by Scanning Electron Microscopy / Energy Dispersive X-Ray Spectroscopy (SEM/EDS). Thermodynamic, kinetic and cycling properties toward hydrogen sorption were determined by the Sieverts method.

Je souhaite également exprimer ma plus grande gratitude à tous les membres de CMTR (Chimie Métallurgie des Terres Rares) pour m'avoir accueilli aimablement dans leur groupe et pour leur soutien. Il ne serait pas juste de mentionner leurs noms, car j'ai reçu de l'aide et le soutien de tous, directement ou indirectement. Je suis également reconnaissant aux gens de l'ICMPE au cours de ces trois années, j'ai rencontré des gens formidables ici.

Gracias al departamento de física de materiales de la Universidad Autónoma de Madrid por su amigable acogida y por la amable ayuda que recibí durante una semana y después por correo electrónico.

I would like to thanks to all my Friends from all the nationalities that I have done during this time for their support and their friendship.

A really special thanks to my family, for always trusting and supporting me in every task of my life, allowing me to achieve great goals such as this dissertation.

I want to express my sincere thanks to all the people who made it possible for me to successfully complete my graduate study.

Finally, I would like to acknowledge the financial support from the Consejo Nacional de Ciencia y Tecnología, y la Secretaria de Energía.

Different methods and phenomena are commonly used for hydrogen storage, such as, i) highpressure gas cylinders (up to 80 MPa) [3], ii) liquid hydrogen in cryogenic tanks (at -253 °C ) [4], iii) adsorbed hydrogen on solid materials with very large specific surface area (at T <-173 °C) [5], iv) absorbed hydrogen at interstitial sites in a host metallic structure (some metal hydrides working at ambient pressure and temperature), v) chemically bonded in covalent and ionic compounds (at ambient pressure) [6]. FigureⅠshows the summary of different technologies mentioned before. It can be seen that metal and complex hydrides (blue area) have volumetric capacities far superior to that of liquid hydrogen.

Figure Ⅰ. Hydrogen storage capacities for different storage media [3] The most common materials for reversible hydrogen storage is metal hydrides MHx. Many metals and alloys can reversibly react with H2 to form hydrides through the general reaction:

2𝑀 + 𝑥𝐻 2 ⇋ 2𝑀𝐻 𝑥
The reverse reaction (i.e. metal M reformation) can be accomplished by either increasing the temperature or reducing the hydrogen pressure. The stored hydrogen can be released by VI reducing the hydrogen pressure in a closed system to a level below the plateau pressure (𝑃 𝑝 ) of the M/MHx equilibria. A hydride stable under a certain temperature and hydrogen pressure will decompose when the temperature is increased to a level where 𝑃 𝑝 is higher than the system pressure. The temperature needed for a hydride to release hydrogen at 1 bar, T (Pp = 1 bar), is a critical parameter as it indicates the minimum working temperature for a hydrogen system based on that material. Unfortunately, the known metals and alloys with T (Pp =1 bar) lower than 80 °C are based on high weight transition metals and rare earths (e.g. LaNi5 and FeTi) and therefore have low gravimetric hydrogen storage capacities (< 3 wt.%).

Because of the low gravimetric capacity for these interstitial metal hydrides (Fig I), the interest of materials with higher capacity has been increasing. This is the case of magnesium hydride which has a storage capacity of 7.6 wt. %. Even though MgH2 has a high T(Pp = 1bar) = 280 °C and sluggish kinetics, research on this hydride has deserved a great interest because of its high abundance and low cost. Many efforts have been done to improve the slow kinetics of MgH2.

One of the most common approaches to improve the hydrogen sorption kinetics of MgH2 is by nanostructuration through mechanical milling methods. Bulk coarse magnesium hydride, with a particle size of ten of micrometers need at least some hours to fully decompose at 350 °C. By milling, the crystal and particle sizes of magnesium hydride reduce at the nanoscale and the surface roughness and area increase. This leads to a reduction in desorption time to some minutes at 350 °C [7].

By combining the last approach with the use of selected additives, further improvement of sorption kinetics can be attained. Hydrogen molecules can easily dissociate at the surface of some catalytic additives. Then, the nascent hydrogen atoms can permeate through them if they exhibit high diffusion coefficients for hydrogen and facilitate hydrogen transport toward the magnesium metal. It is important to find good candidates for magnesium hydride systems, which allows enhancing the kinetics of both absorption and desorption reactions.

Previous studies have demonstrated that Ti is an attractive additive to improve Mg hydrogenation properties [8,9]. It is a lightweight transition metal TM and at relatively low cost. Moreover, the fcc fluorite-type structure of titanium hydride is characterized by a fast diffusion coefficient. The catalytic properties of the TiH2 phase on reversible hydrogenation of MgH2/Mg have already been previously reported by our group at ICMPE/CNRS [10]. In this previous work, the H-cycling properties of a 70MgH2-30TiH2 nanocomposite were characterized and compared to those of Ti-free MgH2 synthesized by the same mechanochemical method (reactive ball milling, RBM). The Ti-containing nanocomposite exhibited outstanding sorption kinetics with absorption/desorption reaction times below 100 seconds at 300 °C. Furthermore, such fast kinetics remain stable over more than 30 cycles. In contrast, pure MgH2 obtained by the same RBM method exhibited sorption reaction times over 60 min and slowed down on H-sorption cycling. This thesis targets three objectives. First one, to find out and understand whether there is an optimum TiH2 amount in the MgH2-TiH2 system. The second aim is to screen other transition metals additives forming hydrides; namely TM = Sc, Y, Zr, V, and Nb. Finally, this thesis aims to design and built-up an experimental system for studying the hydrogen sorption properties of MgH2-TMHx systems over hundreds of cycles.

Chapter 1. Introduction

The European Union has set a target by 2050 of renovating the energetic sector into lowcarbon one, with the intention of reducing greenhouse gas emissions, upholding a strong growth for a global position as well as increasing energy security [1].

The European Commission presented in "EU 2020 Strategy" [2] puts forward three reinforcing priorities: -Smart growth: developing an economy based on knowledge and innovation.

-Sustainable growth: promoting a more resource efficient, greener and more competitive economy.

-Inclusive growth: fostering a high-employment economy delivering social and territorial cohesion.

Both hydrogen and fuel cell technologies are key factors for each point presented above. An innovative economy would involve a new and broad hydrogen market. Meanwhile, a greener economy should be linked with a low-carbon emission society. Finally, with an innovative and green economy, a territorial cohesion work accomplishes the inclusive growth.

According to International Energy Agency (IEA) report scenario by 2050 [3], hydrogen will likely expand the market over the coming decades if three hydrogen costs fall: production, distribution, and end-use (hydrogen storage). Also, effective policies shall be put in place to increase energy efficiency and improve energy security.

Hydrogen as an Energetic carrier

An energy carrier is a substance for delivering mechanical work or heat transfer. Examples of energy carriers include: solid, liquid or gaseous fuels (e.g., biomass, coal, oil, natural gas, hydrogen)

As an energy carrier, hydrogen provides a unique opportunity to contribute to all major European policy objectives in the transport and energy sectors:

• Hydrogen is a nontoxic gas with a high mass-energy density (between 120 and 142 MJ/kg) and is one of the most abundant elements in the biosphere, though scarce in pure form. • It can be produced from water, biomass, biogas, natural gas or from any other fossil fuel, • Used in conjunction with fuel cells, hydrogen provides an efficient decentralized solution for combined heat and power generation, both at an industrial and a domestic level; • Hydrogen is an efficient way to store electricity, especially the renewable electricity generated by intermittent sources (solar, wind, hydroelectricity) • Hydrogen can be used as an alternative fuel for clean mobility using fuel cell electric vehicles, including passenger cars, buses, light-duty and material handling vehicles • Hydrogen offers a significant reduction of greenhouse gases emissions -hydrogen oxidation only produces water at the point of use -and a reduction in air pollution and noise.

Hydrogen can be produced using renewable sources working for clean energy systems. The main interests for hydrogen as an energy vector is due to the flexibility related to both its production and more important its use as a fuel. Since some years, the use of hydrogen as a fuel has been stronger, even with the severe competition with fossil fuels. To be more competitive, hydrogen as to deal with some drawbacks, such as effective and safe store reducing the cost processes.

Fuel cells can generate electricity from various feedstocks and can be used not only in connection with an electricity grid but also as a stand-alone power generator. Fuel cells have broader applications than any other power source currently available. The power produced can be used in many portable, stationary and transport applications and the heat, a byproduct, can also be used for heating and cooling [4].

Hydrogen Generalities

Hydrogen (chemical symbol H) is the lightest element of the periodic table (atomic weight 1.00794 a.m.u), with atomic number Z=1. The origin of its name comes from the combination of the Greek words 'hydor' meaning 'water' and 'geinomai' meaning 'to bring forth', i.e. the element that brings forth water. Hydrogen is the ninth most abundant element on Earth's crust, as well as, the second most abundant element in Earth's sea after oxygen. Under normal conditions of temperature and pressure (at a temperature of 20°C and pressure of 1 atm), it is a colorless, odorless, tasteless, and flammable gas. Hydrogen was discovered by Henry Cavendish in England in 1766 and named by Lavoisier [5,6].

Hydrogen storing methods

The attractiveness of hydrogen is focused on its electron which is accompanied by only one proton (for charge neutrality), i.e. hydrogen has the highest ratio number of valence electrons to protons (and neutrons) of all the elements in the periodic table. [7].

Hydrogen storage technique basically implies the huge volume reduction of the hydrogen gas; 1 kg of hydrogen at ambient temperature and atmospheric pressure takes a volume of 11 m 3 .

To increase the hydrogen density in a storage system, work must either be applied to compress hydrogen, or the temperature should be decreased below its critical temperature.

Different methods and phenomena are commonly used for hydrogen storage: i) high-pressure gas cylinders (up to 80 MPa), ii) liquid hydrogen in cryogenic tanks (at -253 °C), iii)) adsorbed hydrogen on solid materials with a large specific surface area (typically at T <-173 °C), iv) absorbed on interstitial sites in a host metal (usually near ambient pressure and temperature), v) chemically bonded in covalent and ionic compounds (at ambient pressure) [8].

High-pressure gas storage

The most common hydrogen storage solution currently uses pressurized cylinders to store hydrogen. Cylinders lighter than metal are used, typically working at 350 bar, though it exists other thanks working at 700 bar. The construction of cylinder consists in a liner (made from aluminum, steel or polymer) around which carbon fibers are wound and sealed in a polymer resin. The gravimetric storage capacity of the 700 bar tanks is 4.5 %, however, the volumetric capacity is only 0.025 kgH2/l. The energy required for 700 bar compressions is 15 % of lower heating value (LHV) of the hydrogen stored in the vessel.

On-board, compressed gas storage is currently the best compromise. For applications where space is not an issue, for example, stationary storage at a fueling station, compressed gas is a relative cheap solution. For portable applications, compressed gas is too bulky, especially in cars where a significant amount of space must be sacrificed because of the cylindric shape of the tank [9].

Liquid hydrogen

Cryogenic storage of hydrogen is another mature technology. It works at very low temperature (-253 °C) at ambient pressure. Due to the low liquid-vapor critical temperature of hydrogen (-240 °C), liquid hydrogen can only be stored in open systems because there is no liquid phase existing above the critical temperature. The volumetric density of liquid hydrogen is 70.8 kg•m 3 , comparable to that of solid hydrogen (70.6 kg•m 3 ). The challenges of liquid hydrogen storage are the energy-efficient liquefaction process and the thermal insulation of the cryogenic storage vessel to reduce the boil-off of hydrogen. The simplest liquefaction cycle is the Joule-Thompson cycle (Linde cycle). The gas is first compressed and then cooled in a heat exchanger, before it passes through a throttle valve where it undergoes an isenthalpic Joule-Thomson expansion, producing some liquid. The cooled gas is separated from the liquid and returned to the compressor via the heat exchanger [10].

The boil-off rate of hydrogen from a liquid hydrogen storage vessel due to heat leaks is a function of the size, shape and thermal insulation of the vessel. Large-sized spherical containers are expensive because of their manufacturing difficulty. Since boil-off losses due to heat leaks are proportional to the surface to volume ratio, the evaporation rate diminishes drastically as the storage tank size increases. The relatively large amount of energy needed for the liquefaction and the continuous boil-off of hydrogen limit the possible applications for liquid hydrogen storage systems to utilizations where the cost of hydrogen is not an important issue and the hydrogen is consumed in a relatively short time, e.g. air and space applications.

Physisorption of hydrogen

Physical adsorption is the process where hydrogen is stored in molecular form (that means, without dissociating H2 molecule) on the surface of a solid (carbon materials (3.5 wt. %H2), metal-organic frameworks (4 %H2), zeolites (1.5 wt. %H2) [9]. The molecular adsorption of H2 is done through van der Waals forces, between the gas molecules and the atoms at the solid surface.

Adsorption is an exothermic process, where, heat is released during the adsorption of gas molecules on the surface. Owing to the low polarizability of H2 molecule, the interaction is very weak (with an enthalpy of adsorption of between 40 and 10 KJ/mol) These weak interactions mean low temperature is needed to obtain significant amounts of adsorbed gas, so that the hydrogen molecules do not have too much thermal energy which easily overcomes the van der Waals interaction.

In general, porous materials show excellent cyclability as the material itself does not undergo any significant change through the adsorption/desorption process. A cryo-tank is needed as for liquid hydrogen and there are similar issues with boil-off. However, there is a significant energy saving in having to keep the store at -200 °C as opposed to -253 °C. These combined reasons make porous hydrogen storage materials less attractive for automobile on-board storage [11].

Metal hydrides

Metal hydrides are promising materials for both stationary and mobile hydrogen storage applications. These materials are formed by one or several metals (alloys and intermetallic compounds) able to react with hydrogen gas. Hydrogen reacts with many transition metals and their alloys to form those hydrides. The electropositive transition metals (i.e. scandium, yttrium, the lanthanides, the actinides, and the members of the titanium and vanadium groups) form stable hydrides. The binary hydrides of the transition metals are predominantly metallic in character and are usually referred to as metal hydrides. The lattice structure is that of a typical metal with atoms of hydrogen in interstitial sites; for this reason, they are also called interstitial hydrides [12]. Ternary systems ABxHn have an important interest due to the variety of elements which allows tailoring of the hydride properties (see Table 1.

1).

The A element is usually either an alkaline-earth metal (elements group two), a transition or a rare earth metal which tends to form a stable hydride. The B element is often a last transition metal and forms only unstable hydrides. Some well-defined ratios of B to A in the intermetallic compound x = 0.5, 1, 2, 3 and 5 have been found to form hydrides with a hydrogen to metal ratio of up to 2. Metal hydrides are successfully effective for large amounts of hydrogen storing (MgH2 7.6 wt. %), in both a safe and compact way. Up to now, all the reversible hydrides working around ambient temperature and atmospheric pressure consist of transition metals. Therefore, the gravimetric hydrogen density is limited to 3 mass%. It remains a challenge to find out novel lightweight metal hydrides that react reversibly with hydrogen near normal thermodynamic conditions [13].

Complex hydrides

Complex metal hydrides are normally formed by elements of the first and second groups of the periodic table. They are especially interesting because of their light weight and the number of hydrogen atoms per metal atom, which is in a few cases 2. The main difference between the complex hydrides and the metallic hydrides is that the former are ionically and/or covalently bounded. The hydrogen in the complex hydrides is often located in the corners of a tetrahedron with boron or aluminum in the center. The negative charge of the complex anion, [BH4] -and [AlH4] -, is compensated by a cation, e.g. Li + or Na + .

Table 1. 1 The most important hydride-forming intermetallic compounds families including both the prototype and the structure. Mg2Ni, Ti2Ni Mg2NiH4 Cubic, Ti2Ni-type

Intermetallic compound

Chemical reaction with water

Hydrogen can be generated from metals and chemical compounds reacting with water. The most common experiment -shown in many chemistry classes -where a piece of sodium floating on water produces hydrogen, demonstrates such a process. The sodium is transformed into sodium hydroxide in this reaction. The reaction is not directly reversible, but the sodium hydroxide could later be removed and reduced in a solar furnace back to metallic sodium. Two sodium atoms react with two water molecules and produce one hydrogen molecule. The hydrogen molecule produces again a water molecule in the combustion, which can be recycled to generate more hydrogen gas. However, the second water molecule necessary for the oxidation of the two sodium atoms must be added. Therefore, sodium has a gravimetric hydrogen density of 3 mass%. The same process carried out with lithium leads to a gravimetric hydrogen density of 6.3 mass%. The major challenge with this storage method is the reversibility and the control of the thermal reduction process in order to produce the metal in a solar furnace [14].

Metal Hydrides overview

As it was shown above metal hydrides can be successfully employed as efficient and safe storage media of hydrogen gas. In these materials, storage sometimes takes place under moderate pressure and temperature conditions. In that respect, metal hydrides have a significant advantage over classical hydrogen storage methods, which suffer from either extremely high pressures or very low temperatures. Storage in metal hydrides is, therefore, one of the key factors, facilitating e.g. hydrogen-driven fuel cells and Nickel-Metal Hydride batteries [15]. To unravel the hydrogen storage process in metals, a more detailed understanding is, however, essential.

Thermodynamic of metal hydrides

Pressure-composition isotherms are widely used for determining the thermodynamic aspects of hydride formation between hydrogen gas and a solid, here a metal. The most important part of these isotherms is the plateau zone. The amount of hydrogen storage is directly associated with the plateau length. The host metal firstly dissolves hydrogen as a solid solution (α-phase) up to a solubility limit . At that point, increasing the hydrogen pressure will induce the nucleation and growth of a novel hydride phase (β-phase) at the expense of the former α-phase (see Fig. 1.1). The two-phase region expands from  to β, the plateau width decreases with temperature and eventually disappears above a critical point Tc. Beyond this temperature the transition from the α to β phases is continuous. The plateau pressure, Pp, is related to the enthalpy (ΔH) and entropy (ΔS), as a function of temperature T by the van't Hoff equation:

𝐼𝑛𝑃 𝑝 = ∆𝐻 𝑅𝑇 - ∆𝑆 𝑅 Eq. 1.1
where, Pp is the plateau pressure, and R the gas constant. From this equation, the slope of the line in 1/T is equal to the formation enthalpy divided by the gas constant (first term), therefore the intercept at the origin is equal to the formation entropy divided by the gas constant (Figure 1.1). ΔH is a measure of the metal-hydrogen bond energy (and therefore of the hydride stability) meanwhile the entropy change corresponds mostly to the change from molecular hydrogen gas to dissolved solid hydrogen. The latter approximately corresponds to the standard entropy of hydrogen and therefore, ΔSf ≈ -130 J•K -1 mol -1 H2 for most of the metal-hydrogen systems. As for the hydride stability, it is characterized by the enthalpy term. To reach a plateau pressure of 1 bar at 300 K, ΔHf should amount to -39.2 kJ mol -1 H2. The term entropy formation for metal hydrides leads to a significant heat evolution ΔQ = T•ΔSf (exothermic reaction) during hydrogen absorption.

The same heat must be provided to the metal hydride to desorb the hydrogen (endothermic reaction). In the case hydrogen desorbs below RT, this heat can be delivered by the environment. However, if the desorption occurs above RT, the necessary heat has to be delivered from an external source, such as the combustion of hydrogen [16].

Hysteresis

Hysteresis is observed for nearly all metal hydrides: the pressure Pf needed for hydride formation is greater than the hydride decomposition pressure Pd (Figure 1.3). In practical applications, hysteresis represents a loss of efficiency. Another manifestation of this phenomenon is that the maximum hydrogen solubility in the solute phase amax is higher for hydride formation a' than for hydride decomposition a''. A similar behavior takes place for the minimum solubility in hydride phase min. One explanation of hysteresis was proposed by Lacher [17] who suggested that it may result from the effect of particle size (i.e. either small nuclei of the dilute phase alfa in the hydride phase or small nuclei of hydride phase) on the equilibrium hydrogen pressures.

Hysteresis is an important feature for practical applications as it impacts the service pressure of a storage tank. In fact, it represents a loss of efficiency because of the irreversible deformation of hydrogen sorption. Hysteresis should be as small as possible, which can be achieved by element substitution in the alloy and heat treatment.

Reversible Capacity

Reversible capacity Crev is defined as the difference between the hydrogen amount absorbed and the desorbed one, which can be considerably less than the maximum capacity (M/H) max (Figure 1.4). A clear example is vanadium. At T= 25 °C and PH 2 = 10 -7 MPa, it forms a monohydride (VH), whereas a dihydride (VH2) is formed above 1 MPa at the same temperature. Under standard conditions (T= 0°C and P= 0.1 MPa), the dehydrogenation of the monohydride is thermodynamic impossible, so the reversible capacity of vanadium at these conditions is between VH≈1 and VH≈2. Hydrogen capacity, CH, can be reported in different units: i) as the atomic ratio (H/M) between hydrogen and metal atoms; ii) in weight percent (wt. %), denoting the weight of hydrogen stored over the weight of the hydride phase and iii) another way to hydrogen capacity is the volumetric capacity (g/l), expressed by the number of hydrogen atoms per unit volume [18].

Kinetic of metal hydrides

For hydrogen storage in solids, sorption kinetics are of paramount importance. The reaction between metal and hydrogen is a heterogeneous phase transformation. Formation of a metal hydride occurs through several steps taking place in series: H2 transport to the surface, H2 dissociation, H chemisorption, surface-bulk migration, H diffusion and finally, nucleation and growth of the hydride phase. Inverse reactions occur on hydride decomposition. For a reaction in series, the slowest step is defined as the rate-determining step.

Hydrogen reaction kinetics for a given material depends on many sensitive parameters such as activation conditions, structural defects, crystallinity and occurrence of catalytic phases. For example, nanostructured materials including secondary catalytic phases can be synthesized by mechanical ball milling. In this way, nanocomposites material with fast kinetics are obtained thanks to short diffusion lengths for hydrogen diffusion and catalyzed surface reactions [19][20][21][22][23][24][25][26]. Both, the sample temperature and the hydrogen pressure should remain constant. No one of these conditions is easily fulfilled in the case of the metal-hydrogen reaction. Metal particles suffer from fragmentation during absorption and sintering during desorption. Sample temperature changes because of the heat of reaction and, except for constant pressure experiments, the absorption and desorption of hydrogen induce a related decrease or increase of pressure. In order to obtain reliable kinetics data, care should be taken to minimize all these effects [27].

The applicability of a hydrogen storage system is highly dependent on the kinetic parameters (rate and time of full desorption and absorption). The identification of the rate controlling mechanism gives the possibility to understand and thus to improve the kinetic properties. As hydride formation is a multi-step process, the reaction rate is dependent on many individual parameters. Since there are too many parameters describing the hydride formation, usually three different approximates are applied to model the measured α(t) functions [28].

The surface-controlled process assumes that the slowest step of the reaction is the chemisorption. In the case of the contracting volume model, the main assumption is that the initial nucleation on the surface is fast compared to the overall growth kinetics and the nucleation zone is thin compared to the particle diameter. The third model, the Johnson-Mehl-Avrami (JMA), applies to cases where the nucleation and growth of the new phase begin randomly in the bulk and at the surface. From adjusting the measured α functions to the model predicted ones, the rate control mechanisms can be determined [29].

Activation

The metal surface is covered with a native layer of oxide that acts as a hydrogen barrier. It must be overcome to start metal hydrogenation. The time that hydrogen lasts to cross the oxide film is called in literature incubation time [30]. Therefore, the first hydrogenation treatment is normally performed at high temperature and pressure to facilitate hydrogen penetration through the oxide layer. Upon hydrogenation, the metal lattice volume increases significantly while on dehydrogenation, the lattice reverts to its original size. This expansioncontraction breaks the metal particles, exposing fresh metal surface, reducing particle size and enhancing reaction kinetics.

Cycling life

For practical applications, it is of great importance to maintain the sorption properties (capacity, kinetics, reversibility, plateau pressure, etc.) of the hydrogen storage system during the whole life of the device. Typically, it can run from a few hundred to several thousand cycles. Cycling life depends on many parameters such as impurities in the hydrogen gas leading to poisoning, agglomeration of particles, and structural relaxation, etc. Nevertheless, magnesium and magnesium-based compounds could sustain a few thousand cycles without drastic changes in hydrogen sorption properties. A special way to improve cycling stability is by element substitution in the alloy [31].

Magnesium

Considering that metal hydrides could become outstanding energy carries for both mobile and stationary applications, the hydrogen storage mass should be reduced. An important promising candidate is magnesium for this mobile mass sensitive application, Mg-based hydrides have drawn much attention due to Mg low density and cost [32]. MgH2 is an outstanding hydrogen store due to its high hydrogen content both specific (7.6 wt. % H) and volumetric (0.109 kgH2/l).

Metallic magnesium has a hexagonal crystal structure (space group: P63/mmc) crystal structure with lattice parameters a= 3.2094 Å and c=5.2108 Å. it forms a stoichiometric dihydride MgH2 which can crystallize in two polymorphic forms. Stable at ambient conditions, -MgH2 has a tetragonal TiO2 rutile-type (space group: P42/mnm), with lattice parameters a =4.517 Å and c=3.0205 Å. This hydride, at mechanical pressure exceeding 390 MPa transforms into a metastable γ-MgH2 modification that crystallizes with an orthorhombic PbO2-type structure [33].

Hydrogen accumulates within the magnesium body in the form of stoichiometric magnesium hydride MgH2 [34]. The absorption process occurs in several steps, generally defined as follows [35]: 1) adsorption of hydrogen molecules on the surface of a magnesium particle, followed by their dissociation into hydrogen atoms H2 → 2H; 2) diffusion of dissolved hydrogen atoms into the Mg -bulk (α-phase); 3) formation of the magnesium hydride MgH2 (β-phase) at the α-β interface Mg + 2H → MgH2. The desorption process occurs in the reverse order, namely [36]:

1) magnesium hydride decomposes at the α-β interface and hydrogen atoms dissolve in the metallic bulk MgH2 → Mg + 2H; 2) the dissolved hydrogen atoms move to the surface of the particle by diffusion transport; 3) the chain of several surface reactions (association of hydrogen atoms on the surface followed by the release of gaseous molecular hydrogen from the surface of powder particle 2H → H2) provides the surface degassing [START_REF] Drozdov | Modelling and evaluation of hydrogen desorption kinetics controlled by surface reaction and bulk diffusion for magnesium hydride[END_REF] Fernandez and Sanchez [27] found that the rate-determining step for this desorption is hydrogen diffusion through the β phase. The diffusion of H in MgH2 is very slow, and the diffusion coefficient of H in magnesium hydride at 300°C-100°C lies between10 -18 and 10 -24 m 2 /s [START_REF] Xu | Advanced materials for clean energy[END_REF].

Thermodynamic properties of the Mg-H system

Magnesium hydride is a stable hydride, with a heat of formation of -74.5 kJmol -1 H2 and an entropy variation of -135 Jmol -1 H2 K -1 . Figure 1.5 shows the variation of the desorption plateau pressure as a function of temperature for magnesium hydride. For desorbing hydrogen under PH 2 = 1 bar (0.1MPa) (lower blue line), magnesium hydride must be heated to at least 277°C. Similarly, magnesium could not absorb hydrogen below a pressure of 10 bar (upper blue line) if the temperature is higher than 370 °C. Therefore, magnesium hydride reversibly stores hydrogen only at ~300 °C for hydrogen pressures compatible with the feeding of fuel cells (typically around 0.1 MPa). For room temperature applications, the challenge is to reduce the operation temperature of this system while keeping as high as possible its hydrogen storage capacity.

Modification of Mg-H thermodynamics

One solution to destabilize magnesium hydride is to alloy magnesium with another element, usually a late transition metals, besides Sc, Y, and La. One classic example is Mg2Ni where the heat formation of the ternary hydride is reduced to -64.5 kJ mol -1 H2. Unfortunately, hydride destabilization is accompanied by a severe reduction of mass capacity to 3.6 wt.%. Alloying with other elements has been studied, as pointed out by Dornheim et al. [START_REF] Dornheim | Hydrogen storage in magnesium-based hydrides and hydride composites[END_REF]. Several magnesium-based alloys have been investigated for reversible hydrogen storage, and the alloying elements include rare earth metals [START_REF] Darriet | Application of magnesium rich rare-earth alloys to hydrogen storage[END_REF], transition metals [START_REF] Pozzo | Hydrogen dissociation and diffusion on transition metal[END_REF], and other metals such as Al, Sn, In and Pd [20].

Kinetic properties of the Mg-H system

During the last years, many efforts have been done to improve the sluggish kinetics of Mg/MgH2 sorption. A quite effective technique is to use mechanical ball milling, creating clean and enhanced surface areas and reducing the crystallite size to the nanoscale [START_REF] Huot | Structural study and hydrogen sorption kinetics of ball-milled magnesium hydride[END_REF]. Figure 1.6 shows the drastic effect of milling on the hydrogen kinetics of magnesium. Ninety percent of the reaction occurs in 12 min for milled magnesium, meanwhile, for the non-milled one, the absorption time takes more than 10 h. This improvement is attributed to a high density of defects that act as nucleation sites for the hydride phase and by grain boundaries that facilitate hydrogen diffusion to the matrix. Increasing in the specific surface area by milling also plays a significant role. In different investigations, it has been shown that reduction in crystallite size is probably the dominant factor of the enhanced kinetics [21,[START_REF] Wronski | Mechanochemical synthesis of nanostructured chemical hydrides in hydrogen alloying mills[END_REF][START_REF] Liang | Catalytic effect of transition metals on hydrogen sorption in nanocrystalline ball milled MgH2-Tm (Tm=Ti, V, Mn, Fe and Ni) systems[END_REF]. Another way to enhance reaction kinetics is the use of catalysts or additives that facilitate hydrogen transport toward and/or hydride nucleation in Mg matrix.

°C

hydride phase with a maximum hydrogen content of TiH2. Crystal structure data are presented in Table 1.2. The TiH diffusion data were collected by Zhou [START_REF] Zhou | A study of advanced magnesium-based hydride and development of a metal hydride thermal battery system[END_REF] in the temperature range of 25 °C to 400°C, and the pressure is referred as 1 atm (0.1 MPa). The values of hydrogen diffusion (DH) of TiH1.71 are lower than those of TiH1.66 and metallic ht-Ti.

Figure 1.9. Intrinsic diffusion coefficients (DH) for some transition metals, intermetallic, and hydrides [START_REF] Zhou | A study of advanced magnesium-based hydride and development of a metal hydride thermal battery system[END_REF].

PCI curves of the Ti-H system (Fig. 1.10) were measured by Arita [START_REF] Arita | Thermodynamics of the Ti-H system[END_REF], at 300 °C and 1.4x10 -3 MPa, the nH/nTi ratio measured was 1.953, which means that the system shows to be pretty stable at this operation conditions. Mg-Ti binary alloys were studied by Liang & Schulz [START_REF] Liang | Synthesis of Mg-Ti alloy by mechanical alloying[END_REF], they synthesized Mg and Ti by reactive ball milling (RBM) getting a Mg-Ti nanocrystalline alloy with an extended solubility of Ti in Mg.

It was observed that Mg(Ti) solid solution is stable at temperatures below 250 °C and phase decomposition becomes fast after 300 °C. Mg(Ti) solid solution transforms to MgH2 and TiH2 upon hydrogen sorption. Because of such decomposition, a destabilization of MgH2 was not reached. Lu et al. [22] investigated hydrogen sorption properties of MgH2-0.1TiH2 mixtures produced by RBM of MgH2-TiH2 powder mixtures. They observed that this material has reversible hydrogen storage with a cycling hydrogen storage of 6 wt.%. There is no loss of cycling capacity over 80 sorption process. They conclude that both the nano-size and the addition of TiH2 contributed to a significant improvement of the kinetics of both (de)hydrogenation of MgH2. Improvement of the kinetic in Mg-Ti system is also reported by Pasquini et al. [START_REF] Pasquini | Magnesium nanoparticles with transition metal decoration for hydrogen storage[END_REF], they synthesized Mg nanoparticles (size range 100 nm to 1 nm) by inert gas condensation and decorated by titanium clusters via in situ vacuum deposition. They observed a complete hydrogenation as well as good cycling stability working at 330 °C and PH2 = 1.8 MPa. Ponthieu et al. [START_REF] Ponthieu | Structural Properties and Reversible Deuterium Loading of MgD2-TiD2 Nanocomposites[END_REF] investigated the Mg-Ti-H system focusing on MgH2-TiH2 nanocomposites. The samples were prepared by RBM of Mg and Ti mixtures under hydrogen pressure to form (1-x)MgH2+xTiH2 nanocomposites with molar ratio 0 < x < 0.5. They consist of MgH2 (with the coexistence of β-and γ-polymorphs) and ε-TiH2 homogenously distributed with crystallite sizes below 15 nm. A study of the sorption properties showed that only the hydrogen stored in magnesium hydride is reversible under moderate conditions (T <330 °C, PH 2 <1 MPa). The presence of the titanium hydride leads to very fast hydrogen sorption kinetics for the magnesium system. TiH2 inclusions are found homogeneously distributed within the MgH2 matrix. They are claimed to enhance hydrogen mobility in Mg by preserving short diffusion paths for hydrogen. Indeed, TiH2 inclusions are shown to limit grain growth of both Mg and MgH2 phases. Furthermore, TiH2 is expected to favor the diffusivity of hydrogen through the existence of a coherent coupling between TiH2 and Mg/MgH2 [START_REF] Crivello | Mg-based compounds for hydrogen and energy storage[END_REF]. As an additive, Ti-hydride shows remarkable performance in improving the hydrogen storage properties of Mg, especially at moderate temperatures (300 °C). Cuevas et al. found that during milling of Mg-Ti powder mixtures under hydrogen atmosphere (8 MPa), TiH2 phase was firstly formed (t<15 min) before MgH2 [23] They also established that the milled 0.7Mg-0.3Ti composite can keep stable cyclic capacity of 3.7 wt.% H at 300 °C for at least 30 cycles .

Zr-Mg system

Figure 1.11 presents the calculated Mg-Zr phase diagram at the temperature range of 300 to 1000 °C by Hämäläinen [START_REF] Hämäläinen | Thermodynamic evaluation of the Mg-Zr system[END_REF]. Experimental phase diagram data exist in a very limited composition range from 0 to 1 at. % Zr. The melting point of magnesium was assessed at 650 °C and zirconium at 1855 °C. There was evidence of a peritectic reaction near the magnesium-rich end. The maximum solid solubility of 1.042 at.% Zr in (rtMg) is reported. The solid solubilities of magnesium in both (htZr) and (rtZr) are negligible. In figure 1.13, Russel et al. [START_REF]Solid Solution Equilibria in the Zirconium-Hydrogen System1 n[END_REF] measured the solubility of hydrogen in zirconium as a function of hydrogen pressure (range 1 to 760 mm. of Hg) and temperature (range 600 to 900 °C). Inside Figure, the temperature is represented by letters as follows: A, 600 °C; B, 700°C; C, 750°C; D, 800°C; E. 825°C; F, 850°C; G, 875°C; H, 900°C. As it can be noted for the operation conditions ZrH2 system is more stable than TiH2. Korablov et al. [START_REF] Korablov | Mechanism and kinetics of early transition metal hydrides, oxides, and chlorides to enhance hydrogen release and uptake properties of MgH2[END_REF] selected hydrides (TiH2, ZrH2), chlorides (VCl3, ScCl3) and oxides (V2O5) as additives facilitating hydrogen release and uptake for magnesium hydride. Mechanism of hydrogen release and uptake in magnesium hydride system at 270 °C, 10 MPa, were assessed. The weakest kinetic effect for hydrogen desorption is provided by TiH2, whereas the strongest influence was provided by V2O5. Zirconium hydride has an intermediate kinetic effect. A similar analysis of additive effect was also performed at a higher temperature of 320 °C, 10 MPa. The additive ability to accelerate the reaction kinetics follows the sequence: ScCl3 < TiH2 < ZrH2 < VCl3 < V2O5. Czujko et al. [START_REF] Czujko | Investigation of the hydrogen desorption properties of Mg+10wt.%X (X=V, Y, Zr) submicrocrystalline composites[END_REF] studied the effects of metal additives on the hydrogen desorption properties of magnesium hydride (β-MgH2) formed after hydrogenation of the Mg+10wt. %X (X= V, Y, Zr) synthesized by RBM. TGA studies of the powders at 350°C shows that the Mg+10 wt.%Zr composite desorbs hydrogen within the highest and simultaneously narrowest temperature range. Furthermore, the Mg composites with V and Y exhibit a two-step desorption. Finally, hydrogen desorption in a Sieverts-type apparatus at 300, 325 and 350 °C under atmospheric pressure of hydrogen shows that the Mg+10wt.% V composite exhibits the most rapid kinetics.

Sc-Mg system

Mg-Sc phase diagram (Figure 1.15) drawn by Nayed [START_REF] Nayeb-Hashemi | The Mg-Sc (Magnesium-Scandium) system[END_REF] consists of: the liquid, L; the terminal (Mg) solid solution; the high-temperature bcc (htSc) solid solution; the intermediate CsCl-type solid solution designated MgSc rt the terminal low-temperature cph (rtSc) solid solution; a peritectic reaction at 710 °C a peritectoid reaction at 520 °C and a eutectoid reaction at 480 °C. The hydrogen-scandium system is plotted in Figure 1.16. For (ScH2), the upper limit for its existence would most probably be the α<-> β transformation temperature of pure Sc at 1337 °C. However, the range of presently available measurements of equilibrium isotherms for the absorption of H in Sc does not extend as far as this temperature, the highest measured isotherm being that for 1100 °C. Manchester et Pitre [START_REF] Manchester | The h-sc (Hydrogen-Scandium) system[END_REF] drew the PCI curves shown in Figure 1.17, which depicts the high stability of ScH2, the researchers also reported the inexistence of hysteresis for this hydride.

Figure 1.17. Pressure-composition isotherms for H-Sc system [START_REF] Manchester | The h-sc (Hydrogen-Scandium) system[END_REF] The Mg-Sc system was studied by Latroche et al. [START_REF] Latroche | Crystal structure of Mg0.65Sc0.35Dx deuterides studied by Xray and neutron powder diffraction[END_REF], these researchers found that thin films of Mg0.65Sc0.35 exhibited a relatively high reversible hydrogen storage capacity (4.1 wt.%) with improved kinetics. They reported an equilibrium pressure at 300 °C (573 K) close to 0.06 MPa on the plateau. They also mentioned that faster kinetics are achievable due to the change from the rutile structure of MgH2 to the fluorite structure of Mg0.65Sc0.35 hydride, which enhances the mobility of hydrogen. Luo et al. [START_REF] Luo | Hydrogen storage properties of nano-structured 0.65MgH2/0.35ScH2[END_REF] studied the thermodynamic of Mg0.65Sc0.35 alloy at different temperatures 320, 340 and 350 °C founding a plateau pressure in desorption at 0.213, 0.355 and 0.452 MPa respectively. From PCI curves, the reversible hydrogen storage capacity was measured (4.2 ± 0.1 wt.%), close to the theoretical value of 4.4 wt.%. They observed two plateaus, the main one appeared over the capacity range 0.6-3.5 wt.%, while the lower plateau was present when capacity was below 0.6 wt.%. A total of 55 cycles were carried out and the hydrogen storage capacity remained unchanged after the initial 2 cycles, indicating the good cycling stability of the 0.65MgH2/0.35ScH2 system. These same researchers studied the hydrogen storage properties of ball-milled xMgH2/(1-x)ScH2 (x= 0.65-1) samples [START_REF] Luo | Catalytic effect of nano-sized ScH2 on the hydrogen storage of mechanically milled MgH2[END_REF]. They found that a small addition of ScH2 (6 mol.%) effectively improved the kinetics of MgH2 by reducing the activation energy value from 159 ± 1 kJ mol -1 of the asreceived MgH2 to 99 ± 6 kJ mol -1 . It was also observed that the samples in which the ScH2 content ranged from 6 to 12 mol.% and underwent ball milling for a longer duration (20 and 40 h) showed the most improved dehydrogenation kinetics of MgH2 over cycling. They suggested that the nano-sized ScH2 could help the dissociation of H2 as well as functioning as nucleation sites for the growth of Mg during dehydrogenation.

Y-Mg system

The Mg-Y phase diagram (Figure 1.18) was drawn by [START_REF] Okamoto | Mg-y (magnesium-yttrium)[END_REF], in which the phase boundaries of (htY) and (rtY) were estimated due to lack of experimental data, by optimization of thermodynamic and phase diagram data. Hydrogen yttrium phase diagram, Figure 1.19. Khatamian [START_REF] Khatamian | The H-Y (Hydrogen-Yttrium) system[END_REF] refers to a "solid solution phase," i.e., the phase with the pure metal as one boundary and the mixed phase region as the other boundary. For the Y-H system at room temperature this boundary is located at 20 at.% H; in the present evaluation, this phase has been designated α. Greek letters are preferable to lose terminology (such as "the hydride phase") if a general agreement on the choice of Greek letters can be established; e.g., the phase centered around the YH2 composition, which has been variously labeled β phase, ϒ phase, or hydride phase. The present evaluators suggest that this phase is designated δ, to conform to the phase nomenclature used for the Ti-H system. The α phase of the Y-H system retains the cph structure, the δ phase has fcc structure, and the phase above the approximate composition YH3 has a hexagonal structure. Crystal structure data are presented in Table 1.5. Yannopoulos et al. [START_REF] Yannopoulos | The Thermodynamics of the Yttrium-Hydrogen System1[END_REF] published the isothermal equilibrium data for the yttrium-hydrogen system which are represented in Figure 1.20 for the temperature range from 601.4 to 949.4 °C. From the shape of the isotherms, it is inferred for the condensed state the existence of two single-phase regions separated by a two-phase field. In the dilute region of the Y-H system up to 0.1 g.-atom of hydrogen/g.-atom of yttrium. Reversibility of the measurements for the composition range demonstrates that it fits with the isothermal desorption points in the righthand portion and the desorption points (•) in the left of the figure. YHx is commonly used as a switchable mirror, real-time transitions from metallic (YH2) to semiconducting (YH3 ) behavior occur in YHx films during continuous absorption of hydrogen, accompanied by pronounced changes in their optical properties [START_REF] Huiberts | Yttrium and lanthanum hydride films with switchable optical properties[END_REF].

Several thermodynamically unstable transition metal aluminum hydrides with a high hydrogen content like Y(AlH4)3 have been studied because of the relatively high stability of Y(AlH4)3. Y(AlH4)3 was reported to be quite stable and starts to decompose at 50 °C. Y(AlH4)3 has a high theoretical hydrogen content of 6.6 wt%, while until now no further information is known about its hydrogen storage properties and reversibility [START_REF] Cao | Reversible hydrogen storage in yttrium aluminum hydride[END_REF].

V-Mg system

The Mg-V is presented by Smith [START_REF] Smith | Binary Alloy Phase Diagrams[END_REF] in Figure 1.22, who explains that there are no intermediate phases in the system, however, extensive immiscibility was observed in both liquid and solid state. Vanadium, which crystallizes in the bcc cubic structure at 0.1 MPa, absorbs interstitial H up to an atomic H/V ratio of about two. For H/V ratios up to unity, the phase diagram is shown in Figure 1.23. Above 200°C (473 K), the α-phase with bcc structure of vanadium is observed. Only a lattice expansion accompanies the increase of the hydrogen concentration in this cubic phase (solid-solution behavior). At lower temperatures, the β-hydride phase, in which the V atoms form a body-centered tetragonal lattice, appears. Depending on whether the cubic phase has a lower or higher hydrogen concentration than the tetragonal β-phase, it is usually denoted by α or α'-phase, respectively. At room temperature, pure β-phase exists for H/V ratios between 0.45 and 0.9. The lower concentration is in equilibrium with the α-phase containing H/V = 0.05. However, for H/V ratios exceeding 0.9, a new phase, the γ-phase, appears in coexistence with the β-phase. The γ-phase has a fcc arrangement of vanadium atoms and contains a hydrogen concentration of H/V = 2. Crystal structure data are presented in Table 1.6. Figures 1.24 shows a series of pressure-composition isotherms on vanadium films over the whole concentration range V-VH2. Pressure plateaus indicate regions where two hydride phases coexist with the gas phase. The phase diagrams in the inserts [START_REF] Alefeld | Hydrogen in Metals II: Application-Oriented Properties[END_REF] are determined by temperature variations of sealed samples. The hysteresis is negligible for 0.5 < x < 1.0. For 1.0 < x < 1.8, the plateau pressures for hydriding are a factor of 2-3 higher than the dehydriding pressures shown here. Values given against the full curves are in °C. Pressure-composition isotherms for VH [START_REF] Alefeld | Hydrogen in Metals II: Application-Oriented Properties[END_REF].

Several pressure-composition isotherms of the V-H system exists in the region VH0.8 to VH2.0 [START_REF] Reilly | Higher hydrides of vanadium and niobium[END_REF]. The initiation of a pressure plateau at a composition approximating VH0.95 marks the point at which the γ phase appears in the previously existing β phase. At 40 °C, the plateau extends from VH0.95 to VH2.00 at which point the isotherm rises sharply, marking the composition at which the β phase disappears and only the γ phase exists. As indicated in Figure 1.24, hydrogen may dissolve in the γ phase and H: V atomic ratios above 2.00 are possible. It will be noticed that the 67 and 78°C isotherms cross one another in the single-phase (γ) region; this is a consequence of the non-stoichiometry of the system in which the higher temperature isotherm was obtained using a sample which had a greater initial hydrogen content (ie., VH2.08 vs. VH2.01 at room temperature). Liang et al. [START_REF] Liang | Hydrogen storage properties of the mechanically milled MgH2-V nanocomposite[END_REF] have prepared a composite MgH2+5 at.% V by RBM leading to a nanocomposite mixture of β-MgH2 + γ-MgH2 + VH0.8, which can desorb hydrogen at 200 °C under vacuum and re-absorb hydrogen quickly even at room temperature. The rehydrogenation of this nanocomposite is extremely fast. In fact, at PH 2 = 1 MPa, it can absorb 2wt. % in 16 min at 29°C, 4 wt. % in 2 min at 100 °C and 5.6 wt. in 4 min at 200 °C. They concluded that the enthalpy and entropy of the composite are identical to that of pure magnesium. Ren et al. [START_REF] Ren | Hydrogen Storage Properties of Magnesium Hydride with V-Based Additives[END_REF] analyzed the effects of pure vanadium and V-based solid solution alloys (V75Ti5Cr20, V80Ti8Cr12, V80Ti20, and V80Cr20) on the sorption kinetics of magnesium hydride. 5 mol % of V-based additive was used. In all cases, dehydrogenation temperatures of MgH2 with V-based additive were approximately 100 °C, 0.1MPa lower than that of pure MgH2. Among the V-based additives, V75Ti5Cr20 provides the lowest dehydrogenation temperature (217°C). Samples with V75Ti5Cr20, V80Ti8Cr12, and pure V as additives fully dehydrogenate within 6 min. As compared to the pure MgH2 sample, which did not dehydrogenate at all at this low temperature. These results demonstrate a significant improvement in the kinetics of dehydrogenation of MgH2 with V-based alloys additives. The significant hydrogen absorption of MgH2 with V-based additives (1 bar hydrogen pressure and room temperature) suggests a strong effect of V-based additives on the hydrogenation of Mg.

As concerns thermodynamics, they reported that the V-based additives do not affect the thermodynamic stability of magnesium hydride.

Nb-Mg system

The Mg-Nb phase diagram was drawn by Smith [START_REF] Lomnitska | V-Si-Sb phase diagram. Isothermal section at 1070 K[END_REF], which is very similar to Mg-V (Figure 1.22 being more stable Nb than V. It is known that hydrogen is dissolved in large amounts into niobium, the maximum capacity (Cmax  2 H/M) doing it in tetrahedral interstitial sites. The phase diagram of the Nb-H system is depicted in Figure 1.27. At room temperature, the solubility limit max of hydrogen in Nb extends approximately to x= 0.04. In the concentration range of 0.7<x< 1.1, the hydride (βphase) occurs with the hydrogen atoms ordered in chains along a (110) direction. The next higher phase corresponds roughly to the NbH2-phase (δ-phase). Crystal structure data are presented in Table 1.7. Hydrogen is extremely mobile in niobium, in the α-phase the diffusion coefficient D depends only slightly on temperature and at -173 °C it still has a value of some 10 -8 cm 2 /s [START_REF] Welter | Superconducting transition temperature and low temperature resistivity in the niobium-hydrogen system[END_REF].

The Nb-H2 system behaves in a manner similar to the V-H2 system [START_REF] Reilly | Higher hydrides of vanadium and niobium[END_REF][START_REF] Pryde | Phase equilibria and kinetics of evolution of dilute solutions of hydrogen in niobium[END_REF]. In Figure 1.28 is shown several pressure-composition isotherms for the region extending from Nb0.1 to Nb0.8 on the left and NbH0.9 to NbH2.0 on the right. The isotherms rise steeply on the left until a composition of NbH1.1 is reached where the γ-dihydride phase appears. At 40 °C the β+γ mixed phase region extends to NbH1.9 where the isotherm begins a steep ascent indicating the disappearance of the β phase. It may be noted that the lower temperature isotherms indicate an appreciable solubility of hydrogen in the pure γ phase which resulted in hydrogen contents up to a composition corresponding to NbH2.081, niobium dihydride, as are other group Vα hydrides, is a nonstoichiometric compound. Here, too, it is possible for the isotherms to cross one another in the single-phase (γ) region because of the variable initial hydrogen content.

Figure 1.28. Pressure-composition isotherms for NbH0-0.8 (left) and NbH0.9-2.1 (right) [START_REF] Reilly | Higher hydrides of vanadium and niobium[END_REF][START_REF] Pryde | Phase equilibria and kinetics of evolution of dilute solutions of hydrogen in niobium[END_REF].

Some studies have shown that Nb additive can intensively enhance the hydrogen sorption kinetics of magnesium [START_REF] Cui | Ti, Nb, V, Co, Mo or Ni) coreshell like nanostructures: synthesis, hydrogen storage performance and catalytic mechanism[END_REF], [START_REF] Bazzanella | Catalytic effect on hydrogen desorption in Nb-doped microcrystalline MgH2[END_REF]. On the basis of the first principle calculations, Song et al. [START_REF] Song | Influence of selected alloying elements on the stability of magnesium dihydride for hydrogen storage applications: A first-principles investigation[END_REF] confirmed that the alloying element Nb could destabilize magnesium hydride because it makes the bond energies between magnesium and hydrogen weaker from a kinetic viewpoint. Huot et al. [START_REF] Huot | Investigation of dehydrogenation mechanism of MgH2-Nb nanocomposites[END_REF] established that the ball milled MgH2 + 5 atom% Nb composite could absorb 5.2 wt % H2 at 300 °C within 5 min at Pabs= 0.1 MPa and Pdes=0.015MPa. Yavari et al. [START_REF] Yavari | Structural evolution and metastable phase detection in MgH2-5%NbH nanocomposite during in-situ H-desorption in a synchrotron beam[END_REF] also confirmed an increase of hydrogen desorption kinetics when Nb was added to MgH2. De Castro et al. [24] reported that the desorption temperature of the ball milled Mg + 5 atom% Nb nanocomposite was 270 °C, the dehydrogenation behavior was studied by differential scanning calorimetry (DSC). Thus, Nb is considered as an effective additive not only for significantly reinforcing hydrogen sorption kinetics but also reserving the superior hydrogen sorption capacity of the Mg-based composite, due to a result indicating that Mg promoted the hydrogenation of Nb in the same way that Nb promoted the hydrogenation of Mg. With respect to the catalytic mechanism of Nb, Huot et al. [START_REF] Huot | Investigation of dehydrogenation mechanism of MgH2-Nb nanocomposites[END_REF] claimed that the formation of NbHx (x ≈ 0.9) phase during dehydrogenation acted as a gateway to accelerate the hydrogen release from MgH2. Considering the good catalytic effect of Nb on the hydrogen sorption of the milled Mg-Nb composite, Liu et al. [START_REF] Liu | Catalytic Effect of Nb Nanoparticles for Improving the Hydrogen Storage Properties of Mg-Based Nanocomposite[END_REF] intended to fabricate the Mg-Nb nanocomposite by hydrogen plasma-metal reaction, the Mg-7.5 wt % Nb nanocomposite was successfully prepared by this method. Nb nanoparticles of 12 nm are homogeneously decorated on the surfaces of the Mg nanoparticles and effectively prevent the Mg nanoparticles from growing during both the synthesis stage and the hydrogenation/dehydrogenation process. The Mg-Nb nanocomposite showed fast hydrogen absorption/desorption kinetics and high hydrogen storage capacity. The Mg-Nb nanocomposite can quickly uptake 4.0 wt % H2 in 10 min and reach a saturation value of 5.7 wt % H2 in 60 min at 200 °C. Furthermore, it can also release 4.0 wt % H2 in 60 min at 300 °C. The reversible hydrogen storage capacity is as high as 7.0 wt. % at 400 °C. The enhanced hydrogen storage capacities and the reduced apparent activation energies are due to the prominent catalytic effect of Nb nanoparticles and the nanosizing effects of both Mg and Nb.

Chapter 2: Synthesis and characterization techniques 2.1 Materials and samples handling

All the elements involved in this research have been gotten commercially. The sources and forms of all used raw metals are listed in Table 2.1. They have been used as-received without further purification. The rare earth metals used in the automatic cycling device were filed down right before melting to remove the native oxide layer at the surface. The raw elements were carefully weighted to desire stoichiometric proportion with a precision of ± 0.001 mg. All samples handling, weighting, and loading were performed in an argon-filled glove box with a controlled amount of oxygen (<2 ppm) and moisture (H2O < 45 ppm). For the synthesis of hydrides, 6 N Alphagaz (99,9999 % purity) was used.

Synthesis of metal hydrides and alloys

Reactive Ball Milling (RBM)

Mechanochemical synthesis of metal hydrides using ball milling has become a very frequently used method. Different kind of planetary ball mills such as rotational, vibratory or attritor mills is commonly used. The different types of mills differ in their milling efficiency and capacity as well as arrangements for cooling, heating, gas loading etc. Planetary ball mills are used wherever the highest degree of fineness is required. In addition to well-proven mixing and size reduction processes, these mills also meet all technical requirements for colloidal grinding and provide the energy input necessary for mechanical alloying. The extremely high centrifugal forces of a planetary ball mill result in very high pulverization energy and therefore short grinding times. Typically, a few grams of material and balls are placed in a planetary ball mill to give a ball-topowder weight ratio of 10:1-60:1.

Several parameters can be varied for the ball-milling synthesis: milling speed, total milling time, vial and ball composition, powder-to-ball weight ratio, vial diameter, ball diameter and density, milling temperature, milling atmosphere and pressure of the reactive gas. Most planetary mills only allow controlling the speed of the main support disk. The speed of the planets, on which the milling vials are mounted, is usually fixed relative to the speed of the main disk. However, for mills, such as the Fritsch Vario-Planetary Mill Pulverisette 4, both the speed of the support disk and the planets can be varied freely. Thereby, the trajectory of the balls within the vial may be controlled at least when the number of balls is low. Ideally, the milling can be continuously changed from the high-energy mode dominated by high-energy ball-vial impacts to a grinding mode where the balls mainly follow the circumference of the vial [1]. Different MgH2-TMHx nanocomposites had been prepared by reactive ball milling during this work. The systems are listed in Table 2.2. RBM experiments under hydrogen gas (6 N Alphagaz) were performed in a high-pressure milling vial (Evicomagnetics, Germany) equipped with gas pressure and temperature sensors [2]. The commercial vial was connected to a manometric hydrogenation device equipped with calibrated volumes. Using this device, we established the starting hydrogen pressure in the vial (typically 8 MPa) and the volume occupied by the gas in the vial (typically 166 cm 3 ). The milling process was performed in a Fritsch Pulverisette 4 planetary mill at disk and vial rotation speed of 400 rpm and 800 rpm (relative to disk), respectively, with a ball-topowder mass ratio of 60:1. Stainless steel balls of 12 mm in diameter were used. Milling was carried out in two subsequent cycles, 120 min milling time and 120 min rest each.

Electromagnetic processing of materials by induction melting

Electromagnetic processing of material is a technology which works with different interactions between material and electromagnetic field. The high-frequency induction heating and melting are one of the first applications of electromagnetic fields [3]. The alloys are prepared by co-melting according to the technique of high-frequency induction melting, also called levitation in cold crucible. An induction coil is connected to the power supply so that a magnetic field is generated from the current flow. The starting metals with stoichiometric quantities are placed in the copper crucible with inner water-cooling circulation. The crucible is surrounded by a water-cooled alternating current solenoid coil of copper. An aperiodic generator provides a maximum power of 50 kW and delivers a high-frequency alternating current (up to 120 kHz). The LaNi4.66Sn0.34 alloy was prepared by the co-fusion of the elements in a highfrequency induction furnace. A no inductive cooper melting-pot cooled by water circulation was put into a vacuum enclosure incorporated by a silica tube. The magnetic field applied to the sample is created by an external solenoid. The primary and second pumps allow to reach a pressure of 7x10 -10 MPa, the residual pressure is measured by a vacuum gauge, the temperature is measured by an optical pyrometer.

The three high purity metals were cleaned on the surface and weighted in an accuracy of ±0.0001 mg. to produce a 112 g final product. The target was to synthesize more material than calculated before to prevent loss of material during synthesis as well as to have extra material for its characterization. Both nickel and tin were manipulated under air, meanwhile, lanthanum was prepared inside the gloves box due to its high reactivity with the oxygen. The fusion was carried out increasing slowly the temperature to allow the metals to degas. To limit the nickel and tin vaporization, it was performed a reduction with argon. The final product was cracked into several pieces and put inside a silica tube under primary vacuum (1x10 -7 MPa) and annealing was performed at 850 °C for 8 days. Finally, the silica tube was cooled by water quenching. The alloy was recovered inside the glove box and stored for characterization.

Structural, morphological and chemical characterization 2.3.1 X-Ray Diffraction (XRD)

Because of the availability and properties of X-rays, diffraction is one of the most common techniques for phase identification and structural characterization of crystalline materials. This technique allows determining crystalline structures on powdered materials. X-rays wavelengths used for crystallographic studies are typically in the range 0.5-2.5 Å. X-ray is scattered by electrons of the atoms and the important consequences of this are that [4]:

o The atomic scattering factor of elements, 𝑓 𝑖 in Eq. ( 2.3) increases linearly with the atomic number Z; o The 𝑓 𝑖 factors have a significant dependence on the scattering angle, the maximum value (equal to the number of electrons) is at zero scattering angle and decreases with increasing scattering angles; o X-rays interact strongly with matter.

In this work, the XRD analyses were carried out with a Bruker D8 Advance - Diffractometer with Cu-K Radiation (wk1=1.54051 Å, λK2= 1.54433 Å) and graphite rear monochromator. A value of 2θ ranging from 25 to 70 degrees was used during this procedure. Because of the nanocomposites reactivity, a special air-tight sample-holder was used to avoid their exposition with air and moisture. XRD sample-holder was loaded with about 200 mg of powder inside the glove box and hermetically closed with an X-ray transparent lid.

The Rietveld and Loppstra method

X-ray powder diffraction is an efficient method for determining the phase content of a polycrystalline material. Every material exhibits a typical 'X-ray fingerprint', which is stored in databases such as the ICDD PDF2 or PDF4. This fingerprint is utilized in the DIFFRAC.EVA software for reliable phase identification. Once the phases and their structure had been identified, the XRD experimental profile can be fitted to a calculated profile by means of the Rietveld's method [5,6]. For this work, diffraction patterns were analyzed using the Fullprof Suite software [7].

Structure refinements -the Rietveld method

The final step in a structure solution is to fit the experimental diffraction pattern with a calculated pattern while refining the crystallographic parameters of all detected phases. Loopstra and Rietveld [6] suggested using the whole experimental powder diffraction diagram in the refinements by fitting a calculated diagram to the observed data. It should be emphasized that the Rietveld method is a structure refinement method and not a method to solve unknown structures. An adequate structure model must be known in advance.

The minimizing function in the Rietveld method is:

𝑀 = ∑ 𝑤 𝑖 (𝑦 𝑖 𝑜𝑏𝑠 -𝑦 𝑖 𝑐𝑎𝑙𝑐 ) 2 𝑛 𝑖=1 [2.1]
n is the number of measured data points, 𝑦 𝑖 𝑜𝑏𝑠 and 𝑦 𝑖 𝑐𝑎𝑙𝑐 are the observed and calculated intensity at point 𝑖 of the powder diffraction pattern, respectively, and 𝑤 𝑖 is the weight of the point given by 𝑤 𝑖 = 1/𝜎 𝑖 2 where 𝜎 𝑖 is the variance. The calculated intensity (assuming a single-phase sample) can by described by:

𝑦 𝑖 𝑐𝑎𝑙𝑐 = 𝑏 𝑖 + 𝑆 𝑖 ∑ 𝑀 ℎ𝑘𝑙 • 𝐿 𝑖 • 𝑃 𝑖 • 𝐴 ℎ𝑘𝑙 • 𝑇 ℎ𝑘𝑙 • |𝐹 ℎ𝑘𝑙 | 2 • Φ(2𝜃 ℎ𝑘𝑙 -2𝜃 𝑖 ) ℎ𝑘𝑙 [2.2]
𝑏 𝑖 is the background at point 𝑖. It can be either determined manually or by interpolation between selected points or described by a polynomial. The sum covers all reflections contributing to Bragg scattering at point 𝑖. 𝑆 𝑖 is the scale factor, 𝐿 𝑖 is the Lorentz factor that is defined by the geometry diffraction, 𝑃 𝑖 is the polarization factor accounting for the partial polarization of the scattered electromagnetic wave, 𝑀 ℎ𝑘𝑙 is the multiplicity factor accounting for the presence of multiple symmetrically equivalents points in reciprocal space 𝐴 ℎ𝑘𝑙 , 𝑇 ℎ𝑘𝑙 is the preferred orientation factor, Φ is the profile function coming from instrumental and sample broadening in the diffraction pattern, 2𝜃 ℎ𝑘𝑙 is the ideal Bragg angle for reflection ℎ𝑘𝑙, 2𝜃 𝑖 is the scattering angle corresponding to point 𝑖, the structure factor 𝐹 ℎ𝑘𝑙 is the structure factor defined by:

𝐹 ℎ𝑘𝑙 = ∑ 𝑓 𝑖 𝑒𝑥𝑝2𝜋𝑖(ℎ𝑥 𝑗 + 𝑘𝑦 𝑗 + 𝑙𝑧 𝑗 )exp (-𝐵 𝑗 𝑠𝑖𝑛 2 𝜃/𝜆 2 ) 𝑁 𝑗=1 [2.3]
𝑓 𝑖 is the atomic scattering factor that is strongly dependent of both radiation type and elements, 𝑁 is the number of atoms in the unit cell, 𝑥 𝑗 , 𝑦 𝑗 , 𝑧 𝑗 are the atomic fraction coordinates of factor 𝑗 in the unit cell, 𝐵 𝑗 is the isotropic displacement parameter of atom 𝑗 defined as 𝐵 𝑗 = 8𝜋 2 (𝑢 ̅ 2 ) 𝑗 , (𝑢 ̅ 2 ) 𝑗 is the root mean square deviation of atom j from the equilibrium position in Å 2 .

Several factors influence the shape and width of the measured Bragg peaks, the most important are those related with the X-ray apparatus: collimation of the beam, degree of mosaicity of the monochromator, detector resolution; and the related to the nature of the sample: size and in addition to imperfection/mosaicity. The resulting profile shape is a convolution of all these components. The measured profiles show both a Gaussian and Lorentzian contribution. The Rietveld method was developed with a pure Gaussian distribution which well fits data from low-resolution neutron powder diffractometer, but with improved resolution, and X-Ray diffraction in general, a significant Lorentzian contribution is present.

Nowadays the so-called pseudo-Voigt profile function is a commonly used profile function. It is a linear combination of a Gaussian and a Lorentzian function. The peak broadening is measured by the "full-width-at-half-maximum" (FWHM). The Gaussian contribution is given by: 𝐹𝑊𝐻𝑀 𝐺𝑎𝑢𝑠𝑠𝑖𝑎𝑛 = √𝑈 𝑡𝑎𝑛 2 𝜃 + 𝑉 𝑡𝑎𝑛 𝜃 + 𝑊 [2.4] And Lorentzian:

𝐹𝑊𝐻𝑀 𝐿𝑜𝑟𝑒𝑛𝑡𝑧𝑖𝑎𝑛 = 𝑌 cos 𝜃 + 𝑋 tan 𝜃 [2.5]
Table 2.3 are summarized the typical parameters which are refined in the Rietveld method. The starting point for the refinement should be an adequate structure model. Many software programs are available for Rietveld refinements, such as Fullprof [7], GSAS [8] and TOPAS-Academic [9]. The figure of merit to evaluate the quality of the refinements are the weight profile R-factor:

𝑅 𝑤𝑝 = [ ∑ 𝑤 𝑖 (𝑦 𝑖 𝑜𝑏𝑠 -𝑦 𝑖 𝑐𝑎𝑙𝑐 ) 2 𝑛 𝑖=1 ∑ 𝑤 𝑖 (𝑦 𝑖 𝑜𝑏𝑠 ) 2 𝑛 𝑖=1 ] 1/2 [2.6]
And the goodness of fit:

𝑋 2 = ∑ 𝑤 𝑖 (𝑦 𝑖 𝑜𝑏𝑠 -𝑦 𝑖 𝑐𝑎𝑙𝑐 ) 2 𝑛 𝑖=1 𝑛 -𝑝 = [ 𝑅 𝑤𝑝 𝑅 𝑒𝑥𝑝 ] 2 [2.7]
where 𝑝 is the number of free least square parameters and 𝑅 𝑒𝑥𝑝 is the so-called expected profile R-factor.

An important application of the Rietveld method is the determination of relative amounts of different phases in a sample with a so-called quantitative phase analysis. Based on multiphase refinements the molar-, weight-, and volume-fractions can be calculated from the individual scale factors and the unit cells contents and volume.

Finally, the accurate analyses of the profile parameters and their dependence on the scattering angle 2θ can give information about strain and particle sizes in the sample. The term U in Eq.2.4 describes strain contribution to peak broadening, while the term Y in Eq. 2.5. corresponds to the crystal size distribution. The term U in Eq.2.4 describes strain contribution to peak broadening, while the term Y in Eq. 2.5. corresponds to the crystal size distribution [10].

Scanning Electron Microscopy (SEM)

By irradiating a sample with a monochromatic electron beam, scanning electron microscopy (SEM) takes advantage of the diverse interactions between the electrons and the sample to obtain chemical and morphological information on the material.

The images in secondary electrons give information on the surface morphology of the sample. The angular emission of backscattered electrons (BSE) is characteristic of the atomic number on the probe path. Heavy atoms (large atomic number Z) induce stronger elastic scattering between the incident electron and the nucleus. Therefore, BSE images can provide a chemical contrast of the specimen. Qualitative analysis is a routine test on interesting samples to verify chemical compositions. Energy dispersive spectroscopy (EDS) quantitative analysis is performed using a Zeiss Ultra field emission scanning electron microscope equipped with a Bruker SDD EDS analyzer [11].

The morphological characterization of the materials was performed using scanning electron microscopy for imaging. The microscope was provided by the company Carl Zeiss, it is a MERLIN SEM. It has a field effect gun (FEG) to obtain an electronic probe of very small diameter (of the order of one nanometer at the sample), nevertheless carrying a large current (of the order of the nanoampere). This current is constant (variation less than 0.5% per hour), which allows easier and more reliable analyzes. The design of the column and especially of the objective lens of this microscope takes full advantage of this beam by significantly reducing the aberrations of the electromagnetic optical system, even at low electron acceleration voltages. This, together with an annular secondary electron detector located above the objective lens, makes possible to obtain morphological observations with excellent resolution at any acceleration voltage: the guaranteed resolutions are 0.8 nm to 15 kV and 1.4 nm at 1 kV on gold. A second annular detector allows the detection of backscattered electrons with low acceleration voltage, to obtain chemical contrast images with a very good spatial resolution. The secondary and backscattered electron detectors arranged this time in the object chamber, improving the observation capabilities of the microscope. Particle size was measured with the specialized software ImageJ, this software allows to configure the image on a known scale, then to calculate the area of the particles on the image, finally, the size can be measured from the given area [12].

Inductively Coupled Plasma Optical Emission Spectrometry (ICP-OES)

ICP-OES, an abbreviation for Inductively Coupled Plasma, is a method of optical emission spectrometry. ICP/OES is one of the most powerful and popular analytical tools for the determination of trace elements. The technique is based upon the spontaneous emission of photons from atoms and ions that have been excited in a radiofrequency discharge. Liquid and gas samples may be injected directly into the instrument, while solid samples require extraction or acid digestion so that the analyses will be done in solution. The sample solution is converted to an aerosol and directed into the central channel of the plasma. At its core, the inductively coupled plasma (ICP) sustains a temperature of approximately 10 000 °C, so the aerosol is quickly vaporized. Elements are liberated as free atoms in the gaseous state. Further collisional excitation within the plasma imparts additional energy to the atoms, promoting them to excited states. Sufficient energy is often available to convert the atoms to ions and subsequently promote the ions to excited states. Both the atomic and ionic excited state species may then relax to the ground state via the emission of a photon. These photons have characteristic energies that are determined by the quantized energy level structure for the atoms or ions. Thus, the wavelength of the photons can be used to identify the elements from which they originated. The total number of photons is directly proportional to the concentration of the originating element in the sample.

The instrumentation associated with an ICP/OES system is relatively simple. A portion of the photons emitted by the ICP is collected with a lens or a concave mirror. This focusing optic forms an image of the ICP on the entrance aperture of a wavelength selection device such as a monochromator. The wavelength exiting the monochromator is converted to an electrical signal by a photodetector. The signal is amplified and processed by the detector electronics, then displayed and stored on a computer [13].

The analyses were performed with a Varian Vista-Pro ICP-OES. For that purpose, MgH2-TiH2 samples (20 mg) were solubilized with acids (3 mL HCl 37%+ 2 mL HNO3 65% for samples containing Ti, whereas 2 ml HNO3 was used for samples without Ti) and pure water.

Electron Probe Micro-Analysis (EPMA)

The electron microprobe, also known as the electron probe micro-analyzer (EPMA), uses X-ray spectrometry to both identify and measure concentration of elements in microscopic volumes of the specimen. This technique was used to determine the phase composition of induction-melted La(Ni, Sn)5-type alloys used in the automatic cycling device. With EPMA, a finely focused electron beam interacts with the sample to generate back-scattered electrons, secondary electrons, characteristic X-rays and light known as cathodoluminescence. By scanning the electron beam over a surface of the sample, these signals can be used to obtain high resolution scanning electron images, X-ray maps showing spatial distribution of elements, and cathodoluminescence images for phase (element and compound) identification, estimation of phase distribution and proportions, trace element compositional variation and surface textural analysis of multi-phase composites. Characteristic X-rays generated from a microscopic volume in spot mode (i.e., beam not scanning) are utilized to obtain a complete quantitative chemical analysis. Sample preparation is minimal, but a well-polished surface is required for accurate quantitative analysis. The EPMA is mostly used for studying inorganic materials [14]. This is a non-destructive in situ analysis technique that detects all elements from beryllium in a cubic micron volume with a sensitivity of about 100 ppm. For this purpose, it relies on the measurement of the characteristic X-ray intensity emitted by a given element under excitation conditions.

The features of this device are:

Characterization of hydrogenation properties 2.4.1 The Sieverts' method

The Sieverts' method [15] is a volumetric measurement technique that was commonly used to measure the sorption of gaseous hydrogen into materials [16]. In the literature, the detailed description of this measurement method can be found [17]. The amount of hydrogen absorbed by a sample can be calculated using the real gas law:

𝑃𝑉 = 𝑛𝑍𝑅𝑇 [2.8]
where, P is the pressure, V is the volume, n is the value of moles, Z is the gas compressibility, R is the universal gas constant, and T is the temperature.

H-sorption measurements were made in a home-built Sieverts apparatus operated manually. This device is thermalized in water at 25 °C. It is equipped with two reservoirs of different volume each, a supplier hydrogen valve, one release valve to the atmosphere, a primary vacuum pump (1 Pa) and a connection to the sample holder which is heated to the operation temperature by an electric tubular oven.

Pressure-Composition-Isothermal (PCI), kinetic and cyclic measurements

As is has been explained in Section 1.3.1, PCI measurements are used to characterize the thermodynamics of hydrogen sorption/desorption in a solid. Experiments are performed at isothermal conditions while measuring the amount of hydrogen absorbed/desorbed (Eq. 2.8) at an equilibrium hydrogen pressure. PCI measurements on absorption-desorption were carried out for all MgH2-TMHx systems (TM= Ti, Zr, Sc, Y, V, and Nb) with the target of determining their thermodynamics.

Isothermal sorption kinetic measurements were performed at 300 °C in the same system at quasi-constant pressure using large reservoirs. Applied absorption and desorption pressures were set to Pabs = 0.8 and Pdes = 0.03 to provide similar driving forces for absorption and desorption. At 300°C, the equilibrium plateau pressure of MgH2/Mg system is 0.16 MPa. Therefore, driving forces were fixed to ln(Pabs/Pp) = ln (Pp/Pdes)  1.6.

Hydrogen cycling experiments were performed manually for all MgH2-TMHx systems for 20 cycles. Same conditions as for kinetic measurements were applied. In all cases, sorption time both on absorption and desorption sweeps was constrained to 15 min.

Temperature-Programmed Desorption (TPD)

TPD technique is a characterization method for the determination of the kinetic parameters in desorption processes or hydride decomposition reaction [18]. A hydrogenated sample is placed in a closed system and heated up with a temperature program β(t) = dT/dt (with the temperature T usually being a linear function of the time t) at low partial pressure of hydrogen. The amount of hydrogen desorbed from the sample as a function of the temperature is measured by the Sieverts' method. Hydrogen gas desorption experiments were performed at initial Pdes = 0.03 MPa, by heating from RT to 300°C at a heating rate of 2°C/min.

Temperature-programmed absorption (TPA)

The TPA technique is similar to the TPA but with the aim of determining kinetic parameters of the absorption processes. A hydrogen-free sample is placed in a closed system and heated up under a high partial pressure of hydrogen. The amount of hydrogen absorbed by the sample as a function of the temperature is measured by the Sieverts Method. Hydrogen gas absorption experiments were performed at Pabs = 0.8 MPa by heating from RT to 300°C at a heating rate of 2°C/ min.

Effect of TiH2 additive on the hydrogen storage properties of magnesium hydride

In this chapter, the system (1-y)MgH2+yTiH2 has been analyzed with the objective of having a better understanding of the role that TiH2 addition plays during hydrogen sorption process of MgH2. (1-y)MgH2+yTiH2 composites were synthesized by reactive ball milling. Their morphology and structural properties were determined by XRD and SEM in the as-milled state and after hydrogen cycling. Pressure Composition Isotherm (PCI) measurements were done to determine thermodynamics modifications in the Mg-H system by TiH2 addition. Hydrogen cycling properties were determined using Sieverts' type apparatus for twenty absorption and desorption cycles. The evolution of the reversible storage capacity on cycling for a limited reaction time of 15 minutes has been measured with significant variations as a function of the additive quantity. Finally, hydrogenation kinetics were analyzed using available solid-state kinetic models and the rate-limiting step was determined both for absorption and desorption.

Synthesis and chemical characterization of MgH2-TiH2 nanocomposites

(1-y)MgH2+yTiH2 composites with Ti molar ratio y = 0, 0.0125, 0.025, 0.05, 0.1, 0.2, and 0.3 were synthetized by RBM under hydrogen atmosphere as described in chapter 2. For example, Figure 3.1 shows the evolution of hydrogen pressure and vial temperature during the synthesis of the composite with y = 0.3. The synthesis was performed in two cycles of 120 minutes milling and 120 minutes pause between them. At the beginning of milling, the vial temperature increases because of the collisions and frictions between the milling balls and the vial walls. This also leads to an increment of pressure due to gas confinement. Then, the pressure drops down due to the hydrogen absorption by both Mg and Ti metals leading to the formation of MgH2 and TiH2 phases [1], respectively. Once the hydride formation is completed, a pressure plateau is observed. At the end of the first milling cycle, both pressure and temperature decrease to equilibrium during the pause time. The second cycle does not entail any additional hydrogen absorption, being the change in pressure due to thermal variations induced by milling and rest. Thus, the second cycle can be used to calibrate the gas temperature that differs from the measured vial temperature [2]. This allows determining the hydrogen uptake (CH) corresponding to the pressure change observed during the first cycle. Hydrogen contents CH for all synthetized composites are shown in Table 3.1. Experimental CH values concur well with the nominal ones (according to the expected formation of MgH2 and TiH2 di-hydrides) with a maximum discrepancy of 9% for y = 0.2. Hydrogen uptake curves collected during the synthesis of composites y = 0, 0.05 and 0.1 are displayed in Figure 3.2. They were obtained from pressure and temperature data variations of the first milling cycle. For all samples, complete hydrogenation reaction takes less than 80 min. Hydrogenation starts within the first minutes of the milling process and is much faster for Tiadded composites (y = 0.05 and 0.1) than for pure Mg (y=0). The hydrogen uptake at t  20 min corresponds to the Ti dihydride formation. The metal composition of as-milled composites was determined by ICP-OES for y=0 and 0.3. Results are gathered in Table 3.1. Minor Fe contamination ( 0.6 at. %) from milling tools was detected. For y = 0.3, there is a good agreement between the nominal Ti-atomic ratio (30 at.%) and the experimental value (29.2 at.%).

Microstructural characterization

XRD patterns were collected for all as-milled composites to determine their crystallographic properties. Results are shown in Figure 3.3. All patterns exhibit broad diffraction peaks evidencing nanostructured materials. For y = 0, the pattern can be indexed with two phases: the rutile-type β-MgH2 (S.G. P42/mnm, C.P. a= 4.506(3) c= 3.042(4)) phase and high-pressure γ-MgH2 (S.G. Pbcn, C.P. a= 4.541(1), b= 5.444(1), c= 3.973(1)) commonly found in ball milled MgH2 [3]. For the other samples, three different phases are observed: the two previous MgH2 polymorphs and ε-TiH2 (S.G. I4/mmm, C.P. a= 3.171(2), c= 4.420(1)). As expected, the intensity of the diffraction peaks belonging to the TiH2 phase increases with the Ti-molar ratio y. No formation of ternary Mg-Ti-H phases is observed.

Rietveld analysis of XRD patterns is presented in Figure 3.4. Phase amounts and their average crystallite sizes are gathered in Table 3.3. The ratio between γ and β MgH2 polymorphs remains stable (24  3 wt.%) for all composites and the TiH2-content agrees well, though slightly in excess, as compared to nominal values. As concerns crystallite sizes, they are evaluated in the interval of 5-8 nm for both β-MgH2 and γ-MgH2 polymorphs, while they range from 8 to 13 nm for ε-TiH2. Note that the crystal size of both MgH2 phases gradually decreases with TiH2 content. Based on the formation of binary MgH2 and TiH2 phases as well as on their nanostructured state, the synthetized (1-y)MgH2+yTiH2 samples are hereafter named as nanocomposites (NCs). SEM images of three different as-milled NCs (y=0, 0.0125 and 0.3) and their related particle size distribution are shown to the left and right of Figure 3.5, respectively. The particle size has been calculated by ImageJ software [4] and the particle size distribution has been fitted with a log-normal distribution (mean value and standard deviation values are inlets the graph). For pure Mg sample (y=0), mean particles size is 1.2 µm with a standard deviation of 1.9 µm. For the NCs with y=0.0125 and 0.3, the mean particle size is slightly higher (2.3 µm) with same standard deviation. Figure 3.6 shows the morphology of as-milled NCs. The particles are agglomerates formed by tiny (~100 nm) primary particles. They have flake-like shape. For pure Mg (y = 0), the agglomerates have a dense aspect and flat surface. For the Ti-containing NCs, the aspect is more porous, and the surface is rougher.

Hydrogen sorption properties

All hydrogenation experiments were performed using a manometric manual Sieverts type apparatus described in Chapter 2. About 200 mg of each milled powder was loaded into a stainless-steel sample holder placed in a tubular electrical oven. The temperature of the oven was controlled by a PID system, using a type K thermocouple.

a) H-Thermodynamics

Pressure Composition Isotherm (PCI) measurements were carried out at 548, 573 and 598 K for the NC having the highest TiH2 amount (y=0.3) to ascertain possible changes of Mg-H thermodynamics by TiH2 addition. The temperature of the sample was held constant (± 2 K). Results are displayed in Figure 3.7. PCI isotherms are characterized by flat plateau pressures extending from 2 to 5.9 wt.%H and exhibiting low hysteresis. Both (de)hydrogenation enthalpy and entropy were determined from Van't Hoff plots depicted in the inset graph. The hydrogenation and dehydrogenation enthalpies are -73.6±0.9 and 75.5±0.6 kJmolH2 -1 , respectively; while the corresponding entropy values are -133±2 and 135±1 JK -1 molH2 -1 . These data are in agreement with the reported thermodynamic properties of magnesium hydride [5]. Therefore, TiH2 additive does not modify the thermodynamic stability of magnesium toward hydrogen sorption. It should also be noted that 1.9 wt.% H remains trapped in the nanocomposite for pressures below 10 -3 MPa. This is attributed to hydrogen irreversibly stored in the TiH2 phase under the studied thermodynamic conditions. Indeed, at 300 °C and PH 2 = 10 -3 MPa, the equilibrium hydrogen content is as high as 1.95 H/Ti [6]. This implies that reversible hydrogen sorption in NCs nanocomposites (1-y)MgH2+yTiH2 at the used operation conditions is limited to the MgH2/Mg phase transformation.

b) Temperature programmed desorption and absorption

Temperature programmed desorption (TPD) characterization was carried out for assynthetized pure MgH2 and TiH2-containing NCs (Figure 3.8 a). The samples were heated up from room temperature to 350 °C at 2 °C/min in a closed volume. The desorption of pure MgH2 sample starts at 250 °C while for Ti-containing NCs it starts at lower temperatures (< 200 °C). These results indicate that Ti-addition enhances the kinetics of hydrogen desorption as compared to pure MgH2. Moreover, in agreement with the previous thermodynamic results, hydrogen desorption does not cross the thermodynamic limit of the Mg-H system depicted in Fig. 3.8 a by a continuous red line.

Temperature programmed absorption (TPA) curves were monitored for the same samples as described before after TPD measurements (Figure 3.3, b). Experiments were carried out at a hydrogen pressure of 0.8 MPa with a heating rate of 2 °C/min. For the pure Mg sample, hydrogen absorption starts at 100 °C whereas a significant absorption rate occurs for all TiH2containing NCs at room temperature. This result clearly evidences the enhancement of the hydrogen absorption properties of MgH2 with TiH2 addition. In the TPD graph, the thermodynamic limit of the Mg-H system is depicted by a red line.

c) Evolution of the reversible hydrogen capacity with cycling under constrained reaction time

To further characterize the sorption properties of the here studied nanocomposites, manometric cycling measurements over 20 cycles at a constant temperature have been performed in the same Sieverts' type apparatus. The operation temperature was fixed to 300 °C (573 K) as at this temperature the Mg-H system provides a significant plateau pressure (Pp, H 2 = 0.16 MPa) to feed a fuel cell system. In addition, H-cycling sorption experiments have been arbitrarily constrained to 15 min, a characteristic practical time for refilling hydrogen storage tanks [7]. These sorption measurements were performed at quasi-constant pressure by using large reservoirs compared to the sample holder volume. The same driving force, g, was used for absorption and desorption. This was accomplished by adjusting the applied hydrogen pressure in the system to both a higher value (for absorption) and lower (for desorption) than the plateau pressure so that g = ln(Pabs/Pp) = ln (Pp/Pdes)  1.6. For experiments at 300°C, absorption and desorption were set to Pabs = 0.8 and Pdes = 0.03 MPa. Figure 3.9 shows the evolution of the reversible capacity, Crev, on cycling at 300 (Fig. 9a) and 340 °C (Fig. 9b) for representative NCs. The purpose was to determine the influence of sorption temperature on the cycling life. At 300 °C (Fig. 9a), the pure MgH2 sample (i.e. y = 0) exhibits in the first cycle very low reversible capacity, Crev = 0.5 wt.%, with a gradual decrease after 20 cycles: Crev = 0.2 wt.%. For TiH2-containing NCs, Crev is much higher in the first cycle, Crev > 4 wt.%, and keeps rather stable on cycling, especially at high TiH2 contents (y = 0.2 and 0.3). After 20 cycles, the highest reversible H-capacity, Crev = 4.9 wt.%, is observed for the lowest TiH2 amount (y = 0.025). On increasing the temperature to 340 °C (Fig. 9b), the pure MgH2 sample exhibits a slightly higher capacity, Crev = 0.9 wt.%, which decreases down to 0.5 wt.% after 20 cycles. Meanwhile, for Ti-containing NCs the initial capacity is very high at low TiH2 amount: Crev = 5.2 and 5.8 wt.% for y = 0.05 and 0.1, respectively. However, Crev gradually decreases on cycling. In contrast, the cycle-life of NCs with the highest TiH2 amount (y = 0.2 and 0.3) keeps stable on cycling. Therefore, excluding pure MgH2, the reversible capacity increases on lowering the TiH2 amount but the cycle-life deteriorates. This behavior exacerbates on increasing the sorption temperature.

d) Evolution of hydrogen kinetics on cycling

Figure 3.10 displays the hydrogen desorption curves during the 1 st (Fig. 3.10a) and 20 th cycle (Fig. 3.10b) for doped and undoped NCs. At the used desorption pressure, Pdes = 0.03 MPa, hydrogen desorption is exclusively attributed to the MgH2 to Mg transformation. Full desorption (i.e. reacted fraction  = 1) is defined with respect to the hydrogen stored as MgH2 within the nanocomposite. For the 1 st desorption (Figure 3.10a), the slope of the desorption curves at short reaction time increases with the TiH2 amount. This reveals that H-desorption kinetics strongly increases with TiH2-content. At the end of reaction time (tend = 15 min), the MgH2 to Mg transformation is almost completed (𝛼  0.85) and reaction rates slow down for high TiH2-amount (y  0.1). In contrast, for low TiH2-amount (y  0.05) H-desorption rate is still significant at tend and the reacted fraction is lower,  = 0.7. The non-reacted fraction at tend, 1- = 0.3, is attributed to kinetically retained MgH2. Figure 3.10b shows the H-desorption curves at the 20 th cycle. By comparison to the 1 st cycle, the same beneficial effect of TiH2 is observed on H-desorption kinetics. In addition, at a given TiH2 content, H-desorption kinetics accelerate upon cycling leading. Thus, higher reacted fractions at tend are observed for all NCs at cycle 20th as compared to the 2 nd cycle. Figure 3.11 displays the H-absorption curves at the 2 nd and 20 th cycle. Note that 1 st hydrogen absorption is assigned to the RBM sample synthesis. Absorption curves correspond to the reversible transformation of previously reformed Mg into MgH2. The full reacted fraction with α = 1 is defined as the quotient between the amount of hydrogen absorbed in the sweep compared to the total hydrogen stored as MgH2 within the nanocomposite. Thus, values lower than 1 indicate that some Mg remains unreacted. Reaction rates for the 2 nd cycle (Fig. 3.11a) are faster than for the 1 st desorption (Fig. 3.10a) with most of the transformed fraction occurring at t  3 min. However, at longer reaction times, the absorption rate slows down, and α approaches the values of the first desorption. H-absorption kinetics at short time does not change significantly on cycling: H-absorption curves are very similar at the 2 nd and 20 th cycle except for y = 0.1 and 0.2, which hardly saturate at tend = 15 min.

e) Evolution of hydrogen amount on cycling

Based on the previous thermodynamic and kinetic data, the hydrogen content in all NCs has been evaluated for each absorption and desorption sweep and results are displayed in Figure 3.12. The hydrogen content can be split into four contributions (from bottom to top in each graph): the hydrogen irreversibly stored in TiH2, the hydrogen retained as MgH2 (i.e. not desorbed at t  tend), the hydrogen reversibly stored through the MgH2/Mg phase transformation, Crev, and the unreacted Mg (i.e. not rehydrogenated at t  tend). Details for selected compositions are given in the following.

For pure MgH2 sample (y = 0), very little hydrogen is desorbed in the first cycle. Thus, in the desorption state, much of the material (~94 %) remains as MgH2. On the 2 nd absorption, most formed Mg reverts to MgH2 in the absorbed state, but the reversible capacity is extremely low: Crev ≈0.5 wt. %. Then, only a little amount of hydrogen desorbs/absorbs and much of the sample remains as MgH2 on cycling.

For NC y = 0.0125, i.e. with a little amount of TiH2 addition (irreversible part), drastic changes are observed: the reversible capacity Crev increases more than tenfold (≈6 wt %) thanks to a large diminution of retained MgH2. However, a significant amount of unreacted Mg occurs. In addition, it should be noted that the amounts of both unreacted Mg and retained MgH2 gradually increase through cycling and consequently Crev decreases. After 20 cycles, the reversible capacity is 4.7 wt.%. For NCs 0.025 ≤ y ≤ 0.1, during desorption not only some hydrogen from MgH2 has not been desorbed but also part of it is not absorbed to reform MgH2 during cycling. Therefore, there is an increase in unreacted Mg and the amount of MgH2 decreases with cycling. As a result, the Crev is gradually decreasing only during the first cycles due to the increasing amount of unreacted Mg. It is important to mention that Crev decreases also because the TiH2 which is irreversible at this operation conditions.

For NCs with y ≥ 0.2, the amount of both unreacted Mg and retained MgH2 decreases drastically, showing that the Mg/MgH2 phase transformation is practically complete. Furthermore, the reversible capacity Crev keeps constant on cycling. One should note that the contribution of hydrogen stored as irreversible TiH2 is high: 1.2 and 1.7 wt.% for y = 0.2 and 0.3, respectively.

Effect of hydrogen cycling on the microstructure of MgH2-TiH2 nanocomposites a) Crystallographic properties

XRD analysis was performed for all NCs after twenty sorption cycles in absorbed and desorbed states. XRD patterns are shown in Figure 3.13 and relevant crystallographic data from Rietveld analysis are gathered in Table 3.3. Diffraction patterns can be indexed in three phases: β-MgH2, ε-TiH2, and Mg. The phase γ-MgH2, detected in as-synthetized samples, is fully transformed into β-MgH2 after sorption experiments as a result of thermally induced strain relaxation [8]. As concerns ε-TiH2, its amount and crystallinity both in absorbed and desorbed states do not change significantly as compared to as-synthetized NCs. In contrast, MgH2 undergoes drastic changes in content and crystal size. Taking apart TiH2, MgH2 is the main phase (>85 wt.%) in the absorbed state and the minor one (< 15 wt.%) in the desorbed state. Pure MgH2 (y = 0) is not considered as it exhibits extremely low reversibility on cycling staying almost fully hydrogenated in the desorption state. For Ti-containing samples, the MgH2 amount variation between absorption and desorption reflects the reversibility of the MgH2/Mg transformation. The extent of this transformation is shown in Figure 3.14 as determined by XRD analysis as well as from reacted fractions in H-sorption experiments (Fig. 3.11b). A good agreement is found between XRD and kinetic data, which validates the hypothesis that incomplete Mg/MgH2 transformation is due to kinetic effects. As evidenced in Figure 3.14, the reversibility of the MgH2/Mg transformation increases with Ti-amount.

It is also worth to comment on crystallite size changes of MgH2 after H-sorption experiments (Table 3.3). Indeed, significant grain coarsening occurs after 20 cycles. For β-MgH2, the crystallite size enlarges from 6±2 to 45±15 nm between the as-milled and cycled samples. Such enlargement is mainly attributed to thermally driven effects (holding temperature at 300 °C) and not to cycling between Mg and MgH2 phases. Indeed, pure MgH2 sample, which does not desorb hydrogen significantly, undergoes similar grain growth. Also, from this table, it can be noted that, as the amount of TiH2 is increasing, the crystal size of both β-MgH2 and Mg after cycling is decreasing. It is reported that the presence of the TiH2 phase minimizes grain coarsening of both β-MgH2 and Mg phases [1]. a)

b) Particle size and morphological properties

Morphological studies after cycling were carried out by SEM for three representative NCs (y = 0, 0.0125 and 0.3) in both absorbed and desorbed state. SEM images and particle size distribution histograms are shown in Figure 3.15 and 3.16 for desorbed and absorbed states, respectively. For all NCs, particle size distributions obey log-normal distribution: mean and standard variations are shown in histograms insets and gathered in Table 3.4. The size of agglomerates is similar to absorption and desorption as well as compared to the as-milled state and has no significant dependence with TiH2 content. Indeed, the particle size is comprised between 1.2 and 2.9 µm with a standard deviation of 2 µm Figure 3.17 shows the morphology of the NCs after cycling in both absorbed and desorbed states. For y = 0, the agglomerates are formed by primary tiny slate-like particles both in the absorbed and desorbed states, which differs with the as-milled one, for which a flat surface was observed. For Ti-poor NC (y = 0.0125), the agglomerate is dense in the absorbed state with tiny primary particles stick on it. The morphology resembles that of the as-milled one. In contrast, the morphology of the desorbed state shows the formation of protrusions [9]. For Ti-rich NC (y =0.3), both absorbed and desorbed states show the same morphology. It is a dense agglomerate with cracks and small particles stick to it. The morphology resembles that of the as-milled sample but with cracking. This cracking could be attributed to the repeated volume change of the dense agglomerates upon repeated absorption and desorption cycles. For the as-milled composite, Ti and Mg contents are similar at the center and edge of the agglomerate and close to the nominal value (30 at.% Ti). This reflects that there is a homogeneous dispersion of the MgH2 and TiH2 phases at the spatial resolution of EDX analysis ( 1 m). This agrees with microstructural data obtained by Transmission Electron Microscopy by Berti et al. on RMB MgH2-TiH2 NCs [10]. These authors found that MgH2 and TiH2 phases are homogenously distributed at the nanoscale.

For the cycled composite, similar atomic compositions are found at the agglomerate edges in both absorption and desorption states. They concur with the nominal value (30 at. % Ti). However, much lower Ti-contents are observed at the center of agglomerates. This reflects a partial migration of TiH2 from the center to the edge during the cycling process.

Analysis of hydrogen sorption kinetics

As largely shown in this chapter, TiH2 plays a major role in the reaction kinetics of hydrogen sorption in Mg. As shown in Fig. 3.10 and 3.11, isothermal sorption curves at 300 °C have been measured for different TiH2 contents and at different cycling stages. Using these data, this section provides a kinetic analysis aiming to identify which kinetic model fits the best sorption curves to determine the controlling mechanism of the reaction.

A model is a theoretical, mathematical description of what occurs experimentally. In solid-state reactions, a model can describe a particular reaction type and translate that mathematically into a rate equation. Many models have been proposed in solid-state kinetics, and these models have been developed based on certain mechanistic assumptions. Other models are more empirically based, and their mathematics facilitates data analysis with little mechanistic meaning. Therefore, different rate expressions are produced from these models.

In homogeneous kinetics (e.g., gas or solution phases), kinetic studies are usually directed toward obtaining rate constants that can be used to describe the progress of a reaction. Additionally, the reaction mechanism is typically investigated and rate constant changes with temperature, pressure, or reactant/product concentrations are often helpful in uncovering mechanisms.

These mechanisms involve to varying degrees the detailed chemical steps by which a reactant is converted to product. However, in solid-state kinetics, mechanistic interpretations usually involve identifying a reasonable reaction model because information about individual reaction steps is often difficult to obtain [11].

In a typical hydrogenation experiment, the reacted fraction, 𝛼, i.e. the amount of metal that becomes hydride in absorption or the amount of hydride that transforms to metal in desorption is recorded as a function of time, 𝑡, sample temperature, 𝑇, H2 pressure, 𝑃 𝐻2 , and for different sample shape. In most cases, the relationship between all the parameters takes the simple form: being 𝑘 0 a pre-exponential factor, 𝐸 𝑎 the activation energy of the overall reaction and 𝑅 the gas constant. Experimental data are analyzed according to the former expressions. The functional form of function 𝐹() which best fits the data, together with the pressure, temperature and geometry dependence of the reaction rate are used to determine the ratecontrolling step for the reaction [12]. Several reaction models 𝐹() are listed in Table 3.16. Three-dimensional diffusion (Jander equation)

[1 -(1 -𝛼) 1 3 ⁄ ] 2 D4 Three-dimensional diffusion (Ginstling-Braunshtein equation) (1 - 2𝛼 3 ) -(1 -𝛼) 2 3 ⁄ Geometrical contraction models R2 Two-dimensional phase boundary 1 -(1 -𝛼) 1 2 ⁄ R3 Three-dimensional phase boundary 1 -(1 -𝛼) 1 3 ⁄

Nucleation and growth models F1

First-order reaction

-ln (1 -𝛼) A2 Avarami-Erofe'ev [-ln (1 -𝛼)] 1 2 ⁄ A3 Avarami-Erofe'ev [-ln (1 -𝛼)] 1 3 ⁄ A4 Avarami-Erofe'ev [-ln (1 -𝛼)] 1 4 ⁄
Khawam and Flanagan [11] presented kinetic models and their classification based on the graphical shape of their isothermal curves (α vs t or dα/dt vs α) according to their mechanistic assumptions in Table 3.6. Based on their shape, kinetic models can be grouped into deceleratory, or sigmoidal models (Figure 3 The function 𝐹() offering the best fit to the sorption data is considered as that for which the plot 𝐹() vs. t offers the coefficient of determination R 2 closest to 1. Table 3.7 presents the kinetic models that offer the best fits to the data extracted from all kinetic experiments. Comments to these results are given in the following. For the first desorption, the pure MgH2 sample is best fitted with R2 model, but the F1 model is also a good fit with a close R 2 coefficient. For Ti-containing NCs, three different models offer the best fits: A2 for y = 0.0125, R2 for y = 0.05 and 0.2 and F1 for y = 0.1 and 0.3. However, when F1 is not the best fit, this model cannot be discarded as R 2 remains close to 1. Therefore, the F1 model fits reasonably well all desorption data in the first cycle. For the 20 th desorption cycle, very similar results are observed leading to the conclusion that the F1 model is the best candidate to fit all desorption data.

D1 D2 D3 D4 F1 R2 R3 A2 A3 R 2 =0,9815 R 2 =0,9712 R 2 =0,9569 R 2 =0,9667 R 2 =0,9990 R 2 =0,9918 R 2 =0,9947 R 2 =0,8927 R 2 =0
D1 D2 D3 D4 F1 R2 R3 A2 A3 R 2 =0,9810 R 2 =0,9709 R 2 =0,9572 R 2 =0,9665 R 2 =0,9990 R 2 =0,9938 R 2 =0,9960 R 2 =0,9245 R 2 =0,8214 F() Time (seg)
As for absorption data, the F1 model fits the best all sorption curves in the 2 nd cycle for Ti-containing composites. In contrast, for the pure MgH2 sample (y = 0), the D3 model fits better than the others. For the 20 th cycle, absorption curves of the pure MgH2 sample and some Ti-containing NCs are best fitted with model F1 (y = 0.0125, 0.025 and 0.3) while others fit the best to other models: D4, D3 and D1 for y = 0.05, 0.1 and 0.2, respectively. However, the fit of sorption curves to F1 model is always reasonably good.

To summarize, the F1 model offers a good fit for most of the kinetic sorption data obtained in this study. This model corresponds to a first order nucleation and growth model where both nucleation and the subsequent growing of the nuclei, are relevant to understand the kinetic. First, hydride phases are nucleated close to the TiH2 (catalytic sites) (Figure 3.21 (a)) and the nuclei grow (b). As nuclei grow bigger, they overlap one another (c) to form the hydride layer at the surface of a particle (d). When the hydride layer is thickened, the diffusion distance of the hydrogen atom increases and the hydrogen diffusion through the hydride phase gradually affects the hydriding reaction rate. Fernandez and Sanchez [12] explained that, even though there are a great number of papers published about the kinetics of hydrogen-Mg reaction [3,[16][17][18][19], there is not a proper understanding yet of the slowest step of sorption process. Some authors mention that nucleation and growth mechanism should be taken into account for data analyzing [3,12,[20][21][22][23][24], while others have reported that nucleation is insignificant for the kinetics and only the growth of the phase is relevant [16,18]. Zhou et al. studied the MgH2 system with several Tibased metals [20]. They found the F1 model was perfectly fit with kinetic nucleation and growth process in both absorption and desorption reactions. The MgH2-TiH2 system was studied by Li et al [22], they tested several kinetic models concluding that F1 is the most accurate among them to describe the hydrogenation of MgH2-0.1TiH2. Zhang et al [25] exposed that R3 model has a good fit for the dehydrogenation of MgH2/NbHx composites. Fernandez and Sanchez [12] studied the hydrogen kinetic absorption and desorption of Mg, in both cases data seems to obey a rate equation based on nucleation and growth.

Discussion a) MgH2 system

MgH2 hydride powder was successfully synthesized by RBM reaching its nominal capacity of 7.6 wt.% at the end of milling. This is supported by the hydrogen uptake curve where the same storage value was determined. Minor Fe content (0.3 at.%) was detected: a common contamination found in milled samples. Fe is coming from the milling jar, because of the collision of the stainless steel balls with the wall jar. The kinetics of hydrogen absorption in Mg by RBM close to RT is remarkably enhanced as compared to classical solid/gas hydrogenation of Mg at 300°C. The latter takes more than 10 hours while the former need less than one hour using in both cases micrometric Mg. This was presented by Vigeholm et al, they showed that full hydrogenation of 90 m in diameter of Mg particles requires about 1 day at 400 °C and PH 2 = 2 MPa [26].

By XRD it was possible to identify the two phases contained in the sample, the rutile-type β-MgH2 phase, and γ-MgH2. These phases were also identified by Huot et al in ball-milled magnesium hydride [3]. From Rietveld analysis, the abundance of each phase was determined as β-80 wt. % and γ-20 wt. %. This agrees with previous studies. Both phases have a crystal size of ~ 8 nm in close agreement with previous RBM synthesis of MgH2 [1]. The morphology and particle size were determined by SEM. MgH2 forms agglomerates with a mean size of 1.2 µm. The agglomerates have a dense aspect and flat rounded surface.

TPD experiments showed that as-milled MgH2 starts its hydrogen desorption at 280 °C whereas TPA shows that hydrogen absorption starts at 120 °C. These temperatures do not contradict the thermodynamics of the Mg-H system as it was shown in Chapter one.

MgH2 synthetized by RBM exhibits a poor reversible capacity for reaction time limited to 15 min. This is mainly due to its slow desorption kinetics even at the low initial nanometric crystallite size (8 nm). At 300°C and back-pressure of 0.03 MPa, only 7% of its total capacity is desorbed in 15 min.

After the cycling treatment, the MgH2 sample was analyzed by XRD both in the absorbed and desorbed state. Only β-MgH2 phase could be clearly detected. Its crystal size enlarged to 35 ± 5 nm. SEM experiments were done to characterize particle size as well as the morphology. No change in particle size was observed. However, after cycling, surface morphology becomes rougher. The lack of strong changes in particle morphology could be attributed to the poor hydrogen exchange on cycling. In contrast, crystal size growth may explain the increase on surface roughness by the formation of faceted crystals. Crystal growth is attributed to thermal annealing, and not to cycling effects, by holding the sample at 300 °C.

b) MgH2-TiH2 system

MgH2-TiH2 NCs with different Ti-contents were successfully synthesized by reactive ball milling. The measured hydrogen absorption capacity was the same or close to the nominal ones. A slight contamination of Fe (0.6 at.%) was also observed.

From XRD it was found that the synthetized materials consist of immiscible MgH2 and TiH2. No ternary Mg-Ti-H hydrides were found. All patterns can be indexed in three phases: β-MgH2, γ-MgH2, and ε-TiH2. From Rietveld analysis, the crystal size was measured about 8±3 nm for each phase. Both phase abundance and crystal size agree with the published data [1,27,28]. From SEM, it was determined that MgH2-TiH2 NCs consist of agglomerates 2.3 µm in size of ~ 100 nm in size particles and exhibit porous aspect and rough surface.

From PCI measurements, the hydrogenation and dehydrogenation enthalpies of the Mg-MgH2 transformation were measured as -73.6 ± 0.9 and 75.5 ± 0.6 kJmolH2 -1 , respectively. The corresponding entropy values are -133±2 and 135±1 JK -1 molH2 -1 in agreement with literature data for the bulk Mg-H system as was presented in Chapter one. These results can also be compared with previous studies on MgH2-TiH2 NCs. Lu et al. obtained for the system MgH2-0.1TiH2 the hydrogenation enthalpy and entropy values of -68.2 kJmolH2 -1 and -127 JK -1 molH2 - 1 respectively [27]. Patelli et al calculated enthalpy and entropy values of -68.2 kJmolH2 -1 and -119 JK -1 molH2 -1 respectively for the desorption of the MgH2-TiH2 system [29]. Cui et al, also calculated the desorption enthalpy and entropy of the same system (enthalpy -75.1 kJmolH2 - 1 and entropy -136.3 JK -1 molH2 -1 ) [30].

TPD curves showed that doped NCs start desorbing at around 150 °C; i.e. 130°C lower than for pure MgH2. From TPA curves are observed that the hydrogen absorption even starts at room temperature, showing the improvement of hydrogen sorption kinetics in the Mg-H system by TiH2 additives. Nevertheless, both TPD and TPA results do not contradict the thermodynamics of the Mg-H system The reversible capacity measured of doped NC is extremely good. NCs with high (y≥0.1) TiH2 content shown stable Crev through cycling. NCs with low additive amount (y≤)0.05 exhibit fluctuating Crev. Lu et al performed eighty complete absorption-desorption cycles of one hour each, any loss of capacity was observed in the system MgH2-0.1 TiH2 [27].

From the kinetic desorption curves, it was identified that hydrogen desorption is accelerated through cycling. In contrast, absorption kinetics does not improve on cycling and likely slow down since the amount of unreacted Mg increases (Fig. 3.12). As a result, NCs with low TiH2 content are not able to absorb and desorb the initial hydrogen amount during the cycling process having both unreacted Mg and retained MgH2. NCs with high TiH2 content have less unreacted Mg and less retained MgH2. However, for these NCs reach in TiH2 its reversible capacity is not higher than the ones with less additive. Due to if TiH2 amount increases, also the irreversible part does.

After cycling, the XRD results show that there is an increase in the crystal size in the MgH2 and Mg phases for all NC in both absorbed and desorbed states. The crystal size of MgH2 is within the range from 45 to 80 nm, whereas that of Mg is comprised between 65 and 105 nm. It is however remarkable that despite such a significant crystal growth as compared to initial values (8 nm), desorption kinetics do not slow down on cycling. This clearly indicates that H-diffusion in MgH2 or Mg crystallites is not the rate limiting step for hydrogen desorption in MgH2-TiH2 NCs. In contrast, as discussed later, grain coarsening may contribute to the observed increase of unreacted-Mg (absorption sweeps) on cycling (Fig. 3.12). SEM-EDS studies were performed to determine particle sizes, morphology and phase distribution. No significant changes on particle size were observed in NCs on cycling. Therefore, this property is not at the origin of kinetic modifications neither induced by TiH2 content nor by cycling. The morphology of y=0.0125 in absorbed state is like the one as-milled showing a densification which it could be because of the poor cyclability of this. In desorbed state shows the formation of protrusions in its morphology which indicates that even with a poor desorption the morphology is modified. y=0.3, in both absorbed and desorbed state, shows a main cracked particle, this cracking could be attributed to the stress suffered by the particles during the almost complete absorption-desorption cycling.

As concerns phase distribution, SEM-EDS analysis for NC y=0.3 shows homogenous phases dispersion for the as-milled state. After cycling, the TiH2 content at the center of the agglomerates is found to be depleted as compared to the average composition. This is assigned to a partial migration of TiH2 from the core of the agglomerate to the edges on cycling. This migration has been previously observed by Zhou et al. [20,31]. The catalytic additives suffer repeating stresses from the surrounding Mg/MgH2 matrix during the cycling process. The stress might be strong enough so that a "migration" of the catalyst particle is possible (Figure 3.22). Also, the catalyst particle that closely contacted to the Mg/MgH2 matrix might be detached from the agglomerate due to the shrinkage during dehydrogenation.

These migration and detachment phenomena may explain two hydrogenation facts observed in this thesis: i) as for migration, the acceleration of dehydrogenation kinetics on cycling for all NCs may result from the increased amount of TiH2 catalytic sites at the agglomerate surface ii) as for detachment, the diminution of reversible capacity in the first cycles for Ti-poor NCs may result from the detachment of TiH2 additives from the matrix. Thermodynamics of the Mg-H system is not altered by TiH2-addition and only the MgH2 counterpart of the nanocomposite can be reversibly ab(de)sorbed at the working temperature, T = 300 °C, and pressure, 0.01 < PH2 < 1 MPa, conditions (Figure 3.6). This was confirmed by the similar obtained values of enthalpy and entropy formation than for pure MgH2. Indeed, under these thermodynamic conditions, no significant hydrogen desorption from TiH2 phase occurs [6]. TPD curves show that the additive facilitates desorption at lower temperature but without crossing the thermodynamic limit. Also, absorption can occur at much lower temperature, even room temperature, for doped NCs.

It has been reported by some researchers that MgH2 system mixed with additives showed very good stability when cycling at a temperature of 300 °C [32,33], which was confirmed comparing cycle life of the NCs cycled at 300 and 340 ° C. Isothermal cycling for 15 min shows that is possible to get a good reversible capacity adding TiH2 to the MgH2, without this the reversible is practically null. In Figure 3.12, it can see that Ti addition increases the weight of the metallic counterpart of the composite since Ti is heavier than Mg which it can be clearly in NCs with the more TiH2 amount. This is evidenced in Table 3.1, where the Hcontent of (1-y)MgH2+yTiH2 NCs gradually decreases from 7.7 to 6.0 wt.% H for y = 0 and 0.3, respectively.

Absorption and desorption kinetics are both accelerated with the incorporation of TiH2 additive to MgH2 in the system. TiH2 has beneficial kinetic effects on the temperature onset for hydrogen absorption and desorption as displayed in TPA and TPD measurements (Figure 3.8). Also, Ti-free MgH2 sample brings almost negligible desorption at 300°C in 15 min time compared to Ti-doped ones (Figure 3.10). Moreover, sorption rates on absorption (Figure 3.11) and desorption (Figure 3.10) gradually increase with TiH2 content. Consequently, reacted fractions at the end of cycling sweeps increase with TiH2 content both on absorption and desorption. As results of these kinetics improvements, the reversible capacity Crev of MgH2-TiH2 NCs is much higher than that of pure MgH2 system (Figure 3.9).

The kinetics analysis of sorption curves shows that the F1 model fits the best to both absorption and desorption experimental data. This model corresponds to a first order reaction where both nucleation and the subsequent growing of the nuclei describe sorption kinetics. Figure 3.23 shows the nucleation and growth kinetic model proposed by Ponthieu [34] where nucleation would occur at the TiH2/MgH2 interfaces. Moreover, TiH2 if present at the particle surface, can act as a gateway for hydrogen gas facilitating hydrogen dissociation (absorption) and recombination (desorption).

Moreover, TiH2 can also facilitate hydrogen transfer to the Mg phase where MgH2 nucleates. This is supported by the large diffusion coefficient of H in TiH2 [35]. In addition, it is proposed that fast H-absorption is possible in nanocomposites because native MgO barriers are bypassed through the TiH2 phase. This effect seems to be even more important on desorption sweeps (Fig. 3.10): recombination of hydrogen atoms at TiH2 surfaces will enhance hydrogen desorption kinetics in MgH2-TiH2 nanocomposites as compared to the pure MgH2 system. From the analyzed microstructure by XRD, it has been confirmed that crystallite sizes of Mg and MgH2 increase up to 100 nm after the cycling process. In contrast, the TiH2 size remains low and constant around 10 nm. Several authors have proposed that additive inclusions may act as a grain growth inhibitor of Mg and MgH2 phases [20,24,27,31]. Our study shows that significant grain growth exists after 20 sorption cycles. However, this has no detrimental effect on desorption kinetics but may be at the origin of reduced hydrogen absorption on cycling (i.e. increase of unreacted Mg phase). Indeed, as in bulk Mg [12], the growth of MgH2 nuclei during absorption is likely limited by hydrogen diffusion through MgH2 crystallites.

d) Additive contributions

Ti acts as a gateway for hydrogen transport enabling fast hydrogen dissociation/recombination at the TiH2 surface and its diffusion towards the bulk as schematically depicted in Figure 3.24. Different authors have reported that TiH2 can improve the hydrogen sorption kinetics in MgH2. For instance, Patelli et al [29] suggest that TiH2 provides catalytic active sites for surface adsorption and dissociation of H2. Hydrogen diffusion is faster in TiH2 than in MgH2, and therefore the TiH2 nanocrystallites may provide high mobility paths for atomic hydrogen. A similar explanation was given in previous studies performed in our laboratory [1] proposing that TiH2 is the gateway for the transfer from the gas phase to the bulk. In addition, Jangir et al [36] showed that activation energies for hydrogen desorption from MgH2 gradually decrease with TiH2 content. Within the framework of our study, we have demonstrated that TiH2 addition favorably enhances H-sorption kinetics of the Mg/MgH2 reaction and cycling stability. The more TiH2 is added, the more Mg/MgH2 transformation is accomplished in a given time. This can be clearly observed in Figure 3.12: the higher the TiH2 amount, the more hydrogen can be reversibly stored through the Mg/MgH2 reaction. However, as a drawback of TiH2 addition, it decreases the reversible hydrogenation capacity of Mg/MgH2 system because of its higher atomic weight (as compared to Mg) and the formation of irreversible TiH2. Thus, it exists an optimum TiH2 content in MgH2-TiH2 nanocomposites to achieve the highest reversible capacity. It was established here at 0.975MgH2-0.025TiH2, for a limited reaction time of 15 min, leading to a reversible capacity of 4.9 wt.% after 20 sorption cycles.

To summarize, the relative contributions to hydrogen storage in the two-phase (1-y)MgH2+yTiH2 system are presented in Figure 3.25. Starting from pure MgH2 (y = 0; 7.7 wt.%), the total hydrogen capacity decreases with titanium amount due, firstly to its heavier atomic mass (hatched blue area in Figure 3.25). Secondly, as hydrogen in TiH2 is not reversible under the working thermodynamic conditions, the reversible hydrogen stored in MgH2 decreases down to 4.3 wt.% for y = 0.3 (hatched red area in Figure 3.25). From then, the hydrogen that can be reversibly recovered from MgH2 depends on the H-kinetic properties of the NCs. Without Ti (y = 0), no significant H-amount is recovered from pure MgH2. When increasing the Ti-content, such kinetic limitation is minimized, and an optimized value is obtained for the reversible H-capacity of NCs (grey area in Figure 3.25). The fastest kinetics are found at the highest Ti-content (y = 0.3) leading to the highest reversibility in the Mg/MgH2 reaction (white area in Figure 3.25). However, the negligible hydrogen desorption of TiH2 restricts the interest of using high Ti-contents. The best compromise between the detrimental effects of increasing TiH2 content (heavy and irreversible hydride) and its beneficial one (enhancement of Mg/MgH2 kinetics) leads to an optimum for y = 0.025 with a reversible capacity of 4.9 wt.%.

Effect of hydride-forming TM additives (TM = Ti, Zr, Sc, Y, V, and Nb) on sorption properties of MgH2-based nanocomposites

The hydrogen sorption properties of MgH2-TMHx nanocomposites for different transition metals (TM= Ti, Zr, Sc, Y, V, and Nb) are analyzed in this chapter. As shown in the previous chapter, there are two ways to successfully improve the MgH2 kinetics through RBM: by nanostructuration and by additives incorporation. Once that MgH2-TiH2 system was deeply investigated, arose the idea of exploring the effect of other transition metals with properties to close Ti for a better understanding of the additive role. Thus, neighbor elements to Ti in the periodic table were chosen, i.e. metals of 3, 4 and 5 groups and 4 th and 5 th periods. In the previous chapter, we found out an optimum amount of TiH2 (2.5 mol%) to have both high Crev and good cycling in a restricted time of 15 minutes. In this chapter, we decided to fix the TM amount to 5 mol% to maintain fast sorption kinetics, while using a significant amount of additive for facilitating characterization purposes (e.g. XRD-studies).

MgH2 with 4 transition metals Ti and Zr

Synthesis and structural characterization

a) System 0.95MgH2-0.05TiH2 Figure 4.1 shows the RBM synthesis of the MgH2-TiH2 composite. An initial pressure closes to 90 bar (9 MPa) at room temperature were the initial conditions for the RBM. As in the previous chapter, the same 5g total mass was used to perform the RBM. The milling program consists of 2 milling cycles of 120 min each with one pause between them of 120 min. At the beginning of the process, the temperature increases because of the frictions and collisions of the stainless-steel balls with the vial walls reaching a maximum value of 50 °C (323 K). This also leads to an increment of the pressure (95 bar (9.5 MPa)) due to gas confinement. Then, the pressure drops down (65 bar (6.5 MPa)) due to the hydrogen absorption of both Mg and Ti metals during the formation of the hydride phases. Once that hydrogenation is completed, a constant pressure is observed. At the end of the first milling cycle, both pressure and temperature decrease to equilibrium values at ambient conditions during the pause time, P= 56 bar (5.6 MPa) and T= 24 °C (297 K). The second cycle did not show any additional hydrogen absorption, is the change in pressure due to thermal variations induced by milling and the pause Pmax = 63 bar (6.3 MPa) and Tmax= 50°C (323 K). Thus, the second cycle can be used to calibrate the gas temperature that differs from the measured vial temperature [1]. From both pressure and temperature data of the first milling process, the corresponding hydrogen uptake curve was obtained (Figure 4.2). The figure shows the hydrogen uptake during the milling process in both weight percent (wt.%) and hydrogen atoms per formula unit (H/f.u.) scales. The evolution of the hydride phases formation during the first 60 min can be seen then, at the end of 120 min, the maximum uptake capacity is been achieved. Cuevas et al. [2] explained that the formation of the nanocomposite MgH2-TiH2 takes place in two-distinct stages due to the consecutive formation of TiH2 and MgH2, during the fast (0 < tm < 15 min) and long (15 < tm < 50 min) milling times, respectively (Figure 4.2). The maximum value of CH is 1.8 H/f.u. (or 6.8 wt.%) which is close to the one of the targeted material 0.95MgH2+0.05TiH2, i.e. 2.0 H/f.u. or 7.3 wt.% respectively. XRD measurements were used to determine the crystallographic properties of the obtained nanocomposite (NC). Figure 4.3 displays the XRD pattern of as-milled material. The pattern (black line) exhibits broad diffraction peaks evidencing nanostructuration. Three different phases were indexed: two MgH2 polymorphs (β and ) and fluorite-distorted ε-TiH2. The XRD pattern was refined by the Rietveld method [3]. The calculated pattern is shown in Figure 4.3. (red line). Table 4.1 gathers phase amounts and their average crystallite sizes. The ratio between γ and β MgH2 polymorphs is 24  3 wt.%. TiH2-content agrees well as compared to the nominal value. As concerns crystallite sizes, they are evaluated in the interval of 5-7 nm for both β-MgH2 and γ-MgH2 polymorphs, and 8 nm for ε-TiH2. It is important to mention that no ternary Mg-Ti-H phase was detected during the synthesis. Figure 4.5 displays the hydrogen uptake during the milling process in weight percent (wt.%) and hydrogen atoms per formula unit (H/f.u.). The Zr dihydride is formed in less than 20 min. Then, 90 % of the MgH2 phase formation is obtained after 33 min. Finally, the maximum uptake capacity is achieved after 120 min. The maximum values of CH are 1.9 H/f.u. and 6.7 wt.%, which agrees with the nominal ones. The XRD pattern of as-synthetized 0.95MgH2+0.05ZrH2 is shown in Figure 4.6. Broad diffraction peaks are detected evidencing nanostructured materials. Three different phases are observed: the two β-and γ-MgH2 polymorphs and ε-ZrH2. Table 4.2 gathers phase amounts and average crystallite sizes determined by the Rietveld method [3]. The most abundant phase is β-MgH2, followed by γ-MgH2 and finally ε-ZrH2. The crystallite size is identical for all phases (7 nm). No ternary Mg-Zr-H phase was detected. 4.7 shows the PCI isotherms at 300 °C (573 K) of 0.95MgH2-0.05TiH2 (left) and 0.95MgH2-0.05ZrH2 (right) nanocomposites. For Ti, the PCI is characterized by a flat plateau pressure at 0.16 MPa extending from 0.34 to 6.5 wt.%H and exhibiting low hysteresis. For the Zr case, the PCI isotherms show a flat plateau at 0.14 MPa from 0.31 to 6.45 wt.%H. The value of the plateau pressure for both NCs is very close to that of pure MgH2 (0.16 MPa). This MgH2 value is obtained by the reported thermodynamic properties of magnesium hydride the hydrogenation and dehydrogenation enthalpies (74.1 and 75.8 kJ mol H2 -1 ), and entropies (133.2 and 135.1 JK-1molH2 -1 ) [6].

b) Evolution of the reversible hydrogen capacity with cycling under constrained reaction time

MgH2-TiH2 and MgH2-ZrH2 nanocomposites have been cycled under a hydrogen atmosphere at 300°C. H-cycling sorption experiments have been arbitrarily constrained to 15 min, a characteristic practical time for refilling hydrogen storage tanks [7]. These sorption measurements were performed using a concept of quasi-constant pressure and equal driving force for absorption and desorption sweeps. The pressures were fixed to Pabs=8 bar (0.8MPa) and Pdes= 0.03 bar (3x10 -3 MPa) for an equilibrium plateau pressure of Pp = 0.16 MPa. Twenty manual absorption-desorption cycles were done.

c) Hydrogen sorption curves

Figure 4.9 displays the hydrogen sorption curves for MgH2-TiH2 and MgH2-ZrH2 nanocomposites at initial (1 st desorption, 2 nd absorption) and final (20 th absorption and desorption) sorption cycles. Since both TiH2 and ZrH2 are stable hydrides (as it has been explained in Chapter one) under our operation conditions, H-sorption curves are exclusively attributed to MgH2 to Mg transformation (desorption) and the reverse reaction (absorption). Full MgH2-desorption (i.e. reacted fraction, 𝛼 = 1) is defined with respect to the nominal Hcontent stored in MgH2 in the NCs. For the Ti-containing NC, the reaction fraction at the constrained time (15 min) is limited to 60% for the 1 st desorption (i.e. 40% of MgH2 is retained in the NC). On the 2 nd absorption, the absorbed amount is a little more than the half available, meaning that a part of Mg is not reacting to form MgH2. At the 20 th desorption, reaction kinetic is faster than in the initial cycles with α  0.80. Meanwhile, the absorption looks similar, it can be noted that there is a slight acceleration in the kinetic after 20 cycles.

For the Zr-containing NC (right Fig 4.9), the 1 st desorption shows faster kinetics than for the Ti-case with ca. 85% of MgH2 decomposed in 6 min. Moreover, desorption kinetics accelerate on cycling with α = 0.85 in 1.5 min for the 20 th desorption. In contrast, the 2 nd absorption is rather fast (α = 0.85 in 2 min) but deteriorates on cycling for a reacted fraction over α = 0.4. At the end of the 20 th desorption, the reaction fraction attains 60% only and a large amount of Mg should be unreacted.

d) Evolution of hydrogen amount in cycling

Figure 4.10 shows the hydrogen content in MgH2-TiH2 and MgH2-ZrH2 nanocomposites after absorption and desorption sweeps and their evolution on cycling. For the Ti-containing NC, the desorption is not completed during the first cycle (ca. 2 wt.% not desorbed) and some MgH2 is retained. The amount of MgH2 retained gradually decreases on cycling (ca. 1 wt.% in cycle 20). This result concurs with the acceleration of desorption kinetics on cycling (Fig. 4.9). As for the absorption state, it contains more hydrogen in the first cycles, while it gradually decreases on long cycling so that the amount of unreacted Mg steadily increases. Consequently, the amount of reversible hydrogen storage keeps constant on cycling at ca. 5 wt % in agreement with Fig. 4.8. One can notice that the quantity of hydrogen stored in TiH2 is supposed to be constant (ca. 0.34 wt.%).

For the Zr-containing NC (right Fig 4 .10), the hydrogen content strongly decreases for the first four cycles, which concurs with the slowing down of absorption kinetics on cycling (Fig. 4.9. right). Such a decrease leads to a high amount of unreacted Mg. An opposing effect is observed during desorption. The amount of retained MgH2 decreases on cycling due to an acceleration on hydrogen desorption on cycling (Fig. 4.9. right). After 20 cycles, Crev is 3.8 wt.%, which is lower than for the Ti-containing NC, losing the most amount of capacity during the hydrogenation.

Effect of hydrogen cycling on the microstructure of MgH2-TiH2 and MgH2-ZrH2 NCs

XRD patterns of MgH2-TiH2 and MgH2-ZrH2 nanocomposites after 20 sorption cycles both in the absorbed and desorbed state are shown in Figure 4.11 and relevant crystallographic data from Rietveld analysis are gathered in Tables 4.3 and 4.4. XRD patterns of both absorbed and desorbed MgH2-TiH2 can be indexed with three phases: β-MgH2, ε-TiH2, and Mg. Amount and crystallinity of ε-TiH2, both in absorbed and desorbed states, do not change significantly as compared to the as-synthetized NC. In contrast, MgH2 undergoes drastic changes in content and crystal size. Taking apart TiH2, MgH2 is the main phase (83 wt.%) in the absorbed state and the minor one in the desorbed one (9 wt.%). A good agreement is found between MgH2 amount as evaluated from XRD and kinetic data, which validates the hypothesis that incomplete Mg/MgH2 transformation is due to kinetic effects.

MgH2-ZrH2 diffraction patterns, both in absorbed and desorbed states, can be indexed by three phases: β-MgH2, ZrH2, and Mg. In the absorbed state, the main phase is β-MgH2 (51 wt.%) but a significant amount of Mg (34 wt.%) remains unreacted. As it has been shown in the kinetic study, there is a significant amount of unreacted Mg in the sample during absorption. Desorbed state agrees with the kinetic results (Fig. 4.10 right), there is a great amount of Mg (75 wt.%) and a little quantity of retained MgH2 (5 wt.%). There is no change in the amount and crystallinity of ZrH2 on cycling as, at the operation conditions, it is a stable hydride.

MgH2 with group 3 transition metals Sc and Y

Synthesis and structural characterization a) System 0.95MgH2-0.05ScH2

The NC 0.95MgH2-0.05ScH2 was synthesized by RBM. A problem in the acquisition data during milling occurred, so it was not possible to measure the second milling cycle. However, we were able to evaluate the hydrogen uptake from the initial and final hydrogen pressure in the jar. The experimental CH is 7.0 wt.% which agrees with the theoretical one assuming 0.95MgH2-0.05ScH2 formation. As it has been explained in Chapter one, the metallic Sc was received in lumps. It was performed a first hydrogenation and a simultaneous decrepitation of the hydride. Then, it was sized in 20/100 mesh, and the dihydride was mixed with the Mg powder. The XRD pattern of as-synthetized 0.95MgH2-0.05ScH2 is shown in Figure 4.12. The pattern can be indexed in three phases: β-MgH2, γ-MgH2, and ScH2. The crystallographic are presented in Chapter 1. The pattern (black line) exhibits broad diffraction peaks evidencing nanostructuration. Table 4.5 summarizes the crystallographic properties of this nanocomposite as evaluated from Rietveld refinement (red line in Fig. 4.12). The most abundant phases are β-MgH2(69 wt. %) and γ-MgH2 (22 wt. %). ScH2 content (9 wt.%) agrees with the targeted composition of the material 0.95MgH2+0.05ScH2, in agreement with the fact that no ternary Mg-Sc-H phases were detected. Crystal sizes are the same for β-and γ-MgH2 phases (7 nm). Interestingly, crystal size for ScH2 is quite large (23 nm) as compared to previous additives. Several previous studies report this larger crystal size [8,9] . 4.13. The evolution of hydrogen pressure (left) and hydrogen uptake (right) as a function of milling time is displayed. Qualitative evolution of pressure and temperature on milling is like previous synthesis. Final hydrogen uptake is 1.9 H/f.u. and 6.8 wt.% which reasonably agrees with the expected values: 2.05 H/f.u and 6.9 wt.%., respectively.

Hydrogen sorption properties a) Pressure-composition isotherms

Isothermal PCI curves at 300 °C of Sc-containing and Y-containing NC are depicted in Figure 4.15. Their isotherms are very similar. For Sc-containing NC, the isotherm exhibits a flat plateau pressure at 0.14 MPa with a width of 6 wt.%. No significant hysteresis is observed between absorption and desorption isotherms. For the Y-containing NC, the plateau pressure, which exhibits low hysteresis is located at 0.15 MPa. The width of the plateau is also 6 wt. %. The initial Crev capacity for Sc-containing NC is quite small (3.5 wt. %) but improves on cycling to attain 4.8 wt.% after 7 cycles and gradually decrease down to 4.1 wt.% after 20 cycles. For the Y-containing compound, the initial Crev is extremely low (1 wt.%) and it further decreases down to 0.6 wt.% after 20 cycles. A similar behavior has been observed in the NC of pure magnesium (Fig. 3.9 Chapter 3) where y = 0 exhibited in the first cycle very low reversible capacity, Crev = 0.5 wt.%, with a gradual decrease after 20 cycles: Crev = 0.2 wt.%. The Ticontaining composite exhibits higher reversible capacity and more stable cycling behavior as compared to both Sc-and Y-containing NCs.

c) Hydrogen sorption curves

.17 displays the hydrogen sorption curves for MgH2-ScH2 and MgH2-YH3 nanocomposites at initial (1 st desorption, 2 nd absorption) and final (20 th absorption and desorption) sorption cycles. For Sc-containing NC (left), kinetics of the 1 st desorption is sluggish. Only 47% of the reaction fraction takes place in 15 min, leaving 53% of retained MgH2. On the 2 nd absorption, kinetics is, however, faster and 97% of the previously formed Mg is converted into MgH2. At the 20 th desorption, kinetics is faster than for the first desorption sweep but the transformed fraction only attains 64%. At the 20 th absorption, kinetics is very close to the 2 nd sweep but the final reaction fraction is much lower: α = 0.75 and 0.97 for 2 nd and 20 th absorption, respectively.

For Y-containing NC (Fig. 4.17 right), the desorption kinetics are similar and extremely slow at cycles one and twenty, leading to a large amount of retained MgH2. On the other hand, Habsorption of the previously formed Mg is fast: after 15 min., more than 90% of the formed Mg returns to MgH2. Comparing both NCs with TiH2, the kinetics of the latter one is getting faster in both absorption and desorption, this also happened in desorption for Sc-NC. However, the absorption is slowing down for in Sc-NC, as for Y-NC in a slightly way.

d) Evolution of hydrogen amount on cycling

Figure 4.18 shows the evolution of hydrogen content in MgH2-ScH2 and MgH2-YH3 NCs upon cycling both after absorption and desorption sweeps. For the Sc-containing NC, the retained amount of MgH2 in the first desorption is high because of the sluggish desorption kinetic described before (Fig. 4.17 left). Nevertheless, absorption kinetics in the first cycles is fast, and low amount of Mg remains unreacted. As cycling proceeds, desorption kinetics accelerate and the retained MgH2 decreases. In contrast, the amount of unreacted Mg increases on cycling. As a result, the reversible capacity increases on cycling particularly within the first cycles (see also Fig. 4.16). For the Y-containing NC (Figure 4.18 right), the amount of retained MgH2 is extremely high (above 5 wt.%) due to the very slow desorption kinetics of Ycontaining NC (Fig. 4.17 right). In contrast, the absorption kinetics are relatively high and almost full reformation of MgH2 occurs. However, the reversible capacity does not surpass 1 wt. % on cycling. As it has been discussed above the Ti-NC Crev is 5 wt %, for Sc-NC and Y-NC, Crev= 4.4 and 1 wt %, respectively. Ti-NC shows higher value due to the improvement of the kinetic.

Effect of hydrogen cycling on the microstructure of MgH2-ScH2 and MgH2-YH3 NCs

XRD patterns of MgH2-ScH2 and MgH2-YH3 nanocomposites after 20 sorption cycles both in the absorbed and desorbed state are shown in Figure 4.19 and relevant crystallographic data from Rietveld analysis are gathered in Tables 4.7 and 4.8. XRD patterns of both absorbed and desorbed MgH2-ScH2 can be indexed with three phases: β-MgH2, Mg and ScH2. Phase amount and crystallinity of ScH2 in both absorbed and desorbed state remain the same as for the as-synthetized NC. In contrast, MgH2 and Mg change in content and crystal size. MgH2 is the main phase (70 wt.%) in the absorbed state and the minor one (5 wt.%) in the desorbed one showing rather high reversibility for hydrogen storage. A good agreement is found between XRD and kinetic data, which validates the hypothesis that incomplete Mg/MgH2 transformation is due to kinetic limitations. Crystal sizes of MgH2 and Mg are quite close and do not exhibit significant change between absorbed and desorbed states (60±10 nm).

MgH2-YH3 diffraction patterns (Fig. 4.19 bottom) after 20 cycles in absorbed and desorbed sates can be indexed with three phases: β-MgH2, YH3, and Mg. As discussed in the previous kinetic section, the main phase in both states is β-MgH2. Its crystal size is practically the same in both states (ca. 50 nm). Comparing both NCs with Ti-NC, this last one shows an increase of the β-MgH2 phase in the absorbed state as well as, Mg phase in the desorbed state, because of the improve kinetic, that is not observed for Sc-NC and Y-NC. .20, left, shows the synthesis of V-containing NC with typical profiles of both vial temperature and hydrogen pressure. On the right, the hydrogen uptake curve is depicted and consists of two steps. The first uptake, at t < 30 min with ca. 1 wt.% absorption is assigned to VH2 formation. Later, Mg absorbs hydrogen with 90% of the reaction completed at 53 min. The total hydrogen content after 120 min of milling is CH= 7.3 wt.%, 1.9 H/f.u.

(H/f.u.) The XRD pattern of the as-milled V-containing NC is shown in Figure 4.21. It can be indexed with three phases: β-MgH2, γ-MgH2, and VH. As it was explained in Chapter one VH0.8 is where the β1 phase disappears and the β2 (VH0.9-2) phase exists from a pressure of 0.1 MPa [11]. Table 4.9 gathers the main crystallographic properties for all phases as evaluated by the Rietveld method. The most abundant phase is β-MgH2 (80 wt. %), followed by, γ-MgH2 (10 wt. %) and, VH (10 wt. %). The crystal sizes are very small for all phases: 6 nm for both MgH2 polymorphs and 10 nm for VH.

Hydrogen sorption properties

a) Pressure-composition isotherms Figure 4.23 shows isothermal PCI curves at 300 °C for V-containing (left) and Nb-containing (right) NCs. Both nanocomposites exhibit a flat plateau at close hydrogen pressure: Pp= 0.16 and 0.14 MPa for V-and Nb-containing composites, respectively. In the first case, the width of the plateau pressure is ca. 7 wt.%, while in the second extends over ca. 6 wt.% at Pp= 0.16 This difference can be explained by the higher atomic weight of Nb as compared to V. Also, it can be noticed that there is minimal hysteresis between absorption and desorption curves. .24 shows H-cycling sorption experiments for V-and Nb-containing NCs carried at 300°C with 15 min absorption/desorption sweeps. Crev was measured for both NCs and compared to Ti-containing one. For both V-and Nb-NCs the same tendency is observed with a gradual decrease of Crev on cycling: they lose ca. 35 % of the initial capacity on 20 cycles, being initial values 6.1 and 5.7 wt.% for V-and Nb-NCs, respectively. In contrast, Crev for Ticontaining NC is rather stable with a capacity loss of only 7 % in 20 cycles. Crev on cycling at 300 °C for 0.95MgH2-0.05VH and 0.95MgH2-0.05NbH and 0.95MgH2-0.05TiH2 nanocomposites. Reaction time limited to 15 min.

c) Hydrogen sorption curves

.25 displays the hydrogen sorption curves for MgH2-VH and MgH2-NbH nanocomposites at initial (1 st desorption, 2 nd absorption) and final (20 th absorption and desorption) sorption cycles. For V-containing NC (left), kinetics of the 1 st desorption is very fast at short reaction time and 75% of the reacted fraction α is attained in 6 min. However, the reaction slows down at longer time and α only attains 80% after 15 min. After twenty cycles, the desorption kinetic is even faster at short times with α = 0.65 in 3 min but saturating at this reacted fraction for longer reaction time. As concerns absorption kinetics, they are extremely fast in the 2 nd absorption with 85% of the reacted fraction in 3 min, saturating at long reaction time (α = 0.9 at 15 min). For the 20 th cycle, reaction kinetics at short reaction time are as fast as in the 2 nd sweep (α = 0.6 for t < 1min) but then it slows down as reaction proceeds with only 75% of the reaction accomplished after 15 min.

For the Nb-containing NC (Fig. 4.25, right), the first desorption curve resembles that of the V-one with 80% of the reaction rapidly accomplished in 6 min. but slowing down at longer reaction time (α = 0.83 at 15 min.). After 20 cycles, desorption kinetics are slightly faster at short reaction time (t < 2 min) but for long reaction time, the reaction fraction saturates at α = 0.6. Comparing both NCs with Ti-NC, these V-NC and Nb-NC show a faster desorption kinetic through cycling but a reduction in the reacted fraction, on the other hand, Ti-NC exhibits slower desorption kinetic but the reacted fraction increases. For the three NCs, the speed in absorption kinetic is the same in cycles one and twenty, nevertheless, Ti-NC keeps the same reacted fraction, V-NC and Nb-NC reduce this value through cycling.

d) Evolution of hydrogen amount in cycling

For the V-containing NC, the desorption is not completed during the first cycle (ca. 2 wt.% not desorbed) and some MgH2 is retained. The amount of MgH2 retained gradually decreases on cycling (ca. 1 wt.% in cycle 20). This result concurs with the acceleration of desorption kinetics on cycling (Fig. 4.9). As for the absorption state, it contains more hydrogen in the first cycles, while it gradually decreases on long cycling so that the amount of unreacted Mg steadily increases. As a consequence, the amount of reversible hydrogen storage keeps constant on cycling at ca. 5 wt % in agreement with Fig. 4.8. One can notice, that the quantity of hydrogen stored in TiH2 is supposed to be constant (ca. 0.34 wt.%).

Effect of hydrogen cycling on the microstructure of MgH2-VH

XRD patterns of MgH2-VH nanocomposite after 20 sorption cycles both in absorbed and desorbed state are shown in Figure 4.19 and relevant crystallographic data from Rietveld analysis are gathered in Table 4.10. XRD pattern of both absorbed and desorbed MgH2-VH composites can be indexed with three phases: β-MgH2, VH, and Mg. β-MgH2 is the major phase 65 wt.%) in the absorbed state, while Mg is the major one (58 wt.%) in the desorbed state. In both states, the crystal size is the same (50 nm) for MgH2 and Mg. VH0.9 amount and crystal size remain unchanged as compared to the as-synthetized nanocomposite.

Discussion

The following discussion is presented analyzing each MgH2-transition metal system, their main results and their comparison with the literature results. Finally, it is shown a comparison and analysis among them.

MgH2-ZrH2

The system MgH2-ZrH2 shows the faster milling kinetic among all the NCs. In 33 min, 90% of the hydrogenation reaction is achieved. After synthesis, the crystallite size was measured and the three phases γ-MgH2, β-MgH2, and ZrH2 have a uniform size of 7 nm. After 20 sorption cycles, it is observed that Crev decreases through the first cycles then keeps constant. Regarding kinetic, it is observed that desorption improves during cycling and reversibility loss is attributed to the poor absorption kinetic. The same behavior was already shown by Zaluska et al [13]. They studied several additives for improving the MgH2 sorption kinetics. They observed that vanadium markedly improves the kinetics of hydrogen absorption. Even more effective is 5 wt.% of zirconium, which can lead to desorption occurring at lower temperatures than pure magnesium (280°C). However, the best kinetic results are reported for mixtures of additives, for example, Zr-Ti or Mn-Zr. The addition of Mn and Zr was found to be especially effective, leading to hydrogen absorption at much lower temperatures (190 and 170°C). Csujko et al [14] compared the effects of catalytic metal additives on the hydrogen desorption properties of magnesium hydride formed after hydrogenation of the Mg + 10 wt.%X (X = V, Zr, or Y) composites. They found that the desorption is much better for Mg + 10 wt.%V, then Zr and finally Y. The difference is particularly marked at lower temperatures (300 and 325 °C). Other studies have been done with ZrF4 to improve the sorption kinetic [15]. Yamanaka et al [16] studied the thermal properties of zirconium hydride in the range of 25 to 500 °C, detecting that the thermal diffusivity decreased with increasing temperature and was almost independent of the hydrogen content, also observing that the temperature dependence of the heat capacity of the zirconium hydride is small.

After cycling the particle size was measured and it was noted that there is no change in the crystal size for the ZrH2 phase whereas Mg and MgH2 phases have grown larger. This growth was also observed by Csujko et al [14], they explained that during the hydrogenation stage, the powder particles lose their nanostructured character since the particles do not contain nanosized β-MgH2 hydride grains (crystallites) anymore. Therefore, during the desorption stage, the β-MgH2 hydride is already sub-microcrystalline rather than nanocrystalline. The grain growth of MgH2 is also more substantial, apparently due to higher employed activation or hydrogenation temperature (350 °C).

MgH2-ScH2

RBM shows a maximum hydrogen uptake is achieved (7 wt. %). After synthesis, the NC was analyzed by XRD, where three phases are found: γ-MgH2, β-MgH2 and ScH2. By the Rietveld method, the abundance and crystal size were calculated. The most abundant phase is β-MgH2, then, γ-MgH2 and finally ScH2. Both Mg phases showed a crystal size of 7nm and the ScH2 was 23 nm. The ScH2 crystal size is the largest one compared with the other additives, a fact characteristic on this dihydride. Luo et al [9] synthesized by RBM the xMgH2+(1-x)ScH2 (x=0.65-1) system at different times, they calculated a crystal size for ScH2= 40 nm after 40h of milling. In this study, the researchers found than Mg do not react during ball milling for milling time larger than 20 h. They indicate that there is an optimum of 6h milling to reach the full MgH2 formation; also mentioning the good dispersion among the phases observed by SEM analysis. Figure 4.12 shows that the full absorption is carried out in only 40 min having the main phases described above. From this study, an optimum synthesis time could be proposed in 60 min.

After carrying out twenty sorption cycles, the hydrogen reversible capacity is enhanced from the first cycles, then reaches a maximum (Crev= 4.5 wt. %H) and finally fluctuates around this maximum. This good cycling performance was also observed by Luo et al [8], they perform hydrogen sorption cycles of 0.65MgH2+0.35ScH2, in thin films at 290 °C and 2 MPa. After 55 cycles, they observed a good cyclability of the NC, reaching Crev= 4.2 wt% H. They measured a hydrogen absorption of 3.2 wt. % H in 5 min and hydrogen desorption 3 wt. % H in 15 min. As can be seen in Figure 4.16 after 20 cycles, Crev= 4.1 wt% H, which agrees with the previous study exposed above.

From kinetic results it is noticed that kinetic absorption improves during when desorption kinetic is lessening. This could be attributed to ScH2, being a transition metal hydride, only helping the dissociation of H2 to H during absorption. This is in agreement with Ogawa et al [17]. They studied the hydrogen desorption of Mg-10at%ScH2, at several temperatures (350, 375 and 400 °C) at 1MPa founding that even at 400°C the system shows a sluggish desorption kinetic. For the cycling reported by Luo et al [8], after 55 sorption cycles, they conclude a kinetic improvement occurs after 20 hydrogen sorption cycles.

After the cycling process, the crystal size for Mg/MgH2phases reaches 62±7 nm in both absorbed and desorbed state, ScH2 remaining the same. This could be understood if ScH2 is only a gateway, having a fluorite structure allowing a higher mobility of H than MgH2/Mg. Also, the large crystal size of ScH2 limits the growth of the MgH2/Mg phases.

MgH2-YH3

After RBM, three phases were identified by XRD, two polymorphs of MgH2 and YH3. Zlotea et al. [18] reported on the heats of enthalpy formation (ΔHf) for YH2, YH3 and MgH2 with -114, -89 and -37 KJ-1mol -1 H respectively. They explained that the formation of YH2 is the most favorable at moderate pressures, however, at higher pressures, both MgH2 and YH3 are always formed, which agrees with the heats of formation.

Absorption-desorption cycles were carried out in a constraint time of 15 min. The Y-NC shows the poorest Crev value among all the tested materials with Crev only slightly higher than pure MgH2, Ogawa et al [17] tried to study the hydrogen desorption properties of Mg-10at.%Y. However, they explain that they were not successful because of the limited solubility of yttrium in the magnesium host. They observed that the desorption kinetic is very sluggish by comparison with other additives. Borgschulte et al [19] explained that the morphology dependence of the catalytic activity of supported metal clusters in heterogeneous catalysis is attributed to a so-called strong metal-support interaction (SMSI). They demonstrate that the SMSI state of deposited Pd clusters is also the physical origin of the thickness dependence of the hydrogen uptake of switchable mirrors and the same effect is also expected to play a role in hydrogen storage. They investigated the morphology and the electronic structure of Pd clusters grown on oxidized yttrium surfaces, both as deposited and after hydrogenation, founding a strong Pd thickness dependence for yttrium hydrogenation. Chacon et al [20] studied the hydrogen uptake in Pd/Mg1-x-Yx thin films (x= 0-17 at. %). They found that yttrium atoms are substitutionally included into the magnesium lattice giving rise to an extended solution, as compared to bulk alloys. They also found that for low Y content, the hydrogen uptake is like that of pure magnesium. This latter result agrees with the ones explained above. They concluded that the hydrogen uptake can be described with a nucleation and growth model. After cycling, the crystal size was measured, and it was found that there is no change in the sizes while the NC totally absorbs and desorbs. They have similar phase abundance and YH3 does not change its crystallite size. This can be explained by the poor reversibility of the reaction, as a result, the material remains like its initial state, the growth of MgH2 before and after cycling can be attributed to the increase of the temperature.

MgH2-VH

MgH2-VH NC was successfully synthetized by RBM. The mechanically milled MgH2+5 at.%V composite exhibits much better hydrogen sorption kinetics than non-milled and mechanically milled and MgH2, reported previously by Huot et al [21]. A PCI curve was plotted with only one plateau observed, which corresponds to MgH2, and agrees with [22].

Crev was measured during cycling, and a loss of the reversible capacity is observed for the first sorption cycles, then remains constant. This is attributed to the sluggish absorption kinetics. However, desorption kinetic is fast, it is in line with the results of Liang et al. [22] who observed that MgH2-5at.%V is able to desorb hydrogen at 300 °C under vacuum (0.015 MPa) with fast kinetics. Ren et al [23] tested different MgH2-V-based mixtures observing dehydrogenation at 240 °C and 0.001 MPa, the binary system MgH2-5at.% showing the slowest desorption kinetic among all the composites. Sorption kinetic worsening was observed during 20 cycles, which could be related to the migration of catalyst particles to the edge as it was described in Chapter 3. Worsening was also observed by Tan et al [24]. They carried out hundreds of sorption cycles for MgH2-5 at.%V and worsening were observed from the first cycles. They attributed this fact to the growth of the Mg/MgH2 particles through cycling, which increases the hydrogen diffusion path. This disagrees with our results because from XRD analysis, we conclude that Mg/MgH2 growth is limited for the additive phase provoking a coarsening limitation effect [2]. Finally, Liang et al. [25] analyzed the desorption kinetic mechanism. They concluded that it follows a nucleation and two-dimensional growth at temperatures below 300 °C and an interface controlled, two-dimensional growth for temperatures above 300 °C

MgH2-NbH

The synthesis of MgH2-NbH NC was successfully achieved, the hydrogenation during the synthesis reached 90% yield in only 46 minutes, which is shorter than other NCs except for Zr. The PCI curve is very close to the homolog MgH2-VH, the smaller hydrogen capacity is explained by the fact that at 5at. % translates into 10 wt.% vanadium and 15 wt.% niobium NCs. The Crev was measured during cycling. After of twenty cycles, a loss of 2 wt.% H. is observed. The dehydrogenation kinetics of MgH2-NbH is as fast as that of MgH2-VH, and hydrogenation curves have the same shape. This was also observed by Huot et al. [26] that suggested that having the same curve shape, they have the same dehydrogenation ratelimiting step. The capacity loss is seen in absorption, where the kinetic is sluggish. Liu et al [27] carried out absorption of Mg-7.5 wt.% Nb at 4MPa and 200 °C, pointing out that the complete hydrogenation should be done in 60 min. Kinetic worsening through cycling was also observed the, which explains the reduction of Crev. A large hydrogen storage capacity loss (3 wt. %) was also observed by Tan et al [28]. As this system is like MgH2-VH with comparable shape curves, it can be assumed that nucleation and two-dimensional growth is the kinetic model which can better explain the experimental data.

Effect of TM additives on MgH2

MgH2-based NCs with 5 mol% of different additives (Ti, Zr, Sc, Y, V, and Nb) were successfully synthesized by RBM. Under our milling conditions, all additives formed hydrides with the following stoichiometry: TiH2, ZrH2, ScH2, YH3, VH2, and NbH2, in agreement with their thermodynamic properties [11,[29][30][31][32][33][34]. Each NC reached the nominal capacity within the first milling cycle. This was confirmed by XRD analysis with 10 wt % for the additives of period four and 15wt.% for those of the fifth period, according to their atomic number (Table 4.11). For TiH2 and ZrH2, as presented in Figures 4.1 and 4.4, these hydrides are stable at around 8 MPa and temperature close to 60°C (inner vial temperature). ScH2 was hydrogenated before milling as it explained in section 4.2.1 a, so this NC remains stable during the RBM. For YH3, the trihydride formation is explained by Huiberts et al [35] : during hydrogen absorption of YH2, at higher concentration of hydrogen, the material resistivity quickly increases, first due to an increase of octahedral occupation in the fcc dihydride, followed for a further increase due to the transformation of the hcp trihydride. The NC VH2 and NbH2 are obtained after RBM, according to their thermodynamic properties as explained in Chapter one. However, after RBM, the milling vial is unloaded, and the pressure is decreased to 0.03 MPa. At this point, a desorption of the dihydride occurs leading to the formation of the monohydride VH and NbH. Depending on TM, by RBM, a variation of the hydrogenation time is observed. For instance, 90% hydrogenation is completed in 40 min for MgH2-ZrH2, MgH2-ScH2, and MgH2-NbH2. On the other hand, MgH2-TiH2, MgH2-YH3, and MgH2-VH2 need 50 min to reach 90%. These results indicate that the nature of the additive has an influence on the nanocomposites synthesis and, depending on the additive, synthesis time can be improved. The more stable dopants, belonging to the 5 th period, allows a faster hydrogenation than the NCs of period four, with two exceptions: the ones who belong to group 3, Sc because it was previously hydrogenated and Y due to its poor dopant contribution.

From XRD data, it is observed that the main phase after synthesis is β-MgH2, followed by ϒ-MgH2, and finally the additive phase. Formation of any ternary phase Mg-TM-H has not been detected. From XRD analysis, the crystallite size of each phase was determined. The two (β and ) MgH2 phases exhibit a crystallite size of 6 ± 1 nm with all additives. The crystallite size of the additive phases ranges between 7 to 13 nm, except for ScH2 with a larger crystallite size of 23 nm (explained above). Figure 4.28 shows both the absorption and desorption plateau pressures, for each NC compared with pure MgH2 data [36]. The addition of TM leads to almost the same plateau pressure value PP than pure MgH2 for absorption and desorption. This assures that the thermodynamic of MgH2 is not changed and agrees with some other systems already studied. For instance, Cuevas et al. [2] studied the 70MgH2+30TiH2 at 300°C with Pabs=0.8 Mpa, Pdes= 0.03 MPa, there was no evidence for a change in the MgH2 thermodynamic. Liang et al. [30] assessed 95MgH2+5VH2, at 200°C and Pabs= 0.1 Mpa, Pdes= 0.015 MPa, measuring both the enthalpy and entropy of hydride formation which is exactly the same than MgH2. Desorption of 90MgH2+10ScH2 was carried out by Ogawa et al. [17] at 350°C and 1MPa. They calculate the enthalpy and entropy of hydride formation which is like pure MgH2. As for Sc-containing NC, it starts with low Crev (3.2 wt.%), increases up to a maximum after 7 cycles and then gradually decreases on long-cycling, as for the previous additives. A unique behavior is seen for Y-containing system, for which the reversible capacity is very low at the first cycle (1 wt.%) and gradually decreases upon cycling. Finally, the best and most interesting behavior is observed for Ti. It exhibits a high capacity from the first cycle (5.3 wt.%) which remains almost unchanged on cycling (4.9 wt.% after 20 cycles). The low reversible capacity of the MgH2-YH3 composite from the first cycle deserves particular interest despite its poor performances. This Y-NC forms similarly to the other NCs by RBM (Fig. 4.13) and consists as usual of two MgH2 polymorphs in nanostructured state. However, though microstructural properties are favorable, hydrogen desorption is extremely slow and resembles to that of nanometric MgH2 prepared by RBM (Fig. 3.9 Chapter 3). This fact can be attributed to surface barriers, such as formation of yttrium oxides leading to slow recombination of hydrogen atoms at the yttrium hydride surface as compared to the other studied TM hydrides.

Low reversible capacity at the first sorption cycles also occurs for MgH2-ScH2, belonging both to the same group 3. It can be argued from their close chemical properties that reaction barriers may also form at the ScH2 surface limiting hydrogen recombination. Indeed, slow desorption kinetics in the first cycle is detected for MgH2-ScH2 composite (Fig. 4.17). However, contrary to Y-containing NC, the reversible capacity of Sc-NC significantly increases on cycling and therefore both cases differ. Another singular property of Sc-NC is the large crystallite size of ScH2 phase (23 nm), due to its preparation method by the solid-gas reaction, as compared to other TM-hydride additives (ca. 10 nm). Thus, one may tentatively attribute the low reversible capacity of MgH2-ScH2 at first H-cycles to a less efficient dispersion of ScH2 in the MgH2 matrix or at its surface as compared to other TM-hydride additives.

The cases of Zr-, Nb-and V-containing nanocomposites, for which the reversible capacity gradually decreases on cycling, are all characterized by the similar evolution of reaction kinetics on cycling (Figures 4.9 The MgH2-TiH2 composite offers the best behavior in Crev for all NCs with stable reversible capacity on cycling. As it happens for Zr-, Nb-and V-NCs, reaction kinetic accelerates on desorption (Fig. 4.9b) but absorption kinetics after 20 th cycles are better for MgH2-TiH2 than for the other NCs. Indeed, one can notice in the 20 th absorption sweep, that the reacted fraction at which the absorption kinetics slow down is higher for Ti-NC (α = 0.6) than for Zr-, Nb-and V-NCs (α = 0.4).

From the XRD and Rietveld's analysis after the cycling process, some facts must be pointed out: additive phases were identified for the NCs in the absorbed and desorbed state as follows: TiH2, ZrH2, ScH2, VH0.8, NbH0.8, and YH3. In the case of YH3, the phase YH2 was never observed as explained above. The Mg/MgH2crystallite size increases after absorption and desorption in all cases at the end of 20 cycles. The crystal size of Mg and MgH2 increases about 10 times after 20 hydrogen cycles at 300 °C which could be attributed to thermal effects between the Mg metal and MgH2 after cycling process for each NC.

To summarize the hydrogen storage properties of 0.95MgH2-0.05TMHx (TM= Ti, Zr, Sc, Y, V, and Nb) NCs for practical applications, the hydrogen content contributions after twenty cycles of the different counterparts in the NC are displayed in Figure 4.30. The total hydrogen content of the NCs is given by blue triangles assuming the formation of MgH2 TiH2, ZrH2, ScH2, VH2, NbH2 and YH3 during RBM synthesis. It is slightly larger for TM belonging to the period 4 (ca. 7.3 wt.%) than for the period 5 (6.8 wt.%). The available capacity of the nanocomposites is given by red squares assuming that reversible hydrogen storage only occurs for the MgH2 counterpart, i.e. TM hydrides are not reversible at the operating conditions.

Based on the kinetic performance of the nanocomposites, some reversible (grey area) and irreversible (white area) Mg/MgH2 contributions are detected. Their relative extent depends on the additive. Starting from MgH2-YH3 exhibits very poor desorption kinetics and therefore a high irreversible contribution. MgH2-ZrH2 and MgH2-NbH2 have the similar reversible capacity (3.5 wt.%), which value is low due to the sluggish absorption kinetic at high reacted fraction. Like these latter NCs, MgH2-VH2 exhibits slow absorption kinetics but the lighter weight of the TM confers a high reversible capacity (4 wt.%). After 20 sorption cycles, MgH2-ScH2 has a similar behavior to h MgH2-VH2 though it was characterized by sluggish desorption kinetics in the first cycles. Finally, MgH2-TiH2 offers the larger amount of reversibility (4.9 wt.%) thanks not only to its fast kinetics on absorption and desorption but their good stability on cycling.

Chapter 5. Conception of an automatic cycling device to measure the cyclelife of metal hydrides

The target of this chapter is the conception and assembly of an Automatic Hydrogen Cycling device (AC-device) prototype. It has been carried out in different steps which are explained in this chapter. First, the thermodynamic operation conditions have been established. Second, a hydride forming alloy that will act as a source and sink of hydrogen has been selected according to operating conditions. Third, a blueprint has been drawn before the assembly. The conception of hydride reservoirs, pressure, flow and temperature sensors, automatic valves and connectors have been considered in this step. Fourth, a coupling with commercial software has been adapted to the automatic control and data acquisition.

Principle of the Automatic cycling device

Thermochemical hydride compressors

The AC-device will be based on the use of metal hydrides as thermochemical compressors of hydrogen gas. As shown in Figure 5.1, the temperature dependence of the plateau pressure of a hydride forming alloys can be used to compress hydrogen gas. By heating the hydride from a low temperature, TL to a high-temperature TH, its equilibrium pressure increases from a low-pressure plateau, PL, and the high-pressure plateau, PH. Thus, in operation, a reservoir filled with a hydride that is repeatedly heated and cooled between TL and TH will follow the thermodynamic path ABCD (solid line, absorption; dashed line, desorption) shown at Fig. 5.1 (right). The compression ratio depends on the alloy selected, on the temperature range T = TH-TL, and on the overall system dynamics. The thermal energy needed to do the work of compression can be produced from waste heat or solar energy [1].

Development of hydride compressors

In the early 70's, Reilly et al. described one of the first applications of hydrides compressors [2]. The compressor used vanadium hydride. It was built to drive periodically a mercury column acting as a pump for other gases. The first commercial hydride compressor was developed by Golben in 1983 [3]. The device is a four-stage metal/hydrogen compressor that uses hot water at 75 °C as an energy source. Four stages filled with different alloys from the LaNi5 family are assembled in series in two different beds inside a reservoir. Based on this work, in the 80's, two models are commercially proposed by Ergenics, a single electrically powered unit that is able to compress H2 between 0.5 and 3.5 MPa, and a four-stage hot water (75 °C) heated unit that compresses H2 from 0.4 to 4.1 MPa for a weight system of 64 kg [4].

Hydrogen cycling devices

Hydrogen cycling in manual hydrogenation systems is time-consuming and strict manipulations must be considered to keep the purity of the gas. Usually, these experiments are carried out in standard Sieverts' type volumetric devices [5,6]. This kind of equipment is not optimized for performing extended absorption/desorption cycles. Even when automatic devices are designed specifically for cycling, such as PCTPro system from Setaram [7], other problems arise. Typical designs of these devices involve the use of high purity hydrogen as a gas source and a vacuum pump used as a gas sink. The compound under study is held at constant temperature while high-pressure and low-pressure sweeps (i.e. pressure cycling) are imposed from a hydrogen source (usually, 20 MPa tank inlet) and a hydrogen sink (vacuum connected reservoir outlet). Thus, large quantities of high purity gas are unfortunately wasted as cycling proceeds. A second issue is related to gas purity itself. Hydride-forming materials can be poisoned by impurities such as O2, CO2, and water. This gas can be further purified, for example, up to 99.99% grade by means of a hydriding/dehydriding sorption process using a suitable metal hydride. Even in such a case, as the gas is pumped out of the system on every cycle, fresh gas must be constantly added to the system through a gas inlet. This operation increases the risk of leakage and subsequent contamination. In addition, the open sink concept (usually a mechanical vacuum pump) also introduces risk of contamination due, for example, to oil fumes from the pump and gas upstream diffusion. Most of these issues can be solved if decomposition/formation of metal hydrides by thermal cycling is used as a source/sink of hydrogen gas. Purified H2 is only introduced into the equipment's volume at the beginning of the experiment. If all device components are properly sealed, external impurities will not appear and residual ones will likely be dissolved in the metal hydride.

Alternatively, another way to cycle hydride materials is to impose thermal cycles at fixed external pressure. However, this implies using wide temperature ranges to achieve full hydride decomposition and formation. Thus, cycle-life degradation will result not only from hydrogen absorption/desorption processes but also from thermally induced changes. Moreover, to facilitate hydriding/decomposition, the sample could be thermally cycled out from technological working conditions, making the association of results with real applications difficult.

Taking into account last drawbacks, Meyer et al [8]. designed an automated equipment for pressure cycling. The basic idea was to use a single hydride reservoir filled with a hydride forming alloy as a hydrogen source/sink by programmed heating/cooling to generate pressure sweeps. The highest/lowest temperatures are selected so that the generated pressures by a hydride forming alloy are above/below the formation/decomposition pressure of the sample under study, which will be kept at constant temperature. Thus, thermal cycles are applied to the hydrogen source/sink compound and not to the sample under study. The pressure sweeps were used to fill (or empty) a calibrated intermediate volume up to a desired pressure. The gas contained in this volume was then automatically expanded to the sample volume like in a classical Sieverts' apparatus and the amount of hydrogen absorbed in each sweep could be determined. Using a forming hydride material LaNi5, as a source and sink of hydrogen, the equipment eradicated hydrogen waste and ensure high purity gas during the experimentation.

Using this apparatus, 1000 uptake/release cycles were performed in a Mm0.8Ca0.2Ni5 (Mm: 50% cerium, 25% lanthanum, and 15-18% neodymium) sample kept at 25°C. Absorption and desorption pressures were about 5 and 2 MPa. A full absorption-desorption cycle took 12 min. Such a fast cycling was achieved thanks to an effective heating/cooling (internal cartridge heater/external cold-water flow) system.

To study the cycle-life of hydrides, Challet [9] developed a simpler pressure cycling equipment which was able to work autonomously without interruption (Figure 5.2). The hydrogen source/sink was provided by a reservoir filled by a hydride forming alloy: LaNi4.6Sn0.4. Thermal cycling was imposed on this reservoir to generate pressure cycles. To supply hydrogen to the studied sample (absorption), the reservoir is heated so that the plateau pressure PH of LaNi4.6Sn0.4 exceeds that of the studied sample (Ps). The sample then absorbs hydrogen. By cooling the reservoir down to a temperature TL for which the plateau pressure PL of LaNi4.6Sn0.4 drops below that of the sample, the latter desorbs the hydrogen. Cycling absorption/desorption of the sample is therefore carried out by pressure cycling. The studied sample was maintained at constant temperature. With this system, no measurement of the quantity of hydrogen exchanged between the reservoir and the sample was available. It was used for cycling a La0.55Y0.45Ni5 sample at -5°C over 100 cycles in the pressure range from 0.05 to 2 MPa. Each cycle took as much as 5 hours since low-efficient classical tubular oven and ambient air was used as heating/cooling systems.

Principle of AC-device

Taking into consideration the previous works, there is a strong interest to build an ACdevice to measure the cycling-life of hydrides. More specifically, we aimed to determine the cycle-life of MgH2-based NCs which operate at 300°C and possess fast kinetics as shown in Chapters 3 and 4.

The ideal thermodynamic principle of the designed system is schematically described in Figure 5.3. The sample S, a MgH2-TMHx nanocomposite, will be cycled through MgH2 Mg transformation with the help of two equivalent reservoirs filled up with the same hydride forming alloy. These reservoirs will be heated up at different temperatures TH and TL to act as hydrogen source and hydrogen sink, respectively. The reservoir at TH will be initially hydrogenated in an external apparatus at the onset of the β-hydride branch (point 1 in Fig. 5.3). The reservoir at TL will be initially hydrogenated at the end of the α-solution branch (point 4 in Fig. 5.3). Temperatures TH and TL will be set so that the corresponding equilibrium plateau pressures PH and PL of the hydride forming alloy enclose that of the studied MgH2-TMHx sample PS. The sample temperature will be kept constant (typically at 300 °C). Reservoirs and sample holder will be put in contact alternatively. Thus, connection between the hightemperature reservoir and the sample holder will hydrogenate sample PS (arrow from 2 to 3 in Fig. 5.3), whereas connection between the sample holder and the low-temperature reservoir will dehydrogenate it (reverse arrow from 3 to 2 in Fig. 3.5). This results in a complete hydrogenation/dehydrogenation cycle for sample S at constant temperature. Sample mass will be much lower than that of the hydride forming alloy in the sinks as schematically represented in Fig. 3.5 by a large difference in plateau width. It will allow tens of cycling sweeps before all hydrogen from the high-temperature reservoir is transferred to the low temperature one. Once this happens, the temperature between both reservoirs will be inverted.

Setting of the operation conditions

The purpose of the AC-device within the framework of this Thesis is to perform extended hydrogen cycles for MgH2-TMHx NCs. Therefore, it is necessary to consider the thermodynamics of Mg/MgH2. The PCI curves in Chapter 3, Figure 3.7, showed that at 300 °C the Pp of MgH2-TMHx NCs in both absorption and desorption is 0.16MPa. Thus, the pressure and temperature conditions of the AC-device can be fixed considering these thermodynamic data as well as kinetic, safety and thermal management issues. Values are presented in Table 5.1. Thus, absorption (PH) and desorption (PL) pressures were fixed to similar values than those used in chapters 3 and 4 to get high enough driving forces for the reversible MgH2 Mg reaction and ensure fast kinetics. The temperature sink reservoir was fixed at TL = 50°C for using ambient air for thermal management while operating close to RT. The temperature of the source reservoir was fixed at TH ≤ 200 °C for safety issues.

Selection of the hydride forming alloy

There are four important facts to consider for selecting the hydride forming alloy used to fill up the source and sink reservoirs according to the previous operation settings: storage capacity, thermodynamic properties (plateaus), kinetics and cycle-life. First, the alloy must have a high reversible storage capacity to ensure the autonomy of the system. The amount should be much larger than that of the cycled sample. Second, flat plateaus with moderate hysteresis are needed to guarantee a good efficiency of the system and stable driving forces during cycling. Third, alloy kinetics should be as fast as those of the cycled sample for not limiting its absorption/desorption rate and minimizing the time needed for cycling. Forth, the alloy must have a long-cycle life for maintaining the thermodynamic operation conditions of the AC-device over thousands of cycles. An additional fact for this work is a safety issue, because of the device autonomy the hydride material should work at a relatively low temperature (T ≤ 200 °C) and pressure (P ≤ 3 MPa).

Hydride-forming alloys of AB5 and AB2 types are suitable for the compression of hydrogen due to their thermodynamic properties: they exhibit moderate hysteresis and almost flat pressure plateaux. Hydrogen absorption in these intermetallic compounds is characterized by a hydrogen capacity of about one atom per metal atom (H/M).

Van Vutcht et al. [10] reported that LaNi5 could reversible absorb 6 H/f.u. at room temperature under a PP about 0.2 MPa. Thus, LaNi5 forms a too unstable hydride for being used as hydrogen sink at TL = 50°C. To adapt its plateau pressure to these thermodynamic conditions part of B-type atoms (Ni) can be replaced by other transition metals [11,12]. Indeed, plateau pressures of substituted LaNi5-type compounds depend on the nature and amount of the substituting elements [13]. Unfortunately, there is a price to pay for this Percheron-Guégan et al. reported that the partial replacement of Ni for other transition metals and semi-metals generally leads to a lowering of PP [19]. This is clearly evidenced in Fig. 5.5. The extent of this stabilization depends strongly on the type and the rate of substitution. For instance, manganese and aluminum substituted compounds form very stable hydrides: at 40 °C: PP = 5 X 10 -4 and 1.6 X 10 -4 MPa for LaNi4MnH6 and LaNi4AlH4.8, respectively, compared to with PP =0.37 MPa for LaNi5H6.4. Exceptions to this rule occur with platinum and palladium [20,21]. Considering the operation conditions given before (particularly, PL = 0.05 MPa at TL = 50°C) several stabilizing substitutions are of interest: M = Mn, Al, Si, and Sn. To make a choice between these possibilities, capacity, kinetics, and cycle-life should be considered. All the four substituted systems exhibit good hydrogen reversibility and fast kinetics [19]. Nevertheless, the desorption capacity for Al, Si, and Sn is lower than for Mn. On the other hand, Mn substituted alloy has poor cycle-life as compared to the others. Indeed, LaNi4.7Al0.3 has good cycle-life with a loss of 12% in hydrogen uptake after 2500 cycles [22]. Even better, Lambert et al. [23] reported that the hydrogen capacity of LaNi4.8Sn0.2 remains virtually unaffected after 10 000 thermal cycles with a very low hysteresis during the absorption/desorption cycling. Taking the cycle-life as a key factor of choice, the LaNi5-xSnx system has been selected. The substitution amount x is determined in the following from thermodynamic considerations.

Using the Van't Hoff equation 1.1 Chapter 1, the composition of the alloy can be determined from the defined thermodynamic properties as well as enthalpy and entropy values of the reaction. For Pp = 0.05 MPa at T = 50°C, the required formation enthalpy can be calculated as ΔH = -38.11 kJ mol -1 H2 for an entropy of reaction fixed at 112 J K mol -1 H2 [24], which is a typical value for hydrogen absorption in the LaNi5-xSnx system. Lartigue [25] tabulated ΔH values using literature data for different Sn-contents [24,26,27]. They are gathered in Table 5. Following this selection, it is worth noting that Sn-content also reduces the capacity of the alloy as compared to the parent compound. This has been evidenced by Joubert et al. who reported the PCI curves of the LaNi5-xSnx system for 0 ≤ x ≤0.5 [26]. As shown in Fig. 5.6. not only the plateau pressures decrease with Sn-content but also the plateau width reduces in length. Using the calculated thermodynamic data of the selected alloy LaNi4.66Sn0.34 (i.e., ΔH = -38.11 kJ mol -1 H2, ΔS = -112 JK -1 mol -1 H2) we could estimate suitable thermodynamic parameters for the high-temperature hydrogen source reservoir. For a temperature of TH =150 °C, the plateau pressure PH is evaluated as 1.38 MPa, within the range of the thermodynamic operation conditions defined in Table 5.1. Moreover, these conditions allow to obtaining relatively close driving forces (DF) for hydrogen absorption (Ec. As a next step, the amount of LaNi4.66Sn0.34 needed for each reservoir was estimated. This calculation was based on the number of moles, nH,MgH2, needed to hydrogenate a given amount of Mg. This amount was fixed to 100 mg leading to nH,Mg = 0.008. The Mg amount was selected as a compromise between having high enough material for subsequent microstructural characterization after cycling but low enough for minimizing the quantity of hydrogen transferred in each cycle. LaNi4.66Sn0.34 was considered to form LaNi4.66Sn0.34H5 as a hydride (see Fig. 5.7). The needed amount of LaNi4.66Sn0.34 was overestimated by a factor of 25 (nH, LaNi4.66Sn0.34H5 = 0.2) to supply enough hydrogen to hydrogenate the Mg sample 25 times. The estimated quantity of LaNi4.66Sn0.34 was 35 g.

As a final remark, the thermic conductivity of LaNi5 alloys is reported to be low (1.32 Jm -1 s -1 K -1 ) [28,29]. To overcome this issue and based on the studies of Meyer et al. [7] and Suda et al. [30], LaNi4.66Sn0.34 was mixed to Cu (20 wt. %) powders. Cu powder was previously subjected to a reduction treatment (7h at 350°C under Ar/H2 flowing) for getting rid of native surface oxides.

Synthesis of LaNi4.66Sn0.34

In 1972, Buschow and van Mal [31] published the La-Ni phases diagram (Fig 5 .8, left). In the range of 50 to 100 % of nickel atoms, there are six compounds: LaNi, LaNix (x≈1.4), LaNi2, LaNi7, and LaNi5. The binary diagram La-Ni shows a wide homogeneity domain for the phase LaNi5. Also, the ternary La-Ni-Sn system at 400 °C is shown (Figure 5.8 right) [32]. Beside eight ternary compounds (LaNi5Sn, LaNi4Sn2, LaNi2Sn2, La3Ni2Sn6, La3Ni2Sn7, La3Ni8Sn16, LaNiSn, and La5Ni24Sn) the pseudobinary phase La(Ni, Sn)5 with ~ 8 at.% Sn solubility limit exists. The Sncontent in LaNi4.66Sn0.34 is 6 at.%.

c) Thermodynamic properties

The hydrogenation properties of the LaNi4.66Sn0.34 alloy were characterized by PCI isotherms at 50, 100, 125 and 150 °C in the pressure range 10 -3 to 2.5 MPa (Figure 5.12). The maximum capacity is 5.5 H/f.u. with a plateau width around 4 H/u.f. PCI curves exhibit small hysteresis between hydrogen uptake and release with a sloping plateau. From the PP at half-plateau, the van't Hoff plot has been obtained as shown in Figure 5.12. Enthalpy and entropy values both on absorption and desorption are shown. The formation enthalpy, Habs = -33.4 kJ mol -1 H2, is slightly lower (in absolute value) than the targeted one, H = -38.11 kJ mol -1 H2. Note also that a deviation in the same direction occurs for the entropy, which counterbalances this effect. Indeed, the plateau pressures at selected temperatures match well with targeted values (Table 5.1). At TL = 50°C, the plateau pressure on absorption (i.e. to act as a hydrogen sink) is PL = 0.06 MPa. At TH = 50°C, the plateau pressure on desorption (i.e. to act as a hydrogen source) is PH = 1.2 MPa.

Blueprint of AC-device

Once that both hydrogen source and sink hydrides have been chosen, the reservoir or reactor where the hydrides are going to take place it has been selected. Knowing the mass of LaNi4.66Sn0.34 (35 g), and density (8.28 g/cm 3 ) the volume occupied by the hydride forming alloy can be calculated (4.22 cm 3 ) (Cu volume was also considered). With these data and the supplier availability, it was chosen a Swagelok double-ended cylinders, of 50 cm 3 , able to work at pressures up to 34.4 MPa, stainless steel 304L materials resist to intergranular corrosion. The double-ended was considered to set a thermocouple on one side and the other is the out/in gas. The heat provided to the reactor is coming from a heating tight mica collar. The maximum reachable temperature is 375 °C. The supply voltage is 230 V. The electric connection is done by cable isolated of the sink from glass under flexible braid steel. A T-type thermocouple was plugged in the heating collar and connected to a PID controller, for temperature control of the reactor.

According to the operating conditions, two Bronkhorst thermal mass flow meters made on stainless steel 304L, with maximum flow values (1000 mg/h), temperature (70 °C) and pressure (10 MPa), were selected. To work in an automatic way, two Swagelok 3-way solenoid valves were acquired. These valves open when receiving an electrical signal, allowing air to enter the actuator. When the signal stops, the solenoid valves close and vents pressure from the actuator. A pressure indicator is also used in the system likewise a safety valve limiting to a maximum of 25 bar pressure operation into the system. Swagelok stainless steel ball valves (7) has been added to the system for manual operation at the beginning or end of the process. The tubing fitting material is from stainless steel with 4 mm of internal diameter. The volume of the system (out of the two reservoirs) was measured to be 21.24 cm 3 .

A diagram of the AC-device is drawn in Figure 5.13, acronyms are explained from left to right in Table 5 The working steps can be explained with different pictures. For instance, Figure 5.14 depicted in yellow the very first step, vacuum procedure. The vacuum is applied to system or stationary parts. After tens of cycles, the hydrogen source is going to be empty and the hydrogen sink is close to being full. At that time the temperature of both reservoirs is inverted, and the source starts refilling as shown in Figure 5.17

Connections and software of control and measure

Tubbing, electric signal converters, and activators have been fixed on an acrylic surface as can be seen in figure 5.18. An interface box containing the relays controlling each valve and the DC power supplies for the pressure sensors was made. A desktop computer records the pressure data and flux meters. In accordance with these measurements, the computer also controls the opening and closing of the electrically controlled valves. The interface to the computer is performed through a DAQPad-6211 external acquisition board from the company National Instruments, connected to the USB port of the computer.

The control program (Figure 5. 19) is performed with the LabView® 12 software from National Instruments. The operator can assign a name to outputs (valves) an input (pressure gauge). The state of the valves is displayed by indicator lights, text or gauge indicator of pressure, representing the measurements on the screen. The measured data and the state of the valves can be saved on file in real time. The control can be performed either with the mouse or the keyboard.

Measurements

The system was initially tested without any sample to validate the cycling acquisition data. Before starting with the test, it was carried out a leaking checkup, which assured that the system is completely tight and isolated. It was also possible to measure the intern volume of the pipelines and accessories (v= 21.240±0.001 cm 3 ). Then, the cycling simulation was run. During this operation, both flux-meters were saving data through the cycling simulation. At the end, the data were recovered from the Flowplot software and treated in OriginPro® software. The recording information of both flux-meters is shown in Figure 5.21a. The fluxmeter one records the absorption, meanwhile the flux-meter two records the desorption. The Flowplot software gives the time and value data. From that information, a plot with two axes is drawn: the time as x (s) and the flux value y (%), this percent value corresponds at one basis of 1 % for 10 mgH2/second. During this first test, it was possible to simulate more than 40 absorption-desorption cycles. Due to the fact that the sample holder was empty, the hydrogen flux was constant for absorptions, and desorptions. Figure 5.21b, shows the hydrogen flux during the first sorption cycle (black curves) and the integrated flux (red curves). From the integrated data, it is possible to calculate the mole number that went through the flux meters. These values are 0.002 mole for absorption and 0.004 mole for desorption. Taking on account that there is no hydride forming sample for the sorption process, the same quantity is measured for absorption and desorption. After 43 cycles all the hydrogen contained in the source tank was transferred to the sink. Once the system has been tested, an experimentation was carried out with a hydrideforming sample NC: 95MgH2-5TiH2. More than 50 sorption cycles were carried out during this study. The NC used for the experiment was obtained directly form the synthesis, which means that it was already hydrogenated. The sample mass was 0.0995 g. Then, it was introduced into the sample holder (volume= 2.250±0.001 cm 3 ) inside a glove-box (argon atmosphere). Once the sample was inside the sample holder, the free volume was calculated (2.181 cm 3 ). The sample holder was plugged to the AC-device and heated up to 300 °C. Thus, the sample releases a small amount of hydrogen leading to a slight increase in gas pressure. After the sample reached 300°C, and the reservoirs were at their respective temperatures (50 and 150 °C), the cycling setup was started (cycling time 15 min. absorption and 15 min. desorption). During the sorption process, it was necessary to stand-by the process 4 times and to refill the source reservoir with the hydrogen contained in the sink. This was performed when the sink pressure was close to the 95MgH2-5TiH2 plateau pressure (1.6 bar). The refill procedure took in average 3 h. Figure 5.22 shows the first 12 cycles, for both absorption and desorption. As the absorption rate is faster than the desorption one, a peak for desorption is clearly detected. An intensity decrease of the desorption peak is also observed. At the end of the cycling process, measured flux data is obtained from Flowplot software. Then, the hydrogen moles number can be calculated for both absorption and desorption. In Figure 5.23 left, the first cycle is shown in detail for absorption and desorption. The shadowed desorption curve can be integrated to calculate the number of moles that have been desorbed for this first step. Figure 5. 23 (right) shows the integration curve that corresponds to 0.02 moles of hydrogen that have been desorbed by the 95MgH2-5TiH2 NC in 15 min, at 300°C. The same process is repeated through the full cycling process.

Conclusion

From this experiment, it was possible to carry out tens of uptake-release hydrogen cycling with the automatic cycling device designed and assembled in the laboratory. To build this device, a study of the cycling operation conditions has been done. A bibliography research was performed to choose the best hydride that would work as both source and sink, and an AB5-type intermetallic alloy was chosen. The alloy LaNi4.66Sn0.33 was synthetized and was characterized both structurally and thermodynamically. Considering all the main parts of the device (reservoirs, sensors, automatic valves, and connectors), a piping and instrumentation diagram was drawn. Manipulation and safety tests were performed before starting up. The automatization and operation with a commercial software were performed to operate the system from a computer. In the end, this prototype allows performing tens of uptake/release cycles in a safe and automatic way. These cycles are carried out at high hydrogen purity as well as economizing the hydrogen consumption. In comparison with manual cycling process or some commercial cycling machines, this prototype is working with a constant amount of hydrogen, only transferring the gas from the source to the sink passing through the sample of interest. Very first test on the sample 95MgH2-5TiH2 have shown that sorption cycles can be conducted successfully, and quantitative measurements of hydrogen exchanged during the cycling can be obtained from the flux meters. Further works are needed to fully calibrate the device and to perform accurate measurements. However, first trials have shown that such device can efficiently conduct large-scale cycling of metallic hydride with high safety and low cost.

Conclusions and future works

This final chapter aims to point out the main results achieved throughout this Ph.D. Thesis. Different strategies have been implemented through different Mg-based materials to improve the hydrogen sorption properties of Mg. The kinetic properties were enhanced by nanostructuration and addition of early transition metals. Besides kinetics enhancement, hydrogen reversibility and cyclic stability at short time (15 min) of the magnesium-based materials were analyzed and discussed. Additional efforts were dedicated to making a proofof-concept, design and build an automatic hydrogen uptake/release cycling device for future extended cycle-life studies. It is now essential to consider the relevance of these strategies to correctly orientate future investigations.

Conclusions

In this Ph.D. thesis, it was experimentally shown that by reactive ball milling, it is possible to reach hydrogen uptake in microcrystalline Mg in shorter times than for bulk material. The synthetized nanocomposites have a high density of grain boundaries which improves the hydrogen diffusion in the particle. Also, this work demonstrates that hydrideforming early transition metals enhance hydrogen sorption kinetics in Mg, in some cases showing a high hydrogen reversibility for short reaction time (15 min) and stable capacity on cycling. Most significant and detailed results are given in the following.

(1-y)MgH2+yTiH2 system

• Reaction time for MgH2 formation decreases with TiH2 addition during reactive ball milling synthesis, which reveals that TiH2 enhances hydrogen absorption rates of Mg. • PCI measurements show that the hydrogenation and dehydrogenation enthalpies of the Mg-MgH2 transformation (-73.6 ± 0.9 and 75.5 ± 0.6 kJmolH2 -1 ) are not modified by Ti addition. This concurs with the lack of formation of ternary Mg-Ti-H phases. • TPD measurements show that Ti-containing NCs start desorbing at a lower temperature than pure MgH2. TPA measurements demonstrate that hydrogen absorption in Mg even starts at room temperature. • NCs with high TiH2 content (0.1<y<0.3 mol.) showed stable hydrogen reversible capacity through cycling. • Kinetics of the reversible Mg  MgH2 transformation speed with TiH2 addition. The more Ti is added, the faster the reaction takes place. • Morphological SEM studies proved that no significant changes in NCs particle size occur by Ti addition nor in sorption cycling. Therefore, changes in particle size are not at the origin of the observed kinetic modifications. • XRD results show that there is a significant increase in the crystal size both for MgH2 and Mg phases for all NC. • SEM-EDS analysis on cycled NCs revealed migration of TiH2 from the core of the agglomerate particles to the edges upon cycling. • Hydrogen desorption from MgH2 is accelerated through cycling. This is attributed to surface effects: faster hydrogen recombination at NC surface on cycling which is promoted by surface migration of TiH2. Bulk effects are discarded since Mg and MgH2 crystal size increases on cycling which should slow-down reaction kinetics by Hdiffusion if it would be the rate limiting step. • No improvement of hydrogen absorption kinetics occurs on cycling.

• Hydrogen sorption curves both on absorption and desorption can be fitted with a firstorder reaction nucleation and growth model. • TiH2 acts as a gateway for hydrogen transport enabling fast hydrogen dissociation/recombination at the TiH2 surface and its diffusion towards the bulk. • TiH2 enhances the reversible hydrogen uptake of Mg but remains an irreversible hydride at moderate temperatures (300 °C). As a consequence, there is an optimum amount of TiH2 for getting the highest reversible capacity in (1-y)MgH2+yTiH2 NCs. This optimum is set to y = 0.025 for a reaction time of 15 min, with a reversible capacity of 4.9 wt.% at the twentieth cycle.

Effect of TM additives on MgH2

• All early transition metals formed hydride phases during RBM synthesis of MgH2-TMHx NCs. These hydrides have the following stoichiometry: TiH2, ZrH2, ScH2, YH3, VH2, and NbH2. • Depending on the additive, the synthesis time of MgH2-TMHx NCs can be reduced from 50 to 30 min. • The addition of TM leads to equivalent plateau pressures than for pure MgH2, showing that the thermodynamics of MgH2 is not changed. • After sorption cycling, XRD analysis showed that the hydrogen content of transition metal hydrides is the same as in the as-synthetized state except for V and Nb. The composition of these is VH0.8 and NbH0.8. • The order of the hydrogen reversible capacity after 20 cycles is as follows:

Ti>Sc>V>Nb>Zr>Y.

Automatic cycling device

• A closed-loop experimental device has been designed for automatic hydrogen cycling of MgH2-based composites with high-purity hydrogen gas. • The thermodynamic operation parameters aim to characterize reversible hydrogen absorption systems within the range pressure 0.05 ≤ PH2≤ 2 MPa. These parameters were used to define the composition and amount of an AB5 hydride (LaNi4.33Sn0.66) which is used as a hydrogen source and sink in the cycling device. • The cycling device was successfully assembled and operated with a software monitoring. • Many sorption cycles were performed showing the good operation and convenience of the automatic device. • The device allows to perform cycling in a safe way and at high hydrogen purity.

Annex Fitting kinetic models

Function of the reacted fraction vs time and their respective kinetic models in first desorption of (1-y)MgH2+yTiH2. Function of the reacted fraction vs time and their respective kinetic models in twentieth desorption of (1-y)MgH2+yTiH2.

Fig. 1 . 1

 11 Fig. 1.1 Left-hand side: Typical Pressure-Composition isotherms of a Metal-Hydrogen system. Right-hand side: related van't Hoff plot.

Figure 1 . 2

 12 Figure1.2 van't Hoff plots of some selected hydrides[16]

Fig. 1 . 3

 13 Fig. 1.3 Schematic representation of hysteresis for a metal hydride system

Fig. 1 . 4

 14 Fig. 1.4 Schematic representation of reversible capacity for a metal hydride

Fig. 1 . 5

 15 Fig. 1.5 Dependence of the desorption plateau pressure of MgH2 with temperature [39].

Fig . 1 . 6

 16 Fig .1.6 Hydrogen absorption curves of both non-milled (micrometric) (black line) and ballmilled (red line) Mg at 325 °C

Figure 1 . 7

 17 Figure 1.7 Ti-Mg Phase Diagram by Murray[START_REF] Murray | The Mg-Ti (Magnesium-Titanium) system[END_REF]

Figure 1 . 8 .

 18 Figure 1.8. Ti-H diagram phase by San Martin[START_REF] San-Martin | The H-Ti (Hydrogen-Titanium) system[END_REF]

Figure 1 . 10 .

 110 Figure 1.10. Fugacity vs x in TiH2 [49]. 𝑓 𝐻 2 / 𝑃 𝐻 2 = 3.6 𝑥 10 -4

Figure 1 . 11 .

 111 Figure 1.11. Magnesium-Zirconium Binary Phase Diagram [54].

Figure 1 . 12 .

 112 Figure 1.12. The H-Zr System [55].

Figure 1 . 13 .

 113 Figure 1.13. Isothermal pressure-volume curves [60].

Figure 1 . 14 .

 114 Figure 1.14. The temperature dependence of the hydrogen diffusivity in ZrH, measured by PFC NMR [61].

Figure 1 . 15 .

 115 Figure 1.15. Magnesium-Scandium Binary Phase Diagram [64].

Figure 1 . 16 .

 116 Figure 1.16. The H-Sc System [64].

Figure 1 . 18 .

 118 Figure 1.18. Magnesium-Yttrium Binary Phase Diagram [71].

Figure 1 . 19 .

 119 Figure 1.19. Hydrogen-Yttrium Binary Phase Diagram [72].

Figure 1 . 20 .

 120 Figure 1.20. Experimental isotherms for the system yttrium-hydrogen: ○, absorption points; •, desorption points [76]

Figure 1 . 21 .

 121 Figure 1.21. Experimental isothermal curves for the system yttrium-hydrogen. Absorption and desorption reproducible, irreversible curves are designated by the right and left-hand arrows, respectively [76].

Figure 1 . 22 .

 122 Figure 1.22. Magnesium-Vanadium Binary Phase Diagram [79].

Fig. 1 .

 1 Fig. 1.23 V-H phase diagram[START_REF] Smith | Binary Alloy Phase Diagrams[END_REF]

Figure 1 .

 1 Figure 1.24. Pressure-composition isotherms for VH [83].

Figure 1 .

 1 Figure 1.25 shows the diffusion coefficients for H and D in V, Nb, and Ta that have been determined and plotted by Schaumann [85] as a function of the reciprocal temperature. VH has the fastest DH value among the others, then V-H and finally Ta-H.

Figure 1 . 25

 125 Figure 1.25 Logarithm of the diffusion coefficients of H respectively D in V, Nb, and Ta as a function of the reciprocal temperature [85]

Figure 1 . 26 .

 126 Figure 1.26. Magnesium-Niobium Binary Phase Diagram [80].

Fig 1 . 27

 127 Fig 1.27 Nb-H phase diagram [80].

o

  Model: SX 100 CAMEC o Wavelength Dispersive Spectroscopy (WDS) analyzer; o 3 vertical spectrometers with LPC1, LPC2, LPC3, Lalit, LPET and TAP crystals; o 1 inclined spectrometer with LiF, PET, TAP, PC2 crystals; o 1 EDS SDD spectrometer for EDS-WDS coupled analysis.

Figure 1 . 1 .

 11 Figure 1.1. Technical drawing of the Cameca SX-100 EPMA instrument

Figure 3 . 1 .

 31 Figure 3.1. Evolution of the hydrogen pressure (blue line) and the vial temperature (red line) as a function of milling time during the synthesis of the composite 0.7MgH2+0.3TiH2

Figure 3 . 2 .

 32 Figure 3.2. Hydrogen uptake curves during the first milling cycle of powder mixtures (1-y)Mg + yTi with y = 0, 0.05 and 0.1

3 yFigure 3 . 3 .

 333 Figure 3.3. XRD patterns of nanocomposites (1-y)MgH2+yTiH2 after RBM synthesis.

2 Figure 3 . 4

 234 Figure 3.4 Rietveld analysis of XRD patterns of nanocomposites (1-y)MgH2+yTiH2

Figure 3 . 5

 35 Figure 3.5 Left) SEM images of as-milled powder samples. Right) particle size distribution (histograms) and best fit to a log-normal distribution (solid line) for y=0, 0.0125 and 0.3.

Figure 3 . 6 .

 36 Figure 3.6. SEM images showing the particle morphology of as-milled samples for y=0, 0.0125 and 0.3.

Figure 3 . 7 .

 37 Figure 3.7. PCI absorption (full symbols) and desorption (empty symbols) curves for NC y=0.3 at three different temperatures. Van't Hoff plots are given as inset.

Figure 3 . 8 .

 38 Figure 3.8. a) Temperature Programed Desorption (TPD) and b) Temperature Programed Absorption (TPA) for as-milled MgH2 (y = 0) and TiH2-containing NCs.In the TPD graph, the thermodynamic limit of the Mg-H system is depicted by a red line.

Figure 3 . 9 .

 39 Figure 3.9. Reversible capacity on cycling of (1-y)MgH2+yTiH2 NCs at a) 300 °C and b) 340 °C. Reaction time limited to 15 min. Pabs = 0.8 and Pdes = 0.03 MPa.

Figure 3 .

 3 Figure 3.10. H-desorption curves at cycle one (a) and twenty (b) of (1-y)MgH2+yTiH2 nanocomposites at 573 K for reaction time limited to tend = 15 min. Pdes = 0.03 MPa.

Figure 3 .

 3 Figure 3.11. H-absorption curves at cycle two (a) and twenty (b) of (1-y)MgH2+yTiH2 nanocomposites at 573 K for reaction time limited to tend = 15 min. Pabs = 0.8 MPa.

Figure 3 . 12 .

 312 Figure 3.12. Evolution on cycling of the hydrogen content in (1-y)MgH2+yTiH2 nanocomposites after each absorption (full symbols) and desorption (empty symbols)steps. T = 573 K. Reaction time limited to tend = 15 min. Pabs = 0.8 MPa, Pdes = 0.03 MPa.

Figure 3 . 13 .

 313 Figure 3.13. XRD patterns of (1-y)MgH2+yTiH2 nanocomposites after 20 H-sorption cycles: a) absorbed state, b) desorbed state.

Figure 3 . 14 .

 314 Figure 3.14. Reversibility of the MgH2/Mg phase transformation for (1-y)MgH2+yTiH2 nanocomposites determined from H-sorption curves and XRD measurements. The dashed line stands for full agreement between H-sorption and XRD data. Data labels stand for y amount.

Figure 3 .Figure 3 .

 33 Figure 3.15 SEM images showing the particle size distribution in (1-y)MgH2+yTiH2 NCs after 20 cycles at the desorbed state

Figure 3 . 17 .Figure 3 . 18 .

 317318 Figure 3.17. SEM images showing the morphology of (1-y)MgH2+yTiH2 NCs after cycling in both absorbed and desorbed state

2 Where

 2 𝐹(𝛼) = 𝑘(𝑃, 𝑇, 𝑟) * 𝑡 Ec. 3.1 where the function 𝐹 depends of the controlling mechanism of the reaction. The parameter 𝑘 determines the reaction rate which is dependent on 𝑃, 𝑇 and the sample geometry indicated by 𝑟. Also, in simple cases, it is possible to express 𝑘 as a product of three functions: 𝑘(𝑃, 𝑇, 𝑟) = 𝑓(𝑟) * 𝑔 ( 𝑃 𝑃 𝑝 ) * ℎ(𝑇) Ec. 3.𝑓(𝑟) is related with the sample geometry, 𝑔 ( 𝑃 𝑃 𝑒𝑞 ) is the driving force for the reaction meanwhile the function ℎ(𝑇) is given by: ℎ(𝑇) = 𝑘 0 * exp (-𝐸 𝑎 𝑅 * 𝑇 ) Ec. 3.3

Figure 3 . 19 .

 319 Figure 3.19. Isothermal dα/dt time and α time plots for solid-state reaction models (Table3.16); data simulated with a rate constant of 0.049 min -1 : (a-c) deceleratory; (d) sigmoidal[11].

Figure 3 . 20 .

 320 Figure 3.20. Function of the reacted fraction vs time and their respective kinetic models in absorption (left) and desorption (right) in cycle 20 th

Figure 3 . 21 .

 321 Figure 3.21. Schematic diagram of nucleation and growth process

Figure 3 . 22 .

 322 Figure 3.22. Schematic depicting the evolution of microstructure of the catalyzed MgH2particle during cycling[31]

Figure 3 . 23 .

 323 Figure 3.23. Schematic representation of hydrogen absorption and desorption mechanisms in MgH2-TiH2 nanocomposites [34].

Figure 3 . 24 .

 324 Figure 3.24. Schematic shows the TiH2 role as a gateway for hydrogen transfer (yellow arrows) from the surface to the bulk (black arrows) both on desorption and absorption reactions.

Figure 3 . 25 .

 325 Figure 3.25. Reversible and irreversible contributions to hydrogen storage in (1y)MgH2+yTiH2 nanocomposites with reaction time limited to 15 min ( = Total hydrogen capacity, ◼ = H-capacity of MgH2 counterpart, • = experimental H-reversible capacity) .

Figure 4 . 1 .

 41 Figure 4.1. Evolution of the hydrogen pressure (black line) and the vial temperature (red line) as a function of milling time during the synthesis of the composite 0.95MgH2+0.05TiH2

Figure 4 . 2 .

 42 Figure 4.2. Hydrogen uptake curves in weight percent (black line) and hydrogen atoms per formula unit (red line) during the first milling cycle of 0.95MgH2+0.05TiH2.

Figure 4 . 3 .Figure 4 .Figure 4 . 4 .

 43444 Figure 4.3. Rietveld analysis of XRD pattern of as-milled 0.95MgH2+0.05TiH2 NC

Figure 4 . 5 .

 45 Figure 4.5. Hydrogen uptake curves in weight percent (black line) and hydrogen atoms per formula unit (red line) during the first milling cycle of 0.95MgH2+0.05ZrH2

2 Figure 4 . 7 .

 247 Figure 4.7. PCI absorption (full symbols) and desorption (empty symbols) isotherms at 300°C. left) 0.95MgH2-0.05TiH2, right) 0.95MgH2-0.05ZrH2

Figure 4 .

 4 Figure 4.8 shows the evolution of the Crev, upon cycling for MgH2-TiH2 and MgH2-ZrH2 nanocomposites. For the Ti-containing NC, Crev keeps almost constant during cycling; 5.2 ± 0.03 wt.%H. In contrast, for the Zr-containing-NC, Crev is higher for the first cycle (5.6 wt.%) then gradually falls to 3.5 wt.% in 10 cycles and then keeps constant.

2 Figure 4 . 8 .

 248 Figure 4.8. Crev on cycling at 300 °C for 0.95MgH2-0.05TiH2 and 0.95MgH2-0.05ZrH2 nanocomposites. Reaction time limited to 15 min, Pabs=0.8MPa, Pdes= 3 10 -3 MPa.

Figure 4 . 9 .

 49 Figure 4.9. H-sorption curves at 300°C for cycles one, two and twenty. Left) 0.95MgH2-0.05TiH2, right) 0.95MgH2-0.05ZrH2

Figure 4 . 10 .

 410 Figure 4.10. Evolution of hydrogen content with cycling for 0.95MgH2-0.05TiH2,(left) and 0.95MgH2-0.05ZrH2 (right) nanocomposites. H-content is shown after absorption (full symbols) and desorption (empty symbols). Partial contributions from H-stored in TiH2 or ZrH2, irreversibly retained as MgH2 and reversibly stored in Mg (Crev) are shown.

  pressure evolution during RBM of Sc-containing NC (black line) and temperature vial (red line) as a function of milling. Right: hydrogen uptake curves in weight percent (black line) and hydrogen atoms per formula unit (red line).

2 Figure 4 . 12 .

 2412 Figure 4.12. Rietveld analysis of XRD patterns for Sc-NC

Figure 4 . 13 .

 413 Figure 4.13. Left: Hydrogen pressure evolution during RBM of Y-containing NC (black line) and temperature vial (red line) as a function of milling. Right: hydrogen uptake curves in weight percent (black line) and hydrogen atoms per formula unit (red line).

Figure 4 .

 4 Figure 4.14 shows the XRD pattern of the as-synthetized Y-containing NC. The pattern can be indexed with three phases β-MgH2, γ-MgH2, and YH3. No ternary Mg-Y-H phase was detected. Rietveld refinement was carried out and the main crystallographic properties are gathered in Table 4.6. The main phase is β-MgH2 (70 wt. %). Minor phases are γ-MgH2 (20 wt.

%

  ) and YH3 (10 wt. %). The crystal sizes are the same for both MgH2 polymorphs (7 nm) and a little larger for YH3 (13 nm).

Figure 4 . 14 .

 414 Figure 4.14. Rietveld analysis of XRD pattern of as-milled 0.95MgH2+0.05YH3 NC

Figure 4 . 15 .

 415 Figure 4.15. PCI absorption (full symbols) and desorption (empty symbols) isotherms at 300°C. Left) 0.95MgH2-0.05ScH2, right) 0.95MgH2-0.05YH3

Figure 4 . 16 .

 416 Crev on cycling at 300 °C for 0.95MgH2-0.05ScH2 and 0.95MgH2-0.05YH3 and 0.95MgH2-0.05TiH2 nanocomposites. Reaction time limited to 15 min.

Figure 4 . 17 .

 417 Figure 4.17. H-sorption curves at 300°C for cycles one, two and twenty. Left) 0.95MgH2-0.05ScH2, right) 0.95MgH2-0.05YH3

2 Figure 4 . 18 .

 2418 Figure 4.18. Evolution of hydrogen content with cycling for 0.95MgH2-0.05ScH2,(left) and 0.95MgH2-0.05YH3 (right) nanocomposites. H-content is shown after absorption (full symbols) and desorption (empty symbols). Partial contributions from H-stored in ScH2, YH3, irreversibly retained as MgH2 and reversibly stored in Mg (Crev) are shown.

Figure 4 . 19 .

 419 Figure 4.19. Rietveld analysis of XRD patterns of 0.95MgH2-0.05ScH2, (top) and 0.95MgH2-0.05YH3 (bottom) nanocomposites after 20 cycles in both absorbed (left) and desorbed (right) states

Figure 4

 4 Figure 4.20, left, shows the synthesis of V-containing NC with typical profiles of both vial temperature and hydrogen pressure. On the right, the hydrogen uptake curve is depicted and consists of two steps. The first uptake, at t < 30 min with ca. 1 wt.% absorption is assigned to VH2 formation. Later, Mg absorbs hydrogen with 90% of the reaction completed at 53 min. The total hydrogen content after 120 min of milling is CH= 7.3 wt.%, 1.9 H/f.u.

Figure 4 . 20 .

 420 Figure 4.20. Left: Hydrogen pressure evolution and temperature vial (red line) as a function of milling during RBM (black line) of V-containing NC. Right: Hydrogen uptake curves in weight percent (black line) and hydrogen atoms per formula unit (red line).

Figure 4 . 21 .Figure 4 .Figure 4 . 22 .

 4214422 Figure 4.21. Rietveld analysis of XRD patterns for V-NC

Figure 4 . 23 .Figure 4

 4234 Figure 4.23. PCI absorption (full symbols) and desorption (empty symbols) isotherms at 300°C. Left) 0.95MgH2-0.05VH, right) 0.95MgH2-0.05NbH

Figure 4 .

 4 Figure 4.24. Crev on cycling at 300 °C for 0.95MgH2-0.05VH and 0.95MgH2-0.05NbH and 0.95MgH2-0.05TiH2 nanocomposites. Reaction time limited to 15 min.

Figure 4 . 25 .

 425 Figure 4.25. H-sorption curves at 300°C for cycles one, two and twenty. Left) 0.95MgH2-0.05VH, right) 0.95MgH2-0.05NbH

Figure 4 .

 4 Figure 4.26 shows the hydrogen content in MgH2-VH (left) and MgH2-NbH (right) nanocomposites both after absorption and desorption sweeps and their evolution on cycling. For the V-containing NC, Crev decreases gradually due to increasing amount of unreacted Mg on cycling. Based on the previous kinetic measurements (Figure 4.25 b), this is attributed to the sluggish absorption kinetics (i.e. saturation of reacted fraction) at long reaction time. In contrast, there is no significant influence of desorption kinetics since the amount of MgH2 retained remains unchanged during cycling. For the Nb containing-NC (Fig. 4.26, right), a rather similar behavior is observed. Crev decreases due to sluggish absorption kinetics at long reaction time with the expected occurrence of a high amount of unreacted Mg after 20 cycles. In contrast, the amount of retained MgH2 tends to steadily decrease due to the slight enhancement of desorption kinetics on cycling. Doing a comparison of Ti-NC with these NCs; Ti-NC keeps constant Crev through cycling because of the improvement of the kinetic during cycling. On the other hand, V-NC and Nb-NC lose Crev during absorption state due to sluggish kinetic.

Figure 4 . 26 .

 426 Figure 4.26. Evolution of hydrogen content with cycling for 0.95MgH2-0.05VH, (left) and 0.95MgH2-0.05NbH (right) nanocomposites. H-content is shown after absorption (full symbols) and desorption (empty symbols). Partial contributions from H-stored in VH, NbH, irreversibly retained as MgH2 and reversibly stored in Mg (Crev) are shown.

Figure 4 . 27 .

 427 Figure 4.27. Rietveld analysis of XRD patterns for V-NC, in both absorbed and desorbed state

MgH2Figure 4 . 28 .

 428 Figure 4.28. Half-plateau pressures on absorption and desorption for each synthetized NC

Figure 4 . 29 .

 429 Figure 4.29. Crev on cycling at 300 °C for reaction time limited to 15 min

  , 4.25 left and 4.25 right, respectively). For desorption sweeps, kinetics accelerates on cycling in the three cases. In contrast, for absorption sweeps, reaction kinetics slow down on cycling for reacted fraction exceeding ca. 0.4, leading to a large amount of unreacted Mg(Figures 4.10 right and 4.26). This evidences that the gradual diminution of Crev on cycling relates to the absorption process because of the kinetic exposed above.

Figure 4 . 30 .

 430 Figure 4.30. Reversible and irreversible contributions to hydrogen storage in 0.95MgH2-0.05TM NCs after 20 sorption cycles ( = Total hydrogen capacity, ◼ = H-capacity of MgH2 counterpart, • = experimental H-reversible capacity) References

Figure 5 . 1

 51 Figure 5.1 Principle of the hydride compressor, van't Hoff plot (left) and ideal P-c-I [1]

Figure 5 . 2

 52 Figure 5.2 Thermodynamic principle of the device developed by Challet [9]

Figure 5 . 3

 53 Figure 5.3 Schematic operation principle of the AC-device based on the PCT properties of a hydride forming alloy. Temperatures TH and TL correspond to the equilibrium plateau pressures PH and PL of the hydride forming alloy, Ps and Ts correspond to the MgH2-TMHx pressure and temperature sample.

Figure 5 . 5

 55 Figure 5.5 Desorption isotherms at 40 °C for some substituted LaNi5-xMx compounds [19].

Figure 5 . 6 .

 56 Figure 5.6. Formation enthalpy of LaNi5-xSnx hydrides as a function of substitution ratio

Figure 5 . 7

 57 Figure 5.7 PCI curves at 25°C for LaNi5-xSnx compounds[26]

  5.1) and desorption (Ec.5.2) at 300°C in MgH2-TMHx NCs:

Figure 5 . 8 . 5 . 3 . 1 . 1

 585311 Figure 5.8. Binary Phase Diagram of La-Ni (left), ternary Phase Diagram La-Ni-Sn at 400°C (right)

Figure 5 . 9 .Figure 5 . 10

 59510 Figure 5.9. BSE image of LaNi4.66Sn0.34 alloy

Figure 5 .

 5 Figure 5.11 PCI isotherms of LaNi4.66Sn0.34 alloy at several temperatures. Empty and full symbols stand for desorption and absorption, respectively.

5 H 1 ) 2 Figure 5 .

 5125 Figure 5.12 van't Hoff plots corresponding

Figure 5 . 13 .

 513 Figure 5.13. AC-device diagram

Figure 5 . 14 .Figure 5 .Figure 5 . 15 .Figure 5 . 16 .

 5145515516 Figure 5.14. AC-device vacuum process diagram

  .

Figure 5 . 17 .

 517 Figure 5.17. AC-device hydrogen source refilling process diagram

Figure 5 . 18 .

 518 Figure 5.18. AC-device

Figure 5 . 19

 519 Figure 5.19 Control panel in LabView® 12

Figure 5 . 20

 520 Figure 5.20 Flowplot main window by Bronkhorst®

Figure 5 . 21

 521 Figure 5.21 Acquisition data plot (a), first sorption cycle (b)

Figure 5 . 22

 522 Figure 5.22 Hydrogen uptake-release in AC-device

Figure 5 . 23

 523 Figure 5.23 First absorption and desorption steps (left), integral of the desorbed curve (right)

F

  Function of the reacted fraction vs time and their respective kinetic models in second absorption of (1-y)MgH2+yTiH2.

Table 1 . 2 .

 12 Ti-H Crystal structure data

					Density	Volume	Cell	Cell
	Phase	Formula Prototype Pearson	(mg/m 3 )	(nm 3 )	parameter	parameters	Ref.
	label			symbol			(nm)	(°)
		TiH1.97	ThH2	tI6	3.76	0.044	a=0.3177	α=90	[47]
	TiH2 rt			I4/mmm			b=0.3177	β=90
	(ε)						c=0.4363	γ=90
		Ti	W	cI2	4.37	0.0364	a=0.33149	α=90	[47]
	(Ti) ht			Im-3m			b=0.33149	β=90
	(βTi)						c=0.33149	γ=90
		TiH2	CaF2	cF12	3.56	0.0932	a=0.45345	α=90	[47]
	TiH2 ht			Fm-3m			b=0.45345	β=90
	(δ)						c=0.45345	γ=90
		TiH0.17	Mg	hP2	4.54	0.0352	a=0.29472	α=90	[47]
	(Ti) rt			P63/mmc			b=0.29472	β=90
	(αTi)						c=0.46758	γ=120

Table

Table 1 .

 1

	3. Zr-H Crystal structure data
	Phase	Formula	Prototype	Pearson	Density	Volume	Cell parameter	Cell parameters	Ref.
	label			symbol	(mg/m 3 )	(nm 3 )	(nm)	(°)
	(Zr) rt	Zr	Mg	hP2	6.49	0.04665	a=0.32333	α=90	[56]
	α,(αZr)			P63/mmc			b=0.32333	β=90
							c=0.51523	γ=120
	ZrH2	ZrH2	ThH2	tI6	5.63	0.055	a=0.35179	α=90	[57]
	ε			I4/mmm			b=0.35179	β=90
							c=0.4447	γ=90
							T=294 K
	(Zr) ht	Zr	W	cI2	6.42	0.0472	a=0.3614	α=90	[58]
	β,(βZr)			Im-3m	6.42	0.0472	b=0.3614	β=90
							c=0.3614	γ=90
		Zr	W				T=1173 K
				cI2
				Im-3m				α=90
							a=0.3614	β=90
							b=0.3614	γ=90
							c=0.3614
							T=1173 K
	ZrH1.6	ZrH1.8	CaF2	cF12	5.66	0.1091	a=0.4779	α=90	[59]
	δ			Fm-3m			b=0.4779	β=90
							c=0.4779	γ=90

Table 1 .4. Sc

 1

		-H Crystal structure data
	Phase	Formula Prototype Pearson	Density	Volume	Cell	Cell	Ref.
	label		symbol	(mg/m3)	(nm3)	parameter	parameters
							(nm)	(°)
		ScH2	CaF2	cF12	2.85	0.1095 a=0.47844	α=90	[65]
	ScH2		Fm-3m			b=0.47844	β=90
	δ						c=0.47844	γ=90
		ScH0.39	Mg	hP2	2.95	0.05111	a=0.334	α=90	[66]
	(Sc) rt		P63/mmc			b=0.334	β=90
	(αSc)						c=0.529	γ=120

Table 1 .5. Y

 1

		-H Crystal structure data
	Phase	Formula	Prototype	Pearson	Density	Volume	Cell	Cell parameters	Ref.
	label			symbol	(mg/m3)	(nm3)	parameter	(°)
							(nm)
		Y	Mg	hP2	4.47	0.066	a=0.36474	α=90	[73]
	(Y) rt			P63/mmc	4.47	0.066	b=0.36474	β=90
	(αY)	Y	Mg				c=0.57306	γ=120
				hP2
				P63/mmc			a=0.36474	α=90
							b=0.36474	β=90
							c=0.57306	γ=120
		YH3	HoH3	hP24	3.96	0.2313	a=0.63587	α=90	[74]
	YH3			P-3c1			b=0.63587	β=90
	YH3						c=0.66068	γ=120
		YH2	CaF2	cF12	4.28	0.1412	a=0.5207	α=90	[75]
	YH2			Fm-3m			b=0.5207	β=90
	δYH2						c=0.5207	γ=90

Table 1 .6. V

 1

		-H Crystal structure data
	Phase label	Formula Prototype Pearson symbol	Density (mg/m3)	Volume (nm 3 )	Cell parameter	Cell parameters	Ref.
							(nm)	(°)
		V	W	cI2	6.08	0.0278	a=0.303	α=90	[80]
	(V)			Im-3m			b=0.303	β=90
	α						c=0.303	γ=90
		VH2	CaF2	cF12	4.52	0.0779	a=0.4271	α=90	[81]
	VH2			Fm-3m			b=0.4271	β=90
	γ						c=0.4271	γ=90
		V2H1.10	V2H1.1	tI56	5.47	0.25029	a=0.6035	α=90	[82]
	V2H1.1 rt			I41/amd			b=0.6035	β=90
	β						c=0.6872	γ=90

Table 1 .7. Nb-H Crystal structure data

 1

	Phase	Formula Prototype Pearson	Density	Volume	Cell	Cell	Ref.
	label			symbol	(mg/m3)	(nm3)	parameter	parameters
							(nm)	(°)
		NbH	TaH0.5	oS8	7.56	0.08256	a=0.3467	α=90	[81]
	NbH0.85 rt			C222			b=0.4843	β=90
	β						c=0.4917	γ=90
		NbH0.03	W	cI2	8.55	0.0361	a=0.33053	α=90	[89]
	(Nb)			Im-3m			b=0.33053	β=90
	α						c=0.33053	γ=90
						(no data)
	Nb2H lt2
	ζ
						(no data)
	NbH0.72 lt
	ε'

Table 2 .

 2

	1 Characteristics of the starting metallic materials
	Element	Supplier	Purity	State	Mesh
	Mg	Alfa-Aesar	99.8%	Powder	-20+100
	Ti	Alfa-Aesar	99.9%	Powder	-150
	Zr	CERAC	99.7%	Powder	-325
	V	Stream Chemicals	99.5%	Powder	-325
	Nb	Chem Pur	99.9%	Powder	-150
	Sc	China Rare Metals	99.9%	Pieces, irregularly shaped	-100 (post crushed)
	Y	Stream Chemicals	99.9%	Powder	-40
	La	Alfa-Aesar	99.9%	Pieces under mineral oil	-
	Ni	Alfa-Aesar	99.9%	Pieces, irregularly shaped	-
	Sn	Alfa-Aesar	99.9%	Rod	-

Table 2 .

 2

	2 MgH2-TMHx nanocomposites prepared by RBM
	Nanocomposites	Composition (mol %)
	MgH2-TiH2	Ti = 0.0125-0.3
	MgH2-ZrH2	Zr= 0.05
	MgH2-VH2	V= 0.05
	MgH2-NbH2	Nb= 0.05
	MgH2-ScH2	Sc= 0.05
	MgH2-YH2	Y= 0.05

Table 2 .3 Parameters

 2

for Rietveld refinements for a single-phase compound based on constant wavelength X-ray Scale factor (S) Zero-shift Background (when refined as a polynomial function) Lattice parameters: a, b, c, α, β, ϒ Profile parameters: U, V, W, X, Y Coordinates of atoms: xj, yj, zj Isotope displacement parameters: By Occupation parameters: Oj

Table 3 . 1 .

 31 Chemical composition of as-milled nanocomposites. H-content, CH, measured with Evicomagnetics device. Metal content determined by ICP-OES analysis for y=0 and 0.3 (the measured errors are less than 1%).

Table 3 .

 3

		Phase abundance (wt.%)			Crystal size(nm)
					ε-TiH2
	NC (y)	γ-MgH2	γ/(γ+β)-MgH2 β-MgH2	meas./nominal	γ-MgH2 β-MgH2 ε-TiH2 Rwp%
	0		22	78	-	8	8	-	11
	0.0125		22	75	2/2	7	7	12	11
	0.025		24	72	5/5	6	7	12	11
	0.05		27	65	11/9	5	7	8	15
	0.1		26	59	20/17	6	7	13	12
	0.2		23	46	40/32	6	6	9	14
	0.3		22	38	51/45	5	5	9	11

2. Phase amounts, crystallite sizes and Rietveld agreement factors determined by XRD for the phases observed in as-milled NCs (1-y)MgH2+yTiH2.

Table 3 . 3 .

 33 Phase amount, crystal size and weight-profile Rietveld agreement factor of (1-y)MgH2+yTiH2 nanocomposites after 20 H-sorption cycles at the absorbed and desorbed states.

	Absorbed state
	Phase amount (wt. %)		Crystal size(nm)
	Sample	β-MgH2	ε-TiH2	Mg	β-MgH2	ε-TiH2	Mg	Rwp%
	y=0	100	-	-	42	-	-	15
	y=0.012	83	3	14	77	6	-	15
	y=0.025	83	4	13	58	12	-	15
	y=0.05	83	10	7	31	8	-	13
	y=0.1	71	20	9	65	13	-	10
	y=0.2	59	40	1	46	9	-	14
	y=0.3	49	51	1	47	9	-	10
	Desorbed State
	Phase amount (wt. %)		Crystal size(nm)
	Sample	β-MgH2	ε-TiH2	Mg	β-MgH2	ε-TiH2	Mg	Rwp%
	y=0	99	-	1	30	-	-	13
	y=0.012	23	2	75	-	12	105	16
	y=0.025	14	6	80	-	12	105	19
	y=0.05	9	14	77	-	8	87	14
	y=0.1	13	23	64	-	13	69	12
	y=0.2	3	45	53	-	11	94	12
	y=0.3	0.5	56	44	-	9	67	13

Table 3 . 4

 34

	Ps (µm) absorption

. Mean particle size (Ps), of (1-y)MgH2+yTiH2 NCs after 20 H-sorption cycles at the absorbed, desorbed and as-milled states. NC (y) Ps (µm) as-milled Ps (µm) desorption

Table 3 . 5

 35

		As-milled	Desorbed	Absorbed
	Element	Center	Edge	Center	Edge	Center	Edge
	Ti (at.%)	22	25	13	25	10	27
	Mg (at.%)	78	75	87	75	90	73

. Ti and Mg content at the center and particle edge for as-milled and cycled (absorbed and desorbed) NC with y = 0.3.

Table 3 . 6 .

 36 Kinetic models examined in isothermal sorption curves[13][14][15]

	Symbol	Model	Integral F(α) form
		Diffusion models
	D1	One-dimensional diffusion	𝛼 2
	D2	Two-dimensional diffusion	𝛼 + (1 -𝛼)ln (1 -𝛼)
	D3

Table 3 . 7 .

 37 Best fits of H-sorption curves of (1-y)MgH2+yTiH2 NCs to kinetic models (when F1 is not the best fitting model, its value is compared with the best fitted one).

			Desorption			Absorption
		1 st Cycle	20 th Cycle	2 nd Cycle	20 th Cycle
	y	Model	R 2	Model	R 2	Model	R 2	Model	R 2
	0	R2	0.9812	R2	0.9504	D3	0.8911	F1	0.9900
		F1	0.9804	F1	0.9500	F1	0.8714
	0.0125	A2	0.9969	A2	0.9939	F1	0.9980	F1	0.9994
		F1	0.9849	F1	0.9911
	0.025	F1	0.9972	R2	0.9961	F1	0.9951	F1	0.9962
				F1	0.9931
	0.05	R2	0.9972	R2	0.9984	F1	0.9961	D4	0.9913
		F1	0.9945	F1	0.9937			F1	0.9865
	0.1	F1	0.9928	F1	0.9978	F1	0.9990	D3	0.9925
								F1	0.9426
	0.2	R2	0.9970	F1	0.9990	F1	0.9930	D1	0.9931
		F1	0.9927					F1	0.9871
	0.3	F1	0.9992	F1	0.9990	F1	0.9990	F1	0.9990

Table 4 . 1 .

 41 Crystallographic properties of as-milled 0.95MgH2+0.05TiH2 and Rietveld agreement factor.

	Phase abundance (wt.%)	Crystal size(nm)
			ε-TiH2
	γ-MgH2	β-MgH2	meas./nominal	γ-MgH2	β-MgH2	ε-TiH2	Rwp%
	24	65	11/9	5	7	8	15
		S.G.		Cell parameters (Å)
				a= 4.526
	Pbcn	P42/mnm	I4/mmm	(4)	a=4.518(1) a=3.176(1)
				b= 5.427(3) c= 3.028(1) c= 4.403(1)
				c= 4.971 (3)
			Theoretical [4]
	Phase	Prototype	Pearson	Density	Volume	Cell	Cell
	label		symbol	(mg/m 3 )	(nm 3 )	parameter	angles (°)
						(Å)
	TiH2 (δ)	CaF2	cF12	3.56	0.0932	a=4.5345	α=90
			Fm-3m			c=4.5345	β=90
							γ=90

Table 4 . 2 .

 42 Crystallographic properties of as-milled 0.95MgH2+0.05ZrH2 sample and Rietveld agreement factor.

		Phase abundance (wt.%)	Crystal size(nm)
			ZrH2
	γ-MgH2	β-MgH2	meas./nominal	γ-MgH2	β-MgH2	ZrH2	Rwp%
	34	55	11/15	7	7	7	5
		S.G.		Cell parameters (Å)
				a= 4.553
	Pbcn	P42/mnm	I4/mmm	(3)	a=4.530(1) a=3.498 (1)
					c=
				b= 5.405(4)	3.032(1) c= 4.561(2)
				c= 5.030 (6)
			Theoretical [5]
	Phase	Prototype	Pearson symbol	Density	Volume	Cell	Cell
	label			(mg/m 3 )	(nm 3 )	parameter	angles (°)
						(Å)
	ZrH2 ε	ThH2	tI6	5.63	0.055	a=b=3.5179	α=90
			I4/mmm			c=4.4447	β=90
							γ=90

Figure 4.6. Rietveld analysis of XRD pattern of 0.95MgH2+0.05ZrH2 nanocomposite 4.1.

2 Hydrogen sorption properties a) Pressure-composition isotherms Figure

  MgH 2 -TiH 2 -Desorbed

													Yobs
													Ycalc
													Yobs-Ycalc
													Bragg_position
													-MgH 2 Mg
													-TiH 2
			MgH 2 -ZrH 2 -Absorbed			MgH 2 -ZrH 2 -Desorbed
	Intensity (a.u.)									Yobs Ycalc Yobs-Ycalc Bragg_position	(a.u.) Intensity	Yobs Ycalc Yobs-Ycalc Bragg_position
											-MgH 2 Mg ZrH 2		-MgH 2 ZrH 2 Mg
	25	30	35	40	45	50	55	60	65	70	75	25 30 35 40 45 50 55 60 65 70 75
				2 (degrees)				2 (degrees)

Figure 4.11. Rietveld analysis of XRD patterns of 0.95MgH2-0.05TiH2, (top) and 0.95MgH2-0.05ZrH2 (bottom) nanocomposites after 20 cycles in both absorbed (left) and desorbed (right) states

Table 4 .

 4

	3. Crystallographic properties of (des-)absorbed Ti-containing NC after 20 sorption cycles and Rietveld
	agreement factor.
	Absorbed
	Phase abundance (wt.%)	Crystal size(nm)
			TiH2
	Mg	β-MgH2	meas./nominal Mg	β-MgH2	TiH2	Rwp%
	8	83	9/11	-	31	8	13
		S.G.			Cell parameters (Å)
	P63/mmc	P42/mnm	I4/mmm		a=4.518(1)	a=3.169 (1)
					c= 3.021(1)	c= 4.606(3)
	Desorbed
	Phase abundance (wt.%)	Crystal size(nm)
			TiH2
	Mg	β-MgH2	meas./nominal	Mg	β-MgH2	TiH2	Rwp%
	77	9	14/11	87	-	8	14
		S.G.		Cell parameters (Å)
	P63/mmc	P42/mnm	I4/mmm	a=3.212(1)	a=3.169 (1)
				c=
				5.215(2)	c= 4.401(1)

Table 4 . 4 .

 44 Crystallographic properties (des-)absorbed Zr-containing NC after 20 sorptions and Rietveld agreement factor.

	Absorbed
	Phase abundance (wt.%)	Crystal size(nm)
				ZrH2
	Mg	β-MgH2	meas./nominal	Mg	β-MgH2	ZrH2	Rwp%
	34	51		15/15		75	75	11	9
		S.G.					Cell parameters (Å)
	P63/mmc	P42/mnm	I4/mmm	a=3.213(1)	a=4.519(1)	a=3.478 (1)
					c= 5.220(1)	c= 3.022(1)	c= 4.579(1)
	Desorbed
	Phase abundance (wt.%)		Crystal size(nm)
			ZrH2
	Mg	β-MgH2	meas./nominal	Mg	β-MgH2	ZrH2	Rwp%
	75	5	15/15	50	-		8	7
		S.G.				Cell parameters (Å)
	P63/mmc	P42/mnm	I4/mmm	a=3.212(1)	a=3.485 (1)
					c=
					5.216(1)	c= 4.558(1)

Table 4 . 5 .

 45 Crystallographic properties of as-synthetized 0.95MgH2+0.05ScH2 sample and Rietveld agreement factor.

	Phase abundance (wt.%)	Crystal size(nm)
			ScH2
	γ-MgH2	β-MgH2	meas./nominal γ-MgH2	β-MgH2	ScH2	Rwp%
	22	69	9/9	7	7	23	3
		S.G.		Cell parameters (Å)
						a=3.789
	Pbcn	P42/mnm	Fm-3m	a= 4.533 (3)	a=4.519(1)	(1)
				b= 5.420(2)	c= 3.024(1)
				c= 4.958 (3)
			Theoretical [10]
	Phase	Prototype	Pearson symbol	Density	Volume	Cell	Cell
	label			(mg/m 3 )	(nm 3 )	parameter	angles (°)
						(Å)
	ScH2 δ	CaF2	cF12	2.85	0.1095	a=4.7844	α=90
			Fm-3m				β=90
							γ=90

b) System 0.95MgH2-0.05YH3 Synthesis process of Y-containing NC from Mg and Y powders is shown in Figure

Table 4 . 6 .

 46 Crystallographic properties of as-milled 0.95MgH2+0.05YH3 sample and Rietveld agreement factor.

	Phase abundance (wt.%)			Crystal size(nm)
			YH3
	γ-MgH2	β-MgH2	meas./nominal γ-MgH2	β-MgH2	YH3	Rwp%
	20	70	10/16	7	7	13	3
		S.G.		Cell parameters (Å)
	Pbcn	P42/mnm	P-3c1	a= 4.567 (3) a=4.519(1) a=6.3617 (1)
				b= 5.447(3) c= 3.033(1)	c= 6.635(2)
				c= 4.960 (3)
				Theoretical
	Phase label	Prototype	Pearson	Density	Volume (nm 3 )	Cell	Cell angles (°)
			symbol	(mg/m 3 )		parameter
						(Å)
	YH3	HoH3	hP24	3.96	0.2313	a=6.3587	α=90
			P-3c1			c=6.6068	β=90
							γ=120

Table 4 .

 4

	7. Crystallographic properties (des)absorbed Sc-containing NC after 20 sorption cycles and Rietveld
	agreement factor.
	Absorbed
	Phase abundance (wt.%)		Crystal size(nm)
			ScH2
	Mg	β-MgH2	meas./nominal	Mg	β-MgH2	ScH2	Rwp%
	20	70	10/10		70	70	20	3
		S.G.			Cell parameters (Å)
	P63/mmc	P42/mnm	Fm-3m	a= 3.213 (3) a=4.521(1)	a= 4.791 (1)
				c= 4.960 (3) c= 3.023(1)
	Desorbed
	Phase abundance (wt.%)		Crystal size(nm)
			ScH2
	Mg	β-MgH2	meas./nominal	Mg	β-MgH2	ScH2	Rwp%
	85	5	10/10	55		-	20	5
		S.G.			Cell parameters (Å)
	P63/mmc	P42/mnm	Fm-3m	a= 3.211 (1)	a=4.518(1)	a= 4.787 (1)
				c= 5.214 (3)	c= 3.023(1)

Table 4 . 8 .

 48 Crystallographic properties of abs(desorbed) Y-NC after 20 cycles and Rietveld agreement factor.

	Absorbed
		Phase abundance (wt.%)	Crystal size(nm)
			YH3
	Mg	β-MgH2	meas./nominal	Mg	β-MgH2	YH3	Rwp%
	1	84	15/16	-	56	13	3
		S.G.			Cell parameters (Å)
	P63/mmc	P42/mnm	P-3c1		a=4.520(1)	a= 6.355 (1)
					c= 3.023(1)	c= 6.610(1)
	Desorbed
	Phase abundance (wt.%)	Crystal size(nm)
			YH3
	Mg	β-MgH2	meas./nominal	Mg	β-MgH2	YH3	Rwp%
	7	77	16/16	-	50	15	3
		S.G.			Cell parameters (Å)
	P63/mmc	P42/mnm	P-3c1		a=4.519(1)	a= 6.354 (1)
					c= 3.023(1)	c= 6.608(1)

Table 4 . 9 .

 49 Crystallographic properties of as-milled V-containing NC and Rietveld agreement factor.

		Phase abundance (wt.%)		Crystal size(nm)
	γ-MgH2	β-MgH2	VH meas./nominal	γ-MgH2	β-MgH2	VH	Rwp%
	10	80	10/10	6	6	10	3
		S.G.			Cell parameters (Å)
				a= 4.531
	Pbcn	P42/mnm	I4/mmm	(2)	a=4.519(1)	a=3.034 (1)
				b=
				5.425(2) c= 3.020(1)	c= 3.396(2)
				c=
				4.963(2)
				Theoretical [12]
	Phase	Prototype	Pearson symbol	Density	Volume	Cell parameter (Å)	Cell angles
	label			(mg/m 3 )	(nm 3 )		(°)
	VH0.9	-	-	-	0.0729	a=4.46	α=90
						c=3.66	β=90
							γ=120

Table 4 .

 4 10. Crystallographic properties of abs sample V-NC and Rietveld agreement factor.

	Absorption
	Phase abundance (wt.%)				Crystal size(nm)
	Mg	β-MgH2	VH meas./nominal	Mg	β-MgH2	VH	Rwp%
	25	65	10/10	50		50	7	4
		S.G.				Cell parameters (Å)
	P63/mmc P42/mnm	I4/mmm	a= 3.211 (1)	a=4.504(1)	a= 3.119 (1)
					c= 5.215(1)	c= 3.013(1)	c= 3.185(1)
	Desorption
	Phase abundance (wt.%)			Crystal size(nm)
	Mg	β-MgH2	VH meas./nominal	Mg	β-MgH2	VH	Rwp%
	58	32	10/10		50	50	5	5
			S.G.
	P63/mmc	P42/mnm	I4/mmm	a= 3.209 (1)	a=4.804(1)	a= 3.019(2)
					c= 5.211(1)	c= 2.381(1)	c= 3.282(3)

Table 4 .

 4 11. RBM results for doped nanocomposites

	NC	Measured (wt. % )	Nominal (wt. % )	Atomic number (metal)
			(expected)
	TiH2	11	9	22
	ZrH2	11	15	40
	ScH2	9	9	21
	YH3	10	16	39
	VH2	10	10	23
	NbH2	-	15	41

Table 5 . 1

 51 Thermodynamic operation conditions of the AC-device

	Variable	Value
	Sample mass, mS	0.1 g
	Sample temperature, TS	300 °C
	Sample plateau pressure, PS	0.16 MPa
	H2 source pressure, PH	~1 MPa
	H2 sink pressure, PL	~0.05 MPa
	H2 source temperature, TH	≤ 200°C
	H2 sink temperature, TL	50°C
	Sample holder volume	2.25 cm 3

Table 5 . 3

 53 Enthalpy of hydride formation for different alloys in the LaNi5-xSnx system

	x value	ΔH (kJ mol -1 H2)
	0	30
	0.10	33
	0.20	35
	0.32	38
	0.40	40
	0.50	42

Table 5 . 4

 54 Acronyms used for AC-device diagram

	Acronym	Meaning
	Reservoir LT	Reactor at low temperature
	Reservoir HT Reactor at high temperature
	VL	Valve of low-pressure reactor
	VH	Valve of high-pressure reactor
	VPL	Solenoid valve to low-pressure reactor
	VPH	Solenoid valve to high-pressure reactor
	FL	Flux-meter to low-pressure reactor
	FH	Flux-meter to high-pressure reactor
	VCH	Internal recirculation valve
	VCL	External recirculation valve
	RV	Relief valve
	PI	Pressure indicator
	VC	Valve to system
	VS	Valve to sample
	VVc	Valve to vacuum system

  Function of the reacted fraction vs time and their respective kinetic models in twentieth absorption of (1-y)MgH2+yTiH2. y=0

						D1
	1,0					D2 D3
						D4
						F1
	0,8					R2
						R3
						A2
	0,4 0,6 F(r f )					R 2 =0,9341 A3 R 2 =0,9555 R 2 =0,9402 R 2 =0,9212
						R 2 =0,9900
	0,2					R 2 =0,9905 R 2 =0,9904
						R 2 =0,9517
	0,0					R 2 =0,8902
	0		20	40	80	100
				Time (seg)
	0 y=0.1 0,0 0,2 0,4 0,6 0,8 1,0 y=0.3 0 0,0 0,2 0,4 0,6 0,8 1,0 F(r f )	2 2	4 4	6 Time (seg) 8 10 12 14 16 6 8 10 12 14	D1 D2 D3 D4 F1 R2 R3 A2 A3 R 2 =0,9863 R 2 =0,9784 R 2 =0,9670 R 2 =0,9748 R 2 =0,9990 R 2 =0,9891 R 2 =0,9923 R 2 =0,8985 D1 D2 D3 D4 F1 R2 R3 A2 A3 R 2 =0,9803 R 2 =0,9691 R 2 =0,9539 R 2 =0,9646 R 2 =0,9990 R 2 =0,9925 R 2 =0,9950 R 2 =0,8989 R 2 =0,7676 R 2 =0,7732	0 y=0.2 0,0 0,2 0,4 0,6 0,8 1,0	2	4	6	8 10 12 14 16	=0,9953 R 2 =0,9914 R 2 =0,9844 R 2 =0,9894 R 2 =0,9961 R 2 =0,9724 R 2 =0,9776 R 2 =0,8448 R 2 =0,7037 R =0,9914 D1 D2 D3 D4 F1 R2 R3 A2 A3 R =0,9859 R =0,9773 R =0,9833 R =0,9930 R =0,9824 R =0,9865 R =0,8853 R =0,7580	f f f	F(r ) F(r ) F(r )
				Time (seg)						Time (seg)
											Time (seg)

Acknowledgments

First and foremost, I want to say, "thank you", to my supervisors Fermin Cuevas and Michel Latroche, for giving me the opportunity to work on this project and carrying out my Ph.D. Thesis in the ICMPE (Institut de Chimie et des Matériaux Paris-Est) laboratory, and for their expert guidance, supporting, instructions, and encouragement I would like to express my special thanks to Mr. Jean-Louis BOBET (University of Bordeaux), Mr. Salvatore MIRAGLIA (Institute Néel of Grenoble), and Mr. José Francisco FERNÁNDEZ (Universidad Autónoma de Madrid) to have kindly accepted to be the members of the jury for my Ph.D. dissertation.

adaptation. The substituted compounds show smaller capacity than LaNi5. Capacities of several substituted compounds are presented in Table 5.2.

Table 5.2 Hydrogen capacity at 40 °C for various substituted LaNi5-type alloys and their PP [14] Composition

Storage capacity (H/f.u.) PP (MPa) LaNi5 6.24 0.2 LaNi4.7Al0. 3 5.85 0.1 LaNi4.25Co0. [START_REF] Beavis | The Formation and Properties of Rare-Earth and Transition Metal Hydrides[END_REF] 6.23 0.1 LaNi4.6Mn0. 4 6.21 0.1 LaNi4.5Sn0. 5 5.25 0.1

Many studies have been carried out to determine and understand the kinetic properties of AB5 alloys. One of the best advantages of AB5 alloys is the selective oxidation of the A metal, leaving catalytic B species at the surface to promote adsorption and dissociation of the hydrogen molecules [15]. Thus, Ni segregation at the alloy surface provides fast surface reactions and high resistance against poisoning from gas impurities. Schlapbach et al [16] have emphasized the fact that LaNi5 kinetics are extremely fast, they present a diffusion coefficient value for LaNi4.9Cu0.1 of 2x10 -9 cm 2 s -1 at 20 °C, which shows the fast hydrogen mobility in bulk. Also, they pointed out that the thermal conductivity of both the hydride and the LaNi5 in powder form is poor. In many engineering applications, heat transfer between the hydride bed and the reservoir controls the system kinetics.

Cycle-life properties have been investigated for many AB5 alloys. In general, the reduction of the capacity through cycling is attributed to the disproportionation of the alloy upon hydrogen absorption into AHx and B-enriched compounds [17]. This causes losses in reversible storage capacity and modifies the PCT isotherms. This issue is particularly severe in LaNi5 alloy. As shown in Figure 5.4, the partial substitution of Ni by Sn greatly reduces the rate of degradation [18]. Similar improvements have been found for Al, Ge, and Si substitution [15].

Future work

The automatic cycling device design in this Thesis will allow extending the current studies to much longer cycling ranges, which is of high interest for real applications. Considering the materials synthetized in this work, it will be worth to analyze the extended cycle-life of the optimized nanocomposite 0.975MgH2+0.025TiH2 with a reversibility capacity of 4.9 wt.%. Similarly, the extended cycle-life study of the MgH2-ScH2 system, which shows good stability over 20 cycles, deserves interest.

In chapter 3, it was observed by SEM that it exists a partial migration of TiH2 from the core of the agglomerate to the edges on cycling. This phenomenon is relevant to understand the efficiency and role of the additive. A more detailed high-resolution characterization of this migration by Transmission Electron Microscopy as a function of cycling and nature of the additive is required.

The kinetic study of 0.95MgH2-0.05TMHx NCs revealed that the reversible capacity of MgH2-ZrH2, MgH2-VH, MgH2-NbH NCs diminishes at the very first sorption cycles because of their poor absorption kinetic. Future work on the influence of additive amount on kinetic and cycling properties, as well as on microstructural modifications (particle morphology and phase distribution), is of interest to gain a better understanding of additive effects .

The kinetic model here used for the analysis of sorption curves in MgH2-TiH2 NCs can be further pushed to extract rate constants and activation energies using the Arrhenius equation. Results can be compared to pure MgH2 for a better understanding of kinetic properties of MgH2-TMHx NCs,

The performances of the implemented automatic cycling device can be improved. The hydrogen capacity of the reservoirs can be improved by increasing the mass of the hydrideforming alloy. The Labview® program needs improvement to enrich simultaneous data acquisition of flow, pressure and temperature signals as well as further automatization of some manual parts.

Mg/transition-metal nanomaterials for efficient hydrogen storage

Magnesium metal is a prominent element for solid-state hydrogen storage due to its large abundance in earth's crust and its high weight and volumetric hydrogen uptakes. However, hydrogen sorption suffers from sluggish kinetics and the formed hydride is too stable for applications working under ambient conditions. The former issue can be solved by developing composites combining two hydrides, MgH2 and TiH2 at the nanoscale. These materials are synthesized by mechanical milling under reactive atmosphere. By this technique, the formation of nanocomposites and their hydrogenation can be obtained in a single-step. Moreover, these materials can be produced at large scale for application purposes. The work focused on three topics: i) the optimization of the TiH2 content in the (1-y)MgH2+yTiH2 system. This was accomplished by optimizing the titanium content (0.0125≤y≤0.3 mole), while keeping good kinetics, hydrogen reversibility and cycle-life. The data show that y=0.025 is the best compromise to fulfill the most practical properties; ii) the extension to other transition metals for the system 0.95MgH2+0.05TMHx (TM: Sc, Y, Ti, Zr, V and Nb), evaluating the contribution of each additive to kinetics, hydrogen reversibility and cycle-life; iii) the conception of an automatic cycling device able to carry out hundreds of sorption cycles whit the aim of measuring the cycle-life of metal hydrides. The work was done using manifold experimental methods. For synthesis, reactive ball milling under hydrogen atmosphere was primarily used. The crystal structure and the chemical composition of nanomaterials was determined from X-ray diffraction (XRD) analysis. Particle size and morphology were obtained by Scanning Electron Microscopy / Energy Dispersive X-Ray Spectroscopy (SEM/EDS). Thermodynamic, kinetic and cycling properties toward hydrogen sorption were determined by the Sieverts method. Keywords : magnesium hydride, hydrogen storage, nanocomposites, metal hydrides.

Nanomatériaux à base de magnésium et de métaux de transition pour un stockage efficace de l'hydrogène

Le magnésium est un élément de choix pour le stockage de l'hydrogène à l'état solide en raison de sa grande abondance dans la croûte terrestre et de ses fortes capacités de sorption massique et volumétrique de l'hydrogène. Cependant, la réaction de sorption souffre d'une cinétique lente et l'hydrure formé est trop stable pour des applications fonctionnant sous conditions ambiantes. Le premier problème peut être résolu en développant des composites associant deux hydrures, MgH2 et TiH2, à l'échelle nanométrique. Ces matériaux sont synthétisés par broyage mécanique sous atmosphère réactive. Cette technique permet la formation des nanocomposites et leur hydrogénation en une seule étape. De plus, ces matériaux peuvent être produits à grande échelle pour les besoins des applications. Les travaux ont été menés en trois parties : i) l'optimisation de la teneur en TiH2 dans le système (1-y)MgH2+yTiH2. Ceci a été accompli en ajustant la teneur en titane (0,0125 ≤ y ≤ 0,3 mole), tout en conservant une bonne cinétique, une réversibilité de l'hydrogène et une durée de vie utile. Les données montrent que la valeur y = 0,025 offre le meilleur compromis pour développer les propriétés les plus adéquates; ii) l'extension à d'autres métaux de transition pour le système 0,95MgH2 + 0,05TMHx (TM: Sc, Y, Ti, Zr, V et Nb), en évaluant la contribution de chaque additif sur la cinétique, sur la réversibilité de l'hydrogène et sur la durée de vie en cyclage; iii) la conception d'un dispositif de cyclage automatique capable de réaliser des centaines de sorption/désorption dans le but de mesurer la durée de vie des hydrures métalliques. Le travail a été effectué à l'aide de nombreuses méthodes expérimentales. Pour la synthèse, le broyage réactif sous atmosphère d'hydrogène a été principalement utilisé. La structure cristalline et la composition chimique des nanomatériaux ont été obtenues à partir de l'analyse par diffraction des rayons X (DRX). La taille et la morphologie des particules ont été déterminées par microscopie électronique à balayage et spectroscopie de rayons X à dispersion d'énergie (SEM/EDS). Les propriétés thermodynamiques, cinétiques et cycliques de la sorption d'hydrogène ont été déterminées par la méthode de Sieverts. Mots clés : hydrure de magnésium, stockage de l'hydrogène, nanocomposites, hydrures métalliques.