Acknowledgement

As this thesis work is mainly done in the laboratory Femto-ST MN2S, as well as supported by China Scholarship Council and UTBM in the framework of UT-INSA (2013), my thanks firstly goes to all the collaborating parties which made this work possible at all. I would express my greatest gratitude to my thesis supervisors Prof. Omar ELKEDIM, who has been extremely supportive along this journey of scientific research, where difficulties are only the stairs for one to climb up. It is not only the knowledge I could learn from his, but also the confident and cheerful attitude toward obstacles is of great inspiration to me on this road of science in the future. I would also like to say thanks to him for his patience to guide me through all this work. I am very thankful for Mrs. BERTRAND Ghislaine, Mr. HAMZAOUI Rabah and Mr.

Michel DEVEL to take the time to be the juries and help improve this thesis a lot. And my sincere thanks also go to Mr. Chonghe LI who was unfortunately not able to present my thesis defense, but his supports and encouragements is the greatest treasures for me which always inspire me to keep on.

It is also my great pleasure to thank Mr. Mieczyslaw JURCZYK and his friendly group, in particular Mr. Mateusz BALCERZAK. I spent one month with them in Poznan university of technology; it was a great experience working with them and learning from them. His enthusiasm toward science and the passion for life will always be my inspiration and spur me to pursue excellence when doing science.

IV

Thanks also go to Mr. Michel MOLIERE, who is very kind to provide lots of assistances to improve my thesis work on both the methods and technologies.

I have received countless help from many nice working fellows, my colleagues in our lab. Finally, I owe to my parents at the moment, who have been the ones standing firmly behind my back, despite the physical distance between us. Like always, during my study in France, they have been giving me wise advice and are always helping me to be a better person. My sisters deserve my wholehearted thanks as well for their endless love, support and encouragement.

IV

List of figures

INTRODUCTION

Hydrogen is a clean and renewable energy carrier, which can be used to store and transport energy. It is an ideal energy resource as a cold combustion fuel to produce electrical energy, e.g. in fuel cell due to its higher efficiency and zero pollution compare to diesel or gas engines. However, before the arrival of hydrogen-driven economy, some major technical challenges need to be conquered, one of which is finding an efficient and safe hydrogen storage medium with a small volume and low weight. A few currently available technologies permit to store directly hydrogen by modifying its physical state in gaseous or liquid form. But these methods have shortcomings such as: safety concerns, energy loss during operation, cost of recycling and charging infrastructures etc... In addition to these traditional stored form, solid-state hydrogen storage methods bring new possibilities for on-board utilization of hydrogen. A great deal of effort has been made on new hydrogen-storage systems, including metal, chemical or complex hydrides and carbon nanostructures. Metal hydrides are some metals and alloys which can reversibly absorb and desorb hydrogen under moderate temperature and pressure which gives them an important safety advantage over the gas and liquid storage methods. For the vehicle application, depending on the temperature of hydrogen absorption/desorption below or above 150 °C, the alloy hydrides can be classified in high and low temperature materials. The principal disadvantages of alloy hydrides, apart from the cost, are the low hydrogen content at low temperature (e.g. Labased alloys) and the difficulty of reducing desorption temperature and pressure of alloy hydrides having high hydrogen storage capacity and fast rate (e.g. Mg-based materials).

To solve the above mentioned problems, composite materials, partial substitution and modification of composition have been systematically studied for synthesizing novel catalyzed metal hydrides.

Additionally, another notable application area for hydrogen storage material is in Nickel-Metal Hydride batteries. The Ni-MH battery has been widely employed for high-current drawing devices, including mobile phones and laptops, due to its safety, environmental friendliness and overall charge-discharge performance. The metal hydrides used as the negative electrodes of Ni-MH batteries are the key components for battery performance, namely discharge capacity, high-rate capability, initial activation, cycling life, self-discharge rate. In order to compete with the Li-ion battery in emerging electric vehicle (EV) and hybrid electric vehicle (HEV) market, the energy density of the Ni-MH battery needs to be greatly increased. Replacing Li-ion batteries is preferable as these batteries require expensive special safety equipment that prevents over discharging the battery, whereas Ni/MH batteries do not require any special safety equipment. In recent years, a lot of high performance hydrogen storage alloys have been developed. Among them, magnesium based and titanium based alloys have attracted much attention due to their extremely high discharge capacity and low cost.

Theoretically, the hydrogen storage abilities of MgH2 and TiH2 are 7.6 wt.% and 4.0 wt.% respectively. However, due to the irreversible hydrogen absorption reaction of TiH2 at room temperature and the low decomposition-formation kinetics of MgH2, neither of them can be practically applied. In order to improve their hydrogen storage properties, transition metal elements have been introduced as electrocatalysts into the system. For example, amorphous MgNi was measured to process a high electrochemical capacity reaching 500 mAh/g, but with a bad cycling performance. The crystalline Mg2Ni can reversibly absorb and desorb hydrogen at high temperature.

However, the poor kinetics and bad anti-oxidation of Mg2Ni cannot be ignored. TiNi phase was found to possess a capacity of 230 mAh/g and good corrosion resistance in alkaline solution. Unfortunately, the poor absorption/desorption kinetics and complex activation procedure are obstacles for the practical application of TiNi alloys as hydrogen storage material.

These two binary alloys are far from being operational in practical applications due to their mediocre cycling stability and relatively low discharge capacity. Fulfilling all the requirements for optimum performance of hydrogen-based batteries remains thus a real challenge. To overcome this problem, composite hydride materials have been introduced as electrode materials in an attempt to improve the balance of discharge capacity and cycling performances. Alloys pertaining to the Ti-Mg-Ni ternary system are of interest as potential metal hydride since they consist of two hydride forming metals of Ti and Mg. The hydrogen storage properties of these alloys could be upgraded with Ni additions thanks to the high catalytic activity of the latter. A lot of researchers have made efforts to optimize the composition for exploring the balance of discharge capacity and cycling performance. There is therefore a vivid interest for the development of such ternary Ti-Mg-Ni systems in view of hydrogen storage applications.

The objective of the present study is to explore the hydrogen storage properties of Ti-Mg-Ni ternary alloys system and investigate the synergistic effect of different phases on the enhancement of electrochemical performance. Furthermore, this thesis studied the Ti-Ni system with Cu substitution. Meanwhile, the structure of substituted alloys was simulated with first principles calculations in order to have qualitative insight into the variation of their properties. High energy ball milling (HEBM) and subsequent heat treatments have emerged as the technology of choice for their preparation.

In Chapter I, the metal hydride as a hydrogen storage material and the working mechanism of Ni-MH system as a new form of energy storage form is briefly presented.

Afterwards, an introduction on the main varieties of metal hydrides namely Ti-Mg system, Ti-Ni system, Mg-Ni system and Ti-Mg-Ni ternary system are given with an emphasis on recent research on their use as the negative electrode material for Ni-MH batteries.

In Chapter II, the high energy ball milling (HEBM) preparation method as well as its applications in the hydrogen storage material field have been presented. The main instruments for structural characterization and hydrogenation properties measurements used in this study are introduced. This entails of X-Ray Diffraction, Scanning Electron In chapter V, the main points of this study which concern the products preference and hydrogenation properties of Ti-Mg-Ni alloys as well as the experimental and computational results of the influence of Cu substitution on the structural and electrochemical properties of TiNi alloys have been summarized. Prospective views on current research as how it can be continued for further investigation in the search of ideal materials for the negative electrode for Ni-MH batteries have been given.

In appendix I, the publications and attended conferences during PhD studies are listed.

CHAPTER I

Overview of metal hydrides and their applications in Ni-MH batteries

This chapter gives a brief review on hydrogen storage materials and their application as the main materials in the negative electrode of nickel/metal hydride (Ni-MH) secondary batteries. The Ti-Mg-Ni ternary alloys, as the main research object in present study, are introduced from the aspects of Ti-Mg system, Mg-Ni system, Ti-Ni system and Ti-Mg-Ni system. The common ways for manufacturing these hydrogen storage materials are also included.

I.1 Hydrogen storage methods

As energy crisis and environmental pollution are becoming more and more serious, it is very urgent to seek alternative clean energy resources to replace the traditional fossil fuel. Hydrogen is widely regarded as one of the best solutions due to the following reasons [1,2]: (1) Readily Available: Hydrogen is the most abundant element on earth and can never run out theoretically. (2) No Harmful Emissions: Using hydrogen as a fuel is clean and renewable. (3) Environment Friendly: Hydrogen is also non-toxic, which makes it a rarity among fuel sources. (4) Fuel Efficient: Compared to traditional fossil fuel, hydrogen is much more fuel efficient as it can produce more energy per pound of fuel. Especially, since the development of the proton exchange membrane (PEM) fuel cell, which takes in hydrogen gas as fuel and produces only water, hydrogen appears as the most promising energy carrier and has attracted increasing research attentions.

However, a key step for the development of hydrogen technology in application including stationary power and portable power requires overcoming the technical challenge of storage and transport of hydrogen. Hydrogen is very hard to move around compared to traditional energy carriers. Many approaches have been implemented for solving this issue which can be simply classified as physical storage and chemical storage [3]. In this chapter, we will limit our discussion to the hydrogen storage in gaseous form, liquid form and metal hydride, each of these options possesses attractive attributes for hydrogen storage.

I.1.1 Hydrogen storage in gaseous form

Physical storage is the most mature hydrogen storage technology. Hydrogen can be stored physically as either a gas or a liquid. The storage of hydrogen involves a balance and compromise between two different aspects of the storage system; the volumetric and gravimetric capacities. The application scenario of hydrogen results in its preferable storage technology.

Hydrogen has been stored as compressed gas in storage tanks since the beginning of the 20 th century. The tanks can either be made with steel, aluminum or copper alloys that may be encased in fiberglass. The steel tanks are most often used for stationary applications where weight is not a hindrance as the steel tanks tend to be bulky and heavy. High pressure tanks also appear in some test automobiles but safety and space remain significant concerns. The major drawback of gaseous storage for transportation applications is the small amount of hydrogen that may be stored in a reasonable volume.

Even though 10,000 psi (~70 MPa) hydrogen tanks [4] have been successfully employed commercially , the energy content of such hydrogen tanks is significantly less than that for the same volume of gasoline, 4.4 MJ/L (10,000 psi) for hydrogen compared with 31.6 MJ/L for gasoline [5]. Another consideration in the storage of compressed hydrogen requiring special caution is the problem of hydrogen embrittlement, because many of the system parts exposed to hydrogen are metallic. The hydrogen embrittlement of metallic hydrogen tanks occurs under high pressure and high temperature conditions and may result in the cracking of tanks. Thus, all these factors raise a big challenge for the development of tank materials, which should be not only reliable enough but also light of the same time.

I.1.2 Hydrogen storage in liquid form

Hydrogen in liquid form has a considerably higher energy density than in its gaseous form, making it an attractive storage medium. Hydrogen can be stored as a liquid at 21.2 K (-251.95°C) at ambient pressure in cryogenic tanks because the boiling point of hydrogen at one atmosphere pressure is -252.8°C. It is predicted that the energy content of liquid hydrogen is 8.4 MJ/L [6] compared with 4.4 MJ/L for compressed gas (at 10,000 psi). A major drawback of liquid storage is the large quantity of energy required for liquefaction. The cooling and compressing process requires energy, resulting in a net loss of about 30% of the energy stored in the liquid hydrogen.

High-pressure cryogenic tanks are also being explored to alleviate the requirement for very low temperature storage [7]. The main task is to reduce the cost of liquefaction and make it more energy efficient. Besides, the problem of hydrogen boil-off is another unavoidable phenomenon when hydrogen is stored as a liquid form. The boil-off rate of hydrogen from a liquid hydrogen storage vessel is associated with the size, shape and thermal insulation of the vessel. The best shape for vessel is a sphere, however, large-sized spherical containers are expensive due to the manufacturing difficulty. For these reasons, the practical application of liquid hydrogen storage system is quite limited which are mainly in the fields that, the cost of hydrogen is not an important issue and the hydrogen is consumed in a relatively short time, e.g. space applications and military aircrafts [8]. shows the volumetric density of hydrogen inside the cylinder and the ratio of the wall thickness to the outer diameter of the pressure cylinder for stainless steel with a tensile strength of 460 MPa [5]. It is shown that the volumetric density of liquid hydrogen is steady and always higher than that of compressed hydrogen gas. The volumetric density of compressed hydrogen gas increases with pressure and reaches a maximum above 1000 bar, depending on the tensile strength of the material. However, the gravimetric density decreases with increasing pressure and the maximum gravimetric density is found for zero overpressure. Therefore, the increase in volumetric storage density is sacrificed with the reduction of the gravimetric density in pressurized gas systems [8].

I.1.3 Hydrogen storage in metal hydride

Although with the compressed and liquid storage options, hydrogen is accessible for use, these storage methods still cannot meet either the options of DOE targets or practical demands of on-board hydrogen storage system for mobile applications.

Hydrogen storage requires a great technological discovery, that is, chemical storage, where hydrogen is stored in hydrides. Through a chemical reaction, solid metal hydrogen compounds are formed, under hydrogen pressure, and heat is released. For example, 4 kg of hydrogen will be needed to enabled a vehicle with a fuel cell engine to drive a distance of 400 km. Metal hydrides are kind of solid metal-hydrogen compounds which are generated by the reaction of metals, intermetallic compounds and alloys with hydrogen. They can be grouped into three basic types according to the nature of the metal-hydrogen bonds: ionic, volatile covalent and metallic hydrides [8]. Since the first hydrogen absorption phenomenon was found for palladium, various metallic compounds have been proposed as hydrogen storage materials. Actually, almost all of the elements can form binary hydrides with hydrogen, except the inert gas elements.

But due to the problems of capability and reversibility, neither of them can be practically used. The intermetallic hydrides, by contrast, are more attractive because they can adsorb and desorb hydrogen under moderate conditions.

As seen in figure I. 3, which illustrate the schematic model of a metal structure with H atoms in the interstices, charging process can be done using molecular hydrogen gas or hydrogen atoms from an electrolyte [6]. The hydrogen molecule is first absorbed on the surface and then dissociated as strongly bound, individual hydrogen atoms. The hydrogen atoms can diffuse into the metal and form a solid solution of hydrogen atoms dissolved in the metal lattice, which is usually referred to as the α-phase. When the hydrogen pressure increases and the concentration exceeds a certain critical value, a higher concentration hydride phase (β-phase) starts to form.

While the pure β-phase is obtained, α-phase disappears in the mean time, hydrogen enters as the solid solution of β-phase. On the contrary, when the hydrogen needs to be used, it is released from the hydride under certain temperature and pressure conditions following the reversible process. This process can be repeated many times without loss of storage capacity. Intermetallic hydrides contain at least two metals in addition to hydrogen. The general formula of two-metal hydrides are AmBnHx, where metal A has a strong affinity for hydrogen and forms a stable binary hydride while metal B does not interact with hydrogen, such as LaNi5 (CaCu5 structure type), TiNi (CsCl structure type) and Mg2Ni (AlB2 structure type) etc. The hydrides will be formed when the intermetallic expose to hydrogen and follow the reaction:

M + x 2 H 2 ↔ MH x + Q EQN I.1.
Where M represents the metal and Q is the heat of reaction of hydride formation.

The thermodynamic aspects of the hydride formation from gaseous hydrogen can also be described by the means of pressure-composition isotherms (as shown in Figure I. 4). Metal and hydrogen usually form two different kinds of hydrides, α-phase at which only some hydrogen is absorbed and β-phase at which hydride is fully formed.

While the two phases coexist, the isotherms show a flat plateau (in the middle of the curve). The length of this plateau reveals the amount of stored hydrogen.

Accompanying the increase of temperature, the pressure of its plateau increases linearly until the critical point, Tc, above which the phase transition of α-phase to βphase takes place continuously [12]. Where P is the equilibrium pressure and R is the gas constant (8.3145 J mol -1 K -1 ). ΔH and ΔS are respectively the enthalpy and entropy of the formation of the hydride (β-phase). The slope of the line is equal to the dehydrogenation enthalpy divided by the gas constant (ΔH/R), and the intercept is equal to the entropy of formation divided by the gas constant (ΔS/R).

I.

And therefore the overall reaction is represented as:

(Overall) Ni(OH)2 + M = NiOOH + MH ( 3
)
Where M is the hydrogen storage metal/alloy and MH is the hydride of metal M. The key to the whole reaction is in the negative electrode material where metal hydride is formed, and is also where all the research in the field has focused on. For better understanding of the charging and discharging process, the reversible reaction (2) can be decomposed into the following steps:

M + H2O + e¯ ↔ MHads + OH¯ (initial hydrogen adsorption)

MHads ↔ MHabs (hydrogen diffusion)

Habs ↔ MH hydride (metal hydride formation) (6) 2MHads ↔ M + H2 (hydrogen evolution)

MHads + H2O + e¯ ↔ M + H2 + OH¯ (hydrogen evolution)

Where Hads, Habs represent adsorbed hydrogen on the alloy surface, absorbed hydrogen in solid solution. The steps of hydrogen adsorption/desorption, hydrogen diffusion, and formation/decomposition of metal hydride can be illustrated by the reaction Eq.( 4) - (6).

During charge, Hydrogen is firstly produced by Eq. ( 4), known as Volmer reaction. A negative voltage (with respect to the counter electrode) is applied to the metal/metal hydride electrode current collector, and electrons enter the metal through the current collector to neutralize the protons from the splitting of water that occurs at the During discharge, the produced hydrogen can either be absorbed further into the alloy as solid solution, or start the hydrogen evolution reactions, which are the Tafel reaction (Eq. ( 7)) or even Heyrovsky reaction (Eq. ( 8)). Protons in the MH leave the In recent years, a lot of high performance hydrogen storage alloys have been developed. Among them, magnesium based and titanium based alloys have attracted much attention due to their extremely high discharge capacity and low cost.

Theoretically, the hydrogen storage abilities of MgH2 and TiH2 are 7.6 wt. % and 4.0 wt. % respectively. However, due to the irreversible hydrogen absorption reaction of TiH2 at room temperature and the low decomposition-formation kinetics of MgH2, neither of them can be practically applied. In order to optimize the hydrogen storage property of the light weighted metal hydrides, many efforts have been done on the Ti-Mg system. It has been identified as potential materials for hydrogen storage because of their safety, small volume and low weight. It is known that, titanium has a melting point (1943 K) that greatly exceeds the boiling point of magnesium (1363 K). In addition, the equilibrium solid solubility of each metal in the other is very weak and no intermetallic compounds can be found in Ti-Mg binary phase diagram (as shown in Figure I. 6). Therefore, alloying of Mg and Ti by rapid solidification or regular melting method is extremely difficult. Some preparation methods like, physical vapor deposition (PVD), electron-beam and magnetron co-sputter deposition techniques have been employed for manufacturing Ti-Mg alloys [18], but it remains the problem that these techniques can not be scaled-up to industrial level. Instead, Mechanical alloying (MA) is a good candidate which is suitable for fabricating bulk materials as required for Ni-MH battery and fuel cell applications and has been proved to be an excellent technique for extending terminal solid solubility [20].

Many experimental studies have been carried out for investigating the phase transformation of Ti-Mg alloys by means of mechanical alloying. Liang et al. [21] reported the extended dissolution of Ti in Mg up to a maximum of 12.5 at.% upon milling of Mg80Ti20 alloys. Rousselot et al. [22] have synthesized Mg100-xTix (40≤x≤ It is shown that the crystal structure of as-milled Ti-Mg alloys can be very different depending on the milling parameters and the atom ratio of Mg and Ti. The milling parameters can be further subdivided into milling speed, milling duration, grinding media and addition of PCA. Taking Ti50Mg50 as an example, both of Mg and Ti have a hexagonal close packed (HCP) structure, but the phase structures of milling products via different synthetic approaches are BCC and FCC, respectively [24,25]. These authors first synthesized a BCC Mg50Ti50 alloy by ball milling a mixture of 50Mg + 50Ti in a Fritsch P5 planetary ball mill with still balls for 200 h at a rotation speed of 200 rpm. While using zirconia milling balls and pots for ball milling using the same starting material in the same milling conditions, a FCC Mg50Ti50 alloy was obtained.

These variations of phase structures of as-milled samples with the same composition are mainly explained by the plastic deformation process of raw materials during ball milling. The density of stacking faults introduced by milling in the BCC phase is higher than that in the FCC phase because the dynamic energy given by stainless steel balls during milling was higher than that by zirconia ones. The main plastic deformation modes of Ti and Mg are twinning deformation and plane slips, respectively [26]. The different plastic deformation modes caused by various milling conditions further affect the phase structures of milling products. In another investigation of Mg50Ti50 alloys prepared by high-energy mill, twinning deformation was also observed in Ti-rich crystallites at intermediate milling time [START_REF] Anik | Effect of titanium additive element on the discharging behavior of MgNi alloy electrode[END_REF]. They attributed the twinning to the deformation of Ti particles. But they also pointed out that in the Mg-Ti system it might also indicate a strain-induced martensitic transformation of the metastable ω-FCC into BCC. The crystallite boundaries acted as preferential sites for the heterogeneous nucleation of the twins and for the formation of solid solution by release of the lattice strain energy.

By controlling milling conditions and atom ratio of Mg and Ti, Asano et al. [25] have also shown that BCC, FCC or HCP phases could be obtained in the Mg-Ti system by ball milling (as shown in While MgH2 is used instead of Mg as the starting material then, after ball milling a 50MgH2 + 50Ti mixture, the resulting compound is FCC Mg33Ti50H94 and trace amount of MgH2 [27]. Meanwhile, the importance of mechanical effect during milling has also been discussed [28]. It shows that during ball milling of Mg and Ti powders in molar ratio of 1:1, plate-like particles first stick on the surface of the milling pot and balls.

After these plate-like particles fall off from the surface of the milling pot and balls, spherical particles, in which concentric layers of Mg and Ti are disposed, are formed.

These particles have an average diameter of 1 mm. These spherical particles are then crushed into spherical particles with a diameter of around 10 µm by introduction of cracks along the boundaries between Mg and Ti layers. Finally, the Mg50Ti50 BCC phase with a lattice parameter of 0.342(1) nm and a grain size of 3 nm is formed. During milling, Ti acts as an abrasive for Mg which has stuck on the surface of the milling pot and balls. Machio et al. [29] investigated the relationship between the milling time and the thermal stability of Ti-Mg powders which suggested that milling time needs to be optimized for any alloy system since it affects green densities and further processing of the milled powder. Meanwhile, the particle size of starting materials also have significant influence on the milling efficiency. Starting with fine titanium and magnesium particles in ball milling could result in a higher milling yield [24].

Increasing the amount of process-control agent (PCA), could also improve this phenomenon, however, contamination of the milled powder may be a shortcoming to this solution [30]. Kalisvaart et al. [23] reported that the lattice parameter of Ti-Mg FCC phase depended on the composition, i.e., the smaller FCC phase was Ti-rich and the larger one was Mg-rich. It is notable that, all these variations of the lattice parameter and the crystal structure of the milling products could further affect their hydrogen storage property significantly which should be brought to the attention.

Mg-Ti alloys are promising candidates as negative electrode materials for Ni-MH batteries, due to their high hydrogen solubility. Based on the fact that the hydrogen atoms occupy the interstitial sites of crystal lattice in metal hydrides. A BCC structure has more interstitial sites than FCC and HCP structures. A body-centered cubic structure is a coarse packing structure and has more interstitial space than face-centered cubic and hexagonal close-packed structures [31]. Thus, BCC alloys are more attractive candidates to be explored as possible interstitial hydrogen storage materials. The best electrochemical performances, reported by Vermeulen et al. [32], for Ti-Mg based alloys were prepared by mean of electron beam deposition. Discharge capacities were from 750 mAh g -1 in Mg50Ti50 to 1750 mAh g -1 in Mg80Ti20. However, unlike the Ti-Mg thin films, when prepared by mechanical alloying, Mg-Ti materials must be activated for electrochemical application by adding a few at.% of palladium [33].

Without adding Pd as catalyst, the hydrogen discharge capacity of as-milled Mg-Ti materials is close to zero. Rousselot et al. [33] reported a maximum discharge capacity of as-milled Mg50Ti50 alloy with 10 at.% Pd is 400 mAh g -1 after three charge/discharge cycles. The structure of the alloy evolves from a HCP to a FCC structure upon cycling, and does not revert back to its original structure upon dehydrogenation.

A possible explanation for the function of Pd is that the dissolution of Pd changes the electronic structure of Ti-Mg alloy [START_REF] Nambu | Correlation between electronic structure and phase stability of metal hydrides[END_REF][START_REF] Rousselot | Influence of Pd on the structure and electrochemical hydrogen storage properties of Mg50Ti50 alloy prepared by ball milling[END_REF]. The addition of Pd decreases the energy of the hydrogen-sites in the Mg-Ti alloys and has a beneficial effect on their electrochemical hydrogen discharge. Nevertheless, up to now, there is no alternative solution for replacing Pd as the electrochemical catalyst for as-milled Ti-Mg alloy which has become the major disadvantage for its practical application considering the extremely high cost of Pd.

I.3.2 Hydrogenation properties of Ti-Ni system

Ti-Ni based alloys are quite attractive functional materials not only as practical shape memory alloys with high strength and ductility but also as those exhibiting unique physical properties and hydrogen storage performance. The memory behavior of TiNi is due to a reversible structural martensitic transformation (MT) from a high temperature phase with cubic structure and a low temperature phase with monoclinic structure [START_REF] Michal | The structure of TiNi martensite[END_REF]. Besides, there are many phase transformations in Ti-Ni based system, which include not only martensitic transformations, but also diffusional transformations.

Let us note that, the phase diagram of NiTi has been controversial for nearly three decades before some consensus was established. The most reliable Ti-Ni phase diagram, according to the discussion of Otsuka and Ren [START_REF] Otsuka | Physical metallurgy of Ti-Ni-based shape memory alloys[END_REF], is shown in

I.3.2.1 TiNi in Ni-MH system

TiNi phase that has a CsCl type structure (space group pm3m-221), with a lattice constant of 3.15 nm at room temperature, which is also referred to as the high temperature phase or austenitic phase. The other type of TiNi is the martensite phase with a monoclinic B19' structure (space group P21/m) [START_REF] Akiyama | Hydriding combustion synthesis for the production of hydrogen storage alloy[END_REF]. The first investigation of TiNi as a reversible negative electrode for Ni-MH batteries was carried out by Justi et al. [START_REF] Justi | Electrocatalysis in the nickel titanium system[END_REF] in 1970. Afterwards, TiNi hydride received extensive research, thanks to its rather low specific weight, good oxidation resistance and notable discharge capacity: 240 mAh/g [START_REF] Machida | The hydrogen electrode reaction characteristics of thin film electrodes of Ni-based hydrogen storage alloys[END_REF]. Besides, TiNi alloy is also the first alloy studied as a hydrogen storage material, which can absorb hydrogen as much as 1.4 H/f.u. (formula unit) under room temperature and pressure [START_REF] Burch | Absorption of hydrogen by titanium-cobalt and titanium-nickel intermetallic alloys. Part 1.-Experimental results[END_REF].

Lots of attempts have been made in order to improve the hydrogen storage properties of TiNi alloy through substituting element or refining the microstructure.

Various alloying elements have been used to improve the electrochemical capacity and cycle life of TiNi-based ternary alloys by the means of ball milling, e.g. Mg [START_REF] Szajek | Electrochemical and electronic properties of nanocrystalline TiNi1-xMx(M= Mg, Mn, Zr; x= 0, 0.125, 0.25) ternary alloys[END_REF], Mn [START_REF] Drenchev | Influence of alloying and microstructure on the electrochemical hydriding of TiNi-based ternary alloys[END_REF], Zr [START_REF] Drenchev | Influence of alloying and microstructure on the electrochemical hydriding of TiNi-based ternary alloys[END_REF][START_REF] Guiose | Solid-gas and electrochemical hydrogenation properties of pseudo-binary (Ti, Zr) Ni intermetallic compounds[END_REF], Cu [START_REF] Emami | Ti (Ni, Cu) pseudobinary compounds as efficient negative electrodes for Ni-MH batteries[END_REF], Fe [START_REF] Jurczyk | Nanocrystalline titanium-type metal hydride electrodes prepared by mechanical alloying[END_REF] and B [START_REF] Drenchev | Influence of B substitution for Ti and Ni on the electrochemical hydriding of TiNi[END_REF]. Szajek et al. [START_REF] Szajek | Electrochemical and electronic properties of nanocrystalline TiNi1-xMx(M= Mg, Mn, Zr; x= 0, 0.125, 0.25) ternary alloys[END_REF] reported that the respective replacement of Ni in TiNi by Mg, Mn or Zr improved not only the discharge capacity but also the cycle life of these electrodes. They gave an explanation using band structure calculations, which showed that the introduced impurities cause decrease of the width of the valence bands. Emami et al. [START_REF] Emami | Hydrogenation properties of shape memory Ti (Ni, Pd) compounds[END_REF][START_REF] Emami | Electronic and structural influence of Ni by Pd substitution on the hydrogenation properties of TiNi[END_REF] investigated the influence of Pd by partially replacing Ni [163,164]. It is indicated that the unit cell of TiNi is enlarged by the doped Pd, the electrochemical capacity is reduced since the Pd-substitution increases the stability of TiNi intermetallic alloy and decreases the stability of their hydrides according to the electronic calculations. In addition, some other studies on substituting Ni in TiNi with various modifiers revealed that Fe, Co, and Cr increase the electrochemical capacity up to 400 mAh g -1 with good activation performance [START_REF] Young | Effects of Cr, Zr, V, Mn, Fe, and Co to the hydride properties of Laves phase-related body-centered-cubic solid solution alloys[END_REF].

The partial substitution of B for Ti and Ni realized by ball milling showed a formation of solid solutions of B in TiNi after milling, and subsequent annealing yields TiNi and TiNi3 phases, whose discharge capacity was found lower than the initial TiNi alloy [START_REF] Drenchev | Influence of B substitution for Ti and Ni on the electrochemical hydriding of TiNi[END_REF].

Other preparing methods have also been employed for manufacturing TiNi based alloys and shown interesting results. Cuevas et al. [START_REF] Cuevas | Influence of the martensitic transformation on the hydrogenation properties of Ti50-xZrxNi50 alloys[END_REF] have reported a series of Ti50-xZrxNi50 (x = 0, 6, 12, 18 and 24 at. %) alloys with austenite and martensite structure by means of induction melting and melting-spinning, respectively. They also pointed that the type of parental crystal structure has a big influence on their hydrogenation properties. As for austenitic type (CsCl-type structure), the hydride preserves its metal sublattice and forms Ti0.64Zr0.36NiH1.6 (PH2 = 1 MPa and T = 373 K), whereas the monoclinic TiNi-type structure of martensitic Ti0.64Zr0.36Ni is hydrogenated into two coexisting hydrides, Ti0.64Zr0.36NiH and Ti0.64Zr0.36NiH2.2 [START_REF] Cuevas | Influence of polymorphism on the electrochemical properties of (Ti0.64Zr0.36)Ni alloys[END_REF]. Martensitic transformation in TiNi-alloys is promoted by Zr addition, hence can be used to tailor the final product of alloying Ti-Ni-Zr. Martensitic type structure can increase the hydrogenation percentage of TiNi significantly with an obvious PCI plateau, while the austenitic type shows no plateau on the PCI curve. Another study from Emami et al.

I.3.2.2 Ti2Ni in Ni-MH system

Another binary hydride with an fcc structure which is able to store considerable amount of hydrogen in Ti-Ni system is Ti2Ni. The first study of Ti2Ni as hydrogen storage material was done in 1972, from which, Ti2Ni hydride was announced to have a capacity as high as 500 mAh/g in a gaseous hydrogen absorption reaction [START_REF] Buchner | Wasserstoff in intermetallischen phasen am beispiel des systems titan-nickel-wasserstoff[END_REF]. Later studies were proposed by Luan et al. [START_REF] Luan | Mechanism of early capacity loss of Ti2Ni hydrogenstorage alloy electrode[END_REF] using arc melting. They obtained only 160 mAh/g with poor cycling stability. The severe capacity loss during cycling limited the application of Ti2Ni alloy, which was attributed to the formation and accumulation of irreversible Ti2NiH0.5 phase for hydrogen absorption/desorption, based on the results of X-ray diffraction and charge and discharge testing [START_REF] Luan | Effects of potassium-boron addition on the performance of titanium based hydrogen storage alloy electrodes[END_REF]. Zhao et al. [START_REF] Zhao | Effect of mechanical milling on the structure and electrochemical properties of Ti2Ni alloy in an alkaline battery[END_REF] conducted investigations on the performance of amorphous Ti2Ni alloy. They found that amorphous phase of Ti2Ni contribute to the promotion of the cycling performance. The alloy was synthesized by solid-state sintering and then ball milling to obtain an amorphous phase, which had a rather stable discharge capacity that ranges from 100 mAh/g to 125 mAh/g. In comparison, the discharge capacity of non-milled Ti2Ni

(alloyed by solid-state sintering) reached 280 mAh/g in the first cycle, but dropped to less than 100 mAh/g after 50 cycles. Hosni et al. [START_REF] Hosni | Structure and electrochemical hydrogen storage properties of Ti2Ni alloy synthesized by ball milling[END_REF] investigated the influence of the working temperature on the electrochemical performance of Ti2Ni alloys. They indicated that, the electrochemical discharge capacities and cycling suitability of Ti2Ni alloys are increased and decreased respectively with the increase of temperature. In the aspect of Ti2Ni as hydrogen storage material, the highest hydrogen storage capacity obtained so far is 2.1 wt.% reported by Balcerzak et al. [START_REF] Balcerzak | Hydrogenation properties of nanostructured Ti2Ni-based alloys and nanocomposites[END_REF], in which, nanostructured Ti2Ni was prepared by mechanical alloying and annealing with addition of 5 wt.% of Pd.

Elemental substitution of Ti2Ni alloy showed slight improvements in the electrochemical hydrogenation properties. Substitution of Al and B for Ni was carried out by arc melting, but the improvement was not satisfactory [START_REF] Luan | Studies on the performance of Ti2Ni1-xAlx hydrogen storage alloy electrodes[END_REF]. Li et al. [START_REF] Li | Structural characterization and electrochemical hydrogen storage properties of Ti2-xZrxNi (x= 0, 0.1, 0.2) alloys prepared by mechanical alloying[END_REF] studied the effect of Zr substitution on the electrochemical properties of Ti2Ni alloys, in which, the discharge capacity drops drastically with Zr substitution, but with increase of Zr content (from x = 0.1 to x = 0.2), the discharge capacity increases generally, which credits to larger unit-cell-volume caused by Zr substitution. Hiroyuki et al. [START_REF] Takeshita | Hydrogenation characteristics of Ti2Ni and Ti4Ni2X (X= O, N, C)[END_REF] experimented on the hydrogenation process of the ternary compound Ti4Ni2X (X = O, N, C). They found that these alloys all had higher hydrogen desorption pressure compared to Ti2Ni, and presented a hydrogen pressure plateau. Another fabrication method was developed by Zhao et al. [START_REF] Zhao | Electrochemical hydrogen storage properties of a nonequilibrium Ti2Ni alloy[END_REF] for Ti2Ni, which combined induction melting, ball milling and annealing. The product is a mixture of amorphous and nanocrystalline and demonstrates reasonable capacity and good cycle stability.

Moreover, the composite hydrides of TiNi and Ti2Ni also drew others' attentions.

By combining TiNi and Ti2Ni phases, the electrochemical capacity increases up to 320 mAh g -1 which is attributed to the synergetic effect of these two phases [START_REF] Gutjahr | A new type of reversible negative electrode for alkaline storage batteries based on metal alloy hydrides[END_REF]. During hydrogen desorption, TiNi, which has better desorption kinetics, first dehydrides and contributes to the overall electrochemical capacity; then, the hydrogen stored in Ti2Ni

phase is transferred internally into the dehydrided portion of TiNi and discharged through TiNi. Without the assistance of TiNi phase, the hydrogen stored in Ti2Ni phase cannot be released due to the strong metal-hydrogen bond in Ti2Ni [18]. Whereas, in another study, amorphous Ti3Ni2 alloy was prepared by mechanical alloying aiming to combine advantages from TiNi and Ti2Ni, although its capacity was lower than the crystalline form, cycle stability was much improved. The electrochemical properties and the influence of small amounts of added elements (Cr, Co, Al) on the electrode characteristics of the hydrogen storage Ti3Ni2 alloy have been studied in detail as well [START_REF] Xu | The influence of small amounts of added elements on the electrode performance characteristics for Ti3Ni2 hydrogen storage alloy[END_REF].

I.3.3 Hydrogenation properties of Mg-Ni system

Mg-Ni-based alloys are considered to be promising candidates for hydrogen storage applications because of their low cost, light weight and rich mineral resources, as well as high theoretical hydrogen storage capacity (assuming the formation of its hydride is 3.6 mass%, which is approximately 2.7 times that of LaNi5 [START_REF] Iwakura | Hydriding and electrochemical characteristics of a homogeneous amorphous Mg2Ni-Ni composite[END_REF]). The hydrogen storage properties of conventional polycrystalline MgNi are rather poor, and therefore amorphous and nanocrystalline phases of Mg-Ni-based systems are often introduced for improving the hydrogen storage properties.

Amorphous structures, metastable in terms of thermodynamics, differ from the equilibrium, crystalline state essentially in the local configuration of atoms. The lack of a crystallographically defined structural unit results in a distribution of local atomic positions in the structure of the glass. As a consequence, amorphous structure can provide a wide distribution of available sites for hydrogen [START_REF] Eliaz | An overview of hydrogen interaction with amorphous alloys[END_REF][START_REF] Perdew | Generalized gradient approximation made simple[END_REF]. Therefore, the hydrogenation behavior of the amorphous structure is totally different from that of the thermodynamically stable, crystalline material, although both materials may have identical composition. Orimo and Fujii [START_REF] Orimo | Materials science of Mg-Ni-based new hydrides[END_REF] indicated that the dehydriding temperature of amorphous MgNi under argon was lower than that of conventional Mg2NiH4.

Nanocrystalline metals are different in microstructure from both polycrystalline and amorphous phases. Because of the reduction of crystallite size, a large number of interfaces and grain boundaries are produced, which can enhance the solubility and provide easier channel for the diffusion of hydrogen atoms, avoiding the long-range diffusion of hydrogen through an already formed hydride. It has been reported that the introduction of nanocrystalline Mg2Ni decreased the dehydriding temperature of Mg2NiH4 [START_REF] Huang | Structural characterization and electrochemical hydrogen storage properties of Mg2Ni1-xMnx (x= 0, 0.125, 0.25, 0.375) alloys prepared by mechanical alloying[END_REF].

It is worthy to point out that, amorphous and nanocrystalline phases often coexist and the nanocrystallites are embedded in the amorphous matrix. This microstructure characteristic is favorable for the hydrogen storage [START_REF] Huang | Mg2-xTixNi (x= 0, 0.5) alloys prepared by mechanical alloying for electrochemical hydrogen storage: Experiments and first-principles calculations[END_REF], which may be explained by an interaction between the amorphous regions of grain boundaries and the nano-grains.

The amorphous region is hydrogenated first, expands and so puts the encased nanograins under pressure, thus showing synergistic effects on the hydrogen storage properties of Mg2Ni-type alloy. Orimo and Fujii [START_REF] Orimo | Materials science of Mg-Ni-based new hydrides[END_REF] reported that the maximum hydrogen concentrations of the three regions had been experimentally determined to be 0.3 wt.% in the grain region of nanocrystalline Mg2Ni, 4.0% in the interface between Mg2Ni grains and 2.2% in the amorphous region.

Partial element substitution could change the composition of milled alloys and even introduce new phases. Therefore, this method has been widely used to improve hydrogen storage properties of Mg-Ni-based alloys. The partial substitution of Al for Mg could result in the formation of Al2O3 protective film on the surface of matrix alloy.

The Al2O3 plays a role of an inhibitor against further corrosion of the alloy. Therefore, the partial substitution of A1 for Mg can prolong the cycle life of MgNi alloy [START_REF] Jianshe | Electrochemical characteristics of Al-substituted Mg2Ni As negative electrode[END_REF].

Wang et al [START_REF] Wang | An electrochemical investigation of Mg=AlxNi (0≤x≤0.6) hydrogen storage alloys[END_REF] synthesized Mg2-xAlxNi (0 < x < 0. Moreover, MgNi amorphous phase could be prepared by ball milling with a variety of starting materials. Varin et al. [START_REF] Varin | The effect of MgNi2 intermetallic compound on nanostructurization and amorphization of Mg-Ni alloys processed by controlled mechanical milling[END_REF] synthesized amorphous MgNi with the starting material of MgNi2 and Mg2Ni in less than 10 h and indicated that the presence of hard MgNi2 phase helps to reduce crystallite size of Mg2Ni phase and thus facilitates amorphization. Nohara et al. [START_REF] Nohara | Electrochemical characteristics of a homogeneous amorphous alloy prepared by ball-milling Mg2Ni with Ni[END_REF] reported that amorphous MgNi prepared by ballmilling of Mg2Ni and Ni powders exhibits a high discharge capacity. Rongeat and Roué [START_REF] Rongeat | Effect of particle size on the electrode performance of MgNi hydrogen storage alloy[END_REF] investigated the influence of the amorphous MgNi particle size on its electrochemical behavior and indicated that, the initial discharge capacity, the cycling stability and the high-rate discharge ability are improved by increasing the MgNi particle size. It seems to be related to a lower sensitivity to oxidation and to a better pulverization resistance of the electrode as the MgNi particle size increases [START_REF] Anik | Electrochemical hydrogen storage capacities of Mg2Ni and MgNi alloys synthesized by mechanical alloying[END_REF].

Nevertheless, the main shortcoming of MgNi alloys is its rapid degradation upon charge and discharge cycling. The capacity decay of amorphous MgNi reaches 70% after only 20 charge/discharge cycles [33]. It is common understanding that the degradation of amorphous MgNi electrodes is due to the formation of Mg(OH)2 on the material surface. In fact, magnesium based alloys are easily oxidized to Mg(OH)2 by reacting with the KOH electrolyte. The oxidation process can be further accentuated by the formation of fresh surfaces due to the particle pulverization through volume expansion during cycling [START_REF] Ruggeri | Correlation between charge input and cycle life of MgNi electrode for Ni-MH batteries[END_REF]. Therefore, many studies have been conducted in order to improve the cycle life of MgNi alloy, for example by coating the alloy surface [START_REF] Gray | Protective coatings on magnesium and its alloys-a critical review[END_REF],

modifying the electrolyte composition [START_REF] Kim | Effects of F-treatment on degradation of Mg2Ni electrode fabricated by mechanical alloying[END_REF], the elaboration of composite materials [START_REF] Souza | Structural and electrochemical properties of MgNi-based alloys with Ti, Pt and Pd additives[END_REF],

the modification of the bulk composition of the alloy [START_REF] Iwakura | Effects of surface and bulk modifications on electrochemical and physicochemical characteristics of MgNi alloys[END_REF] and powder sieving [START_REF] Anik | Electrochemical hydrogen storage capacities of Mg2Ni and MgNi alloys synthesized by mechanical alloying[END_REF].

Among these, partial substitution of MgNi seems to be the most promising solution, which will be mainly discussed in the following section. The choice of the substituting elements and the substitution ratio as well as the substituted element (Mg and/or Ni) leads to a large number of possible combinations [START_REF] Rongeat | Evaluation of different approaches for improving the cycle life of MgNi-based electrodes for Ni-MH batteries[END_REF].

Several works have shown that Ti as partial substitute for Mg in MgNi alloys is effective in improving the electrode cycle life. Han et al. [START_REF] Han | Effects of Ti on the cycle life of amorphous MgNi-based alloy prepared by ball milling[END_REF] attributed the reason to the formation of Ti oxide. The exposure of the bare surface of the Ti-substituted MgNi alloy to the KOH electrolyte is suppressed during the pulverization of the alloy. Ti substitution is found to be very effective for improving the cycle life while Ti surface coating proves to be hardly effective. Similar results was obtained by Anik et al. [START_REF] Anik | Effect of Al, B, Ti and Zr additive elements on the electrochemical hydrogen storage performance of MgNi alloy[END_REF],

in which, the positive effect of titanium substitution in MgNi alloy is estimated to create pathways for hydrogen penetration. Huang et al. [START_REF] Huang | Influences of Ti substitution on the structure and electrochemical properties of Mg1-xTixNi (x=0, 0.1, 0.2, and 0.3) hydrogen storage alloys[END_REF] indicated that, he substitution of Ti for Mg helps to improve the anti-oxidation/corrosion ability of the MgNi alloy but demolishes the electrochemical kinetics of hydrogenation/dehydrogenation.

Rongeat et al. [START_REF] Rongeat | Synergetic effect between Ti and Al on the cycling stability of MgNi-based metal hydride electrodes[END_REF] investigated the synergetic effect between Ti and Al on the cycling stability of MgNi alloys. It is shown that, the dopants of Ti and Al for Mg improves the cycling stability of the MgNi-based metal hydride electrode significantly.

The formation of TiO2 and Al2O3 appear to be very efficient at preventing the accumulation of Mg(OH)2 onto the particles upon cycling. Further study indicated that the electrode pulverization of amorphous MgNi alloy is significantly lower by the substitution of Ti and Al, which is mainly attributed to the lower porosity of the substituted material, limiting the accumulation of H2 bubbles in the agglomerate pores [START_REF] Etiemble | Effect of Ti and Al on the pulverization resistance of MgNibased metal hydride electrodes evaluated by acoustic emission[END_REF]. Moreover, the synergetic effect of Ti and various other elements on the MgNibased materials have also been extensively investigated, such as Zr, Cr and V [START_REF] Santos | Effect of transition metal additions on the electrochemical properties of a MgNi-based alloy[END_REF]. The synergetic effect of polyelement substitution is assumed to be associated with the formation of a more compact and more protective composite oxide layer on the electrode surface than with Ti alone.

I.3.4 Hydrogenation properties of Ti-Mg-Ni system

As we have discussed above, Mg-Ni and Ti-Ni systems have been investigated by researchers for several decades due to their potential gravimetric and volumetric assets.

However, these two binary alloys are far from being operational in practical applications due to their mediocre cycling stability and relatively low discharge capacity [START_REF] Rongeat | On the cycle life improvement of amorphous MgNi-based alloy for Ni-MH batteries[END_REF][START_REF] Zhang | The phase transformation and electrochemical properties of TiNi alloys with Cu substitution: Experiments and first-principle calculations[END_REF]. Fulfilling all the requirements for optimum performance of hydrogenbased batteries remains thus a real challenge. To overcome this problem, composite hydride materials have been introduced as electrode materials in an attempt to improve the balance of discharge capacity and cycling performances. Alloys pertaining to the Ti-Mg-Ni ternary system are of interest as potential metal hydride since they consist of two hydride forming metals of Ti and Mg. The hydrogen storage properties of these alloys could be upgraded with Ni additions thanks to the high catalytic activity of the latter. There is therefore an interest for the development of such ternary Ti-Mg-Ni systems in view of hydrogen storage applications.

Alloys that belong to the Ti-Mg-Ni system are of considerable interest for hydrogen storage since they contain two hydride forming metals, i.e. Mg and Ti, which combined with the high catalytic activity of Ni, may result in certain compositions exhibiting outstanding hydrogen storage properties. Most of research on the Ti-Mg-Ni system was focused on metastable phases (e.g., thin films [START_REF] Gremaud | Hydrogenography: An Optical Combinatorial Method To Find New Light-Weight Hydrogen-Storage Materials[END_REF] or mechanical alloys).

Combinatorial search for equilibrium phases followed by screening of hydrogen storage properties of prepared alloys in Mg-Ti-Ni system has been attempted. It is found that, the addition of Ti content improved the cycling stability of Mg50-xTixNi50 alloys at the expense of reduction of the discharge capacity [START_REF] Ölmez | Combinatorial search for hydrogen storage alloys: Mg-Ni and Mg-Ni-Ti[END_REF]. Surface coating of alloy particles with Ni was also tested for improving the cycling stability of the ternary alloys and found beneficial for cycling stability and bad for discharge capacity. Zhang et al. By synthesizing series of Ti-Mg-Ni alloys by mechanical alloying, Anik [START_REF] Anik | Effect of titanium additive element on the discharging behavior of MgNi alloy electrode[END_REF] pointed that, the initial discharge capacity was observed to depend on Mg/Ni atomic ratio rather than Ti/Mg atomic ratio in alloys. The alloy charge transfer resistances decreased with the increase of Ti/Mg atomic ratio. The reason is assumed to be partial selective dissolution of the surface Ti/Ti-oxides and thus the limited enrichment of the surface by the electro-catalytic Ni [START_REF] Anik | Effect of titanium additive element on the discharging behavior of MgNi alloy electrode[END_REF].

Besides, in order to achieve equilibrium phases, other manufactory methods have also been employed. Borissova et al. formation [START_REF] Han | Synthesis of composite metal hydride alloy of A2B and AB type by mechanical alloying[END_REF]. The alloy structure variation with the ball milling parameters is also a key factor in order to achieve high electrode cycle life improvement. This can explain the contradictory results from materials having similar nominal composition but dissimilar structure. Thus, it is of interest to explore phase transformation and related hydrogen storage properties of Ti-Mg-Ni ternary alloys system, as well as the synergistic effect of different phases on the enhancement of electrochemical performance.

In this chapter, we present a review of synthesis methods, and describe the experimental techniques and characterization used in this study, including the synthesis method (MA), microstructure characterization (XRD, SEM, TEM), hydrogenation properties measurement (solid gas and electrochemical approaches) and a brief introduction of the DFT (Density Functional Theory).

II.1 Material preparationmechanical alloying (MA)

Mechanical alloying is normally a dry, high-energy ball milling technique and has been employed to produce a variety of commercially useful and scientifically interesting materials [1]. A wide range of intermetallic alloys could be synthesized by MA. One of the advantages of MA is the possibility to synthesize alloys whose constituent elements have very different melting points at room temperature.

MA starts with mixing of the powders with desired proportion of different powders.

The loading of powder mix into the mill can be done under an inert gas atmosphere if necessary, during when the grinding balls and PCA (process control agent) are also put in the vial. The essential factor that plays often a decisive role is the energy input to the milling, which is controlled by several milling conditions, such as B/P (ball to powder ratio), speed, time etc. A milling should be considered to be complete until it reached a steady state when particles share the same composition.

The important events involved in mechanical alloying are repeated welding, fracturing, and rewelding of powder particles. MA is a complex process and hence involves optimization of a number of variables to achieve the desired product phase and/or microstructure. Some of the important parameters that have an effect on the final constitution of the powder are: type of mill, milling container, milling speed, milling time, type, size, and size distribution of the grinding medium, ball-to-powder weight ratio (BPR), extent of filling the vial, milling atmosphere, amount of process control agent, and temperature of milling.

In general, the energy input into the milling system is highly dependent on the parameters settings. A higher value of milling speed, time or B/P ratio gives more energy. Other parameters also play crucial or even decisive roles in a milling setup. For examples, Nouri et al. [3] studied the influence of different types of PCA on milling product. Milling atmosphere is usually either inert gas or reactive gas added purposely.

The former prevents the reaction between the milling powder and air, whereas the latter is expected to react with the milling powder, which is also called "reactive ball milling (RBM)". Extent of filling the vial is suggested to be around 50% to have optimal efficiency. Material of the vial and ball should be picked so that it does least contamination to the milling product, and it was reported to affect the final product of the milling as well [4]. There have been some studies on the influence of temperature of milling, in which it showed that temperature has a significant effect in an alloy system as expected. Therefore, temperature rise induced by long time of milling should not be neglected.

Besides, Heating of the sample may be suppressed by using short periods of milling separated by breaks during which intrinsic heat produced in the grain boundaries can be dissipated and the sample can thermally equilibrate. Therefore, not only the total milling time is important for obtaining the desired compound, but breaks within the period of milling is in some cases crucial, which possibly also reduce agglomeration of the powder on the vial walls and balls [5].

In

II.2 Structural characterization

Structural characterization used in this study includes XRD (X-Ray diffraction), SEM (Scanning electron microscopy) and TEM (Transmission electron microscopy).

Chemical analysis was done by the Energy-dispersive X-ray spectroscopy method.

II.2.1 X-Ray diffraction

X-ray powder diffraction (XRD) is a rapid analytical technique primarily used for phase identification of a crystalline material and can provide information on unit cell dimensions. It is adopted in this work to determine the phase and to verify the purity of the obtained compounds. This technique allows, in the case of ceramics, to determine whether the desired phase was obtained and whether the reaction occurred as expected.

It provides information on the purity, crystallinity and crystallographic parameter values.

Crystals are regular arrays of atoms, and X-rays can be considered waves of electromagnetic radiation. An X-ray striking an electron produces secondary spherical waves emanating from the electron. This phenomenon is known as elastic scattering. A regular array of scatterer produces a regular array of spherical waves. Although these waves cancel one another out in most directions through destructive interference, they add constructively in a few specific directions, determined by Bragg's law:

𝑛𝜆 = 2𝑑 sin 𝜃 EQN II.1
Where d is the distance between atomic layers in a crystal, λ is the wavelength in metres of the incident beam and θ is the angle of incidence. n is an integer representing the number of wavelengths required for constructive interference to occur. At the smallest angle of incidence (θ) for a maxima n = 1, at the next smallest angle n = 2, etc.

Each particular structure has specific directions which is called the high intensity diffraction position as well as the characteristic peaks of the compounds.

The equipment used in this study is branded of Bruker AXS D8 FOCUS. It is equipped with a curved counter, and allows measurements at a fixed incidence of 4°.

The device is equipped with a cobalt anticathode (λ = 0,17889nm).

II.2.3 Transmission electron microscopy

Transmission Electron Microscopy (TEM) uses electrons that have passed through a thin specimen and are focused to create the image with a spatial resolution, which is significantly higher than light microscopes. Provided the prepared sample is thin enough, all electrons are expected to be transmitted through. However, the electrons are scattered by the atoms of the specimen. In any part of the field of view where no specimen is present, and the electrons remain unscattered, it appears bright relative to the specimen. And this results in a bright-field image. Considering the wave nature of the incident electron beams, TEM can also be used to obtain diffraction images of a crystalline material. In this case, it is very similar to the XRD technique, and it also features the Bragg's law: nλ = 2dsinθ. Take Figure II. 5 as an example, the relationship between the radius R of a given diffraction ring and scattering angle θ is given by: R = L tan θ, where L is the distance from where the diffraction pattern, were recorded to the specimen. • thermal electron emission gun type, LaB6.

• accelerating voltage, 80kV, 200kV.

• HTP polepiece, point resolution 0.25 nm, lattice 0.14 nm.

• Energy Dispersive X-Rays Spectroscopy (EDS JEOL JED 2300T 30 mm 2 , resolution 133 eV on Mn K)

• Gatan Erlangshen ES500W camera (1350 x 1040), port 35mm

• Gatan slow-scan camera USC 1000 (2000 x 2000)

Its capability includes: crystallographic and chemical analysis, nanomaterials (nanotubes, nano-ribbons), catalysis nanoparticles, steels, alloys and compounds. In the initial state, V 1 is filled with hydrogen and the pressure is Pi. When V2 is open for hydrogen access, the sample is expected to absorb hydrogen till a stable final pressure is reached at Pf. The amount of hydrogen absorbed is given by the formula [6]:

∆n = 𝑃 𝑖 𝑉 𝑖 𝑍 𝑖,𝑇 𝑅𝑇 - 𝑃 𝑓 (𝑉 1 +𝑉 2 )
𝑍 𝑓,𝑇 𝑅𝑇

EQN II.3

In which, Zi,T and Zf,T are the hydrogen compressibility considering the real gas properties rather than ideal gas, which is realized using Hemmes method [7]. To continue the test in an iteration manner, the final state can be the initial state for the next step. Eventually a PCI (pressure-composition-isotherm) curve can be obtained.

In current study, the hydriding/dehydriding behavior was measured by a Sievert's type apparatus (PCTPro, Setaram), which is shown in Figure II. 9. Each sample with a mass of 1 g was introduced in a stainless steel container. For the first activation, samples were evacuated at 673 K for 1 h, and then hydrogen with a pressure of 3.5 MPa was introduced into the reactor for 2 h. After absorption, the samples were evacuated at 673 The constant current applied to the electrode causes reactants at the electrode surface to react. The potential of the working electrode changes to values characteristic of the electrochemical reaction that is occurring. Galvanostatic experiments provide a means to gain information about the charge and discharge capacities of electrode materials. For example, as the discharge current is constant in time, the multiplication of the applied current and the charging time will yield the exact amount of charge transferred during a particular potential response of the working electrode. This allows for the fact that, for example, the hydrogen concentration in a hydride-forming compound can be precisely tuned, as ideally one electron is transferred per hydrogen atom stored.

In the present study, the electrochemical charge and discharge measurements were

II.4 Density functional theory

Density functional theory (DFT) is a computational quantum mechanical modelling method used in physics, chemistry and materials science to investigate the electronic structure (principally the ground state) of many-body systems, in particular atoms, molecules, and the condensed phases. In many cases the results of DFT calculations for solid-state systems agree quite satisfactorily with experimental data.

Computational costs are relatively low when compared to traditional methods, such as Hartree-Fock theory and its descendants based on the complex many-electron wave function.

II.4.1 Fundamental Principles

DFT methods model all the electrons in a system, and the zero energy corresponds to having all nuclei and electrons at an infinite distance from one another. Quantum mechanics (QM) is the correct mathematical description of the behavior of electrons and thus of chemistry. In theory, QM can predict any property of an individual atom or molecule exactly. In practice, the QM equations have only been solved exactly for one electron systems. These approximations can be very useful, but this requires an amount of sophistication on the part of the researcher to know when each approximation is valid and how accurate the results are likely to be.

The Schrödinger equation, shown in EQN II.4, was devised as an equivalent formulation of QM and it is the basis for nearly all computational chemistry methods:

𝐻 ̂𝛹 = 𝐸𝛹 EQN II.4
where 𝐻 ̂is the Hamiltonian operator, Ψ a wave function, and E the energy. The wave functionΨis a function of the electrons and nuclei positions. As the name implies, this was the description of an electron as a wave. The Hamiltonian operator 𝐻 ̂ is, in general : The eigenvalues of equation cannot be calculated accurately except in the case of hydrogen-like systems. In the case of a multi-electronic system, because of electronic interactions, there is no exact analytical solution of Schrödinger equation. In order to find acceptable solution of the equation, we need to make approximations of Hamiltonian operator. The first approximation is the Born-Oppenheimer approximation [8].

𝐻 ̂= -∑ { ℏ 2 2𝑀 𝛼 ∇ 𝛼 2 } 𝑁 𝛼=1 -∑ { ℏ 2 2𝑚 𝑖 ∇ 𝑖 2 } 𝑛 𝑖=1 -∑ 𝑍 𝛼 𝑒 2

II.4.2 Born-Oppenheimer approximation

All methods of solving the Schrödinger equation based on this approximation, use the Born and Oppenheimer assumption that there is a time scale difference between the electronic and nuclear motions due to a big difference in mass between the nuclei and electrons. Therefore, it is possible to decouple the movement of the nuclei and the electrons, then to write the wave function into two wave functions: nuclei and electrons.

𝛹(𝑟 ⃗, 𝑅 ⃗⃗ ) = 𝛹 𝛼 (𝑅 ⃗⃗ ) × 𝛹 𝑒 (𝑟 ⃗, 𝑅 ⃗⃗ ) EQN II.7

Where 𝛹 𝛼 (𝑅 ⃗⃗ ) is the wave function associated to the nuclei and 𝛹 𝑒 (𝑟 ⃗, 𝑅 ⃗⃗ ) is the wave function associated to the electrons when the nuclei are fixed at the positions of 𝑅 ⃗⃗ . The total energy is then written as the sum of a nuclear and electronic contribution.

𝐸 = 𝐸 𝛼 (𝑅 ⃗⃗ ) + 𝐸 𝑒 (𝑅 ⃗⃗ ) EQN II.8
This approximation is known as the Born-Oppenheimer adiabatic approximation.

The position of the nuclei in the field is supposedly fixed. The operators of nuclei kinetic energies in Hamiltonian operator can be omitted.

(𝑇 ̂𝑒 + 𝑉 ̂𝛼𝛼 + 𝑉 ̂𝛼𝑒 + 𝑉 ̂𝑒𝑒 )𝛹 𝑒 (𝑟 ⃗, 𝑅 ⃗⃗ ) = 𝐸 𝑒 (𝑅 ⃗⃗ )𝛹 𝑒 (𝑟 ⃗, 𝑅 ⃗⃗ ) EQN II.9

Several methods are employed for the solution of this equation. The premier effective method is that of Hartree-Fock based on the assumption of free electrons. This method is quite useful to quantum chemistry to treat the atoms and molecules, but it is less accurate for calculation in solids. The DFT is proved probably to be a more modern and powerful method due to the two theorems by Hohenberg and Kohn [9] established in 1964. These authors have demonstrated that all aspects of the electronic structure of a system in a non-degenerate ground state is also completely determined by using its electron density instead of its wave function.

II.4.3 The Kohn-Sham Formulation

The Kohn-Sham formulation centers on mapping the full interacting system with the real potential, onto a fictitious non-interacting system whereby the electrons move within an effective ''Kohn-Sham'' single-particle potential υKS(r). The Kohn-Sham method is still exact since it yields the same ground state density as the real system, but greatly facilitates the calculation.

The Kohn-Sham potential υKS(r) is given by: The ground state density is obtained in practice by solving the one-electron Schrödinger equations:

[- ℏ 2 2𝑚 𝑖 ∇ 𝑖 2 + υ 𝑒𝑥𝑡 (r) + υ 𝐻 (r) + υ 𝑥𝑐 (r)] 𝜓 𝑖 (𝑟 ⃗) = 𝐸 𝑖 𝜓 𝑖 (𝑟 ⃗), 𝑖 = 1, … 𝑛 EQN II.13
Determining the ground state of the system, then returning to solve selfconsistently, the set of equations (EQN II.20) is known as Kohn-Sham equation. The sum of three υ 𝑒𝑥𝑡 (r) + υ 𝐻 (r) + υ 𝑥𝑐 (r) is an effective potential υKS(r) which can be characterized as it depends only on the parameter r. This method is formally correct, but for the practical calculation, the exchange and correlation energy, which is a density functional, requires the introduction of some approximations.

II.4.4 The exchange-correlation functions

The fact that the DFT gives no information on the form of the exchange-correlation functional 𝐸 𝑥𝑐 [𝜌(𝑟)], hence approximations have to be introduced for its determination to be applicable to different systems. The actual calculation of the energy and the potential for exchange-correlation is based on a number of approximations. Two types of approximation are discussed in this part.

 Local-density approximation (LDA)

The local-density approximation (LDA) relies on the assumption that the electron density varies slowly in space and therefore the exchange-correlation terms depend only on the local value of 𝜌(𝑟); that is to say, it deals with a non-homogeneous system as being locally homogeneous. In this regard, LDA is generally synonymous with functional based on the homogeneous electron gas (HEG) approximation, which are then applied to realistic systems (molecules and solids).

In general, for a spin-unpolarized system, a local-density approximation for the exchange-correlation energy is written as:

𝐸 𝑥𝑐 𝐿𝐷𝐴 [𝜌(𝑟 ⃗)] = ∫ 𝜌(𝑟 ⃗)𝜀 𝑥𝑐 [𝜌(𝑟 ⃗)] 𝑑𝑟 ⃗ EQN II.14
Where 𝜀 𝑥𝑐 [𝜌(𝑟 ⃗)] is the exchange-correlation energy per particle of a homogeneous electron gas of charge density 𝜌(𝑟 ⃗).

The exchange-correlation energy is decomposed into exchange and correlation terms linearly:

𝐸 𝑥𝑐 𝐿𝐷𝐴 [𝜌(𝑟 ⃗)] = 𝐸 𝑥 𝐿𝐷𝐴 [𝜌(𝑟 ⃗)] + 𝐸 𝑐 𝐿𝐷𝐴 [𝜌(𝑟 ⃗)] EQN II.15
The exchange-energy density of a HEG is known analytically. The LDA for exchange employs this expression under the approximation that the exchange-energy in a system where the density is not homogeneous, is obtained by applying the HEG results pointwise, yielding the expression:

𝐸 𝑥 𝐿𝐷𝐴 [𝜌(𝑟 ⃗)] = - 3 4 ( 3 𝜋 ) 1 3 ⁄
∫ 𝜌(𝑟 ⃗) [11]; Cole-Perdew (CP) [12]; Perdew-Wang (PW92) [13].

The LDA can only describe the ground state of electronic systems but not the excited states. The wide band gap of semiconductors and insulators are underestimated in this approximation. For certain strong correlation effects systems, LDA approximation does not properly describe the system properties. In particular, the transition metal compounds of Mott-Hubbard insulating type are predicted metal.

The extension of density functional to spin-polarized systems is straightforward for exchange, where the exact spin-scaling is known, but for correlation further approximations must be employed. A spin polarized system in DFT employs two spin- The Generalized gradient approximations bring an improvement compared to the LDA. In the local approximation, the exchange and correlation potential depends only on the density, whereas in the GGA approximation, the potential is expressed as a function of local electron density and its gradient:

𝐸 𝑥𝑐 𝐺𝐺𝐴 [𝜌(𝑟 ⃗)] = ∫ 𝜌(𝑟 ⃗)𝑓[𝜌(𝑟 ⃗), ∇𝜌(𝑟 ⃗)] 𝑑𝑟 ⃗ EQN II.18
Where, 𝑓[𝜌(𝑟 ⃗), ∇𝜌(𝑟 ⃗)] dissolved the .exchange energy as a function of both electron density and its gradient.

There are several versions of GGA most frequently used which are introduced by Perdew and Wang [13].

In many cases, the GGA approximation provides better results than LDA for total energy, cohesive energies, the volumes to balance and incompressibility modules.

However, the systems with strong correlations are poorly described.

 Hubbard-corrected energy functionals (LDA+U, GGA+U)

The basic idea behind DFT+U is to treat the strong on-site Coulomb interaction of localized electrons, which is not correctly described by LDA or GGA, with an additional Hubbard-like term. The on-site Coulomb interactions are particularly strong for localized d and f electrons, but can also be important for p localized orbitals. The strength of the on-site interactions is usually described by parameters U (on site Coulomb) and J (on site exchange). Hubbard correction is added to the energy functional in order to overcome the imprecision of the strong correlations systems calculation:

𝐸 𝐷𝐹𝑇+𝑈 = 𝐸 𝐷𝐹𝑇 + ∑ 𝑈-𝐽 2 (𝜌 𝑎 - 𝑎 𝜌 𝑎 2 ) EQN II.19
where 𝜌 𝑎 is the atomic orbital occupation matrix.

II.4.5 Vienna Ab initio Simulation Package (VASP)

The Vienna Ab initio Simulation Package (VASP) [14] is a computer program for atomic scale materials modelling, e.g. electronic structure calculations and quantummechanical molecular dynamics, from first principles. It computes an approximate solution to the many-body Schrödinger equation, within density functional theory (DFT), by solving the Kohn-Sham equations.

In current study, VASP was adopted to study the structural and electronic properties, including the Fermi energy, the density of states, band structure etc., of the Sr, Ge doped apatite-type lanthanum silicates LSO. This code is known for good parallelization which authorizes the calculation on more than one hundred atoms systems. It provides accuracy and efficiency for ab initio computer simulation and for these reasons it has become a widely used code.

The wave functions used in the VASP are developed on a plane wave basis adapted to the periodic calculation and have the advantage of being easily implemented with calculations using USPP of Vanderbilt and PAW. Blöchl and Kresse et al. [15] showed that the PAW method is especially good for the transition metal oxides, lanthanides, actinides and magnetic systems. So we used the PAW method for all of our work.

The wave functions are expended on basis of reciprocal space so that they are defined for an infinite number of k points in the Brillouin zone, and each point contributes to the electronic potential. In order to numerically evaluate the integrals on the occupied states of each point, the calculation is performed with a limit grid of k points. To automatically generate this grid points k in VASP, the method of Monkhorst and Pack [16] is used. The integration of occupied states may be achieved by the method of the tetrahedra (Jepsen et al. [17]), with the corrections of Blöchl et al. [18]. The VASP code uses conjugate gradient type algorithms to relax the structures. In practice, to relax an atomic structure, we minimize Hellmann-Feynman forces until the residual force on each atom fall below 0.01 eV/Å.

II.5 References

In this chapter, a series of Ti-Mg-Ni alloys have been synthesized by mechanical

III.1 Experimental methodology

The samples of Ti-Mg-Ni system were prepared by mechanical alloying (MA).

The phase structures of studied samples were characterized by X-ray diffraction (XRD)

with Co Kα radiation (λ = 1.789 Å). The powder morphologies were determined using scanning electron microscope (SEM) and transmission electron microscopy (TEM)

with Energy-dispersive X-ray spectroscopy support (EDS). The electrochemical charge and discharge measurements were performed under galvanostatic conditions at room temperature using Multi-channel Battery Interface ATLAS 0461. The Pressurecomposition-isotherm (PCI) tests were done with a Sievert's type apparatus (PCTPro, Setaram).

III.1.1 Preparation of material

III.1.2 Optimization of milling conditions

Mechanical alloying is a powerful powder processing method, which has been used to synthesize various equilibrium and non-equilibrium alloy powders with extremely fine microstructures. It is also a complex process which involves a number of variables. As discussed in chapter 2, almost every parameter has an effect on the final constitution of the products, and these process parameters are not completely independent. For example, the metastable Ti50Mg50 alloy with BCC structure were synthesized by Rousselot et al. [1] and Asano et al. [2] using high-energy ball milling with different milling speed and milling time. While the zirconia milling balls and pots were used instead of steel ones for ball milling the same starting materials in the same milling conditions, a Mg50Ti50 alloy with FCC structure was obtained [3]. Thus, the milling parameters are the key factors to achieve a desired phase or microstructure.

At the beginning of this study, the optimization of milling parameters was carried out for the equimolar mixture of titanium and magnesium powders. The main variables are milling ball size, process control agent (PCA), milling speed and milling time in consideration of the milling efficiency and contamination. All of the milling processes were conducted under protective atmosphere (Ar). an influence on the milling efficiency which further results in a variation of milling products. When using 4 balls of 14 mm as milling medium, the peaks of elemental Ti are still present after 60 h of milling. On the contrary, when using 2 balls of 21 mm, a fully amorphous phase can be obtained. The total weight of 2 balls of 21 mm are heavier than that of 4 balls of 14 mm. Generally speaking, a large size (and high density) of the grinding medium is useful since the larger weight of the balls will transfer more impact energy to the powder particles. It has also been reported elsewhere that the final constitution of the powder is dependent upon the size of the grinding medium used [4,5]. Thus, in our study, using 2 balls of 21 mm diameter is preferable for the formation of amorphous phase since they possess higher milling energy and efficiency than the other type of balls. synthesized TiMg with BCC phase by mechanical alloying with the speed of 200 RPM.

It is easy to realize that the faster the mill rotates the higher would be the energy input into the powder. Hence, a higher milling speed may be a disadvantage because the increased temperature accelerates the transformation process and results in the decomposition of supersaturated solid solutions or other metastable phases formed during milling [6]. In our cases, the initial elemental peaks can still be identified after 100 h of milling, which means the milling process are far from completion. When the milling speed was set at 400 RPM, an amorphous phase can be formed after 60 h of milling. Higher rotating speed can provide more milling energy which can result in a reinforcement of milling efficiency. Furthermore, no contamination or phase decomposition were found in the mill products. Thus, in this study, the milling speed is set as 400 RPM. These results are is inconsistent with the previous conclusion reported by Rousselot et al. [8], in which, the formation of BCC MgTi was inferred at the very beginning of the milling process in both cases of MgTi and MgTi0.5Ni0.5 alloys. Whereas, other reports relatively sharp while that of Ti was significantly broadened. The same phenomenon had also been reported, which is related to the investigation of milling production of pure Ti and pure Mg [2]. According to the index results of phase, the broad feature of Ti centered peak is attributed to the presence of the overlapped diffraction peaks of Ti2Ni at 47.8°. It is notable that both of Ti and Mg were abundant in this composition, but only the Ti-Ni phase was detected after milling. The formation of Ti2Ni instead of Mg2Ni may be due to two following reasons. Firstly, the atomic radii of Mg and Ti are 0.160 nm and 0.147 nm, respectively, which means Ti possesses a relatively higher diffusivity than Mg. Secondly, the enthalpy of formation of Mg2Ni and Ti2Ni are -51.9 kJ/mol [11] and -71.4 kJ/mol [12] respectively. All these evidences indicate that, the combination of Ti and Ni at the atomic scale is easier than that of Mg and Ni. Besides, While the observed phases, so far, were mainly binary phase with equal atom ratio. According to the stoichiometric proportion, these may lead to two possible assumptions. One is a high degree of excess dispersion of magnesium or/and titanium, while the other one is the existence of another Mg-rich or/and Ti-rich binary phase. As no obvious elemental peak was observed in the XRD pattern, the formation of nanocrystals Mg2Ni and Ti2Ni cannot be excluded. In fact, the recrystallization of MgNi amorphous phase during prolongation milling was report elsewhere [15,16]. The crystallization of MgNi alloy into a mixture of nanocrystals Mg2Ni and MgNi2 took place after a certain time of milling which may differ from the various mill mode. In addition, an observation reported by Guzmá n et al [17] suggested that Mg2Ni and All XRD patterns are dominated by the diffraction peaks of Ni which is not surprising. Firstly, there are no full rings. Secondly, the XRD pattern of TiMgNi milled 60h does

 not give much information about the crystallographic phases which could be present; These absences induce numerous possibilities of existing crystallographic phases (Mg, Ti, Ni, Mg2Ni, MgNi2, TiNi…) which make the indexation of the appearing rings impossible. Nevertheless, correlations with XRD results could be done. that, the particles with the same elemental distribution of zone 1 which consists of the same atom ratio of Ti and Ni and relatively few amount of Mg, have been widely observed in our TEM test. This particle is assumed to be constituted by a mixture of amorphous Ti-Ni and/or Mg-Ni rather that a homogeneous amorphous Ti-Mg-Ni phase.

Because the stoichiometry of this particle is inconsistent with the only stable ternary Ti-Mg-Ni phase (Mg3TiNi2). In the case of zone 2 and 3, the existence of Mg and amorphous Mg-Ni phase are indicated obviously. Moreover, the content of Mg in selected area is less than that in nominal composition. It is not surprising because of the lower hardness of magnesium compared to either titanium or nickel which rise a tendency to adhere to the walls of milling balls and vial. Although the mill procedure was interrupted several time to scrap this adherence. Despite a homogeneous area, regional elemental enrichments were also observed. All these phenomena are in good agreement with the XRD results.

III.2.3 Hydrogenation properties of TiMgNi x alloys

It has been shown that the as-milled TiMgNix alloys, consisting of nanocrystallites and amorphous phase, possess good structural stabilities in alkaline solution at room temperature. However, the thermal stability of Ti-Mg-Ni amorphous as well as its phase According to the intensities of peaks, the ratio of phases are varied with the variation of composition. The formation of TiNi3 is promoted with the increase of Ni content.

The consumption of Ni by TiNi3 suppresses the formation of hydrogen absorbing phase, such as Mg2Ni, which results in the degradation of its hydrogen storage capacity.

Whereas, the formation of TiNi3 phase is not observed in as-milled samples, which rise a problem, that whether the TiNi3 phase is formed during the activation process (heat treatment) or the hydrogenation process. Thus, comparative tests were carried out for TiMgNi alloy.

                    (c) (d) (b) Intensity (a.u.) 2 (degree)
as-milled after annealing hydrogenated dehydrogenated (a)

 Mg MgH2 TiH2 Mg2NiH4 TiNi3 Mg2Ni TiNi     
absorption, TiH2 and TiNi3 are still present, whereas, MgH2 and Mg2NiH4 were formed instead of Mg and Mg2Ni. In fact, the structural variation upon hydrogen absorption is not surprising. The existences of TiH2 and TiNi3 are due to their stability in hydrogen exposing environment. The hydrogen desorbed temperature of TiH2 is 450°C -600 °C [21] which is much higher than the condition in our experiment. TiNi3 alloy is an inactive phase which cannot store hydrogen. The working hydriding phases are therefore Mg and Mg2Ni, which were correspondingly converted to MgH2 and Mg2NiH4 after hydrogen absorption. These results are also coincide with the characteristics of PCT curves which show two steps of pressure plateau during absorption.

The nanocrystallites and amorphous phases prove significantly well crystallized after the PCI the tests. A previous study suggested that the equilibrium phase of the Mg33.3Ti33.3Ni33.3 composition at 500 °C consists of TiNi, Mg and MgO [22] which is in agreement with our results of the activated sample. However, the characteristics of PCI curve of Ti-Ni alloy [23][24][25] are in contradiction with that of our studied TiMgNi alloy which possesses distinct absorption plateau. Thus, it can be concluded that the asmilled TiMgNi alloy is constructed by the Mg, Amorphous TiNi and Mg-Ni phases.

The formation of equilibrium phase of Mg2Ni and TiNi3 are catalyzed by hydrogen under high temperature. Whereas, the mechanism of deformation of TiNi phase during hydrogenation cycling remains unclear. Further studies are in progress to clarify the effect of hydrogen and temperature on the phase transformation during the hydrogenation process. With increasing the atom ratio of Ni/Ti, the broad peak shifts towards the high angles as well as the intensity of free Mg is decreased. It is notable that, the obtained phase structures of MgTi1-xNix (x = 0.4 and 0.6) are inconsistent with the reported results with a similar composition corresponding to the ball milling of MgTi0.5Ni0.5 alloy using a vibratory type mill machine (Spex 8000 M) [8]. No obvious free Mg was found in their 20 h milled sample, instead, the broad peak was obtained which is believed to be associated with the presence of MgNi and TiNi. The possible reason for

(a.u.) 2 (degree) (a)    Mg2Ni  TiNi  Mg TiNi3   (d) (c) (b) Intensity
this difference of as-milled products are due to the milling type and the morphology of the Mg powder. The deformation of raw materials during ball milling is due to the compression of them by high energy collisions of moving balls and the rotating milling pot [5]. The vibratory type mill can provide higher milling energy than planetary highenergy ball mill. To reinforce the milling energy and efficiency, bigger size of milling balls are employed. Whereas, the Mg powder used in this study, as shown in capabilities of studied samples were ameliorated, among which, sample of MgTi0.4Ni0.6 exhibits the best discharge capacity.

Meanwhile, the characteristics of the discharge capacities curve of studied samples can be sorted in two types which are named as "activation needed" and "activated" samples. In the case of MgTi0.6Ni0.4, whose atom ratio of Ni/Ti less than 1:1, its discharge capacity increases linearly with cycling until reaching the maximum capacity.

On the contrary, when the atom ratio of Ni/Ti greater than 1:1, electrode samples perform the their best discharge capacities in the first cycle but drop immediately in subsequent cycles. Taking XRD results into consideration, it is found that all the "activated" samples are constructed by crystalline Mg2Ni, which is not obtained in "activation needed" samples. Thus, it can be inferred that the improved activation properties of MgTi0.4Ni0.6 and MgTi0.8Ni0.2 are attributed to the formation of crystalline Mg2Ni. discharge capacity is virtually nil. The maximum discharge capacity of TiMgxNi1-x increase linearly with increase of the atom ratio of Ni/Mg. The improvement of discharge capacity is attributed to the formation of Ti-Ni phase in such Ti-rich ternary system. As discussed in last section, as-milled TiMg0.2Ni0.8 consist of amorphous Ti2Ni and TiNi phases [29]. Trace amount of Mg is assumed to be dissolved into Ti-Ni alloys.

All the alloys exhibit good cycling stability after activation process, which is in agreement with the characteristic of Ti-Ni alloys. boundaries and defects. The nanostructured phase can avoid the long-range diffusion of hydrogen atoms through the already formed hydride phase [31]. Quantities of microstrains introduced by the defects and grain boundaries can accelerate the hydrogen absorbing/desorbing process [32]. In the meantime, the co-existence of nanocrystalline and amorphous composite of Mg2Ni may further ameliorate the electrochemical properties. Nanocrystalline possesses a great number of interfaces, and the amorphous phase has larger capacity of hydrogen than nanograins. In addition, the amorphous region around nanocrytallites provides an easier access of hydrogen into the nanograins [33]. Therefore, the structure of as-milled sample which co-exised nanocrystalline and amorphous phase of Mg2Ni is very beneficial for its discharge capability. On the other hand, the grain refining weakens the anti-corrosion capability of the MA alloy due to the inevitable inter-nanocrystalline corrosion, which consequently accelerates the corrosion of MA alloys in alkaline solution and result in the formation of Mg(OH)2.

The layers of Mg(OH)2 are not transparent to hydrogen so that these layers prevent hydrogen from further penetrating into the alloys [18]. As a result, the complete amorphization can ameliorate the discharge capacity of Mg-based alloys but may deteriorate their cycling properties as well.

Even though the discharge capacity of MgTi0.4Ni0.6 alloy has been dramatically improved by separate milling method, but it is still lower than that the results with a similar composition (TiMg2Ni) reported by Rousselot et al. [8]. Two reasons are proposed for the lower discharge capacity of our samples. Firstly, the oxidation of Mg due to its high affiliation with oxygen. Even though all the operational process were carried out under the protection of Ar, the separate milling process may aggravate the risk of oxidation. The formation of MgO will deplete the reacting powder of Mg required for synthesizing Mg-based hydrogen storage alloys. Secondly, the biggest difference from our prepared sample is that, 10 wt.% of Pd powder was added as additive in Rousselot's case. Pd as an effective electrochemical catalyst can ameliorate the discharge capability of electrode significantly. Especially for TiMg alloy, Pd is essential which can greatly facilitates the charge-transfer reaction and destabilizes the Mg-Ti hydride. However, it is also a barrier for the practical utilization of Ti-Mg based alloys due to the extremely high price of Pd. In our case, Pd-free MgTi0.6Ni0.4 also exhibits moderate discharge capacity. In the meantime, the synthesis method used in ref. [8] is a vibratory type mill machine which can not be employed in large-scale. Thus, the samples prepared by separate milling process in present study with a moderate discharge capacity are still promising which could be further improved.

III.7 Conclusion

Series of Mg-Ti-Ni alloys are prepared by mechanical alloying using a planetary high-energy ball mill. In the case of MgTiNix (with x = 0.2, 0.4, 0.6, 0.

IV.1 Experimental methodology

The samples of TiNiCu system were prepared by mechanical alloying (MA). The phase structures of the samples obtained were characterized by X-ray diffraction (XRD)

with Cu Kα radiation (λ = 1.5418 Å) at room temperature. The powder morphologies were determined using a scanning electron microscope (SEM) with Energy-dispersive X-ray spectroscopy support (EDS). The electrochemical charge and discharge measurements were performed under galvanostatic conditions at room temperature using Multi-channel Battery Interface ATLAS 0461. The calculation is based on the density functional theory (DFT) with the code VASP.

IV.1.1 Preparation of samples

The starting materials were pure elemental powders of Ti (purity 99.5%, particle size ≤ 150 µm, GoodFellow), Ni (purity 99.5%, particle size ≤ 250 µm, GoodFellow)

and Cu (purity 99.5%, particle size ≤ 100 µm, Gerac) according to the following stoichiometries: TiNi1-xCux (x = 0, 0.1, 0.2, 0.3). The nominal compositions of mixtures were poured into cylindrical stainless steel vials (capacity 50 ml) with two stainless steel balls in an argon protected glove box. The mechanical alloying process was performed using Retsch PM 400 for 20 h at room temperature with the speed of 400 RPM. The weight ratio of balls (diameter 20 mm) to powder (BPR) was 10:1. In order to dissipate the heat and reduce the cold welding of powder particles during high-energy milling, the milling was interrupted for 30 minutes after every 1 hour of operation. After every 10 h of milling, the ball milling was stopped for scraping the powder adhered to meV/atom for TiNi unit cell. The convergence criteria was set at 5.0 × 10 -6 eV/atom for energy change, 0.01 eV/Å for maximum force, 0.02 GPa for maximum stress and 5.0 × 10 -4 Å for maximum displacement.

In this study, the enthalpy of formation of a compound was defined as the difference between its total energy and the energies of its constituent elements in their stable states (referred to the elementary substances). The zero-point energy (ZPE) contributions are significant in reactions where hydrogen molecules are absorbed or desorbed [5]. The reactions investigated in this work do not include hydrogen. Thus, the ZPE contributions are not considered in this study. Therefore, for a lattice containing

x Ti atoms, y Ni atoms and z Cu atoms, the enthalpy of formation is given by: and EDS studies carried out on small amount of material recovered from walls confirmed the presence of a mixture of elements with a high content of Ni (not shown here). Besides the fact that, the abundance of Ti further results in the generation of secondary phases will be demonstrated by the following XRD analysis. phase. It is in agreement with our observation from the EDS results, that all of the annealed samples were constituted with a higher atom ratio of Ti than that of stoichiometry composition. The A2B-type phases yielded from such Ti-rich ternary alloys is corresponding to the ternary Ti-Ni-Cu phase diagram [9]. Similar results were reported in previously cited Ref. [10], where, the XRD measurements of homogeneous TiNiCu alloys were performed at 473K.

IV.3 Computational results

In order to better understand the influence of Cu substitution for Ni on the crystal structure of TiNiCu alloys, first principle calculations were employed for investigation.

The original lattice parameter and coordinates of TiNi B2 phase were based on the experimentally confirmed structure. The replacements of Cu atoms were carried out gradually and randomly in Ni site corresponding to the experimental composition.

Geometry optimizations were performed with full optimization before energy calculation. The computational results are listed in Table IV.2 and compared with the experimental data. The average differences between calculated lattice parameters and experimental values are below ± 0.1%, which indicates that the present calculation models are reasonable and in good agreement with the experimental results. What is important, the calculated lattice parameters also increase linearly with the increasing Cu substitution which concur with the tendency observed in experimental part. It further confirmed that the substituted Cu atoms were dissolved into the TiNi lattice structure successfully which resulted in the little peak deviation of B2 phase in XRD patterns.

Therefore, the subsequent calculations of formation enthalpy were carried out based on this configuration.

The calculated formation enthalpy is also shown in Table IV.2 which represents for evaluating the structural stability of the composition. It can be seen that, with the substitution of Cu, the enthalpy of formation decreased gradually. Among them, TiNi0.7Cu0.3 possesses the most negative enthalpy of formation which is equal to -29.3518 eV. Generally, a lower enthalpy of formation results in a higher stability of product. The Cu substituted pseudo-binary Ti(Ni,Cu) phase, that possesses higher stability than the binary TiNi phase, is preferably formed in an equilibrium state.

Additionally, the electronic structure of TiNi as well as its modification by Cu substitution also focused our attention. The calculated densities of electronic states are presented in Figure IV.4. Based on the previous investigations [11], the significant broadening of valence band for the nanocrystalline TiNi based alloys could be explained by a strong decrystallization of the nanocrystals. In comparison with the DOS of TiNi, the new structure referring to the Cu atoms in substituted compound enlarges the valence band towards the lower energy. Hence this band becomes much broader and thus stabilizes the intermetallic compound. The calculation for TiNi0.7Cu0.3 alloy gives a width of the valence band wider by about 0.5 eV than that for TiNi alloy. This phenomenon also corresponds to the XRD analysis, with the increase of Cu contents, the intensity of (101) reflections of TiNi B2 phase decreased significantly. Furthermore, the DOS at the Fermi level decreases with Cu substitution and also favors the compound stabilization [12]. The behavior explains the decrease of ∆H 𝑓 and the increasing stability of substituted alloys with increasing Cu content. The equilibrium sites of hydrogen on the surface were determined by minimizing the total energy with respect to the distance from hydrogen to the surface layer of the relaxed surface which can be expressed by the following equation (EQN IV.2)

Ead = EH/substr -Esubstr -EH EQN IV.2
Where Ead is the adsorption energy, EH/substr and Esubstr mean the total energy of TiNi1-xCuxH system and total energy of clean TiNi1-xCux super cell, respectively. EH is the energy of free hydrogen atom which is determined as -12.422 eV by calculation. Maximum discharge capacity obtained for unmodified TiNi alloy equals to 154 mAh/g. It is in agreement with our calculation results that, the hydrogen absorption capabilities of TiNi1-xCux alloys decrease linearly with Cu substitution. Unfortunately, during cycling, material lost that capacity to reach only 82% of capacity retaining rate after 20 cycles. However, it is still more optimistic than that obtained by Drenchev [14] and Jankowska et al. [16] with the similar preparing procedure. This may be ascribed to the higher heating treatment temperature we employed in this study, by which, better crystalline samples were obtained. After substituting Cu for Ni, the cycling stability of TiNi1-xCux alloys was highly improved, but the discharge capacity decreased as an expense.

To facilitate the analysis of electrochemical properties, the cycle stabilities of different samples were evaluated by capacity retaining rate (Rh) after 20 cycles, and shown in Table IV changed their electrochemical properties. The discharge capacities of TiNi1-xCux alloys decreased gradually with the increase of Cu substitution, on the contrary, the retaining rate increased. All of the substituted samples exhibit at least 97% retaining rate after 20 cycles which is very impressive. It can be explained as following: on the one hand, as we have demonstrated, the lattice parameters of TiNi B2 phase increased gradually with increasing of Cu substitution. The cell volume expansion makes the electrode more exposed to the electrolyte and therefore the corrosion becomes more accentuated. On the other hand, as we found in XRD patterns, with increasing Cu substitution the reduction of diffraction peaks intensities may further aggravate the degradation of their capacities. Similar hydrogenation properties of Cu substituted TiNi alloys with B2 phase manufactured by induction melting had also been reported in Ref. [1,10], where were attributed to the electronic effects. An experimental rule proposed by Cuevas et.al [12,17] revealed that, electronic effects rather than geometric factors dictate the thermodynamic hydrogenation properties of TiNi-based alloys. As a result, even though the electrode cycle stability of Cu substituted TiNi alloys has been significantly ameliorated by mechanical alloying, the improvement of their discharge capacity are not satisfying.

IV.5 Conclusion

In this chapter, TiNi1-xCux (x = 0, 0.1, 0.2, 0.3) alloys were prepared by mechanical alloying using a planetary high-energy ball mill and subsequent heat treatment at 750 °C for 0. All of the substituted samples exhibit at least 97% retaining rate after 20 cycles.

In the meantime, there are several issues that are particularly interesting in the current study, which can be further investigated:

• As discussed in the thesis, the preferential reactivity of Ni to Ti is higher than that to Mg in Ti-Mg-Ni system, which lead to the remaining of Mg after mechanical alloying. Thus, improving the milling efficiency of Mg in the Ti-Mg-Ni system is a key technical issue. Other type of mill or combined synthesis method (like functional characteristics which will contribute to the enhancement of the experimental efficiency.

  They are: Soichi FUKUHARA, Xianda LI, Djafar CHABANE, Xinkun SUO, Taikai LIU, Xiaohua FENG, Duo YI, Fouad ELTOUMI, Bilel HOSNI, Yu WU, Yi FANG et al. and peers from different sections: Marc Cescutti et al. Especially our "scientific interest group": Shuo YIN, Chaoyue CHEN, Yangzhou MA and Yingchun XIE, thank you for giving me some academic discussions, yielding us steady progress.

Figure I. 1 10 Figure I. 2 13 Figure I. 3 14 Figure I. 4 19 Figure I. 6 25 Figure I. 8 28 Figure I. 9 38 Figure II. 1 56 Figure II. 3 59 Figure II. 5 63 Figure II. 8 67 Figure III. 1

 1102133144196258289381563595638671 Figure I.1 Volumetric density of compressed hydrogen gas as a function of gas pressure, including the ideal gas and liquid hydrogen. [5]. ................................................... 10

Figure III. 9

 9 Figure III.9 XRD patterns of TiMgNi mixture after different milling time: (a) 20 h, (b) h. .........................................................................................................................................

Figure III. 10

 10 Figure III.10 SEM micrographs of (a) TiMgNi0.2, (b) TiMgNi0.4, (c) TiMgNi0.6, (d) TiMgNi0.8, (e) TiMgNi milled 40h ....................................................................................................

Figure III. 11

 11 Figure III.11 Discharge capacities as a function of charge/discharge cycles of TiMgNix alloys. .........................................................................................................................................

Figure III. 12

 12 Figure III.12 XRD patterns of electrode, after 21 cycles of electrochemical studies (a) TiMgNi0.2, (b) TiMgNi0.4, (c) TiMgNi0.6, (d) TiMgNi0.8, (e) TiMgNi. ............................

Figure III. 13

 13 Figure III.13 Transmission electron microscopy image (a) for the as-milled TiMgNi alloy and its corresponding selected area electron diffraction (SAED) (b) and EDS map (c) ...... 100

Figure IV. 5

 5 Figure IV.5 The structural illustration of TiNi1-xCuxH system. .............................................

Figure IV. 6

 6 Figure IV.6 The discharge capacities as a function of charge/discharge cycles of annealed Ti-Ni and Ti-Ni-Cu alloys. .................................................................................................

  Microscopy, Transmission Electron Microscopy, PCI test and electrochemical properties under galvanostatic conditions. At last, the principle of Density functional theory (DFT) calculation is explained with some words on the application CASTEP (Cambridge Serial Total Energy Package) that we used subsequently to study the Cu substituted Ti-Ni alloys. In Chapter III, a series of Ti-Mg-Ni alloys have been synthesized by mechanical alloying using a planetary high-energy ball mill. The effect of Ni addition in TiMg, partial substitution of Ni for Ti in TiMg and partial substitution of Ni for Mg in TiMg have been studied through the compared analysis of TiMgNix, MgTi1-xNix and TiMg1-xNix. The effect of the variation of composition and milling time on the microstructure and hydrogenation properties of Ti-Mg-Ni alloys are investigated. Additionally, an approach used to improve the electrochemical properties of Ti-Mg-Ni alloys is developed. In Chapter IV, Cu substituted TiNi alloys have been investigated as hydrogen storage material for Ni-MH batteries by experiments and first principles calculations. The phase transformations and electrochemical properties of TiNiCu alloys are characterized. The influence of Cu substitution for Ni on their electronic structure and hydrogen adsorption energies are investigated by first principles calculations.

Figure I. 1

 1 Figure I.1 Volumetric density of compressed hydrogen gas as a function of gas pressure, including the ideal gas and liquid hydrogen. The ratio of the wall thickness to the outer diameter of the pressure cylinder is shown on the right side for steel with a tensile strength of 460 MPA. A schematic drawing of the pressure cylinder is shown as an inset [5].

  Figure I. 2. gives a visual comparison of volumetric scale for storing 4 kg of hydrogen in different ways. It can be found that, the volumetric densities of chemical hydrides (i.e., Mg2NiH4 and LaNi5H6) are much higher than that of compressed hydrogen and liquid hydrogen. Chemical hydrides, therefore, are mainly considered as the most potential technology for storing hydrogen.

Figure I. 2

 2 Figure I.2 Volume of 4 kg of hydrogen compacted in different ways, with size relative to the size of a car [6].

Figure I. 3

 3 Figure I.3 Schematic model of a metal structure with H atoms in the interstices between the metal atoms, and H2 molecules at the surface [6].

Figure I. 4

 4 Figure I.4 Schematic PCT diagram at different temperature (T1<T2<T3<T4<Tc) (left side), and the Van't Hoff plot (right side) [13].

  metal/electrolyte interface (Figure I. 5(a)).

Figure I. 5

 5 Figure I.5 Schematics showing the electrochemical reactions between water and metal hydride during charge (a) and discharge (b) [18].

  surface and recombine with OH¯ in the alkaline electrolyte to form H2O, and charge neutrality pushes the electrons out of the MH through the current collector, performing electrical work in the attached circuitry (Figure I. 5(b)).I.3 Ti-Mg-Ni system I.3.1 Hydrogenation properties of Ti-Mg system

Figure I. 6

 6 Figure I.6 Phase diagram of Ti-Mg system. [19]

  Figure I. 7). The HCP phase was formed by solution of Ti into Mg and the BCC phase by solution of Mg into Ti. While, the FCC phase was stabilized by introduction of stacking faults in Mg and Ti which have a HCP structure.

Figure I. 7

 7 Figure I.7 XRD patterns of MgxTi100-x (x=35,50,65 and 80) alloys milled for 200h using stainless steel milling balls and pots (left) and using zirconia milling balls and pots (right) (Cu Kα radiation) [25].

Figure I. 8 .

 8 It can be found that, the stable phase structure exist in Ti-Ni phase diagram are TiNi, Ti2Ni and TiNi3. For hydrogen storage, it is known that both TiNi and Ti2Ni in the Ti-Ni system absorb hydrogen at room temperature. The electrochemical properties such as the discharge capacity, and the discharge kinetics of Ti-Ni alloys were studied based on these two compositions by synthesizing the alloys with raw materials of Ti and Ni with high purity.

Figure I. 8

 8 Figure I.8 Phase diagram of Ni-Ti binary alloy [37].

[ 45 ]

 45 revealed the differences of hydrogen storage performance between austenitic and martensitic phase by substituting Cu in TiNi alloys. Pseudo-binary TiNi1-xCux (x ≤ 0.5) compounds have been synthesized by induction melting. The XRD results in Figure I.9show, that the products consist of B2 structure above 350 K and either of B19' (x < 0.1) or B19 (0.2 ≤ x ≤ 0.5) at room temperature. The electrochemical discharge capacity increases with Cu content from 150 mAh g -1 for TiNi up to 300 mAh g -1 for TiNi0.8Cu0.2 and then decreases again for larger Cu amounts.

Figure I. 9

 9 Figure I.9 Diffraction patterns of Ti1.01Ni0.99-xCux alloys a) at 473 K and b) at room temperature (Cu Kα radiation) [45].

Figure I. 10

 10 Figure I.10 Mg-Ni binary phase diagram[START_REF] Massalski | Binary alloy phase diagrams: American Society for Metals Metals[END_REF]

Figure I. 11

 11 Figure I.11 Models of the unit cell of Mg2Ni [67].

  6) series alloys by interdiffusing Mg, Ni and Al powders together and investigated the effects of partial substitution of Al for Mg on the structure and electrochemical performance. They found that a new compound, Mg3AlNi2 was formed and the electrochemical capacity and cycle life of alloy electrodes increased markedly with increasing amount of this new compound in alloys.Partial substitution of Mg by Ti could reduce the activation energy of desorption from 69 kJ mol -1 for nanocrystalline Mg2Ni to 59 kJ mol -1 for Mg1.9Ti0.1Ni and therefore destabilize the hydride phase[START_REF] Guanglie | Study on the phase composition of Mg2-xMxNi (M= Al, Ti) alloys[END_REF]. Additionally, Ti substitution for Mg could improve cycle stability of Mg-Ni-Ti ternary alloys due to the formation of TiO2 that would limit Mg(OH)2 formation[START_REF] Ruggeri | Properties of mechanically alloyed Mg-Ni-Ti ternary hydrogen storage alloys for Ni-MH batteries[END_REF]. Finally, the partial substitution of Mg by Ti in Mg-Ni-based system can introduce new phases, such as NiTi, NiTi2 and Ni3Ti. Composite phase structure is favorable for improving the hydrogen storage properties of Mg-Ni-based system.I.3.3.2 MgNi in Ni-MH systemAmorphous MgNi prepared by mechanical alloying can be charged/discharged in a secondary battery system at room temperature[START_REF] Lei | Electrochemical Behaviour of Some Mechanically Alloyed Mg-Ni-Based Amorphous Hydrogen Storage Alloys[END_REF]. It has an initial discharge capacity close to 500 mAh/g[1] compared to that of 300 mAh/g for LaNi5-type alloys used in commercial Ni-MH batteries which give rise to a new possibility for Mg-based materials. Many approaches have been explored for synthesizing MgNi amorphous phase. The influence of ball milling duration on the characteristic of the Mg:Ni (1:1) compounds was investigated by Ruggeri et al. [78] and shown in Figure I. 12. It is indicated that ball milling of MgNi mixture leads to an amorphous MgNi alloy after only 10 h of milling. Further milling results in a crystallization of amorphous MgNi into nanocrystalline MgNi2 and Mg2Ni which decrease electrode performance significantly. It is in coincidence with the study of Rojas et al. [79] concerning mechanically alloyed MgNi alloy, in which, the sequence of phase transformations during milling leading to amorphization was proposed as: c-Mg + c-Ni → nc-Mg + c-Ni → amorphous + nc-Ni + nc-Mg → amorphous + nc-Ni + nc-Mg2Ni → nc-Mg2NiWhere c and nc represent crystalline and nanocrystalline state, respectively. The first step shows the fact that grain refinement in nickel is slower than in magnesium. It has been demonstrated that the grain size attainable by milling depends on the crystal structure of the material being milled[START_REF] Fecht | Nanocrystalline metals prepared by high-energy ball milling[END_REF]. Usually, BCC materials tend to reach the smallest sizes, HCP materials somewhat larger grain sizes, and FCC materials tend to produce the largest grain sizes. Since the crystal structure of Mg is HCP and that of Ni is FCC, such a difference in the grain size after milling is expected.

Figure I. 12

 12 Figure I.12 XRD pattern evolution of Mg:Ni (1:1) with milling duration (Mg •, Ni ■, Mg2Ni ○, MgNi ▲, amorphous MgNi +) (Cu Kα radiation) [78].

  [START_REF] Zhang | The effect of Ni content on the electrochemical and surface characteristics of Mg90-xTi10Nix (x= 50, 55, 60) ternary hydrogen storage electrode alloys[END_REF] synthesized the Mg-based ternary Mg90-xTi10Nix (x=50, 55, 60) alloys with amorphous phase by mechanical alloying. The electrochemical test indicates that the cycling stability of the electrodes made of these alloys improves with increasing Ni content, while the initial discharge capacity decreases dramatically. Han et al.[START_REF] Han | Synthesis of composite metal hydride alloy of A2B and AB type by mechanical alloying[END_REF] have shown an important cycle life improvement for a composite material of Mg2Ni and TiNi synthesized from elemental Mg, Ni and Ti powders, with a charge retention of 55% after 150 cycles but with a very low maximal capacity (130 mAh g -1 ).

  [START_REF] Borissova | Synthesis and hydrogen sorption properties of Mg-Ti-Ni composite, produced by mechanical milling at liquid nitrogen temperature[END_REF] synthesized Mg-Ti-Ni composite with nominal composition (Mg0.8Ti0.2)2Ni by mechanical milling at liquid nitrogen temperature and subsequent annealing at 663 K for 3 h. It is found that the Mg2Ni is the main phase for the cryomilled powder, while the annealed composite consists of Mg2Ni, TiNi3 and trace amount of Mg3TiNi2. The reaction occurring during the hydrogenation of Ti-Mg-Ni alloy at 373 K is assumed to be as following: Mg2Ni + Ni -Ti(Mg) + xH2 → amorphous hydride + Mg2NiH0.3. Denys et al. [105] investigated several phases equilibria of Mg-Ti-Ni system at 500 °C by high energy ball milling and annealing heat treatment. One ternary intermetallic compound Mg3TiNi2 with the ordered Ti2Ni-type structure has been identified. CsCl-type structured MgxTi1-2/3xNi1-1/3x (0<x<0.5) ternary phases can be formed by dissolving Mg in TiNi compounds. The phenomena of hydrogen-induced decomposition of annealed equilibrium phases are widely observed in Ti-Mg-Ni system. Moreover, several studies of Mg-Ti-Ni ternary system from the group of Prof. Roué et al. are especially interesting and significant [76]. In their works, series of alloys with the composition of MgNiTix, Mg1-xTixNi and MgNi1-xTix (with x varying from 0 to 0.5) alloys have been prepared by high energy ball milling and tested as hydrogen storage electrodes. It is shown that, the addition of Ti improved the cycling properties of Mg-Ni-Ti ternary alloys compared to Mg-Ni binary alloys with the same Mg/Ni atomic ratio. Mg0.5Ti0.5Ni electrode possesses the best cycle life in all studied materials which retains 75% of initial capacity after 10 cycles in comparison to 39% for MgNi electrodes. On the basis of the XRD patterns, the MgTi0.5Ni0.5 and Mg0.5Ni0.5Ti alloys may be composed of a mixture of MgTi + MgNi and MgTi + NiTi phases, respectively, or made of a single phase that contains all the elements [106]. In general, the ternary Mg-Ni-Ti electrodes had a much higher discharge rate capability than the binary MgTi electrode. It can be concluded that, the ternary Mg-Ti-Ni electrodes displayed a much better cycle life than the binary Mg-Ti or Mg-Ni electrode. The cycle life improvement of the Mg-Ti-Ni electrode can be related to the formation of TiO2 which limits Mg(OH)2

  Figure II. 1 exhibits the schematic depicting the ball motion inside the ball mill which contains two types of pressing the submitted materials undergo during milling: clashing of two balls (Figure II. 2(a)) or clashing of the ball and the vial (Figure II. 3(b)). The vials and the supporting disc rotate in opposite directions and the centrifugal forces alternately act in like and oppositedirections. This causes the milling balls to run down the internal wall of the vial --the friction effect, followed by the material being milled and milling balls lifting off and traveling freely through the inner chamber of the vial and colliding against the opposing inside wall[2].

Figure II. 1

 1 Figure II.1 Schematic depicting of the ball motion inside the ball mill [2].

  the present study, a planetary type mill Retsch PM 400 is used with steel vials and steel balls. The details of the milling setup are illustrated in Figure II. 2., where two steel balls (d = 20 mm) are placed in 50 ml volume steel vials. The filling of the starting powder was carried out in a glove box filled with argon as the protective gas. The rotation speed was fixed at 400 rpm. In order to dissipate heat, all millings were scheduled a 30 min. pause every 30 min.

Figure II. 2

 2 Figure II.2 Retsch PM 400, vials and balls and present planetary mill setup.

  Figure II. 3. shows the equipment in which the sample stage, and the detector/counter lies at the two sides of the stage. It consists of: (1) goniometer, (2) X ray tube, (3) primary optics, (4) sample platform (5) secondary optics, (6) detector. Note that the focus of the X-ray tube, the specimen, and the detector slit are located on a single circle.

Figure II. 3

 3 Figure II.3 Photography and X-ray diffraction mechanism of Bruker-AXS D8 FOCUS Advanced Bragg-Brentano X-ray Powder Diffractometer.

Figure II. 4

 4 Figure II.4 JSM-7800F field emission scanning electron microscope.

Figure II. 5

 5 Figure II.5 Example of TEM diffraction image.

Figure II. 6

 6 Figure II.6 TEM model JEOL JEM-2100.

Figure II. 7 2 Figure II. 8

 728 Figure II.7 Equipped TEM sample preparation platform.

K for 8 h

 8 to ensure the full dehydrogenation of the alloys. After the third absorptiondesorption cycle, PCI measurements were carried out at different temperatures. The temperature of the sample was controlled using a furnace with ± 2 K accuracy. The measurement conditions were set as: delay time 30 s, maximum pressure 3.5 MPa, minimum pressure 2 kPa.

Figure II. 9

 9 Figure II.9 PCTPro Sievert's type apparatus.

Figure II. 10

 10 Figure II.10 Simplified block diagram of a chronopotentiometric measurement device.The working-counter-and reference electrode are denoted as WE, CE and RE, respectively.

  performed using a Multi-channel Battery Interface ATLAS 0461, as seen in Figure II. 11. The negative electrodes (working electrodes) consist of as-milled powder, carbonyl Ni powder with the ratio of 90:10. The mixtures were compressed between two nickel dies by means of cold isostatic pressing under a pressure of 500 MPa to obtain a small pellet. NiOOH/Ni(OH)2 and Hg/HgO were respectively employed as counter electrode and reference electrode. The electrolyte was 6 M KOH aqueous solution. In order to eliminate adverse oxidative effects, the electrodes were soaked in the electrolyte 24 h at room temperature and 1h at 100°C for activation, before electrochemical tests of each sample. The current density for charging and discharging was 40 mA/g. The cut-off potential was -0.7 V vs. the reference electrode.

Figure II. 11

 11 Figure II.11 Appearance of Muti-channel Battery test platform.

υ

  𝐾𝑆 (r) = υ 𝑒𝑥𝑡 (r) + υ 𝐻 (r) + υ 𝑥𝑐 (r) EQN II.10 with the Hartree potential υ 𝐻 (r), υ 𝐻 (r) = 𝜕𝐸 𝐻 [𝜌(𝑟)] 𝜕𝜌(𝑟) = ∫ 𝜌(𝑟 ′ ) |𝑟-𝑟 ′ | 𝑑𝑟 ′ EQN II.11and the exchange-correlation potential υ 𝑥𝑐 (

  densities, 𝜌 𝛼 (𝑟 ⃗) and 𝜌 𝛽 (𝑟 ⃗) with 𝜌(𝑟 ⃗) = 𝜌 𝛼 (𝑟 ⃗) + 𝜌 𝛽 (𝑟 ⃗), and the form of the localspin-density approximation (LSDA) is: 𝐸 𝑥𝑐 𝐿𝑆𝐷𝐴 [𝜌 𝛼 (𝑟 ⃗), 𝜌 𝛽 (𝑟 ⃗)] = ∫ 𝜌(𝑟 ⃗)𝜀 𝑥𝑐 [𝜌 𝛼 (𝑟 ⃗), 𝜌 𝛽 (𝑟 ⃗)] 𝑑𝑟 ⃗ EQN II.17



  Generalized gradient approximations (GGA)

  alloying using a planetary high-energy ball mill. The effect of Ni addition in TiMg, partial substitution of Ni for Ti in TiMg and partial substitution of Ni for Mg in TiMg have been studied through the compared analysis of TiMgNix, MgTi1-xNix and TiMg1-xNix,. The effects of the variation of composition and milling time on microstructure and hydrogenation properties of Ti-Mg-Ni alloys are investigated. Additionally, a systematic approach has been employed for improving the electrochemical properties of Ti-Mg-Ni alloys.

A

  planetary type mill machine (Retsch PM 400) was used in this study with steel vials and steel balls. Elemental Ti (hcp, size ≤ 150 µm), Mg (hcp, size ≤ 100 µm) and Ni (fcc, size ≤ 150 µm) powders purchased from GoodFellow with a minimum purity 99.5% were used as starting materials. The morphologies of elemental Ti, Mg and Ni were characterized by SEM and shown in Figure III.1. As seen in Figure III. 1, the elemental Ti powders are of irregular blocky shape with particle size ranging 20-100 µm . The elemental Mg powders mainly consist of flake structure with size ranging 10-50 µm . The elemental Ni powders are spheres with particle size ranging 20-100 µm .

Figure III. 1

 1 Figure III.1 The morphologies of starting materials: (a) Ti, (b) Mg and (c) Ni.

Figure III. 2

 2 Figure III.2 shows the XRD results of Ti50Mg50 alloys ball milled for 60 h by

Figure III. 2

 2 Figure III.2 XRD results of Ti50Mg50 alloys ball milled for 60 h by (a) 4 balls of 14 mm and (b) 2 balls of 21 mm.

Figure III. 3

 3 Figure III.3 XRD results of Ti50Mg50 alloys ball milled at (a) 200 RPM for 100 h and (b) 400 RPM for 60 h.

Figure III. 4

 4 Figure III.4 XRD results of Ti50Mg50 alloys ball milled for 60 h (a) with PCA and (b) without PCA.

  MgOindicated that, a sufficiently long time of ball milling is necessary for the synthesis of MgTi BCC phase. Asano et al.[2] synthesized BCC TiMg by mechanical alloying for 200 h with the speed of 200 RPM. Thus, it seems that our milling process is efficient and reliable.

Figure III. 5

 5 Figure III.5 XRD results of Ti50Mg50 alloys ball milled for (a) 5 min, (b) 10 h, (c) 20 h, (d) 40 h and (e) 60 h.

Figure III. 6 

 6 Figure III.6 XRD patterns of (a) TiMgNi0.2, (b) TiMgNi0.4, (c) TiMgNi0.6, (d) TiMgNi0.8, (e) TiMgNi milled 40h.

Figure III. 7 

 7 Figure III.7 XRD patterns of TiMgNi0.2 milled for (a) 20 h, (b) 40 h and XRD pattern of TiMgNi0.1 milled 40h (c)

Figure III. 8

 8 Figure III.8 XRD patterns of TiMgNi0.4 mixture after different milling time: (a) 20 h, (b) 40 h.

  Figure III.9 (a) and Figure III.9 (b). Particularly, the diffraction pattern shown in Figure

Figure III. 9

 9 Figure III.9 XRD patterns of TiMgNi mixture after different milling time: (a) 20 h, (b) 40 h.

  MgNi2 nanocrystals detected by high-resolution microscopy bedded on amorphousMgNi may also lead to the absence of Mg-Ni phase in the diffraction pattern. And it is considered that the formation of Mg2Ni is reasonable compared with that of MgNi2 in such a superstoichiometric magnesium environment. The SEM images of different samples after 40 h of milling are shown in Figure III.10. The morphology of all samples consists of micrometric particles. In comparison with the different micrographs, with increasing the Ni content, the particles are getting more irregular which have a tendency to agglomerate. Accompanying the increasing of the Ni content, more fine particles smaller than 5 micrometer and agglomerated particles bigger than 20 micrometers were found. This reflects that, the repeated fracturing and the cold-welding events are occurring together during the ball milling process. Ni as a positive factor which can improve the milling efficiency of the alloys.

Figure

  Figure III.10 SEM micrographs of (a) TiMgNi0.2, (b) TiMgNi0.4, (c) TiMgNi0.6, (d) TiMgNi0.8, (e) TiMgNi milled 40h

Figure III. 11

 11 Figure III.11 Discharge capacities as a function of charge/discharge cycles of TiMgNix alloys.

  10 wt.% of carbonyl nickel powder were added as electrode additive to improve the conductivity before electrochemical tests. The presence of nickel in electrode material can increase electrical conductivity, catalyze electrode reaction and surface hydrogen dissociation, prevent oxygen from diffusing to internal part of the alloy[19]. Except the peaks of Ni, the broad peak in the range of 40°-50° is present for each diffraction pattern which is consistent with the broad peak found for the as-milled samples. It can be inferred that the amorphous phase does not undergo decomposition or crystallization process during cycling, which implies that the TiMgNix alloys achieved a good corrosion resistance in alkaline electrolyte. The other peaks observed in the XRD patterns of Fig. III.12 (b) and (e) at 2θ = 45°are assumed to correspond to the formation of Mg(OH)2 [20], whereas, they are not found in Fig. III.12 (c) and (d). This will be used to reveal the structural modification of the active electrode material as it goes through the successive charge/discharge cycles. It was however difficult to estimate the content of Mg(OH)2 due to their relatively low intensities and overlapped peaks with amorphous phase/nanocrystallites. The formation of Mg(OH)2 can cause a capacity deterioration of Mg based hydrogen storage alloys. According to the electrochemical results, the cycling stability of TiMgNix alloys increases with adding Ni ratio until x = 0.8, which possesses the best cycling property. It is in agreement with the XRD characterization of electrodes after cycling. A longer cycling test should be carried out in the future to examine the suppressing effect of Ni addition on the formation of Mg(OH)2.

Figure III. 12

 12 Figure III.12 XRD patterns of electrode, after 21 cycles of electrochemical studies (a) TiMgNi0.2, (b) TiMgNi0.4, (c) TiMgNi0.6, (d) TiMgNi0.8, (e) TiMgNi.

Figure III. 13

 13 Figure III.13 Transmission electron microscopy image (a) for the as-milled TiMgNi alloy and its corresponding selected area electron diffraction (SAED) (b) and EDS map (c)

Figure III. 14

 14 Figure III.14 suggests the presence of an amorphous material or a "poly-

Figure III. 14

 14 Figure III.14 Transmission electron microscopy image for the as-milled TiMgNi alloy and its corresponding selected area electron diffraction.

Figure III. 15

 15 Figure III.15 Dark field image of as-milled TiMgNi alloy (a) and its corresponding selected area (SAED) images with points of measurements (b).

Figure III. 16

 16 Figure III.16 HRTEM image of the nanocrystallites.

  to high temperature with hydrogen also drew our attention. Therefore, PCT measurements of TiMgNix alloys were carried out at 325 °C for investigating its solid-gas hydrogenation properties. The experimental methodology of PCT measurements are described in chapter II.

Figure III. 17 Figure

 17 Figure III.17 Pressure-composition-isotherm (PCI) curves of TiMgNix alloys (x = 0.2, 0.4, 0.6, 0.8 and 1) measured at 325 °C

  The XRD of as-milled TiMgNi alloys and activated TiMgNi alloys (heat treated at 400 °C for 1 h) are shown in Figure III.19 (a) and (b). It is found that, the activated TiMgNi alloys is constructed by crystalline Mg and TiNi phase with broaden feature. Amorphous Mg-Ni phase may overlapped in the broaden feature of TiNi phase. It is in agreement with our XRD analyses of as-milled TiMgNi samples. The crystallization of alloy is promoted by heat treatment process. Whereas, the formation of TiNi3 or any other phase is not observed in the activated sample.

Figure III. 19

 19 Figure III.19 XRD patterns of TiMgNi after (a) mechanical alloying (b) activation (c) hydrogenation (b) dehydrogenation

III. 3

 3 High energy ball milled MgTi 1-x Ni x alloys III.3.1 Microstructure of as-milled MgTi 1-x Ni x alloys A series of MgTi1-xNix alloys (x = 0.2, 0.4, 0.6 and 0.8) were prepared by ball mill.

Figure III. 20

 20 Figure III.20 shows the X-ray diffraction patterns of different composition samples with

Figure III. 20

 20 Figure III.20 XRD patterns of (a) MgTi0.8Ni0.2, (b) MgTi0.6Ni0.4, (c) MgTi0.4Ni0.6 and (d) MgTi0.2Ni0.8 milled for 40 h.

  Figure III.1 (b), are flake structure with size ranging 10-50 µm. Due to the softness of Mg, the plate-like Mg tend to stack on the surface of the milling pot which results in the hindering of the alloying of Mg. The residues found on the surface of milling pot after milling process further consist this conclusion which mainly consist of Mg. Moreover, The TEM micrograph of MgTi0.4Ni0.6 and its corresponding EDS results are shown in Figure III.21. It is confirmed from the selected area electron diffraction patterns (SAED) of the composite matrix that only a halo pattern is present (as seen at the top of Figure III.21 (a)), which indicates that no obvious long-range ordered structure of the composite is obtained after ball milling. The MgTi0.4Ni0.6 alloy mainly consists of amorphous structure. The corresponding EDS mapping analysis (as seen in Figure III.21 (b), (c) and (d)) reveal the elemental distribution of Mg, Ti and Ni in the selected area. It is found that Mg is dispersed in the entire area, whereas Ti and Ni are concentrated in the same area where is in the center of the selected area. Energy dispersive X-rays analysis data shows the following distribution in the central area of the image: 44 at.% Mg, 26 at.% Ti and 30 at.% Ni. It also suggests that the Ti and Ni are equally distributed which is in accordance with the EDS result of TiMgNi in the last section. It indicates that a Ti-Ni-(Mg) could exist as a disorder phase in the composite.

Figure III. 21

 21 Figure III.21 TEM micrograph of MgTi0.4Ni0.6 (a) and corresponding EDS mapping analysis of (b) Mg, (c) Ti and (d) Ni.

Figure III. 22

 22 Figure III.22 Discharge capacities as a function of charge/discharge cycles of MgTixNi1-x alloys.

  peak corresponding to Mg is superimposed on this broad feature. Similar phase structure of TiMg0.8Ni0.2 has also been obtained in the previous XRD result of asmilled TiMgNi0.2 (Figure III.6 (a)), which indicated that no matter the addition of Ni or the substitution of Ni will decrease the solution of Mg in Ti and further result in hindering of the formation of TiMg binary alloy. With the decrease the atom ratio of Mg/Ni, the diffraction peaks of Mg disappear from the XRD patterns while those of broad peak are always present and shift towards the high angles. It is believed that the broad peaks are associated with the formation of Ti-Ni and Mg-Ni binary phases. To illustrate the phase transformation during milling process, the XRD patterns of TiMg0.2Ni0.8 with different milling time are shown in Figure III.24.

Figure III. 23

 23 Figure III.23 XRD patterns of (a) TiMg0.8Ni0.2, (b) TiMg0.6Ni0.4, (c) TiMg0.4Ni0.6 and (d) TiMg0.2Ni0.8 milled for 40 h.



  Figure III.24 XRD patterns of TiMg0.2Ni0.8 milled for (a) 10 h, (b) 20 h and (c) 40 h.

Figure III. 25

 25 Figure III.25 Transmission electron microscopy image for the as-milled TiMg0.2Ni0.8 alloy (a) and its corresponding selected area electron diffraction (b).

Figure III. 26

 26 Figure III.26 Discharge capacities as a function of charge/discharge cycles of TiMgxNi1-x alloys.

Figure III. 27

 27 Figure III.27 Illustration of discharge capacities as a function of ternary alloys in Ti-Mg-Ni system.

Figure III. 28

 28 Figure III.28 XRD results of TiMgNi milled for (a) 10 h, (b) 20 h, MgTi0.4Ni0.6 milled for (c) 10 h, (d) 20 h and TiMg0.4Ni0.6 milled for (e) 10 h, (f) 20 h.

Figure III. 29

 29 Figure III.29 Discharge capacities as a function of charge/discharge cycles of TiMgNi milled for (a) 10 h, (b) 20 h, MgTi0.4Ni0.6 milled for (c) 10 h, (d) 20 h and TiMg0.4Ni0.6 milled for (e) 10 h, (f) 20 h.

8 , 1 )

 81 , all the asmilled TiMgNix alloys consist of a mixture of Mg and amorphous Ti-Ni and Mg-Ni phases. Within all studied samples, the TiMgNi composition shows the highest discharge capacity. The discharge capacities and activation properties of TiMgNix alloys increase linearly when the Ni content increases. TEM analysis shows that nanocrystallites and an amorphous phase co-exist in the as-milled TiMgNix alloys. The MgTiNi0.8 alloy possesses the best cycling properties which is consistent with its XRD diagram after the electrochemical test. The decay of the discharge capacity during cycling is mainly attributed to the formation of Mg(OH)2. No decomposition or crystallization of the main alloy components are found after 21 consecutive charge/discharge cycles which implies that the as-milled TiMgNix alloys boasts a good corrosion resistance in alkaline electrolytes. PCI tests were carried out on TiMgNix alloy at 598 K. The reversible hydrogen storage capacities of TiMgNix alloys are decreased with increasing of Ni content. All of the studied materials are constructed by the phase composition which are Mg, Mg2Ni, TiH and TiNi3. The formation of equilibrium phase of Mg2Ni and TiNi3 are catalyzed by hydrogen under high temperature. XRD analyses show that the dehydrogenated TiMgNi sample consists of TiH2, Mg, Mg2Ni and TiNi3, in which Mg and Mg2Ni phases appears as the major actors of the hydrogen absorption process.In the case of MgTixNi1-x alloys (x = 0.2, 0.4, 0.6 and 0.8), all the as-milled samples consist of a mixture of Mg and amorphous Ti-Ni and Mg-Ni phase, except MgTi0.2Ni0.8 alloy, which is mainly constructed of Mg2Ni, TiNi, TiNi3 and slight amount of Mg residue. TEM results further confirm the existence of TiNi phase with equal atom ratio in amorphous phase. It is indicated that, the crystallization of MgTixNi1-x alloys tend to be reinforced with the increase of atom ratio of Mg/Ni. Among all of the electrochemical studied materials, MgTi0.4Ni0.6 possesses the best discharge capacity of 96 mAh/g in the first charge/discharge cycle. The good electrochemical performance is considered to attribute to the formation of Mg2Ni amorphous phase. the average capacity decay of MgTi0.4Ni0.6 was slow as 1%, which is due to the coexisted TiNi and Mg2Ni. The synergistic effects of TiNi and Mg2Ni MA alloys on the electrochemical performance has been proposed.In the case of TiMgxNi1-x alloys (x = 0.2, 0.4, 0.6, 0.8) alloys, all of the XRD patterns exhibit a broad peak located around ca. 2θ = 48° except TiMg0.8Ni0.2, in which a sharp peak corresponding to Mg is superimposed on this broad feature. the sequence of phase transformations during ball milling of TiMg0.2Ni0.8 is proposed as: TiMg0.2Ni0.8 → Ti2Ni → Ti2Ni + TiNi → amorphous.With the increase the atom ratio of Mg/Ni, the diffraction peaks of Mg disappear from the XRD patterns while those of broad peak are always present and shift towards the high angles. The maximum discharge capacity of TiMg1-xNix increase linearly with increase of the atom ratio of Ni/Mg. The improvement of discharge capacity is attributed to the formation of Ti-Ni phase in such Ti-rich ternary system. All the alloys exhibit good cycling stability after activation process, which is in agreement with the characteristic of Ti-Ni alloys.By comparing analysis of different series of samples, it is found that the amorphous phases existed in most of as-milled samples are mainly constituted by the Ti-Ni because of their high dependence on the atom ratio of Ti/Ni. The existence of Mg is assumed to be dispersed in the Ti-Ni amorphous phase. Within all the studied alloys, TiMg0.2Ni0.8 which possesses the least amount of Mg but exhibits the highest discharge capacity. The reason of degradation of the rest of alloys are attributed to the unreacted Mg during milling process. In order to improve the milling efficiency of Mg in Ti-Mg-Ni alloys, two-steps milling have been carried out for selected composition. It is found that, the two-steps milling process can promote the amorphization of Ti-Mg-Ni alloys, which results in a significant amelioration of their discharge capacities.TiNi is as known for its outstanding shape memory properties which are based on a reversible martensitic transformation. In the meantime, TiNi exhibits electrochemical activity and good resistance to corrosion in alkaline media which can be used as negative electrode material in Ni-MH batteries. Our previous study has shown that TiNi based alloys can be easily synthesized by mechanical alloying and exhibits moderately good discharge capacity. Moreover, a recent study arose our interest on this system: In that paper, Cu substituted TiNi alloys with Martensite phase have been synthesized by induction melting and showed ameliorative electrochemical performance at room temperature. Thus, it is of interest to investigate this system by other approaches.Meanwhile, it is known that the electronic structure plays an important role in understanding physical and chemical processes on metallic surfaces. In order to understand the adsorption of hydrogen on TiNi surface at the microscopic level it is necessary to investigate the electronic structure of clean TiNi surfaces and its change under the influence of hydrogen.In this chapter, Cu substituted TiNi alloys have been investigated as hydrogen storage material for Ni-MH batteries by experiments and first principle calculations. The phase transformations and electrochemical properties of TiNiCu alloys are characterized. The influence of Cu substitution for Ni in TiNi on the electronic structure of the resulting alloys and the hydrogen adsorption energies are investigated by first principle calculation.

ΔHf= 1 where

 1 Figure IV.1 shows the XRD patterns of TiNi1-xCux alloys after 20 h of milling. All

Figure IV. 1

 1 Figure IV.1 XRD patterns of Ti-Ni and Ti-Ni-Cu alloys after 20 h of milling: (a) TiNi, (b) TiNi0.9Cu0.1, (c) TiNi0.8Cu0.2, (d) TiNi0.7Cu0.3.

Figure IV. 2

 2 Figure IV.2 SEM micrographs of annealed samples: (a) TiNi, (b) TiNi0.9Cu0.1, (c) TiNi0.8Cu0.2, (d) TiNi0.7Cu0.3.

The

  XRD patterns of annealed TiNi1-xCux alloys are depicted in Figure IV.3. It was found that all samples presented a multi-phase structure. As denoted, the main phases can be indexed in terms of TiNi B2 phases. Comparing to as-milled samples, the diffraction peaks of B2 phases became narrower and their intensities increased. The secondary phase of TiNi alloys (in Figure IV.3(a)) was indexed as Ti2Ni phase. With the substitution of Cu, Ti2(Ni,Cu) phase was observed which is coexisted with Ti2Ni

Figure IV. 3

 3 Figure IV.3 XRD patterns of annealed Ti-Ni and Ti-Ni-Cu alloys at 750 °C for 0.5 h: (a) TiNi, (b) TiNi0.9Cu0.1, (c) TiNi0.8Cu0.2, (d) TiNi0.7Cu0.3.

Figure IV. 4

 4 Figure IV.4 The total DOS of TiNi1-xCux alloys. The Fermi level is chosen as the origin of the energy.

Figure IV. 5

 5 Figure IV.5 The structural illustration of TiNi1-xCuxH system.

Figure IV. 6

 6 Figure IV.6 The discharge capacities as a function of charge/discharge cycles of annealed Ti-Ni and Ti-Ni-Cu alloys.

. 4 .

 4 It is shown that partial substitution of Ni by Cu completely

•••

  5 h. The structures of Cu substituted TiNi alloys were simulated by first principle calculation. Based on the XRD analysis, as-milled samples possessed single phases with Pm-3m cubic structure. After heat treatment, studied materials consisted of Ti(Ni,Cu) B2 phase as main phase, Ti2Ni and Ti2(Ni,Cu) as secondary phases. The lattice parameters of B2 phase increased linearly with increasing Cu substitution. The computational results concurred with the experimental results and indicated that TiNi0.7Cu0.3 possessed the most negative enthalpy of formation. The stability of TiNi1-xCux phase increased with the increasing Cu content. The width of valence bands was enlarged by substituting Cu atoms. Meanwhile, With the substitution of Cu for Ni, the hydrogen adsorption energy of TiNi1-xCux alloys decrease linearly. In the electrochemical tests, annealed TiNi showed the highest discharge capacity among all of studied materials which is equal to 154 mAh/g. Accompanying the substitution of Ni by Cu, The discharge capacity and cycling abilities of Cu substituted TiNi alloys decline and increase, respectively. All of the substituted samples exhibit at least 97% retaining rate after 20 cycles. As energy crisis and environmental pollution are becoming more and more serious, hydrogen is widely regarded as the most promising energy resource to replace the traditional fossil fuel. A safe hydrogen storage method with high storage efficiency and low cost is the key issue for the utilization of hydrogen as an energy carrier. Metal hydrides are some metals and alloys which can reversibly absorb and desorb hydrogen under moderate temperature and pressure. This gives them an important safety advantage over other traditional hydrogen storage methods (like in gaseous or liquid forms). Besides, another notable application area for metal hydride is in Ni-MH batteries. The materials used for the negative electrode of Ni-MH batteries are key to battery performance, namely discharge capacity, high-rate capability (both charging and discharging), initial activation, cycle life, self-discharge rate. Among all metal hydrides candidates, Mg-Ni and Ti-Ni systems have being investigated by researchers for several decades due to their potential gravimetric and volumetric assets. However, these two binary alloys are far from practical application due to their mediocre cycling stability and relatively low discharge capacity. Fulfilling all the requirements for optimum performance of hydrogen-based batteries thus remains a real challenge. To overcome this problem, composite hydride materials have been introduced as electrode materials as a way to improve the balance of discharge capacity and cycling performances. Alloys pertaining to the Ti-Mg-Ni ternary system are of interest as potential metal hydride since they consist of two hydride forming metals (Ti and Mg). The hydrogen storage properties of these alloys should be improved with Ni additions thanks to the high catalytic activity of the latter. There is therefore a vivid interest for the development of such ternary Ti-Mg-Ni systems in view of hydrogen storage applications, especially by means of high energy ball milling (HEBM). In this thesis, series of Ti-Mg-Ni based alloys which are used as the negative electrode materials in Ni-MH batteries have been synthesized by mechanical alloying. The structural transformation and hydrogenation properties of prepared samples have been characterized and examined. Different approaches have been applied for improving the discharge capacities of selected system. Additionally, the influence of Cu substitution on the phase transformation and electrochemical properties of TiNi alloys have been investigated. The experimental study is also added and complemented by first principle calculation, which further demonstrates the structural results of the experiments. Important findings and conclusions by this research are summarized as the followings: In the case of MgTiNix (with x = 0.2, 0.4, 0.6, 0.8, 1), all the as-milled TiMgNix alloys consist of a mixture of Mg and amorphous Ti-Ni and Mg-Ni phases. Within all studied samples, the TiMgNi composition shows the highest discharge capacity. The discharge capacities and activation properties of TiMgNix alloys increase linearly when the Ni content increases. TEM analysis show that nanocrystallites and an amorphous phase co-exist in the as-milled TiMgNix alloys. The MgTiNi0.8 alloy possesses the best cycling properties which is consistent with its XRD diagram after the electrochemical test. The decay of the discharge capacity during cycling is mainly attributed to the formation of Mg(OH)2. No decomposition or crystallization of the main alloy components are found after 21 consecutive the amorphization of Ti-Mg-Ni alloys, which results in a significant amelioration of their discharge capacities. In the case of TiNi1-xCux (x = 0, 0.1, 0.2, 0.3) alloys, as-milled samples possessed single phases with Pm-3m cubic structure. After heat treatment, studied materials consisted of Ti(Ni,Cu) B2 phase as main phase, Ti2Ni and Ti2(Ni,Cu) as secondary phases. The lattice parameters of B2 phase increased linearly with increasing Cu substitution. The computational results concurred with the experimental results and indicated that TiNi0.7Cu0.3 possessed the most negative enthalpy of formation. The stability of TiNi1-xCux phase increased with the increasing Cu content. The width of the valence band was enlarged by substituting Cu atoms. Annealed TiNi showed the highest discharge capacity among all of studied materials which is equal to 154 mAh/g. Accompanying the substitution of Ni by Cu, The discharge capacity and cycling abilities of Cu substituted TiNi alloys decline and increase, respectively.

Table I

 I Especially, the volumetric and gravimetric capacities are considered along with cost and efficiency[10]. DOE analyses are most readily available for physical methods of hydrogen storage because they are the most traditional storage methods and so were the first to be investigated. Unfortunately, no physical hydrogen storage technology currently meets DOE standards[11], because hydrogen is diffuse and buoyant and with extremely low density. Although it has been stored in the gaseous form for centuries and as a liquid since the 1940's, traditional mechanical means of

	Storage system parameter	Original 2010 target Revised 2010 target 2017 target Ultimate target
	Gravimetric capacity kgH2 /kg	6%	4.5%	5.5%	5.5%
	Volumetric capacity g H2/L	45	28	40	70
	Operational cycle life	1000	1000	1500	1500
	Filling time (min, for 5 kg)	3	4.2	3.3	2.5
	Fuel purity	99.99%	99.97%	99.97%	99.97%

.1 A subset of original and revised DOE targets for on-board hydrogen storage systems for light-duty vehicles

[9]

.

As storing hydrogen is a key aspect of a hydrogen energy infrastructure, the US Department of Energy (DOE) released targets for hydrogen storage systems in 2003.

Whereas, the targets were revised in 2009. The DOE targets consider volumetric and gravimetric capacities, refilling time, costs, cycle life and loss of usable hydrogen. The original and revised DOE targets are shown in Table

I

. 1. It is indicated that the standard of revised 2010 targets are lowered somewhat in comparison to the original 2010 targets with regard to system gravimetric capacity, system volumetric capacity, and refilling time, which means the hydrogen storage for onboard application is not that easy to accomplish. hydrogen storage (compression and liquefaction) are very energetically and financially expensive. Therefore, research is constantly underway to improve these storage techniques and develop new storage mechanisms. Many alternative storage methods are being extensively investigated, in which, metal hydride is one of the most potential candidates for solving this technical issue.

  Table III.1. It is found that TiMgNi0.8 possess the best cycling properties whose retaining rates are more than 85% after 21 cycles. It can

be inferred that Ni atom can ameliorate the cycling performance of TiMgNix alloys, but may not be the only factor which accounts for the capacity decay.

  Table III.1 Discharge capacities and retaining rates of as-milled samples

	Composition	TiMgNi0.4 TiMgNi0.6 TiMgNi0.8 TiMgNi
	Cmax (mAh/g)	60	95	110	162
	C21 (mAh/g)	33	81	97	119
	Retaining rate (R21)	55%	85%	88%	74%
	To explain the capacities decay during charge-discharge cycles, the structure of
	electrodes after 21th cycles was characterized by XRD and are presented in Fig. III.12.

Table III

 III

	.2 Quantitative analysis on the element compositions of zone 1 in Figure
		III.12 (C)
	Spectrum	Content (at.%)
		Ti	Mg	Ni
	Zone 1	37.1	22.9	40
	Zone 2	4.3	90	5.7
	Zone 3	3.1	42.4	54.5

Table IV .

 IV 1 Average particle size and chemical composition of annealed samples.

	Sample	Average particle size (SEM) [µm]	Ti	Atom ratio Ni	Cu	Chemical composition
	TiNi	44	54.9 45.1	-	TiNi0.82
	TiNi0.9Cu0.1	36	53.5 42.9 3.6	TiNi0.8Cu0.07
	TiNi0.8Cu0.2	27	52.5 39.6 7.9	TiNi0.75Cu0.15
	TiNi0.7Cu0.3	29	53.2 33.2 13.6	TiNi0.62Cu0.25

  Table IV.2 Experimental and calculated lattice parameters of annealed samples.

		Experimental		Calculated
	Sample	Lattice parameter	Cell volume	Lattice parameter	Cell volume	Formation enthalpy
		[Å]	[Å 3 ]	[Å]	[Å 3 ]	[eV]
	TiNi	3.014	27.380	3.016	27.434	-13.0948
	TiNi0.9Cu0.1	3.016	27.456	3.017	27.459	-15.7643
	TiNi0.8Cu0.2	3.020	27.546	3.021	27.592	-21.3593
	TiNi0.7Cu0.3	3.031	27.866	3.027	27.753	-29.3518

Table IV .

 IV 4 Discharge capacities and capacity retaining rate of annealed samples.

	Composition	TiNi	TiNi0.9Cu0.1 TiNi0.8Cu0.2 TiNi0.7Cu0.3
	Cmax [mAh/g]	154	116	93	81
	C20 [mAh/g]	127	113	91	79
	Retaining rate [R20] 82%	97%	97%	98%

List of tables

CHAPTER II

Fundamentals of Preparation,

Characterization and Calculation.

CHAPTER III

Synthesis, characterization and hydrogenation properties of Ti-Mg-Ni alloys

It is known that Mg2Ni phase possesses a high discharge capacity but also rapid decay with cycling [27]. Unlike that, the average capacity decay of MgTi0.4Ni0.6 was lower than 1%, which is attributed to the co-existence of Mg2Ni and Ti-Ni phases. From the TEM results of MgTi0.4Ni0.6, the existence of Ti-Ni-(Mg) phase has been confirmed. the synergistic effects of TiNi and Mg2Ni MA alloys on the electrochemical performance was also reported by Huang et al. [28]. The characteristic of Ti-Ni phase which possesses a good corrosion resistance in alkaline solution but mediocre activation property is beneficial for the electrochemical property of composite.

Therefore, amorphous TiNi and Mg2Ni cooperate to improve the discharge capacity.

In contrast to MgTi0.4Ni0.6 alloy, MgTi0.2Ni0.8 alloy which also consisted of TiNi and MgNi amorphous phases possessed a lower discharge capacity but better cycling performance comparing with MgTi0.4Ni0.6. This is mainly due to the formation of TiNi3. TiNi3, as one of the main phase achieved in the as-milled powder is an inactive phase which can not store hydrogen [20]

CHAPTER IV

The experimental and ab initio study of

Cu substituted TiNi alloys

the milling chamber walls and grinding balls. All of these operations were carried out under a high pure argon atmosphere.

Annealing process was applied subsequently for all of the obtained materials to recrystallize TiNi cubic structure. Drenchev et.al. [1] reported that the crystallization of ball-milled TiNi alloys can not be achieved sufficiently by annealing at 700 °C. A higher temperature or a longer holding time should be required. Thus, in our studies, all of the samples were heat treated at 750 °C for 0.5 h under argon protected environment.

IV.1.2 Computational model and method

In order to investigate the influence of Cu substitution on the crystal structure evolution of TiNi alloys, the cubic TiNi phase was constructed. The unit cell of TiNi BCC phase belongs to the space group Pm-3m with lattice parameters a = b = c = 3.015 Å [2]. All first principle calculations in this work were performed with the code VASP.

Perdew Burke Ernzerhof (PBE) [3] and generalized gradient approximation (GGA)

were used as approximate function for the exchange and correlation potential. • By comparing analysis of different series of samples, it is found that the amorphous phases that existed in most of as-milled samples are mainly constituted by the Ti-Ni because of their high dependence on the atom ratio of Ti/Ni. The Mg is assumed to be dispersed in the Ti-Ni amorphous phase. Within all the studied alloys, TiMg0.2Ni0.8 possesses the least amount of Mg but exhibits the highest discharge capacity. The reason of degradation of the rest of alloys are attributed to the unreacted Mg during milling process. In order to improve the milling efficiency of Mg in Ti-Mg-Ni alloys, two-steps milling have been carried out for selected composition. It is found that, the two-steps milling process can promote combination of melting and mechanical alloying) should be carried out.

• We have demonstrated that the composite TiMg2Ni prepared by two-steps milling process exhibits the best discharge capacities in our studied samples. It is attributed to the co-existence of Mg-Ni and Ti-Ni phases. Thus, it is worthy to continue optimizing the composition in Ti-Mg-Ni system. Meanwhile, the synergistic effect of Mg-Ni and Ti-Ni phases is still of interest to be investigated.

• It is reported that the TiNi martensite phase is a potential hydrogen absorbing phase which possesses more sites for the absorption of hydrogen than austenite TiNi phase. Whereas in our study, the martensite TiNi phase could not be obtained by substituting Cu for Ni using mechanical alloying at room temperature. Hence it would be worthy to find out the milling condition or mechanism of the production of the martensite TiNi phase. Once it can be obtained by mechanical alloying, its hydrogen storage properties should be studied.

• The calculation results for the TiNiCu system obtained in this thesis are in good agreement with the experimental observations. It is revealed that the variation of the electronic structure influences the hydrogen adsorption energy of the system which may further result in the modification of its hydrogenation property.

However since the investigation of surface adsorption in this study is not sufficient, this inference based on empiricism is uncertain. More computational studies concerning hydrogen adsorption on other surfaces and/or in the different position should be launched. Meanwhile, these calculation process should be employed in other elemental substituted TiNi ternary system for predicting their structural and Publications: Abstract:

The storage of hydrogen is one of the biggest technical problems that restrict the practical application of hydrogen. Metal hydrides are mainly regarded as the solution for facing this issue since they can reversibly absorb and desorb big amount of hydrogen under moderate temperature and pressure. Meanwhile, metal hydrides used as negative electrodes of Ni-MH batteries are also key components for battery performance.

In this thesis, the metallic composites TiMgNix, MgTi1-xNix and TiMg1-xNix were synthesized by mechanical alloying from elemental powder. The microstructure and phase transformation of prepared samples were characterized by XRD, SEM and TEM (with EDS support). The hydrogenation properties were measured by hydrogen solid-gas reaction and electrochemical tests.

Ré sumé :

Le stockage de l'hydrogè ne est l'un des plus grands problè mes techniques qui restreignent l'application pratique de l'hydrogè ne. Les hydrures mé talliques sont considé ré s comme la solution principale à ce problè me puisqu'ils peuvent absorber et dé sorber de faç on ré versible une grande quantité d'hydrogè ne sous une tempé rature et une pression modé ré es. Par ailleurs, les hydrures mé talliques utilisé s comme é lectrodes né gatives dans les accumulateurs Nickel-Mé tal Hydrure (Ni-MH) sont é galement les composants clé s des performances de ces derniers.

Dans cette thè se, les alliages mé talliques TiMgNix, MgTi1-xNix et TiMg1-xNix ont é té synthé tisé s par broyage mé canique à partir de poudres é lé mentaires. La microstructure et les transformations de phase des é chantillons pré paré s ont é té caracté risé es par DRX, MEB et MET (avec microanalyse EDS). Les proprié té s d'hydrogé nation ont é té mesurées par réaction d'hydrogène solide-gaz et par des essais é lectrochimiques. Un diagramme de composition-capacité 3D a é té é tabli sur la base du diagramme de phase ternaire Ti-Mg-Ni. Un procé dé de broyage en deux étapes a été mis en oeuvre pour améliorer les performances é lectrochimiques des alliages Ti-Mg-Ni.

De plus, les alliages TiNi1-xCux ont é té synthé tisé s par broyage mé canique et ensuite recuits. L 'influence de la substitution du nickel par le cuivre sur la structure et les proprié té s é lectrochimiques est é tudié e en utilisant une double approche: expé rimentale et par simulation.

Les ré sultats obtenus par la thé orie de la fonctionnelle de la densité (DFT) en utilisant le programme VASP montrent que l'enthalpie de formation et l'énergie d'adsorption de l'hydrogène de la phase pseudo-binaire Ti(Ni, Cu) sont en bon accord avec les ré sultats expé rimentaux.

Mots