Raphaël Janot

Benoît Fleutot

Matthieu Courty

Theo Kudu

Xavier Cédric

Viki Aman

Garima

Alexis Yinhue

Marianita

Nour, Carlito, Josesito, Kiril, Gaspard, Caro Vito Tu

Roberto

Arina, Sorina, Alexis, François, Tristan, Chabert, Clément Teo

LRCS, Amiens) for all the thermal analysis measurements and Jean-Noël Chotard (LRCS, Amiens) for teaching me how to analyze XRD patterns. I would like to thank also Yaroslav Filinchuk (Université catholique de Louvain) for his help on crystal structure resolution.

General introduction

In a world consuming more and more electricity (cf. Figure 1a), 1 rechargeable batteries are playing a key role. 2 They are used in many applications such as: electrical vehicles (EVs), mobile electronics and stationary energy storage. Rechargeable batteries have been mainly used until now for portable electronics but a growing demand is coming from other applications such as electrical vehicles (cf. Figure 1b). 3 EVs are expected to replace the majority of the vehicles using internal combustion engines in the coming years. The production of stationary batteries could also start to grow in the context of renewable energies. Indeed, renewable energies are intermittent and need to be stored to ensure a continuous electricity supply. b) Li-ion batteries sales per year for different applications. 1,3 Nowadays, the most common rechargeable batteries are based on Li-ion technology. This leadership is mainly due to their energy density that is the highest of all existing technologies (up to 200 Wh.kg -1 , cf. Figure 2). 4 Their large impact on the society has been rewarded with the Nobel Prize of Chemistry 2019 to John B. Goodenough, Stanley Whittingham and Akira Yoshino for the development of rechargeable Li-ion batteries. This technology is so versatile that its massive utilization starts to create problem of sustainability. Elements such as Co, which is used in the cathode material of most of the commercial Li-ion batteries, are rare and present in only few regions of the world. For these reasons, their price and availability are unstable. 5 Solutions are studied to substitute these elements by more abundant ones or to replace them totally by using post Li-ion technologies such as Li-S batteries. Beyond Li-based batteries, other technologies are studied. For instance, systems based on Na or K ions, two elements which are much more abundant than lithium, are developed. Then, systems based on multivalent Mg 2+ , Ca 2+ , Zn 2+ or Al 3+ ions are studied to try to increase the energy density of the batteries. 6 Among these multivalent ions, Mg 2+ has been the one attracting the more interest during the last two decades. Rechargeable Mg batteries have been studied since the early 90's but it is only in 2000 that a functional cell has been assembled by Aurbach et al.. Nevertheless, this first cell could not deliver [9] large energy amounts. Thus, important efforts have been put by the scientific community to design better cells but none of them can deliver high capacity and/or high voltage for many cycles.

The main problems lay in finding cathode materials allowing fast insertion of Mg ions and electrolytes with good ionic conductivities and wide electrochemical stability windows. To improve rechargeable Mg batteries, our approach was to study the possibility of using solid-state electrolytes. Indeed, they could offer benefits such as better chemical and electrochemical stabilities and better safety. The candidates as solid electrolyte materials were selected among the existing inorganic solid Mg 2+ ionic conductors (i.e. no mixed electronic/ionic conductors such as electrode materials). Few materials are reported as Mg 2+ is a strongly polarizing cation and, thus, it has a sluggish mobility in solids. Most of the inorganic solid state Mg 2+ ionic conductors are listed in Chapter I. Three families can be found: the phosphates, the sulphides and the complex hydrides. The phosphates, based on Mg-Zr-PO4, have high ionic conductivities but only at high temperatures, typically ~10 -3 S.cm -1 at 800°C. Then, the sulphides, based on MgS-P2S5 glasses or glass/ceramics, have too low ionic conductivities (~10 -7 S.cm -1 at 200°C) to be interesting as solid electrolytes. Finally, the complex hydrides, based on Mg(BH4)2, have the best ionic conductivities at low temperatures (~10 -5 -10 -6 S.cm -1 below 150°C). Furthermore, they are the only ones that have been characterized regarding possible applications in Mg batteries (e.g. Mg deposition/stripping, electrochemical stability, etc.). It was decided to start our study with Mg(BH4)(NH2) that was previously reported as a promising Mg 2+ solid ionic conductor. 7 Mg(BH4)(NH2) was described for the first time by Noritake et al. in 2013. 8,9 They synthesized it from the reaction of Mg(BH4)2 with Mg(NH2)2. Unfortunately, the synthesis parameters were not reported (ball-milling speed, annealing time or temperature, etc.).

The first part of this thesis is dedicated to the determination of the appropriate synthesis parameters needed to obtain Mg(BH4)(NH2) (cf. Chapter II). After several attempts with different synthesis parameters, crystalline Mg(BH4)(NH2) could be obtained. It appeared however by 11 B MAS-NMR spectroscopy that none of our samples was completely pure as they all contain an additional phase. Interestingly, this additional phase had an important effect on the ionic conductivity values.

The next part of the thesis was the exploration of the Mg(BH4)2-Mg(NH2)2 binary phases diagram (cf. Chapter III). A series of syntheses with different stoichiometries has been performed: it revealed the existence of three other crystalline phases. Two rich in amide that we labeled α and β-phase, and one rich in borohydride labeled the γ-phase. These phases were studied for the first time as no other reported work describes these compounds. Their crystal structures and ionic [10] conductivities were especially studied. It appeared that the ionic conductivity follows a trend as a function of the composition and that new compounds with a high Mg 2+ mobility can be obtained.

Once it was demonstrated that good inorganic solid Mg 2+ ionic conductors could be obtained, these materials had to be further characterized to assess their potential use as solid electrolytes.

The final part of this PhD thesis is dedicated to the characterization of the chemical and electrochemical stability of the newly discovered Mg 2+ ionic conductors (cf. Chapter IV). The

I-1. Batteries: short reminder of their principle and main characteristics

A battery is the assembly of one or several cells used to store electrical energy under chemical one through electrochemical reactions. Each cell is made of two electrodes separated by an electrolyte (cf. Figure I. 1). Batteries can be separated into two categories: the primary cells and the secondary cells. The primary cells are not rechargeable and secondary cells can be recharged.

If a cell is charged, connecting the two electrodes will allow the electrical current to pass from one electrode to the other and the battery will discharge. Once a cell is discharged, it can be charged by applying a reverse current. Redox reactions are occurring simultaneously at both electrodes: at the anode and at the cathode. In electrochemistry, the oxidation happens at the anode (negative electrode) while the reduction happens at the cathode (positive electrode). In batteries, it is commonly admitted to call "anode" the electrode with the lower redox potential and "cathode" the one with the higher redox potential. The cells are characterized by their working potential Vcell and their capacity Qcell.

The working potential Vcell is the difference between the potential of each electrode. The capacity Q is the quantity of charges that can be released or stored by the electrodes and it can be calculated by the Faraday's law: 𝑄 = 𝑛. 𝑛 𝑒 . 𝐹 Equation I. 1

with n: the number of exchange species (mol) ne-: the number of electrons exchanged F: the Faraday's constant (96500 C.mol -1 , and 1 C = 1 A.s) [16] Equation I.1 gives the capacity Q in Coulomb but the unit commonly used for batteries is the Ah.

The specific and volumetric capacity can be calculated modifying Equation I. 1. The specific capacity is: From Vcell and Qcell it is possible to determine the energy Ecell (usually in Wh) that can be delivered by the cell:

𝑄 𝑠𝑝𝑒 = 1
𝐸 𝑐𝑒𝑙𝑙 = 𝑉 𝑐𝑒𝑙𝑙 . 𝑄 𝑐𝑒𝑙𝑙 Equation I. 6 The power Pcell (in W) of the cell is calculated by dividing the energy by the time:

𝑃 𝑐𝑒𝑙𝑙 = 𝐸 𝑐𝑒𝑙𝑙 𝑡 Equation I. 7
with t: the time of charge/discharge

To ensure the charge balance of the redox reactions, it is necessary to have an electrolyte between the two electrodes. The electrolyte is an ionic conductive medium that allows ions to pass from an electrode to the other; whereas it is an electronic insulator avoiding short-circuits between the electrodes. It is characterized by different intrinsic properties:

Chapter I [17] -the transference number, -the viscosity, -the ionic conductivity, -the electrochemical stability window.

For solid electrolytes, the important parameters are the ionic conductivity and the electrochemical stability window. The transference number in inorganic solid electrolytes is usually 1 as only one ion is the charge carrier, the counter anions/cations are fixed in the structure and the electronic conductivity must be close to zero.

I-2. Rechargeable Mg batteries

I-2a. From Li-ion to post Li-ion batteries

Nowadays, the most commonly used rechargeable batteries are Li-ion batteries (LIBs). They have many advantages in terms of energy density, lifetime and can be used at different scales: portable electronics, vehicles and stationary energy storage systems. Nevertheless, the massive production of LIBs needed to meet the ever-growing demand in electricity has arisen new problematics such as the shortage of raw materials because of their low relative abundance and geopolitical issues. For instance, there is a mismatch between some of the most abundant element on earth which are the constituents of the biomass (H, O, S, P, Na, K, Ca, Mg, Mn, Fe, Ti) and main constituents of LIBs (for example, Li, Co, Ni, Cu, F). In 2017, the European Commission has classified Li and Co as critical raw materials for the European Union (EU) because of risks of supply shortage. 1 To anticipate this situation, researchers are looking for post Li-ion technologies based on more abundant elements like Na, K, Mg, Ca, Zn or Al. Among these new technologies Na-ion batteries (NIBs) are the most advanced one, achieving high voltage (>3 V) and good capacity retention. 2 Mg batteries are another option to replace LIBs. On the contrary to NIBs for which the knowledge transfer from LIBs helped their development, the rechargeable Mg batteries need specific electrolytes and electrode materials. The particular chemistry of Mg open the way for new advantages but also new challenges that will be described in the following subsections. 3 Chapter I -State-of-the-art: Rechargeable Mg batteries and Mg 2+ inorganic solid electrolytes [18]

I-2b. Magnesium metal anode

The intrinsic properties of magnesium present many advantages that make it an excellent candidate as anode material for batteries. First of all, it has a highly negative standard potential of -2.356 V vs. NHE close to the one of lithium (-3.040 V vs NHE, cf. Figure I. 2b). The second interesting feature of magnesium as anode material is its divalency as a cation. Its redox reaction involves two electrons and its theoretical gravimetric capacity is 2205 mAh.g -1 . This value is lower than the one of lithium (3829 mAh.g -1 ) but taking into account the higher density of magnesium (1.74 g.cm -3 ), its volumetric capacity overpasses the one of lithium (3833 mAh.cm -3 vs. 2046 mAh.cm -3 , cf. Chapter I [19] Magnesium has also interesting technical properties. It is solid at room temperature (Tmelting = 650°C). It can be manipulated in ambient atmosphere. It has a large abundance in the Earth crust of 21,100 ppm (vs. 10 ppm for Lithium) 4 but it is considered as a critical raw materials by the European Comission as Li and Co because of geopolitical issues. Actually, 94 % of Mg consumed in the EU comes from China. In terms of electrochemistry, its deposition has been considered for many years to be dendrites free. [5][6][7] In fact, many studies have proved that the electrodeposition of magnesium leads to the formation of smooth films (cf. Figure I. 3a). 8 Nevertheless, this consideration has been questioned and a recent paper from Davidson et al.

warns on the possibility of dendrites formation on magnesium electrodes 9 . They show that magnesium, as any other electrodeposited metals, can form dendrites under certain conditions.

They concluded that local inhomogeneities can result in an electrochemical reaction rate that overpasses self-diffusion rate and thus leads to the formation of dendrites (cf.

I-2c. Prototypes of rechargeable Mg batteries

Here, we describe how the rechargeable Mg batteries have evolved during the past 30 years and give a non-exhaustive list of examples of rechargeable prototypes. Thomas D. Gregory et al. were the first to consider the feasibility of rechargeable Mg batteries in 1990. 10 Trying different intercalation materials as cathode and different electrolytes, they assembled the first rechargeable Mg battery. Their cell was made of Mg for the anode and Co3O4 for the cathode with

Chapter I -State-of-the-art: Rechargeable Mg batteries and Mg 2+ inorganic solid electrolytes [20] an electrolyte made of 0. It has been demonstrated that other MXn electrodes (M = transition metal and X = S, Se) can allow reversible Mg intercalation. 13 Among them, TiS2 has been reported as an interesting insertion electrode for rechargeable magnesium batteries. Sun et al. showed that at 60°C and at C/10 it was possible to obtain a stable capacity of 115 mAh.g -1 (cf. Figure I. 6). 14 This capacity is lower than the one obtained for the first discharge (250 mAh.g -1 ). The authors explain this capacity

Chapter I -State-of-the-art: Rechargeable Mg batteries and Mg 2+ inorganic solid electrolytes [22] drop as a consequence of Mg 2+ entrapment in the crystalline structure owing to complex ordering and kinetic effects that were not studied in that paper. Furthermore, the working potential is quite low (~1 V vs. Mg 2+ /Mg), so TiS2 as cathode material is limited in terms of energy density. Still, it is an interesting material to study the Mg diffusion in solid phases. This can provide guidance in the search for new Mg positive electrode materials. It can also be considered as a standard intercalation material for evaluating new electrolytes. There is a need to find high capacity cathode materials for magnesium batteries to achieve high energy densities. A promising cathode for rechargeable Mg batteries is sulphur. Using S as the cathode allows to reach very high capacities and to develop batteries with high energy density.

With a theoretical voltage of 1.77 V vs Mg 2+ /Mg, Mg-S batteries would surpass the volumetric

Chapter I [23] energy of Li-S batteries, 3221 Wh.L -1 and 2856 Wh.L -1 respectivley. 15 Zhao-Karger et al. did a comprehensive study of Mg-S batteries. They used an electrolyte based on Mg-HMDS (HMDS = N(SiMe3)2 -= hexamethyldisilazide) and have obtained a high capacity of 800 mAh.gS -1 during the first discharge. 16 They could also demonstrated that the reaction pathway of Mg with S is similar to the one of Li-S batteries forming different MgSx species (8 > x > 1, cf. Figure I. 7). More works have been done to improve Mg-S batteries but, as Li-S batteries, suffer from fast capacity fading due to side reactions as the dissolution of polysulphides in the electrolyte. [START_REF] Zhao-Karger | Beyond intercalation chemistry for rechargeable Mg batteries: A short review and perspective[END_REF] CuS has also been reported as high capacity cathode (560 mAh.g -1 in theory). This conversion electrode suffers however from poor cyclability because of important volume changes during the magnesiation and demagnesiation. , in tetraglyme (denoted as TEG in red) using CMC or PVDF as binder. 16 To reach high energy densities, the working potential of cathode materials for Mg batteries has to be increased. In this way, an interesting material as cathode is V2O5 because of its large interstitial vacancies able to accommodate Mg cations and its working potential of about 2.4 V vs Mg 2+ /Mg. However, the electrolytes stable at this potential are scarce and generally very corrosive, thus, leading to corrosion issues with the steal of the cells and aluminum of the current collectors. To study V2O5, Gershinsky et al. used an electrolyte made of 0.1 M of MgTFSI2 in acetonitrile but Mg metal cannot be used with this electrolyte as a passivation layer would be [24] formed at its surface. To solve this problem, they used activated carbon as the negative electrode on which Mg ions are reversibly adsorbed and desorbed. They could cycle a thin film of V2O5 and get a relatively stable capacity of 150 mAh g -1 with a good coulombic efficiency (cf. Figure I. 8b). [START_REF] Dolci | Magnesium imide: Synthesis and structure determination of an unconventional alkaline earth imide from decomposition of magnesium amide[END_REF] Fu et al. used a similar approach to cycle V2O5 nano-tubes, but instead of activated carbon as the negative electrode, they used MgxMo6S8. They did in-situ measurements on their V2O5 electrode and observed the reversible magnesiation/demagnesiation. Nevertheless, their electrode had a poor cyclability with an important capacity fading. 20 Figure I. 8. a) Discharge and charge profile of V2O5 thin film b) its capacity and coulombic efficiency upon cycling. [START_REF] Dolci | Magnesium imide: Synthesis and structure determination of an unconventional alkaline earth imide from decomposition of magnesium amide[END_REF] These studies show that electrodes enabling reversible magnesiation/demagnesiation exist.

They also show that for some materials the downsizing is important to promote the magnesium diffusion. Nevertheless, there are few systems with high voltage, high capacity and high cycling capability. The next section describes the main limitation of these systems focusing on the problems of the electrolytes.

I-2d. Limitations of rechargeable Mg batteries

In the paper of Gregory et al. of 1990 cited previously, it was said: "Over the past 25 years, a considerable amount of attention has been paid to the use of alkali metals, particularly Li, in

Chapter I [25] non-aqueous secondary batteries, but comparatively little work has been done with the alkaline earth metals". 10 At that time, the first commercial rechargeable Li-ion battery was arriving on the market. In the case of rechargeable Mg batteries, after more than 25 years of work on this topic, no commercial device exists. There are two main reasons explaining why the scientific community is still struggling to find a viable system for rechargeable Mg batteries. The first one is the difficulty to find electrodes enabling reversible magnesiation/demagnesiation with good performances. This is due to the sluggish diffusion of Mg 2+ into solids: the double charge of the cation increases the electrostatic interaction with the surrounding ions. The other reason is the difficulty to find a good electrolyte for Mg-based systems. 13 Contrarily to Li-ion batteries, there is no standard liquid electrolyte working for Mg batteries yet. For instance, the electrolytes using carbonate-based solvents, which are the best for LIBs, cannot be used with the Mg metal anode. The carbonates are reduced at the surface of the Mg metal forming the well-known solid electrolyte interface (SEI). This interface is, in the Mg case, very detrimental because it blocks the Mg ions migration. 6 To replace carbonate-based electrolytes, liquid electrolytes compatible with Mg metal were needed. Grignard reagents-based electrolytes are stable versus Mg metal and allow a very good Mg deposition/stripping. They are known since the beginning of 20 th century as electrolyte for Mg electrodeposition. It is only in the 90's that they started to be used as electrolytes for Mg batteries. 21 The best ones have a good oxidative stability: typically 3.7 V vs Mg 2+ /Mg using a Pt working electrode for (Mg2(μCl)3•6THF)(B(C6F5)3Ph). 22 Nevertheless, they are very air sensitive and strongly corrosive because of the presence of chlorine. Indeed, they cannot be used at potentials higher than 2.2 V vs. Mg 2+ /Mg if the cell contains stainless steel as a consequence of corrosion. stripping with an excellent coulombic efficiency (> 99%). It has also a high electrochemical stability (~3.8 V vs. Mg 2+ /Mg using a stainless steel working electrode). Nevertheless, these salts have a low solubility so the conduction properties of these liquid electrolytes are still limited. 24 More recent studies from Zhao-Karger et al. show the good properties of Cl-free electrolytes using Mg(B(hfip)4)2 (hfip = hexafluoroisopropyloxy) dissolved in DME. 25,26 These electrolytes have high ionic conductivities (~10 -3 mS.cm -1 ). They also have a large electrochemical stability window of 3.5 [26] V vs. Mg 2+ /Mg (using a Pt working electrode). They are used in Mg-S batteries for which the dissolution of polysulphides in the electrolyte remains problematic.

Consequently, there is a real need of new electrolytes for rechargeable Mg batteries. These new electrolytes should:

-be compatible with Mg anode and cathode materials, -be non-corrosive,

-have good transport properties, -have a large electrochemical stability window.

Beside the liquid electrolytes, some groups have been working on gels, polymers, ionic liquid electrolytes or MOF-based electrolytes (MOF = metal-organic framework) 27 but these works are out of this thesis scope. Our approach was to look for inorganic solid electrolytes and to know if they can fulfill the criteria described above.

I-3. Interest for solid electrolytes

To improve existing rechargeable magnesium batteries, it would be interesting to use a solid electrolyte. Two main possibilities exist to use solid electrolytes in batteries:

-as bulk electrolyte in an all-solid-state battery (cf. Figure I. 9a and b) -as coating for the electrode materials (cf. figure 9b andc)

The all-solid-state battery is an old concept developed at the beginning of the 20 th century with the work of Tubandt and Lorentz on α-AgI as solid-state electrolyte. The interest for solid electrolytes grew in the 60's with the discovery of the superionic conduction of Na + in -Al2O3. This solid electrolyte enabled the development of efficient Na-S batteries working between 270°C and 350°C. Nowadays, a new hype is growing around all-solid-state batteries as many private and academic groups are focused on the discovery of functional solid electrolytes at room temperature. This interest comes from the advantages that all-solid-state batteries could offer. They are expected to be safer as no liquid electrolyte is used, avoiding problems of thermal decomposition that can lead to the release of toxic and flammable gases. Then, without liquid electrolyte, an increase of the range of workable temperatures is expected. The assembly of an all-solid-state battery could also be simpler than the one of a battery using a liquid electrolyte. Instead of assembling several cells, an all-solid-state battery could be made simply by successive stacks of all-solid-state cells as the one on Figure I. 9a or b. 28 We have to bear in mind that solid electrolytes can be used as bulk electrolytes but also as coating materials for electrodes. Actually, some solid electrolytes have a poor ionic conductivity but have really good electrochemical and chemical stabilities. These stable materials can be used

Chapter I [27] as interface between the electrodes and the electrolyte. Of course, their thickness has to be minimized (typically few nanometers) to ensure a good ionic diffusion. The coating of electrodes with solid electrolytes can be done with a solution of dissolved solid electrolyte in which the electrode is dipped before subsequent evaporation of the solvent and recrystallisation of the solid electrolyte. 29,30 More sophisticated techniques can also be used as physical vapor deposition. [START_REF] Takahashi | All-solid-state lithium battery with LiBH4 solid electrolyte[END_REF] The following section reviews the main families of inorganic solid state Mg 2+ conductors known at the time of this PhD thesis.

I-4. Mg 2+ inorganic solid electrolytes

I-4a. Phosphates

The first solid Mg 2+ inorganic electrolytes were reported in 1987. That year, the team of Kaname Ito from the Nagoya Technological Institute was the first one to propose the use of solid electrolytes in Mg batteries. [START_REF] Ikeda | Solid electrolytes with multivalent cation conduction. 1. Conducting species in Mg-Zr-PO4 system[END_REF] As explained in their work, the research in monovalent solid ionic conductors was progressing fast but in the meantime almost nothing was done on multivalent ion

Chapter I -State-of-the-art: Rechargeable Mg batteries and Mg 2+ inorganic solid electrolytes [28] conductors. They foresaw that the discovery of Mg 2+ solid electrolytes would "not only be academically interesting but also be highly expected to compose higher energy density electrochemical devices using such materials". They studied different NASICON-like (NASICON = Na super ionic conductor) Mg 2+ conductors in the system Mg-Zr-PO4. A standard ceramic synthesis route at high temperature (1200°C) was used to synthesize materials with different Mg:Zr:PO4 molar ratio. Then, they made pellets from the as-obtained materials by sintering at 1200°C for 24h. It is known that oxides are hard to compact and in their study, the pellets have an important porosity, even after sintering (about 20-30%). The pellets are used to measure the ionic conductivity by AC (alternative current) impedance spectroscopy. The ionic conductivity of these materials follows an Arrhenius law (cf. MgZr4(PO4)6 was synthesized via a sol-gel method. 34 They could obtain pellets with better compactness after sintering at 1300°C for 24 h (99 % of the theoretical density). This improvement of compactness leads to a slight increase of the conductivity value reaching 7.23x10 NMR spectroscopy. They proved that, in these materials, the hopping of Mg 2+ from one site to another was related to a large deformation of the coordinating PO4 tetrahedra. Thus, they could link the Ea determined from the relaxation time measured by [START_REF] Takahashi | All-solid-state lithium battery with LiBH4 solid electrolyte[END_REF] P NMR spectroscopy to the Ea determined by impedance spectroscopy and confirmed the fast motion of Mg 2+ . [START_REF] Kawamura | High temperature 31 P NMR study on Mg 2+ ion conductors[END_REF] So, MgZr4(PO4)6based materials allow to reach high ionic conductivities (~10 -3 S.cm -1 ) but only at high temperatures (~800°C). The only application reported for these solid electrolytes is as Mg 2+

conductor in Cl2 gas sensors. [START_REF] Imanaka | A chlorine gas sensor based on the combination of Mg 2+ cation conducting and O 2anion conducting solid electrolytes with lanthanum oxychloride as an auxiliary electrode[END_REF] The very high operating temperatures make the application of these materials in solid-state batteries difficult.

It has been demonstrated recently by Su et al. that magnesium phosphates, MgxPyOz thin films (~100 nm) can be deposited by ALD. [START_REF] Su | Atomic layer deposition of a magnesium phosphate solid electrolyte[END_REF] These films are glasses and can have different networks depending on the deposition parameters. Such films have really low ionic conductivities, 1.6x10 -7 S.cm -1 at 500°C and an activation energy of 1.37 eV. The ionic conductivity should be improved for use in batteries. The authors proposed that further work on the nitridation of the films could be done to increase the ionic conductivity, equivalently to what has already be done with LiPON solid electrolytes. [START_REF] Kozen | Atomic layer deposition of the solid electrolyte LiPON[END_REF] We can give the example of a patent from Panasonic describing nitridated magnesium silicate glasses, MgxSiOyNz, thin films (< 20 µm) but their ionic conductivities are also low, around 1x10 -8 S.cm -1 at 800°C. [START_REF] Fujinoki | Solid electrolyte having magnesium ion conductivity and magnesium secondary battery using the same[END_REF] [30]

I-4b. Sulphides

The group of Masahiro Tatsumisago from Osaka Prefecture Univeristy has a great experience in the synthesis of sulphide-based glasses (e.g. Li2S-P2S5 and Na2S-P2S5) and their application as solid electrolytes. [START_REF] Hayashi | Development of sulfide solid electrolytes and interface formation processes for bulk-type all-solid-state Li and Na batteries[END_REF] They discovered that after annealing of these glasses, they could obtain glass-ceramic materials. These materials are interesting solid electrolytes as the coexistence of a crystalline and an amorphous phase results in superionic conductivity at room temperature. [START_REF] Mizuno | highly ion-conductive crystals precipitated from Li2S-P2S5 glasses[END_REF][START_REF] Tsukasaki | Direct observation of a non-crystalline state of Li2S-P2S5 solid electrolytes[END_REF] In 2014, they studied new glasses based on MgS-P2S5. [START_REF] Yamanaka | Preparation of magnesium ion conducting MgS-P2S5-MgI2 glasses by a mechanochemical technique[END_REF] They synthesized first the mother glass 60MgS-40P2S5. After annealing of the glass at 450°C for 1h, they could observe some diffraction peaks indexed as Mg2P2S6 crystallizing in the space group C2/m (cf. Figure I. 11a and b). They confirmed this result by Raman spectroscopy as only the vibration bands of the P2S6 4- groups remained after annealing. Hence, they claimed that they synthesized a glass-ceramic material. The conductivity at 200°C of the 60MgS-40P2S5 solid electrolyte jumps from 2x10 -10 S.cm -1 for the glass to 1x10 -7 S.cm -1 for the glass-ceramic (cf. Figure I. 11c). They also added MgI2 to the material with different (100-x)(60MgS-40P2S5)-xMgI2 compositions to increase the number of charge carriers. For various compositions, new unexplained weak peaks appeared by X-ray diffraction while the ones of Mg2P2S6 remained. The effect of MgI2 on the ionic conductivity is Chapter I [31] negligible for the glass-ceramic but clearly improves the one of the glass (i.e. by two orders of magnitude, cf. Figure I. 11c). For the glass-ceramic samples, the ionic conductivity at 200°C increases from 1x10 -7 S.cm -1 to 2x10 -7 S.cm -1 for x=0 and x=20, respectively. The ionic conductivity of these sulphides is too low and needs to be improved for solid electrolyte application.

I-4c. Complex hydrides

Another class of materials has been studied as solid state electrolyte for Mg batteries: the complex hydrides. This thesis is exclusively based on the study of complex hydrides as Mg 2+ ionic conductors. Hence, extra details will be given in this subsection to understand better why these materials are interesting as solid electrolytes.

In 2006, the group of Shinichi Orimo at the Tohoku University in Japan revealed for the first time the possibility to use complex hydrides as solid state ionic conductor. While studying LiBH4 as hydrogen storage material, they measured how the material behaves when it is heated up by [32] microwaves. [START_REF] Nakamori | Dehydriding reaction of metal hydrides and alkali borohydrides enhanced by microwave irradiation[END_REF] They observed that the behavior of LiBH4 is similar to other complex hydrides at low temperature but drastically change around 110°C (cf. Figure I. 12a). This is due to the phase transition that undergoes LiBH4 with a reorganization of the (BH4) -anions in the structure (cf.

Figure I. 12c

). The crystal structure changes from a low temperature (LT) orthorhombic phase (space group: Pnma) to a high temperature (HT) hexagonal phase (space group: P63mc). They actually discovered that this structural transition affects the electrical properties of the material.

Performing dielectric constant measurements, they could determine that LiBH4 was passing from an insulator to a conductor behavior (cf. Figure I. 12b). They predicted that the charge carrier would be Li + cation as B and H are strongly covalently bonded in the BH4 units. They measured the ionic conductivity by AC impedance spectroscopy and as predicted, the conductivity exceptionally increases with the structural transition at 110°C from ~10 -7 S.cm -1 to ~10 -3 S.cm -1 (cf. Figure I.

12d). [START_REF] Matsuo | Lithium superionic conduction in lithium borohydride accompanied by structural transition[END_REF] Thus, they demonstrated the superionic conductivity of the hexagonal HT-LiBH4 phase. The activation energies were found by the Arrhenius law to be 0.69 eV and 0.53 eV for the orthorhombic LT and hexagonal HT LiBH4 phases, respectively. They confirmed that Li + was the mobile specie by 7 Li solid state MAS-NMR spectroscopy. A great narrowing of the NMR peak can be observed after the phase transition from LT-to HT-phase, indicating the increased Li + mobility (cf. Figure I. 12e). Then, they measured the spin-lattice relaxation time (T1) as a function of the temperature from which they could extract an activation energy of 0.56 eV which is in agreement with the one determined by impedance spectroscopy (0.53 eV).

Chapter I

[33] 7 Li solid state NMR spectra of LiBH4 at different temperatures. [START_REF] Matsuo | Lithium superionic conduction in lithium borohydride accompanied by structural transition[END_REF] Later, the electrochemical stability of LiBH4 was determined by cyclic voltammetry done on HT-LiBH4 at 120°C. [START_REF] Matsuo | Lithium fast-ionic conduction in complex hydrides: Review and prospects[END_REF] A pellet was sandwiched between a lithium and a molybdenum electrode and the potential scanned at a speed of 5 mV.s -1 . Only the lithium deposition/stripping (around 0 V vs Li + /Li) and no oxidation of the electrolyte were observed (cf. Figure I. 13a). The electrochemical stability window of LiBH4 was estimated between 0 and 5 V vs Li + /Li which is really interesting for all-solid-state batteries. Nowadays, we know that the characterization of solid electrolyte with planar counter electrodes (here molybdenum) can hinder the decomposition process. 51 This electrochemical stability window was overestimated. Recent studies have demonstrated that the actual stability of LiBH4 is about 2.1 V vs Li + /Li. [START_REF] Gulino | Phase stability and fast ion conductivity in the hexagonal LiBH4-LiBr-LiCl solid solution[END_REF][START_REF] Asakura | Electrochemical oxidative stability of hydroborate-based solid-state electrolytes[END_REF] Nevertheless, in 2014, Takahashi et al.

assembled the first all-solid-state Li battery using LiBH4 as solid electrolyte and working at 120°C. [START_REF] Takahashi | All-solid-state lithium battery with LiBH4 solid electrolyte[END_REF] They used Li as anode and LiCoO2 as cathode. When bare LiCoO2 was used an important interface resistance was observed after cycling. By Raman spectroscopy, they showed that this interfacial resistance was due to the reaction of LiBH4 with charged Li1-xCoO2. Species as Co3O4 and CoO(OH)

were formed. To prevent the reaction of the borohydride with the high potential cathode, a LiCoO2 cathode thin-film coated with Li3PO4 solid electrolyte was prepared (cf. Figure I. 13b). Li3PO4 is used here as a protective layer. This strategy worked as no degradation of the electrode was observed after cycling and the interfacial resistance with the electrolyte remained stable (cf.

Figure I. 13c

). This first all-solid-state battery using a borohydride showed the feasibility of such systems but also shed the light on the reactivity of these materials with high potential cathodes.

Another study made by Unemoto et al. was done on a full-cell using a low potential cathode material, TiS2 (cf. Figure I. 13d). [START_REF] Unemoto | Stable interface formation between TiS2 and LiBH4 in bulk-type all-solidstate lithium batteries[END_REF] This time, no coating of the cathode was necessary to achieve good capacity retention. They observed that a stable interface was formed between TiS2 and LiBH4. This interface could be made of closo-borane as Li2B12H12 and appeared to be stable, ensuring a good capacity retention after 300 cycles (cf. ). These mechanical characteristics are of major interest for the assembling of all solid-state batteries. The low density of borohydride (< 1 g.cm -3 ) is also an advantage as, for the same volume, these solid electrolytes would represent less mass than other solid electrolytes (e.g. sulphides or oxides) so the energy density of the cell would increase.

Other all-solid-state batteries have been assembled using LiBH4 as solid electrolyte [START_REF] Unemoto | Complex hydrides for electrochemical energy storage[END_REF][START_REF] Guzik | Lightweight complex metal hydrides for Li-, Na-, and Mg-based batteries[END_REF] but an important issue remained: the high conductivity of LiBH4 can only be reached at elevated temperatures (> 120°C). Different approaches were developed to obtain good conductivity at room temperature with LiBH4-based materials. Among these approaches, the mixture of the lithium borohydride with lithium halides can be cited. It has been proved that a mixture with LiX (X = Cl, Br and I) can form a solid solution where HT-LiBH4 is stabilized at lower temperatures. [START_REF] Maekawa | Halide-stabilized LiBH4, a room-temperature lithium fast-ion conductor[END_REF] The best mixture with LiI resulted in superionic conductivity at room temperature (cf Figure I. 14a).

The HT-LiBH4 polymorph can also be stabilized by nano-confinement into mesoporous silica scaffolds (cf. (black). [START_REF] De Jongh | Complex hydrides as room-temperature solid electrolytes for rechargeable batteries[END_REF] Another approach was to use compounds from the LiBH4-LiNH2 binary system that have already been studied as hydrogen storage materials. LiNH2 was also known as a fast Li + ionic conductor since 1979. [START_REF] Boukamp | Ionic conductivity in lithium imide[END_REF] Its ionic conductivity is 3x10 -4 S.cm -1 at 25°C but its poor electrochemical stability (0.7 V vs Li + /Li) does not allow a practical use as solid electrolyte. Two crystalline defined compounds have been reported in the LiBH4-LiNH2 binary phases diagram: Li2(BH4)(NH2) (space group: R-3) and Li4(BH4)(NH2)3 (space group: I213) (cf. Figure I. 15a). 61 Li2(BH4)(NH2) has an ionic conductivity of 1×10 -4 S.cm -1 at room temperature with Ea = 0.56 eV. Li4(BH4)(NH2)3 has a conductivity of about 2×10 -4 S.cm -1 at room temperature and its activation energy is two times lower (0.26 eV). [START_REF] Matsuo | Complex hydrides with (BH4)and (NH2)anions as new lithium fast-ion conductors[END_REF] These materials melt at elevated temperatures which induces a decrease of the activation energy and an increase of the ionic conductivity (cf. Figure I. 15b). Their conductivities

Chapter I [37] reach 6×10 -2 S.cm -1 at 105°C for Li2(BH4)(NH2) and 2×10 -1 S.cm -1 at 240°C for Li4(BH4)(NH2)3. Later, it was also shown that the ionic conductivity of Li4(BH4)(NH2)3 can be enhanced by changing the stoichiometry. With a 2:3 Li(BH4):Li(NH2) ratio, the conductivity reaches 6.4 × 10 -3 S cm -1 at 40°C. [START_REF] Yan | A lithium amide-borohydride solid-state electrolyte with lithium-ion conductivities comparable to liquid electrolytes[END_REF] Yan et al. even demonstrated that this material could be used as solid electrolyte. A cell with Li as the negative electrode and LTO (LTO = Li4Ti5O12) as the positive was assembled and showed excellent rate and stable cycling capability (140 mAh.g -1 for 400 cycles at 1C at 40°C). An analogous strategy was used to find solid-state Na + ionic conductors. The NaBH4-NaNH2-NaI ternary system was explored (cf. Figure I. 16a). [START_REF] Matsuo | Sodium ionic conduction in complex hydrides with [BH4]and [NH2]anions[END_REF] On the NaBH4-NaI side, a solid solution is formed with different (1-x)NaBH4-xNaI compositions. This solid solution has a slightly better ionic conductivity than the raw materials. Then, on the NaNH2-NaI side, a defined compound is obtained with the composition Na(NH2)0.5I0.5 with also an improved ionic conductivity. Finally, on the NaBH4-NaNH2 side, the compound Na(BH4)0.5(NH2)0.5 was synthesized. This compound has the best ionic conductivity from all the phases described previously: 2x10 -6 S.cm -1 at room temperature. The better ionic conductivity is assumed to come from its anti-perovskite structure where Na ions can migrate through vacancies. This material is known to have two crystalline allotropes with a structural transition at 100°C. In this study, they declared to observe only the high temperature one. They have also characterized the electrochemical stability of Na2(BH4)0.5(NH2)0.5 by sandwiching this compound between Na and Mo electrodes and applying a scan rate of 1 mV.s -1 (cf. Chapter I [39] Now, coming back to the case of Mg 2+ ionic conductors, a material inspired from the discoveries introduced previously was studied as solid electrolyte for Mg batteries. This material is Mg(BH4)(NH2). As the above complex hydrides, it was first synthesized and studied as hydrogen storage material. [START_REF] Noritake | Dehydrogenation properties and crystal structure analysis of Mg(BH4)(NH2)[END_REF] Its structure was solved by Noritake et al. in a tetragonal crystal system (space group: I41) with a= 5.814(1) Å, c = 20.450(4) Å, V = 691.3(2) Å and Z = 8. The material has a low density of 0.9966 g.cm -3 . [START_REF] Noritake | Synthesis and crystal structure analysis of complex hydride Mg(BH4)(NH2)[END_REF] This low density is very interesting for a solid-state electrolyte as it allows obtaining good ionic percolation in composite electrodes with a relatively low weight of electrolyte, ensuring reasonable electrochemical gravimetric capacities for the composite electrodes. In 2014, a team from the Toyota research center in Japan studied Mg(BH4)(NH2) as solid electrolyte. [START_REF] Higashi | A novel inorganic solid state ion conductor for rechargeable Mg batteries[END_REF] A band gap of 4.9 eV was determined by first principle calculations so this material is an electronic insulator. It was also calculated that the bonds between Mg and the anions are ionic while the B-H and N-H bonds are covalent. From this first theoretical work, it appeared that this material could be an ionic conductor. They measured an ionic conductivity of 1x10 -6 S.cm -1 at 150°C by AC impedance spectroscopy and an activation energy of 1.31 eV (cf.

Figure I. 17a

). This conductivity was the highest ever reported for a Mg 2+ inorganic ionic conductor at low temperatures so far. It is three orders of magnitude higher than one measured for Mg(BH4)2 (~10 -9 S.cm -1 at 150°C). This improved ionic conductivity might be due to the shortest distance Unfortunately, they were unable to cycle these cells and only observed a potential decay while applying current. Even so, the stack was done well enough to have the expected OCV for each cathode tried. It was the most advanced work reported so far on all-solid-state magnesium batteries.

The last ionic conductor mentioned in this state of the art chapter is not a fully inorganic material but rather a hybrid material combining Mg(BH4)2 with organic molecules. Indeed, Roedern et al. discovered that it was possible to coordinate ethylene diamine molecules with Mg(BH4)2. [START_REF] Roedern | Magnesium ethylenediamine borohydride as solid-state electrolyte for magnesium batteries[END_REF] They studied two different mixtures prepared by ball-milling: Mg(en)3(BH4)2 and Mg(en)1(BH4)2 (en = ethylene diamine).The best conductivity was obtained for Mg(en)1(BH4)2 with 6 × 10 -5 Scm -1 at 70°C (cf. Figure I. 18a). It is a crystalline compound but the structure remains [40] unknown yet. Infrared and Raman spectroscopy measurements unveil that the chelating cis configuration is favored. This material as a thermal stability up to 75°C before to undergo an endothermic phase transition then it starts to decompose above 100°C. They proved that reversible Mg deposition/stripping was possible at 60°C by cyclic voltammetry with this material as solid electrolyte (cf.

Figure I. 17. a) Thermal evolution of the ionic conductivity of Mg(BH4)2 and Mg(BH4)(NH2) compared to MgZr4(PO4)6 and Mg1.4Zr4P6O24.4+0.4Zr2O(PO4)2. b) Cyclic voltammogram measured at 150°C of Mg(BH4)(NH2) sandwiched between magnesium and platinum electrodes. c) Discharge curves at 150°C of Mg-S, Mg-FeS and Mg-Ag2S cells using Mg(BH4)(NH2) as solid electrolyte. Inset shows the photo of a cell. 67

Chapter I

[41]

I-5. Aims of the PhD thesis

Reviewing the state-of-the-art of rechargeable Mg batteries shows that some systems are very interesting in terms of energy density but suffer from limitations. One of the main limitations is the difficulty to find viable liquid electrolytes. Hence, our approach is to study the possibility of using a solid-electrolyte in Mg rechargeable batteries and to determine if it would be beneficial or not. The second part of the state-of-the-art chapter was on inorganic solid-state Mg 2+ ionic conductors and it appears that few of them exist. This can be explained by the important electrostatic interaction of the double positive charge of Mg 2+ with its surrounding atoms making its diffusion in solids very sluggish. Each of the three families of materials reported as solid electrolytes has different range of ionic conductivities. The phosphates have high ionic conductivities but only at high temperatures, typically ~10 -3 S.cm -1 at 800°C. This temperature is too high for using these materials in practical energy storage devices. Then, sulphides have too low ionic conductivities (~10 -7 S.cm -1 at 200°C) to be interesting as solid electrolytes. The third and most recent family of materials is the one of complex hydrides. These materials have the best ionic conductivities at low temperatures (~10 -5 -10 -6 S.cm -1 below 150°C). Furthermore, they are the only ones that have been really characterized regarding possible applications in Mg batteries (e.g. Mg deposition/stripping, electrochemical stability, etc.). To assemble a magnesium all-solid-state battery, it is paramount to have a good solid electrolyte in terms of ionic conductivity, electrochemical stability and ease of implementation (i.e. good electrode/electrolyte interface, chemical stability with the electrode components, etc.).

Nevertheless, to be used as a solid electrolyte in batteries, an ionic conductor must have an ionic conductivity higher than 10 -4 S.cm -1 as close as possible to room temperature. Diffusion of Mg 2+ is sluggish in solids at low temperatures so we first target materials having good ionic conductivities below 200°C. Few Mg 2+ inorganic solid electrolytes exist and none of them present an ionic conductivity reaching our target. So the goals of this PhD thesis was: Chapter I [43] To achieve these goals, the most promising starting material was Mg(BH4)(NH2). Indeed, it was the best Mg 2+ inorganic conductor reported in the literature at the beginning of this thesis and it was demonstrated that it allows reversible Mg deposition/stripping at 150°C. Then Mg(BH4)(NH2) [50] Chapter II

[51]

II-1. Introduction

My PhD work started with the study of Mg(BH4)(NH2) that was previously reported as a promising Mg 2+ solid ionic conductor. 1 It has been described for the first time by Noritake et al. in 2013. 2,3 They synthesized this compound in the scope of finding new solid state hydrogen storage materials. Through a series of synchrotron X-ray diffraction measurements at different temperatures, they could obtain this phase pure and solve its crystal structure. The synthesis was done from the reaction of Mg(BH4)2 with Mg(NH2)2 (cf. Equation II. 1):

Mg(BH4)2 + Mg(NH2)2 → 2 Mg(BH4)(NH2) Equation II. 1
Magnesium borohydride has been extensively studied as hydrogen storage material because of its high gravimetric hydrogen density (14.5 wt%) and its capability of dehydrogenation at moderate temperatures (250-400°C). 4 It can also be used as salt in liquid electrolytes for Mg batteries, 5 as described in chapter I, but its ionic conductivity as solid electrolyte is very low (~10 -12 S.cm -1 at 70°C and ~10 -9 S.cm -1 at 150°C). 1,6 Mg(BH4)2 is known to have several crystal structures which will be described in the following section.

Magnesium amide is also mainly known in the literature in the field of hydrogen storage materials, especially in the mixture Mg(NH2)2-2LiH, which can reversibly store 5.6 wt.% of hydrogen with suitable thermodynamics according to the following reaction: [7][8][9] Mg(NH2)2 + 2 LiH  Li2Mg(NH)2 + 2 H2 Equation II. 2

This chapter will depict the precursors used for the synthesis of Mg(BH4)(NH2) i.e.

commercial Mg(BH4)2 and homemade Mg(NH2)2. Then, it will describe the optimization of the synthesis parameters of Mg(BH4)(NH2) and how its purity influences its ionic conductivity.

II-2. Magnesium borohydride: Mg(BH4)2 II-2a. Polymorphs of Mg(BH4)2

The crystalline polymorphs of Mg(BH4)2 that are described the most in the literature are the α-, β-, γ-, and δ-phase. Their structures and cell parameters are presented in Figure II. 1. In the literature, the existence of other phases is also discussed (e.g. amorphous phase, β', ε, etc.). 4 The α-Mg(BH4)2 polymorph is stable at room temperature and undergoes a phase transition at about 220°C forming the β-phase. 10 After cooling, the structure of β-Mg(BH4)2 remains, it is considered Mg(BH4)(NH2) [52] as a metastable phase at room temperature. γ-Mg(BH4)2 has an extremely porous structure and a very low density (0.5431 g.cm -3 ). The pores of its structure can accommodate gas molecules such as H2 or N2. 11 It has two phase transitions: around 135-165°C to the ε-phase and around 165-215°C

to the β'-phase. 12 Finally, δ-Mg(BH4)2 is obtained under high pressures (> 2.1 GPa) from γ-Mg(BH4)2 and it is stable at room temperature. 11

II-2b. Commercially available Mg(BH4)2 powders

The first magnesium borohydride used for the synthesis of Mg(BH4)(NH2) was the commercial γ-Mg(BH4)2 (> 95 %, Sigma-Aldrich). X-ray diffraction (XRD) measurement and Le Bail profile matching indicate that it is pure crystalline γ-Mg(BH4)2 (cf. Chapter II [55] Infrared spectroscopy was also used to characterize these magnesium borohydrides. Only the characteristic vibrations bands of Mg(BH4)2 appear on the infrared spectra of both products:

the stretching bands around 2300 cm -1 and the bending bands around 1200 cm -1 (cf. Figure II. 4).

No -OH impurities are visible (no bands around 3500 cm -1 ). 14 The differential scanning calorimetry (DSC) measurement shows the phase transition that undergoes γ-Mg(BH4)2 as described previously (cf. Figure II. 5a). 12 At 146°C, it starts to transform into the ε-phase and at 185°C into the β'-phase. For the product made of γ and α-Mg(BH4)2 the result is different as the thermal phenomena linked to the two different phases are observed 10,12 :

γ → ε → β' at 146°C and 185°C then α → β around 219°C (cf.

II-3. Magnesium amide: Mg(NH2)2

Mg(NH2)2 is not commercially available and had to be synthesized during this thesis. This compound has been obtained from the reaction of magnesium hydride with gaseous ammonia as described in other studies. 9,15 The synthesis process is described on Figure II. 6 and the reaction is written in Equation II. 3: Commercial MgH2 (> 98%, Alfa Aesar), is first ball-milled to increase its reactivity. In an argon-filled glovebox, 2 g of MgH2 is introduced in a 50 cm 3 stainless steal jar with a balls to powder weight ratio of 53:1 and then ball-milled at 600 rpm using a Retsch PM100 miller for 2 h with a 2 min break every 30 min. By X-ray diffraction, MgH2 appears well amorphized as all its reflections are much weaker and broader after ball-milling (cf. Figure II. 7b). A color change of the powder was noticed, before and after ball milling, from grey to brown. This is due to the partial transition from the α-phase (s.g.: P42/mnm, crystalline phase of MgH2 at ambient conditions) to the γ-phase (s.g.: Pca21) which is formed under high pressure (13.9 GPa). 16 Here, weak reflections of γ-MgH2 appear at 30 and 36°.

MgH2(s) + 2 NH3(g) → Mg(NH2)2(s) + 2 H2(g) Equation II. 3 [57]
[58] Then, a jar with an inlet valve is used to perform the milling of MgH2 under NH3 pressure (99.9%, Messer). In an argon-filled glovebox, 600 mg of milled MgH2 is placed in a 25 cm 3 stainless steel jar equipped with a gas inlet. 23 stainless steel balls of 8 mm diameter are added (balls to powder weight ratio is 77:1). The jar is tightly closed and the following three steps are performed: After step 3, the process is repeated from step 1 twice to achieve a total of three cycles.

The repetition allows to evacuate the formed H2 and to ensure that all of the MgH2 has reacted with NH3. Once this ball milling is finished, the product is amorphous (cf. Figure II. 7c). The last step of the synthesis is an annealing at 290°C under 4 bar of NH3 to crystallize Mg(NH2)2 and to make sure that the reaction is completed. After annealing, pure crystalline Mg(NH2)2 is obtained

Chapter II [59] (cf. The IR spectrum exhibits the typical vibration bands of Mg(NH2)2: sharp asymmetric and symmetric stretching bands at respectively ~3324 and ~3270 cm -1 and bending bands at 1564 cm - 1 (cf. Figure II. 9). [START_REF] Zhao-Karger | Beyond intercalation chemistry for rechargeable Mg batteries: A short review and perspective[END_REF]18 Note that there is no trace of impurities as -OH groups around 3500 cm -1 . Chapter II [61] The DSC shows that Mg(NH2)2 starts to decompose into MgNH around 350°C (cf. 11). Here, the decomposition involves two endothermic-steps which might be due to a difference in the crystallites size. Indeed, after annealing, the product always consists in two different parts:

a fine powder and a heap of agglomerate powder, in which the crystallites could have different sizes. Similar multi-step decompositions have been observed in the literature and they are also attributed to inhomogeneities in the sample. [START_REF] Dolci | Magnesium imide: Synthesis and structure determination of an unconventional alkaline earth imide from decomposition of magnesium amide[END_REF] At higher temperatures, MgNH decomposes into Mg3N2 (around 500°C). The complete decomposition of Mg(NH2)2 is described according to the following reaction 20 : no parameters are given. 1 The third and last paper reporting the synthesis of Mg(BH4)(NH2) comes from Yang et al.. 21 The formation of Mg(BH4)(NH2) was observed while heating up a mixture of Mg(BH4)2.2NH3 and MgH2. At 100°C, pure Mg(BH4)(NH2) was detected by XRD.

Mg(NH2)2 → MgNH + NH3 → 1/3 Mg3N2 + 4/3 NH3 Equation II. 4
Owing to this literature, we decided to synthesize Mg(BH4)(NH2) by ball-milling and subsequent annealing as described by Noritake et al.. The main problematic of this synthesis was to find the correct ball-milling and annealing parameters as none of them are reported in detail.

X-ray diffraction was measured for each set of parameters to determine which one were optimal to obtain purely crystalline Mg(BH4)(NH2). The procedure of the synthesis optimization is summarized in Figure II. 12, the different synthesis parameters tested are described in Table II Mg(BH4)(NH2) [64] Table II. 1. Synthesis parameters investigated for the Mg(BH4)(NH2) synthesis.

[65] ). The product was still not pure but similar to sample 3. So a temperature of 120°C seems to be high enough to obtain Mg(BH4)(NH2) but a simple mixing in an agate mortar does not allow an intimate mixture of the two reactants and a complete reaction.

We tried to optimize the ball-milling parameters starting with softer conditions than those used previously. The balls to powder weight ratio was decreased from 105:1 to 38:1 and the ballmilling speed reduced to 100 rpm for 1h. In this case, the ball-milling does not degrade Mg(NH2)2 [67] into MgNH as no bands corresponding to the imide appear on the infrared spectrum (cf. 13). We noticed also that this sample appears foamy after the annealing treatment. We considered that the additional phase could be a decomposition product of the samples during annealing under Ar. Subsequently, the annealing atmosphere was changed from 1 bar of Ar to 10 bar of H2 to avoid any gas evolution during the synthesis. With this new annealing atmosphere, the product was not foamy anymore but it was impure and the additional phase was still present ). The annealing time of 72h was chosen as optimal parameter to make sure that the reaction is completed.

Later during the PhD, γ-Mg(BH4)2 stopped to be commercialized. We started to do the synthesis from the magnesium borohydride consisting in a mixture of the γ and α-Mg(BH4)2. We first used the same synthesis parameters than for sample C, i.e. a ball milling at 400 rpm for 1h with a balls to powder weight ratio of 38:1 and a subsequent annealing at 120°C for 72h under 10 bar of H2. The as-obtained product was mainly Mg(BH4)(NH2) but it was not pure (sample 12 on In the following sections, the samples labeled A, B, C and D will be compared due to their different compositions, especially regarding the content of additional phase as observed by XRD (cf. Figure II. 13). Their purity will be further analyzed by 11 B solid-state MAS NMR spectroscopy.

The thermal stability of the samples will be evaluated and the influence of the additional phase on the ionic conductivity will be investigated by electrochemical impedance spectroscopy (EIS).

II-4b. MAS-NMR spectroscopy

To verify the purity of our materials and the local environment of boron in Mg(BH4)(NH2), [70] between -37 ppm and -40 ppm, which can be unambiguously attributed to (BH4) -anions. 12,25 Samples A and B show actually two peaks in this chemical shift area: a first one at about -37.6 ppm and a second one at about -39.6 ppm (cf. Figure II. 17b). We can consider that one of the NMR peak is due to (BH4) -in Mg(BH4)(NH2) and the other one to (BH4) -in a slightly different environment as in the additional phase detected by XRD. To assign the two peaks, the spectra were fitted with two components and the percentage of boron atoms contributing to each environment (cf. Figure ppm decreases from sample A to sample B, with the decrease of the additional phase content. We conclude therefore that the peak at -37.6 ppm is due to the (BH4) -anions in Mg(BH4)(NH2) and that the second peak at -39.6 ppm is due to (BH4) -present in the additional phase.

In the case of sample C and D, for which only Mg(BH4)(NH2) was observed by XRD (cf. Figure Interestingly, the bending vibration bands of BH4 groups have changed compared to Mg(BH4)2. In the borohydride Mg(BH4)2, there are several bending bands around 1200 cm -1 (cf. Figure II. 4). In the as-synthesized Mg(BH4)(NH2), these bands are split into four distinct bands at 1109, 1230, 1439 and 1638 cm -1 . This is explained by the deformation of the BH4 tetrahedra with two B-H lengths shorter than the two others (cf. Table II. 2). The vibration due to NH2 has also changed compared to the one of Mg(NH2)2. The asymmetric stretching at 3326 cm -1 has a higher amplitude than the symmetric one at 3270 cm -1 which is the opposite of the Mg(NH2)2 amide (cf. Figure II. 9). The bending band at 1541 cm -1 has lower amplitude than in the amide. These amplitude changes might be caused by the change in the N-H bond length (cf. Table II.

2) and the change in

Chapter II

[73] the NH2 environment. Note that few imide has been formed during the synthesis (small absorption bands around 3200 cm -1 ) and that no OH groups are present (no absorption bands around 3500 cm -1 ).

Mg(BH4)(NH2)

B1-H1 1.17 [78]

The conduction phenomena at high and middle frequencies can be separated by the equivalent circuit described previously and described by their intrinsic characteristics: their activation energy (Ea) and their real capacitances. These characteristics helped us to validate or not our model. To extract the activation energy of each phenomenon, ln(σ.T) is plotted as a function of 1/T for each sample as shown on Figure II. 23a. For the samples A, B and C, a clear difference of behavior can be observed for the two phenomena. The Ea for the middle frequency phenomenon remains constant around 1.3 eV whereas the one at high frequency decreases as the sample is getting higher purity in Mg(BH4)(NH2), evolving from 1.3 eV to 1.0 eV, which confirms our attribution of Mg(BH4)(NH2) as the grain in our model. These Ea values are similar to those reported in the literature for the same material (1.31 eV) and other Mg 2+ ionic conductors. 6,27 For sample D, the activation energy is about 1.0 eV for the high frequency phenomenon and 0.9 eV for the middle frequency one.

To evaluate our fit, model and physical meaning, the real capacitance values of each conduction phenomenon is extracted for the R//CPE couple. The calculated capacitances for each phenomenon are plotted versus the temperature on Figure II. 23b. They remain constant as a function of temperature, which is scientifically sound as the capacitance is only dependent on the electrical permittivity of the phases and their geometrical dimensions, these parameters being independent of the temperature. For the phenomenon at high frequency, a real capacitance of the order of 10 -10 F.cm -2 is obtained for all the samples whereas for the one at middle frequency, the real capacitance is around 10 -9 F.cm -2 for sample A, B and C but around 10 -8 F.cm -2 for sample D. Capacitance values around 10 -9 -10 -10 F.cm -2 are typical of ionic conduction phenomena 28 and, hence, confirm our assumption that the high and middle frequencies phenomena are due to the ionic conduction in both phases. The value of 10 -8 F.cm -2 corresponds to an interface between the electrode and the electrolyte. So, for sample D, the fitted middle frequency phenomenon is not due to the additional phase. The additional phase is in too low quantity in this sample to form grain boundaries detectable by EIS. For the other samples (A,B and C), the interface did not need to be modeled even though they were measured with the same electrodes. Thus, the equivalent circuit with two components was not appropriate for sample D. The data were fitted with a second equivalent circuit made of only one R//CPE component (cf. Figure II. 24). This model fits very well the data. The calculated activation energy (1.0 eV) and real capacitance (10 -10 F.cm -2 ) are the same than the one found for the phenomenon at high frequency while using a model with two R//CPE.

So, an equivalent circuit with one R//CPE is more appropriate to model the ionic conduction in sample D.

Chapter II

[79] [80] Once the models were validated, the ionic conductivity for each sample was calculated. For samples A, B and C the ionic conductivity is calculated from R1 and R2 as both were determined to be linked to ionic conduction phenomena. The equation used is as follow:

𝜎 = 1 (𝑅 1 + 𝑅 2 )
. 𝑙 𝑆 Equation II. 5

For sample D, only R1 was used in Equation II. This value of ionic conductivity is one of the highest reported at this temperature, getting close to the one obtained by Roedern et al. for Mg(en)1(BH4)2 (en = ethylene-diamine). 6 These results could be explained by the presence of the additional phase. Especially, once this phase gets amorphous as shown by 11 B MAS-NMR spectroscopy, it would create a glass-ceramic-like composite material with improved ionic conduction, as previously described by Tsukasaki et al. on Li2S-P2S5 glassceramics. 29 In their study, nano-crystals of Li3PS4 were observed by TEM within an amorphous matrix in Li2S-P2S5 glass-ceramics explaining the improvement in ionic conductivity compared to the glasses. This assumption is also supported in our study by the fact that the ionic conductivity of the pellet decreases with the increase of the proportion of crystalline additional phase. conditions. The ball-milling speed must be not too strongly energetic in order to avoid the decomposition of the reactants but strong enough to make an intimate mixture of the two reactants. Besides the ball milling, the annealing has to be properly performed as well. The annealing temperature should not be too high, otherwise the product will decompose, but the temperature should be high enough to crystallize Mg(BH4)(NH2). We also found out that a back pressure of H2 helps to keep the product stable during the synthesis and that the time of annealing has to be long enough to get a complete reaction. To obtain a compound with no other crystalline phase than Mg(BH4)(NH2) using pure γ-Mg(BH4)2 as precursor, a ball-milling at 400 rpm and a subsequent annealing at 120°C for 72h under 10 bar of H2 were performed. Whereas with a borohydride made of the γ and α-phases, a more energetic ball-milling at 500 rpm and the same subsequent annealing were needed. 11 B MAS-NMR spectroscopy revealed that even if by X-ray diffraction only crystalline Mg(BH4)(NH2) appears in samples C and D, an additional amorphous phase was still present. It was observed that this additional phase could be crystallized in samples for which the synthesis parameters were not optimized.

The content in additional phase present in the samples strongly influences the ionic conductivity. If the sample contains a large amount of the additional phase, its ionic conductivity will be lower than if it contains less of it. An ionic conductivity of 1.10 -5 S.cm -1 was obtained for sample D that is made of 80 molar % of crystalline Mg(BH4)(NH2) and 20 molar % of the amorphous additional phase. This conductivity is four orders of magnitude higher than the one previously reported by Higashi et al.. The activation energy is also lower than the one reported, 1.0 eV vs.

1.31 eV, indicating an improved Mg 2+ mobility. The ionic conductivity values become especially high once this additional phase is amorphous forming a glass-ceramic-like composite material.

This could explain why the ionic conductivity values measured are several orders of magnitude higher than the one reported in the literature. We can also consider that the additional phase is actually detrimental for the ionic conductivity as we observe that the more it is present in the sample, the lower is the ionic conductivity. However, it is necessary to know if the compound of Higashi et al. contained the additional phase or not to be able to conclude if it has a benefic or detrimental effect on the ionic conductivity.

III-1. Introduction

In the previous chapter, the formation of unknown phases has been observed during the synthesis of Mg(BH4)(NH2). It was determined that these phases were not decomposition products but were rather obtained from non-optimized synthesis. Thus, the hypothesis of the existence of obtained such as Li4(BH4)(NH2)3 and Li3(BH4)(NH2)2. [1][2][3][4][5][6][7][8] The Li + ionic conduction properties of these compounds are higher compared to Li2(BH4)(NH2) so it was important to know if similar phases do exist or not in the Mg(BH4)2-Mg(NH2)2 binary phases diagram and if so, to characterize their Mg 2+ ionic conduction properties.

In this chapter, the exploration of the Mg(BH4)2-Mg(NH2)2 binary phases diagram is presented. Looking for new Mg(BH4)2x(NH2)2(1-x) compounds, a series of synthesis with different stoichiometries has been performed. It actually confirmed the existence of new phases. Then, the synthesis parameters were optimized to try to improve the purity of these new phases and to characterize them more precisely.

[93]

III-3a. Structural analysis of the α-phase

Here above, a solid solution where the (BH4) -replace the (NH2) -anions in the crystal structure of Mg(NH2)2 was obtained for x = 0.20 (i.e. α-Mg(BH4)0.40(NH2)1.60). A new synthesis for this stoichiometry was performed but this time the ball-milling speed was increased to 600 rpm to make sure that a maximum amount of Mg(NH2 [95]

To make sure that the α-phase is a solid solution, another synthesis was performed for x = 0.17 with a ball-milling speed of 600 rpm Compared to the α-phase at x = 0.20, the one at x = 0.17 has smaller cell parameters. Surprisingly, the ball-milling at 600 rpm seems to help the stabilization of the α-phase for x = 0.25 instead of the β-phase that was obtained with a ball-milling at 400 rpm (cf. Figure III. 1). The α-phase obtained for x = 0.25 has larger cell parameters; so the cell parameters increase with x (cf. Figure III. 5). These observation confirms that the α-phase is a solid solution crystallizing in the same structure as Mg(NH2)2, for which the (BH4) -anions are substituting the (NH2) -leading to the expansion of the unit cell. The cell parameters are not following a Vegard's law as they are not changing linearly with x.

In order to determine where the (BH4) -anions were exactly located in the crystal structure, a Rietveld refinement was performed on each compound using the structure of Mg(NH2)2. A change in one of the N sites occupancy could mean that it is the actual location of the B. The H atoms are not taken into account due to their very low scattering effect. Before that the occupancies of the N sites were fitted, the structure of the different compounds could already be well fitted with exactly the same structure as Mg(NH2)2 (cf. Figure III. 4). This result can be explained by the fact that (NH2) -and (BH4) -anions have exactly the same number of electrons, 9 electrons each. Indeed, X-rays interact with the electronic cloud of the atoms so in our case it is difficult to have enough contrast between these two anions using standard powder X-ray diffraction (PXRD). After fitting, the N sites occupancies vary around their initial values but without important variation (cf. Table III. 3a). The variations are too low to be significant considering that only light elements are present in the structure and that an error is added because of the intensity correction from the Be window. Nevertheless, for x =0.25, the occupancy of the site of the atom N1 decreased more than the others. Another fit was realized exchanging the N1 atom on this site by a B atom (cf. Table III. 3b). It resulted in occupancies values higher than the initial one (0.5) meaning that this site is not occupied only by B atoms. This result confirmed that our PXRD data do not allow a complete resolution of the structure. It is assumed that the BH4 randomly substitute the NH2 in the structure.

To go further in the determination of the structure of the solid solution, neutron diffraction was considered as B and N atoms have a large difference of neutron absorption. However, for this kind of compound, it is really challenging to perform neutron diffraction because of the presence of incoherently scattering H and strongly absorbing B. A test was done on the sample x = 0.25 but no neutron diffraction peaks were observed (cf. Annex II). To avoid these negative effects, the solution would be to enrich the sample in 2 D and 11 B. [97] [99]

III-3b. Ionic conductivity of the α-phase

The ionic conductivity was measured by EIS (Electrochemical Impedance Spectroscopy) using 6 mm diameter pellets of about 0.9 mm thickness (cf. Figure III. 8). In the case of Mg(NH2)2,

the Nyquist plot at 100°C displays only a straight line generated by the purely capacitive behavior of the sample. Mg(NH2)2 does not conduct ions nor electrons. For x = 0.17, 0.20 and 0.25, a semicircle starts to appear around 90°C and it is followed by a straight line at lower frequencies.

The semicircle is attributed to the bulk ionic conduction of the materials. The straight line is the capacitance formed by the carbon blocking electrodes on the two sides of the solid electrolyte.

The data are fitted with one R//CPE to model the bulk ionic conduction in series with a resistance for the electrical resistance of the cell and a CPE for the capacitance at low frequency. The resistances below 80°C for x = 0.25 and below 90°C for x = 0.17 and 0.20 are too high to be measured precisely so the ionic conductivity cannot be determined below these temperatures.

The values obtained at 100°C for x = 0.17, 0.20 and 0.25 are 4x10 -9 , 5x10 -9 and 2x10 -8 S.cm -1 , respectively. High activation energies of about 1.6-1.7 eV were calculated for these compounds, revealing the relatively low Mg 2+ mobility in these -phases. The ionic conductivity values are

b) Ionic conductivity as a function of 1/T. c) ln(σ.T) as a function of 1/T. d) Ionic conductivity at 100°C as a function of x.

[102]

The improvement of the ionic conductivity can be explained by the evolution of the structure. By XRD, it was shown that the cell parameters were increasing with x. The expansion of the unit cell can help the mobility of the Mg 2+ ions forming a larger pathway through the structure for the Mg 2+ cations. The 11 B NMR spectroscopy showed also that there might be an increase of the BH4 rotational mobility with x. It could promote the Mg 2+ mobility through the paddle wheel mechanism, well described in the literature, where the rotating anions shovels the cations. 10

III-3c. Thermal analysis and α to β phase transition

DSC measurements at 5 °C/min were performed on the samples obtained after ball-milling and before the annealing of the compounds (cf. Figure III. 9). All the samples undergo an endothermic phenomenon at low temperatures (65-80°C). This endothermic peak has a low enthalpy; between 16 and 19 J.g -1 (~0.8-1 kJ/mol, cf. Table III. 4) and its onset temperature

Table III. 5. Enthalpy and temperatures values of the endothermic phenomenon recorded by DSC for different α-Mg(BH4)2x(NH2)2(1-x) compositions.

Then, the thermal stability of the samples obtained after annealing, i.e. the -phases, was measured by TGA coupled with DTA (cf. Figure III. 10). The endothermic phenomenon observed before annealing has disappeared so it can be attributed to the formation of the α-phase. All the annealed samples undergo an endothermic phenomenon at about 150°C. The enthalpy of this phenomenon is approximately 105 J.g -1 (~5.9 kJ.mol -1 , cf. Table III. 5) and there is no mass loss as observed on the TGA curves.

To determine the nature of this endothermic phenomenon, a DSC measurement was performed on the sample for x = 0.25 upon heating, until to 170°C, and cooling. This phenomenon is irreversible. Then, the sample with x = 0.25 has been annealed at 170°C for 12h under 1 bar of Ar to see what happens to the α-phase. The crystalline β-phase was formed. So, the α-phase turns irreversibly into the β-phase at about 150°C through an endothermic structural transition caused by the reorganization of the structure. As the β-phase remains stable at room temperature, it can be considered that the α-phase with x = 0.25 is actually metastable. This metastability is highlighted by the fact that only a powerful ball-milling can lead to its synthesis. As a matter of fact, for x = 0.25, a ball-milling at 400 rpm leads to the formation of the β-phase (cf. whereas the α-phase is obtained at 600 rpm.

Chapter III

[105]

III-4a. Crystal structure of the β-phase

In the previous paragraph, the β-phase was obtained for x = 0.25 by annealing of the α-phase or by ball-milling at lower speed (400 rpm). Another synthesis was performed for x = 0.33 with a ball-milling at 600 rpm followed by an annealing at 120°C for 72h under 10 bar of H2. The as-obtained sample is very amorphous and only weak reflections of the β-phase appear (cf. Figure III. 12). The low crystallinity of the sample might be due to a low stability of the phase with this stoichiometry. It is interesting to compare XRD pattern of the β-phase to the one of sample A which was obtained for a non-optimized synthesis of Mg(BH4)(NH2) and contains mainly a crystalline additional phase. Indeed, the reflections of the β-phase are similar to the ones of the additional phase so it can be assumed that it is actually the same phase but with slightly different unit cell parameters. Therefore, the β-phase seems to behave like a solid solution as the α-phase

III-4b. Ionic conductivity of the β-phase

The ionic conductivity of the β-phase was measured by EIS using 6 mm diameter pellets of about 0.9 mm thickness. The results were fitted with the same model as the one used previously for the α-phase. For the sample obtained after annealing of the α-phase, the resistances were too high to be measured at low temperatures (cf. Figure III. 13). The value of ionic conductivity at

100°C is approximately the same as the α-phase, 2.10 -8 S.cm -1 . However, the activation energy increased, reaching a high value of about 2 eV. The activation energy is similar to the one of the α-phase, 1.6 eV. The conductivity in the β-phase seems to increase with x as observed in the α-phase. Nevertheless, the two samples were not synthesized with the same route so the comparison remains difficult and not fully conclusive.

III-4c. Thermal analysis of the β-phase

Thermal analyses were performed on the sample with x = 0.33 synthesized at 600 rpm. The DSC performed on the mixture after ball-milling and before annealing is similar to what was obtained for the compositions x = 0.17, x= 0.20 and x =0.25 (cf. Then, TGA coupled with DTA was performed on the sample obtained for x = 0.33 after annealing. It shows that no thermal phenomenon happens below 180°C but, after this temperature, a slight mass loss is noticed which is related to the beginning of the material decomposition.

Chapter III

[109] The structure of the γ-phase can be indexed in a tetragonal unit cell with space group I-4

(No.: 82) (cf. Figure III. 17). The fitted unit cell parameters reveal that the increase of the amount of (BH4) -anion leads to an expansion of the cell (cf. Figure III. 18). This expansion is similar to the one observed in the α-phase so the γ-phase is also a solid solution in which (NH2) -anions can be partially substituted by (BH4) -ones. It is also interesting to note that this cell has a tetragonal symmetry as other compounds in the Mg(BH4)2:Mg(NH2)2 binary system: Mg(NH2)2 (s.g.: I41/acd), Mg(BH4)(NH2) (s.g.: I41) and δ-Mg(BH4)2 (s.g.:I41/acd).

Chapter III

[111] [112] III . 8). This structure still needs to be confirmed and optimized by DFT calculations, especially for the H atom positions. Better XRD data would be also useful.

[113]

III-5b. Ionic conductivity of the γ-phase

The ionic conductivities of the samples containing the -phases were measured by impedance spectroscopy using 6 mm diameter pellets of about 1.1 mm thickness (cf. 21). The Nyquist plots show only one semicircle at high frequency that is attributed to the ionic conduction in the materials. The semicircle is followed by a straight line that is the capacitance formed by the carbon blocking electrodes placed on both sides of the samples. In the same way than for the α and the β-phases, the data are fitted with one R//CPE for the bulk ionic conduction in series with a resistance for the cell and with a CPE for the capacitance formed by the electrical contacts on both sides of the electrolyte. The ionic conductivity values extracted from these fits are presented as a function of the temperature on Figure III. 21b. At 100°C, the ionic conductivity is 0.8x10 -5 , 2x10 -5 and 4x10 -5 S.cm -1 for x = 0.75, 0.71 and 0.67, respectively (cf. temperature. The ionic conductivity is increasing as x is decreasing in the γ-phase. This trend is opposite to what was observed for the α and β-phases.

The activation energies determined from these measurements are low with values equal to about 0.8, 0.9 and 1.0 eV for x = 0.67, 0.71 and 0.75, respectively (cf. Figure III. 21c). These activation energies are low for an inorganic Mg 2+ solid ionic conductor, indicating a high mobility of the Mg 2+ cations in the γ-phase. This mobility is increasing as x decreases which is also confirmed by the ionic conductivity values.

Thanks to the good mobility of Mg 2+ in the γ-phase, it was possible, for the sample x = 0.71, to record 25 Mg solid state MAS NMR spectra with a good signal in a short time without isotopic enrichment in 25 Mg. A good spectrum could be recorded in 5h whereas, for Mg(BH4)(NH2), an acquisition of about 15h was necessary. This improvement of the signal allowed us to do measurement in temperature and to try to obtain more information about the Mg 2+ mobility in this -phase (cf. Figure III. 22a). Tries were attempted to measure the spin-lattice relaxation time T1 at different temperatures to calculate an activation energy using the Bloembergen-Purcell-Pound formula. Indeed, from this formula, a Mg 2+ motion correlation time τ can be obtained. This correlation time should follow an Arrhenius behavior from which the activation energy can be calculated. Unfortunately, the variation of T1 was too small and in the range of the signal-to-noise ratio of the measurement. So, it was not possible to determine the activation energy with this technique.

Still, with these measurements in temperature, a clear decrease of the FWHM of the peak was observed, indicating an increase of the Mg 2+ mobility with the temperature. Looking at the 1 H solid state MAS-NMR spectra recorded in the same conditions, no changes of the shape of the peak can be observed. These two measurements prove that the mobile species in this phase are Mg 2+ cations and definitively not H.

Chapter III

[117]

III-5c. Thermal stability of the γ-phase

DSC was measured on samples for x = 0.67, 0.71 and 0.75 just after the ball-milling step and before the annealing (cf. Figure III. 23). As the α-phase, these compounds undergo an endothermic phenomenon at low temperature, around 60°C. Here, the enthalpy is higher with values ranging from 25 to 37 J.g -1 (1.4 to 2.0 kj.mol-1, cf. Table III. 9). The enthalpy and the onset temperature of this phenomenon decrease with x. It could be due to a rearrangement of the crystalline structure.

At about 110°C, an exothermic phenomenon occurs with a higher enthalpy, about -160 J.g -1 (~9 kJ.mol -1 ). This exothermic phenomenon is related to the crystallization of the γ-phase. At 200°C, there is another endothermic phenomenon followed by the degradation of the product. This last phenomenon might be the melting of the products as observed for Li2(BH4)(NH2) and Li4(BH4)(NH2)3. 6 [119]

III.6. Comparative study of the Mg(BH4)2-Mg(NH2)2 binary phases diagram

From the data presented here above and the ones obtained in the chapter II, a pseudobinary Mg(BH4)2-Mg(NH2)2 phases diagram is proposed between room temperature and 200°C (cf. and Mg(NH2)2 as they are very bad ionic conductor. The ionic conductivity of Mg(NH2)2 was too low to be measured and the one of Mg(BH4)2 is reported to be about 10 -9 S.cm -1 at 150°C. 11 So it is interesting to see that the phases based on a mixture of this two compound show better ionic conductivities. For instance, the substitution of (NH2) -anions by (BH4) -in the α-phase improves the ionic conductivity, reaching 2x10 -8 S.cm -1 for x = 0.25. This increasing trend is followed by the β-phase as well with an ionic conductivity of 2x10 -7 S.cm-1 for x = 0.33. In the center, Mg(BH4)(NH2) has a good ionic conductivity of 1x10 -5 S.cm -1 and a low activation energy of 1.0 eV. Finally, the phase rich in borohydride, the γ-phase, is the one having the highest ionic conductivity and the lowest activation energy with 4x10 -5 S.cm -1 and 0.8 eV for x = 0.67. In this phase, the ionic conductivity decreases as x increases. Overall, even if the crystalline structure changes from one compound to another, a trend appears in terms of ionic conductivities. Indeed, the maximum was measured for the γ-phase at x = 0.67. From this point, the ionic conductivity gradually decreases and it is true from one compound to the other. This trend is not random and is concomitant to the evolution of the structures and of the different substitutions of (NH2) -anions by (BH4) -.

Figure III. 26. Ionic conductivity at 100°C and activation energy as a function of x in Mg(BH4)2x(NH2)2(1-x).

Chapter

IV-1. Introduction

During this PhD thesis, new solid-state Mg 2+ conductors have been discovered but before to use them as solid electrolytes, it is crucial to know more about their chemical and electrochemical stabilities. Indeed, a solid electrolyte needs, of course, to be electronically insulative and to have a good ionic conductivity but other parameters are also very important. The chemical stability is important to make sure that the electrolyte will not react with one of the active or inactive material of the battery. Any unwanted reaction can hinder or make impossible the cycling of the electrochemical cell. The compatibility between the solid electrolyte and the electrodes is fundamental for a rechargeable device. A bad interface would lead to a quick fading of the device. Finally, the electrochemical stability window (ESW) of a solid electrolyte has to be as large as possible to avoid its reduction at low potential or its oxidation at high potential. A wide ESW would provide a larger choice in terms of usable electrodes for rechargeable Mg batteries.

In this chapter, Mg(BH4)(NH2) and γ-Mg3(BH4)4(NH2)2 have been characterized as solid electrolytes. It was shown in the previous chapters that they have both good ionic conductivities (~10 -5 S.cm-1 at 100°C). First, different active and inactive materials that can be used in a rechargeable Mg battery have been selected. Then, the chemical stability of the ionic conductors was tested by annealing mixtures of electrode materials with the solid electrolytes. The electrochemical stability is evaluated by polarization and cyclic voltammetry measurements. The polarization is especially used to characterize the interface between the electrolyte and the anode. The cyclic voltammetry enables the determination of the ESW.

[130]

IV-2. Chemical stability

IV-2a. Selection of electrode materials

First, it was needed to know if our solid electrolytes would react with components of a rechargeable all-solid-state Mg battery. Two different materials have been chosen for the anode:

Mg and Bi. The Chevrel phase Mo6S8 has been selected as the cathode material and carbon C45 has been used as electronic conduction additive. Each material is described here after.

Magnesium metal anode

One of the main advantages of using a solid electrolyte should be the possibility of using a metallic anode. In the first chapter, the different properties of magnesium metal as negative electrode for batteries have already been described. As a reminder, its redox standard potential is -2.375 V vs. NHE and it has a gravimetric capacity of 2205 mAh.g -1 . Commercial magnesium powder (Alfa Aesar, 99.8 % purity) was used. Magnesium metal is usually covered by a very stable layer of MgO. So, the powder was ball-milled under argon to break this oxide surface layer and make the Mg particles more reactive. The ball-milling was performed under Ar as followed: 2 g of commercial Mg powder was placed in a 50 cm 3 WC airtight jar and planetary ball-milled at 600 rpm for 2 h with a 2 min break every 30min and with a balls to powder weights ratio of 53:

Bismuth anode

Bismuth has been considered as a second solution for the anode in case of reactivity and/or instability of the interface between Mg and the solid electrolytes. The Mg-Bi binary phases diagram indicates that no solid solution exists between these two elements and that only the intermetallic compound Mg3Bi2 can be formed. The magnesiation of Bi leads to the formation of a biphasic mixture of Bi and Mg3Bi2 as follows: 1,2 3x Mg 2+ + 2 Bi +6x e-→ x Mg3Bi2 + (2 -2x) Bi Equation IV. 1

This reaction happens at ~ 0.27 V vs Mg 2+ /Mg and has a specific capacity of 385 mAh.g -1 .

Bi powder obtained after ball-milling of Bi ingots (Prolabo, 99.999 % purity) was used. The ballmilling parameters were the same than the ones described previously for Mg. After ball-milling, the powder was sieved to separate large aggregates from the fine powder. Then it was stored in an Ar filled glovebox. The powder does not show any traces of impurities as analyzed by XRD. The only seen peaks are from Bi crystallizing in a hexagonal structure with the space group R-3m and the cell parameters a = 4.5438(4) Å and c= 11.854(1) Å (cf. Figure IV. 1b).

Mo6S8 Chevrel phase as cathode

For the cathode, the Chevrel phase Mo6S8, 3 which is commonly used for Mg intercalation, 4,5 has been selected. It has been synthesized at the ICGM in Montpellier from Cu2.5Mo6S8 by acid leaching in a solution of HCl to remove Cu. The reaction with Mg is as follows:

x Mg 2+ + Mo6S8 + 2x e-→ MgxMo6S8 Equation IV. 2

The intercalation occurs in two steps: a first one at 1.2 V vs. Mg 2+ /Mg and a second one at 1.0 V vs Mg 2+ /Mg. Each step corresponds to the intercalation of Mg into two different crystallographic sites. The maximum quantity of Mg that can be intercalated is x = 2 for a theoretical specific capacity of 122 mAh.g -1 . This maximum capacity is hard to reach as, upon charge, the ionic conductivity decays as the concentration of Mg ions decreases, leading to the entrapment of Mg ions. From our XRD measurement, pure crystalline Mo6S8 is obtained with the hexagonal space group R-3 and the cell parameters a = 9.1837(8) Å, c = 10.869(1) Å and

V = 793.8(1) Å 3 (cf. Figure IV. 1c).
[132]

C45 carbon black

To enhance the electronic conductivity of the electrodes, it is necessary to use a carbon

IV-2b. Chemical stabilities of Mg(BH4)(NH2) and γ-Mg3(BH4)4(NH2)2

To characterize the chemical stability of our solid electrolytes, Mg(BH4)(NH2) (sample D in

Chapter II) and γ-Mg3(BH4)4(NH2)2 were mixed with the electrode materials described above in a 1:1 mass ratio. The mixtures were done in an Ar filled glovebox by hand with an agate mortar.

They were subsequently annealed under Ar for 12h at 100°C for Mg(BH4)(NH2) and 150°C for γ-Mg3(BH4)4(NH2)2 as it was observed in Chapter III that this compound is thermally stable at least up to 180°C. The powders are then recovered after annealing and analyzed by XRD.

For the mixture of the solid electrolytes with the anode materials: Mg and Bi (cf. Nevertheless, the volume of cell is similar to the one of the raw materials. So, no bulk reaction has occurred. This is an interesting result, because it indicates that the solid electrolytes do not react with the anode materials. Indeed, some liquid electrolytes are not stable at low potential and form a passivation layer on the anode. This passivation layer is blocking the Mg 2+ ions diffusion and thus makes impossible the cycling of the cell. 6 With these solid electrolytes, the interface with the anode would be chemically stable.

On the contrary, the mixture with Mo6S8 led to a clear reaction between the solid electrolyte and the electrode material. The magnesiation of the Chevrel phase occurs forming MgxMo6S8. A profile matching was performed on the XRD pattern of the mixture after annealing. The refined unit cell parameters of MgxMo6S8 are a = 9.7693(4) Å and c= 10.3734(8) Å (V = 857.31 Å 3 ) for the mixture with Mg(BH4)(NH2) and a = 9.7549(8) Å and c= 10.362(2) Å (V = 853.8(2) Å 3 ) for the one with γ-Mg3(BH4)4(NH2)2. These parameters correspond to the Chevrel phase with x = 2 so it has been completely magnesiated (the cell parameters of Mo6S8 are a = 9.1837(8) Å, c = 10.869(1) Å and V = 793.8(1) Å 3 ). The solid electrolytes are therefore oxidized in contact with the cathode. This oxidation forms amorphous products that are not detected by XRD. It even leads to a complete amorphization of the solid electrolyte for γ-Mg3(BH4)4(NH2)2. This kind of reaction was also observed in the study of Unemoto et al. where the lithiation of TiS2 in contact with LiBH4 is observed. 7 Titanium disulfide has a working potential of about + 2.3 V vs. Li + /Li and recent studies showed that the actual stability of LiBH4 is + 2.0 V vs. Li + /Li. 8,9 So, LiBH4 can be spontaneously oxidized by TiS2. Similarly, our results indicate that the oxidation potential of our solid electrolytes is lower than the working potential of Mo6S8.

[134] [136]

IV-3. Electrochemical stability

III-3a. Anode/electrolyte compatibility

To determine the compatibility between the electrolyte and the anode, galvanostatic polarization measurements were performed on a symmetrical cell. The cell was simply made of a parallel stack of two anodes (Mg or Bi) separated by the solid electrolyte. Ideally, the anodes are pressed onto the electrolyte to form a self-standing stack and to get good electrical contacts.

Unfortunately, Mg is a relatively hard metal and it cannot be stacked onto the electrolyte. Mg has a Brinell hardness (HB) of 260 at room temperature which is close to the HB of Al (245) for instance. 10 By comparison, alkaline elements like Na or K have HB < 1. This means that Mg is not a soft metal and it does not easily form good contact with other materials. So, instead, our stack was made with two magnesium discs made from cold pressed powder and a pellet of solid electrolyte. This assembly was not optimum for good contacts but it was the only way to use Mg powder as anode in a symmetrical cell.

The results for Mg(BH4)(NH2) are presented on Figure IV. 6. A current of ±5 µA.cm -2 was applied for 2h and each polarization was separated by 5 min of rest. During the first cycles, the potential evolved slowly before reaching its steady-state equilibrium with a plateau at about 0.5 V. This evolution of the potential can be due to the slow diffusion of Mg 2+ in the solid electrolyte.

If the ionic conductivity of our pellet is 1x10 -5 S.cm -1 at 100°C, the expected polarization potential should be around 25 mV. The measured polarization was 20 times higher. The bad contacts between the electrolyte and the electrodes might be the cause of this over potential, creating a high interfacial resistance. The interfacial resistance was estimated to be about 120 kΩ (152 kΩ.cm -2 ). As the cycling continued, the interfacial resistance increased and, after 60 h of cycling, the polarization did not reach a steady-state anymore but reached a maximum before slowly decreasing. This new evolution of the potential can be due to side reactions happening, as inhomogeneous deposition of Mg (dendrites, foam, etc.) 11 or decomposition of the solid electrolyte. Therefore, this symmetrical cell measurement indicates that the interface between Mg(BH4)(NH2) and Mg is electrochemically unstable for such stack. A different anode was then used, the magnesiated form of Bi, i.e. Mg3Bi2. With this anode, no metal deposition/stripping process occurs anymore but a biphasic reaction instead as described in Equation IV. 1. For polarization measurements, the same redox reaction is needed at both electrodes to be able to reach equilibrium and to be sure that the polarization evolution is only due to the interface evolution and not to electrochemical reactions occurring at the electrodes. Thus, Mg3Bi2 cannot be used directly otherwise on one side would occur the biphasic reaction described in Equation IV. 1 and, on the other side, Mg deposition:

x Mg 2+ + 2x e-→ x Mg Equation IV. 3

[138]

An electrode was prepared by ball-milling Mg and Bi with the objective that only half of the Bi was magnesiated. In an Ar filled glovebox, 0.177g of Mg were mixed with 1.840g of Bi. The mixture was placed in a stainless steel jar with 12 stainless steel balls of 8 mm diameter for a balls to powder weights ratio of 12:1. The ball-milling is performed using a Spex mill for 5h. The XRD pattern of the as-obtained product is presented on Finally, to use this material, 10 wt% of C45 was added to ensure a good electronic conductivity as Bi is a bad electronic conductor. The electrodes were cold pressed into discs and used in the same configuration as with the Mg electrodes.

Even if Bi has a lower hardness than Mg (HBBi = 94.2), it is still too hard to stick to the solid electrolyte. A low current density had to be used for the galvanostatic measurements to avoid a too important polarization. Already, with a current density of 1 µA.cm -2 , the polarization potential was about 1 V (cf. Figure IV. 8) so the interfacial resistance was about 1.2 MΩ (1.5 MΩ.cm -2 ). This huge resistance was not due to chemical reactivity as shown in the previous section. The interface was inhomogeneous as the electrode surface was made of Bi, Mg3Bi2 and C45 and this inhomogeneity could explain the initial bad interface. The first polarization cycles were stable and featured the formation of an arc before to reach a steady-state. Rapidly, the polarization turned into an unstable state and the potential increased. The complex shape of the polarization observed after several cycles might be due to a detrimental evolution of the contact between the electrode and the solid electrolyte because of the volumetric expansion of Bi upon cycling. The volume increases of 73 % when Bi transforms into Mg3Bi2. Actually, important volume changes are very problematic in all-solid-state batteries because it cannot be compensated by any change of shape or mechanical flexibility. It can also be assumed that upon cycling some of the Bi or Mg3Bi2 particles could lose their electrical contact with the carbon. So, the change of Mg metal for Bi anode unfortunately did not improve the compatibility of the anode with the solid electrolyte. [140]

III-3a. Electrochemical stability window

One of the most important parameter to know about our solid electrolytes is their electrochemical stability window (ESW). To determine the ESW of our materials, cyclic voltammetry was performed using a cell made of a Mg anode versus a composite positive electrode made of the solid electrolyte mixed with 10 wt% C45. Both electrodes were separated by the solid electrolyte. The cyclic voltammetry was performed with a slow scanning rate of 0.05 mV.s -1 at 100°C starting by the reduction process. the oxidation potential is lower than that reported by Higashi et al.. 12 In their work, they assumed that Mg(BH4)(NH2) would not be oxidized before 3V vs Mg 2+ /Mg. This value was clearly overestimated as they performed their cyclic voltammetry using a flat Pt electrode as counter electrode and a fast scanning rate of 1 mV.s -1 . The use of a flat counter electrode does not allow a good contact with the solid electrolyte and can hinder the oxidation reaction as only a few amount of the electrolyte would react and the measured current is therefore very low. The measured current will be even lower if the scanning rate is fast (kinetic limitation). A recent study showed that using carbon with different specific surface area changed the kinetic of oxidation reaction and changed the observed oxidation potential. 9 The most accurate measurements were obtained for composite electrodes using high specific surface area carbon blacks. The high contact surface area allows a good electronic percolation through the solid electrolyte and hence ensures a precise measurement of the oxidation potential. This highlights once more the importance of the contact surface area between the counter electrode and the solid electrolyte. In our case, the fact that a composite counter electrode was used instead of a flat electrode explains the So, on one hand, these materials are unstable at potentials higher than 1.0 V vs. Mg 2+ /Mg and would be oxidized by cathodes such as Mo6S8. This oxidation leads to the detrimental formation of non conducting compounds. On the other hand, they are stable chemically and electrochemically versus the anode materials. Unfortunately, it is hard to obtain a good interface between the anode and the solid electrolytes due to the important hardness of Mg metal. This point is critical to develop all-solid-state Mg batteries and might be improved by using foldable Mg thin foils or very fine powder of the anode material that could be dispersed in a composite electrode with the solid electrolyte.

General conclusion

The main aims of this PhD thesis were to enhance the ionic conductivity of known Mg 2+ inorganic ionic conductors, to find new ones and to characterize them electrochemically in order to know if they could be used as solid electrolytes in Mg batteries. The first studied material was mainly the additional phase (73 %) has an ionic conductivity of 8.10 -8 S.cm -1 at 100°C with an activation energy of 1.3 eV; while the sample containing less additional phase (20 %) has an ionic conductivity of 1x10 -5 S.cm -1 at 100°C with an activation energy of 1.0 eV. This last ionic conductivity value is four order of magnitude higher than the one previously reported by Higashi et al. (~10 -9 S.cm -1 at 100°C and Ea = 1.31 eV). 2 Two hypotheses can be made to explain this discrepancy. In the first hypothesis, the material of Higashi et al. is considered as 100 % pure Mg(BH4)(NH2), then the presence of the additional phase in our composite samples would enhance the ionic conductivity, especially, once it becomes amorphous forming a glass/ceramiclike material with enhanced Mg 2+ mobility. 3 For the second hypothesis, the additional phase has actually a negative effect on Mg 2+ mobility. Indeed, it was observed in our study that the lowest is its content; the highest is the ionic conductivity. In that case, the material of Higashi et al. would contain an important amount of amorphous additional phase, lowering its ionic conductivity. To validate one of these two hypotheses, a precise description of the synthesis parameters used by Higashi et al. and 11 B MAS-NMR spectroscopy measurements on their sample would have been very useful. However, in both cases, the additional phase has clearly an impact on the ionic

General conclusion

[148] conductivity and it was important to try to obtain it pure. As it contains (BH4) -anions, as-probed by 11 B MAS-NMR spectroscopy, that it was not a decomposition product and that it was obtained for non-optimized synthesis, the assumption was made that it was actually a never reported Mg(BH4)2x(NH2)2(1-x) compound.

Then, to try to identify this additional phase and to find new Mg 2+ ionic conductors, the Mg(BH4)2-Mg(NH2)2 binary phases diagram has been explored (cf. Chapter III). Three new Mg(BH4)2x(NH2)2(1-x) crystalline phases were discovered. Two phases with a composition rich in amide were named the α-phase and the β-phase. The α-phase crystallizes with the same crystal structure as Mg(NH2)2 in which large (BH4) -anions are substituting the (NH2) -. This substitution leads to an expansion of the unit cell and an improvement of the ionic conductivity, increasing from 4x10 -9 S.cm -1 for x = 0.17 to 1x10 -8 S.cm -1 for x = 0.25. Upon heating, the α-phase undergoes an irreversible endothermic structural transition, forming the β-phase. The β-phase crystallizes in a new unknown crystal structure. This phase transition has a little impact on the ionic conductivity of the material, σ100°C = 2x10 -8 S.cm -1 , and the activation energy of the ionic conductivity increases from 1.7 eV to 2 eV. For x = 0.33, with a ball-milling at 600 rpm, poorly crystallized β-phase was obtained with an amorphous background visible by XRD. The ionic conductivity at 100°C of this last compound is improved, as it reaches 2x10 -7 S.cm -1 and its activation energy is 1.6 eV. The βphase might be the additional phase present in the composite materials described in Chapter II during the Mg(BH4)(NH2) synthesis. Indeed, for the samples with the crystalline additional phase, the X-ray diffraction patterns and ionic conductivity values are similar. This observation would confirm the second hypothesis we have made, considering that the presence of additional phase decreases the ionic conductivity of the composite materials.

Another phase crystallizes in the borohydride-rich region of the Mg(BH4)2-Mg(NH2)2 phases diagram: it was named the γ-phase. It has a tetragonal structure with the I-4 space group. Similarly to the α-phase, adding more (BH4) -anions into this compound increases the unit cell volume but, here, it decreases the ionic conductivity. The best ionic conductor is γ-Mg(BH4)1.33(NH2)0.67 (which can be also written γ-Mg3(BH4)4(NH2)2) with an ionic conductivity at 100°C of 4x10 -5 S.cm -1 and an activation energy of 0.8 eV. These values mean that it features one of the best Mg 2+ mobility in an inorganic solid ionic conductor reported so far. The good mobility of Mg 2+ was confirmed by 25 Mg solid state MAS-NMR spectroscopy with a significant narrowing of the NMR peak upon heating. was not affected by the composition of the phase studied. Unfortunately, it is hard to obtain a good interface between the anode and the solid electrolytes due to the important hardness of Mg metal. This point is critical to develop all-solid-state Mg batteries and might be improved by using foldable Mg thin foils or very fine powder of the anode material that could be dispersed in a composite electrode with the solid electrolyte.

To summarize, through the exploration of the Mg(BH4)2-Mg(NH2)2 phases diagram, we have discovered new solid inorganic Mg 2+ ionic conductors. A trend appears between the ionic conductivity and the composition of these materials. One of them reaches one of the highest ionic conductivity ever reported for an inorganic Mg 2+ ionic conductor (4x10 -5 S.cm -1 at 100°C) and has also one of the lowest activation energy (0.8 eV). Deeper structural (synchrotron X-ray diffraction)

and theoretical studies (DFT) could help to understand this trend. It might depend on the crystalline structure (Mg 2+ mobility pathway) and the different substitution of (NH2) -anions by

[150] (BH4) -(i.e. interaction of Mg 2+ with its surrounding). Understanding this trend in this binary phases diagram could allow the design of better Mg 2+ ionic conductor based on other compounds. For instance, the substitution of (NH2) -and (BH4) -by other anions such as (B12H12) 2-(as demonstrated with Na3(B12H12)(NH2)) 4 or (CB11H12) -could lead to the discovery of better ionic conductors. To obtain this kind of compounds, important effort on the synthesis process would be needed as there is no report of pure magnesium closo-borane or carborane synthesis. They are stabilized only with solvent molecules or in carbon scaffolds. 5,6 Synthesis process could be imagined to react Mg(BH4)(NH2) with solid or liquid B10H14 for instance, similarly to what is done to get Li2B12H12 and Na2B12H12. 7 In terms of electrochemistry, it is hard to imagine using these materials directly as bulk solid electrolytes. However, their chemical and electrochemical stability vs. the Mg anode makes them interesting as coating materials. Indeed, it is possible to think about a system where the electrodes are coated by a thin layer of stable electrolytes and separated by a good solid or liquid ionic conductor. Then, of course, the contact with Mg metal will have to be optimized and our materials will have to be chemically stable with the bulk electrolyte. If a coating is necessary at the cathode, materials such as phosphates should be stable. 8 Other compounds such as magnesium halides are studied by Canepa's group to form stable interfaces that are crucial for rechargeable batteries. 9 For bulk solid state electrolyte, more studies could also be done on Mg2P2S6. Indeed, this compound has been reported only once in the work of Yamanaka et al., where it was obtained as a side product studying MgS-P2S5-MgI2 glass-ceramics. 10 A recent study showed that Zn2P2S6 features a good Zn 2+ mobility with an ionic conductivity of about 10 -7 S cm -1 at 60°C and more interestingly, a very low activation energy of 351 ± 99 meV. 11 So it is possible that Mg2P2S6 also has good Mg 2+ mobility because it has a very similar crystal structure as that of Zn2P2S6

Annex I

[155]

AI.1 -Note on the air sensitivity of the materials

The studied materials are extremely moisture and O2 sensitive so all the experimental techniques have been adapted to be able to keep the materials protected. The storage and all the handling of the different products were done in an Ar filled glovebox. All the different syntheses and measurements were performed using airtight set-ups. In contact with H2O or O2, the hydrides, borohydrides and the amides would form very exothermically oxides and hydroxides and release an important quantity of gases such as H2 and NH3. This can be dangerous for instance if an important quantity of material is exposed to H2O, especially if the material is powdery. In cases where the materials have to be exposed to air, like for the cleaning of the ball milling jar, it is important to neutralize them quickly after exposition to air and if possible under a fume hood. For the neutralization, it is better to use ethanol as the production of gases will be slower than with water and thus avoid explosions (cf.

AI.2 -Synthesis methods

AI.2a -Ball-milling

Ball-milling was used to perform mechanochemical syntheses and also to reduce the grains size of powders. Two kinds of ball-milling were used during this work: the planetary ball-milling and the Spex ball-milling. Planetary ball milling will produce frictions while the balls rolls on the walls of the jar and compression when the balls are projected from one side of the jar to the other.

The Spex ball milling will mainly shake the balls and the powder describing a 3D 8-shape rotation.

The planetary ball-milling was performed with a Retsch PM100 mill and a 50cm 3 tungsten carbide (WC) jar with 10 mm diameter WC balls. The ball-milling jar was sealed by a rubber o-ring encased in its cap which is tight together with the jar with a clamping bracket. The spex ball-milling was performed with a Spex 8000M mill. To perform the Spex ball-milling under NH3 atmosphere, a 25 cm 3 stainless steel Spex jar equipped with a gas inlet was used (cf. Figure AI. 2). The atmosphere inside the jar was changed using the same montage than described here after for the annealing reactor. For Spex ball-milling under Ar, a jar without gas inlet was used. 8 mm diameter stainless steel balls were used. The Spex jars were airtight thanks to a rubber o-ring encased in the jar and tight with the cap using screws. Annex I

[157]

AI.2b -Annealing under controlled atmosphere

Annealing was needed to crystallize samples or to promote chemical reactions. A stainless steel airtight reactor was used (cf. Figure AI. 3). The sealing of the reactor is ensured by a copper ring that is strongly tightened between the body of the reactor and the cap. The cap was equipped with a gas valve using graphite seals. With such reactor, the annealing can be performed at high temperature and under several bar of gas pressure. During this thesis, three different atmosphere of annealing had to be used: Ar, H2 and NH3. To do annealing under Ar atmosphere, the atmosphere of the Ar filled glovebox in which the samples were loaded in the reactor was used.

Then to place the reactor under different atmospheres, a montage was assembled. It was made of a vacuum pump and a gas inlet (H2 or NH3), both connected to a series of pipes and valves. The montage allowed to vacuum the reactor up to 1x10 -5 bar. Then, by closing the valve of the vacuum and opening the one of the gas inlet, the reactor could be refilled without any air contamination.

AI.3 -Structural analysis

AI.3a -Powder X-ray diffraction

Powder X-ray diffraction (PXRD) was performed on Bruker D8 diffractometers using Co Kα1-Kα2 or Cu Kα1-Kα2 radiation in a θ-θ configuration. For the measurements, the LeRiCh'S cell that has been developed in the lab was used. 1 It is airtight and uses a X-ray-transparent Be window.

The Fullprof suite was used to analyze our data. Le Bail method 2 (profile matching) was used to determine the unit cell parameters. Rietveld refinement 3 was used to get information on the complete crystalline structure of the unit cell and also to quantify the amount of phases present in some samples. For unknown structures, the software DICVOL 4 was used to find the symmetry of the unit cell and the software checkcell helped us to find the best space group based on the extinction rule.

To perform correctly the Rietveld refinements, the measured intensity had to be corrected from the absorption of the X-ray by the Be window. The correction was done with the Beer Lambert law: 𝐼 = 𝐼 0 . exp (-µ. 𝑥) Equation AI. 1

with I: the X-ray intensity after absorption I0: the X-ray intensity before absorption µ: the mass absorption coefficient of Be (m -1 )

x: the distance travelled by the X-ray in the Be window (m)

The mass absorption coefficient does not follow a linear behavior as a function of the photons wavelength so it has to be calculated for each radiation (cf. [160]

AI.3b -Solid state MAS NMR spectroscopy

The solid state MAS NMR spectroscopy measurements were performed and analyzed at the CEMHTI in Orléans by Elodie Salager and Pierre Florian.

11 B, 15 N, 25 Mg and 15 N, 25 Mg and 1 H solid-state MAS-NMR spectroscopy measurements were performed using a 4 mm CP-MAS probe.

The 15 N experiments were recorded at a magic angle spinning rate of 14 kHz under pure nitrogen at 278 K using a 1 H-15 N cross-polarization (CP) experiment with a contact time of 1.4 ms.

Spinal-64 1 H heteronuclear decoupling was applied at 50 kHz during the acquisition period (15 ms).

The recycle delay was 1 s, for a total experiment time of approximately 15 h. All spectra are referenced to a (NH4)(NO3) sample set at 22 ppm and processed with a 50 Hz exponential broadening.

The 25 Mg and 1 H variable temperature experiments were recorded at a magic angle spinning rate of 10 kHz under pure nitrogen for temperatures ranging from -25°C to +100°C.

The 1 H spectra were recorded with a rotor-synchronized Hahn echo (π/2-tau-π-tau) using a RF power of 110 kHz and a tau delay of 3 rotor periods. The recycle delay was 2 s, for a total experiment time of approximately 30 s. All spectra were referenced to pure TMS set at 0 ppm.

The 25 Mg signal was enhanced through a double-frequency sweep (DFS) followed by a rotor-synchronized Hahn echo. The Hahn echo pulses were performed using a RF power of 360 W (25 kHz for MgCl2 solution). Spinal-64 1 H heteronuclear decoupling was applied at 100 kHz during evolution (200µs) and acquisition (8 ms). The recycle delay was 1.5 s, for a total experiment time of approximately 5 hours. All spectra are referenced to a 1 mol.L -1 solution of MgCl2 set at 0 ppm.

Annex I

[161]

Baseline correction and removal of the probe background were performed using a homemade Matlab program. The spectra were fitted with the DMfit software. 6 NMR spectroscopy was used to identify the phases present in our compounds. From gave us an information on the crystallinity of the present phases as a broad peak suggest that the phase is disordered and might be amorphous. Then 15 N NMR was performed to have more information on the structure of our phases. Finally, 25 Mg and 1 H NMR were performed in temperature to confirm that Mg 2+ was the mobile specie and not H + .

AI.3c -Infrared spectroscopy

Fourrier transform infrared (FTIR) spectroscopy was performed in transmission mode. In an argon filled glovebox, the samples were mixed with dried KBr powder and pelletized into thin transparent 13 mm diameter pellets. KBr is used as IR-transparent matrix to support our sample.

The pellet is placed in an airtight sample holder that has two thick IR-transparent KCl windows.

The measurements were performed on Nicolet iS10 spectrometer (Thermo Scientific company) from 4000 to 600 cm -1 with 100 spectra recorded by measurement.

The infrared spectroscopy gave us information on the B-H and N-H bonds in the different samples. It also indicates if some impurities such as -OH groups are present in our samples. These information are given by the vibration that are observed and by their wavenumbers.

AI.4 -Thermal analysis

The thermal analysis measurements were performed by Matthieu Courty (LRCS, Amiens).

AI.4a -Differential scanning calorimetry

The differential scanning calorimetry (DSC) measurements were carried out on a Netzsch DSC 204 F1 heat flux differential calorimeter at a heating rate of 5°C.min -1 under a constant argon flow of 50 mL.min -1 . The samples were placed in aluminum crucibles which were sealed in a Ar filled glovebox and their lids were pierced under argon flow prior to the measurement to allow gas release upon heating. The onset and starting temperatures of the heat flow peaks were determined with the software Proteus Analysis. The same software was used to measure the enthalpy of each thermal phenomenon.

The DSC tells us if the samples will undergo exothermic or endothermic thermal phenomena. To determine at which temperature a thermal phenomenon occurs, three values can be defined: the starting temperature (Tonset), the temperature at the maximum of the peak (Tmax)

and the finishing temperature (Toffset). The DSC also gives the enthalpy of each phenomenon by integration of the area under each peak. This measurement was performed under Ar atmosphere in a glovebox. The mass spectrometry (MS) was coupled to the TGA measurements to identify any gas molecule that would be released.

At the entrance of the device, the gases are ionized and then separated by their mass-to-charge ratio (m/z).

plungers. The pressure in the cell is maintained by a spring placed at the back of one of the plunger.

The glass part of the cell is placed in a homemade tubular oven that heats only the middle of the cell and hence avoids the warming up of the teflon rings that could be detrimental to the sealing.

For the polarization measurements, a 10 mm diameter cold pressed pellet of a thickness of about 0. The polarization measurements with symmetrical cells were used to determine the compatibility between the anode and the solid electrolyte. A stable interface should result in a stable polarization. If it is not the case, the polarization potential will increase and evolve in

Annex I

[169] different ways depending on the phenomena occurring at the interface (dendrite or foam formation, degradation of the solid electrolyte, lost of contacts, etc.).

AI.5c -Cyclic voltammetry

The cyclic voltammetry (CV) measurements were performed with a VSP potentiostat (Biologic company) with a scanning rate of 0.05 mV.s -1 at 100°C using the high temperature electrochemical cell presented in the previous subsection. The samples were prepared in an argon filled glovebox where 20 mg of solid electrolyte were poured in a 6 mm diameter die, a first low pressure of 0.5 ton was applied only with the plunger of the die to tamp the powder. Then, the plunger was taken out and 10 mg of a composite electrode made of the solid electrolyte and carbon black (90:10 wt%) mixed by hand in an agate mortar was poured on the top of the tamped solid electrolyte in the die. A pressure of 2 tons was applied to obtain a dense pellet made on one side of the solid electrolyte and on the other side of the composite electrode (cf. Figure AI. 7). The fact to tamp the solid electrolyte in the die before to introduce the composite electrode ensure that there will be no carbon of the composite electrode passing on the other side of the pellet which would cause a short circuit. The side of the pellet made only of solid electrolyte is placed on a disc made from cold pressed Mg powder. They are stacked in the electrochemical cell where the pressure applied by a spring ensures the contact between them. As for polarization measurements, the CV measurements were hard to realize so the parameters as the voltage window or current range had to be modified depending on the cell (see chapter IV).

Cyclic voltammetry measurements were performed to measure the electrochemical stability window of our solid electrolytes. By measuring anodic and cathodic currents, it is possible to know at which potential our solid electrolytes are oxidized and/or reduced. If a measured electrochemical reaction is reversible, its coulombic efficiency can be calculated from a plot where the current is plotted as a function of time. On this plot, the area under the peak of oxidation or reduction gives the capacity of the electrochemical reaction in Ah. The coulombic efficiency of a reaction is the ratio between the measured areas of the reduction and oxidation during the same cycle.

Annex I -Experimental methods

[170]

Résumé

Cette thèse porte sur l'identification de nouveaux conducteurs d'ions Mg 2+ inorganiques pour les utiliser en tant qu'électrolyte solide dans des batteries Mg. Le choix a été fait d'étudier les composés du diagramme de phases binaire Mg(BH4)2-Mg(NH2)2. Dans ce diagramme, seuls les composés Mg(BH4)2, Mg(NH2)2 et Mg(BH4)(NH2) étaient connus. Mg(BH4)(NH2) avait été rapporté comme un bon conducteur ionique mais ses paramètres de synthèse n'étaient pas rapportés. Dans un premier temps, il a donc fallu déterminer les paramètres de synthèse nécessaires pour l'obtention de Mg(BH4)(NH2). Après avoir tester différentes vitesses de broyage, différents temps, atmosphères et températures de recuit, Mg(BH4)(NH2) a pu être obtenu. Au cours des différentes optimisations des paramètres de synthèse, une phase additionnelle inconnue fut identifiée. Cette phase a une influence importante sur la conductivité ionique des échantillons. L'hypothèse fut faite que cette phase était en fait un composé du type Mg(BH4)2x(NH2)2(1-x). A partir des paramètres de synthèse déterminés précédemment, de nouvelles synthèses furent réalisées avec différentes stoechiométries. Trois nouvelles phases cristallines du type Mg(BH4)2x(NH2)2(1-x) purent être identifiées et une tendance fut observée entre la stoechiométrie et la conductivité ionique des phases. Un maximum est atteint pour x = 0.67 avec une conductivité de 4x10 -5 S.cm -1 à 100°C, ce qui représente l'une des meilleures conductivités ioniques pour Mg 2+ jamais mesurée dans un solide inorganique. Les propriétés d'électrolyte solide des meilleurs conducteurs ioniques identifiés lors de ces travaux furent caractérisées. Il apparut que ces composés ne sont pas stables à hauts potentiels. Ils commencent à s'oxyder aux alentours de + 1.0 V vs. Mg 2+ /Mg. Néanmoins, ils sont stables vis-à-vis de Mg et permettent son dépôt et sa dissolution électrochimique.

Mots clés : batteries magnésium, électrolyte solide, hydrures complexes, conducteurs ioniques solides, stockage de l'énergie

Abstract

This PhD thesis concerns the identification of new Mg 2+ inorganic solid ionic conductors for possible use as solid electrolytes in Mg batteries. The choice was made to study the compounds from the Mg(BH4)2-Mg(NH2)2 binary phases diagram. In this phases diagram, only Mg(BH4)2, Mg(NH2)2 and Mg(BH4)(NH2) were known. Mg(BH4)(NH2) was reported previously as a good ionic conductor but its synthesis parameters were unknown. First, it was necessary to determine the synthesis parameters of Mg(BH4)(NH2). After testing different ball milling speeds, and different annealing times, atmospheres and temperatures, Mg(BH4)(NH2) could be obtained. During the optimization of the synthesis parameters, an unknown additional phase was identified. This phase has a strong influence on the ionic conductivity of the samples. The hypothesis was made that it was actually a Mg(BH4)2x(NH2)2(1-x)type compound. From the synthesis parameters determined previously, new synthesis were performed with different stoichiometries. Three new crystalline Mg(BH4)2x(NH2)2(1-x) phases were identified and a trend was observed between the composition of the products and their ionic conductivity. A maximum is reached for x = 0.67 with an ionic conductivity of 4x10 -5 S.cm -1 at 100°C, which is one of the best Mg 2+ ionic conductivity ever reported in an inorganic solid. The solid electrolyte properties of the best ionic conductors identified during this work were characterized. It appeared that these compounds are not stable at high potentials. They start to get oxidized around + 1.0 V vs. Mg 2+ /Mg. Nevertheless, they are stable against Mg and allow its electrochemical deposition and stripping.

Keywords: magnesium batteries, solid electrolyte, complex hydrides, solid ionic conductors, energy storage

Figure 1 .

 1 Figure 1. a) World electricity consumptions per year for the different regions of the world.b) Li-ion batteries sales per year for different applications.1,3

Figure 2 .

 2 Figure 2. Comparison of different battery technologies in terms of volumetric and gravimetric energy densities.4

1 .I- 3 .I- 4 .I- 5 .

 1345 chemical stability was characterized by annealing a mixture of our solid electrolytes with electrode materials that can be used in rechargeable Mg batteries (e.g. Mg metal as the anode electrode and the Chevrel phase Mo6S8 as the cathode). Then, the electrochemical stability was characterized by polarization and cyclic voltammetry experiments. Batteries: short reminder of their principle and main characteristics ....... I-2. Rechargeable Mg batteries ..................................................................... I-2a. From Li-ion to post Li-ion batteries ............................................................... I-2b. Magnesium metal anode ............................................................................. I-2c. Prototypes of rechargeable Mg batteries ...................................................... I-2d. Limitations of rechargeable Mg batteries ..................................................... Interest for solid electrolytes .................................................................. Mg 2+ inorganic solid electrolytes ............................................................. I-4a. Phosphates ................................................................................................. I-4b. Sulphides..................................................................................................... I-4c. Complex hydrides ......................................................................................... Aims of the PhD thesis ............................................................................ References .................................................................................................

Figure I. 1 .

 1 Figure I. 1. Scheme of the principle of a battery upon discharge (a) and charge (b).

  molar mass of the electrode The volumetric capacity is obtained in a similar way with the equation II.3: the density of the electrode In a full cell, the capacity is limited by the electrode having the lowest capacity. To avoid dead weight of electrode materials, the battery must be balanced to have the same capacity at both cathode and anode. The specific and volumetric capacities are expressed as: 𝑄 𝑠𝑝𝑒 𝑐𝑒𝑙𝑙 = 𝑄 𝑠𝑝𝑒 𝑐𝑎𝑡ℎ𝑜𝑑𝑒 . 𝑄 𝑠𝑝𝑒 𝑎𝑛𝑜𝑑𝑒 𝑄 𝑠𝑝𝑒 𝑐𝑎𝑡ℎ𝑜𝑑𝑒 + 𝑄 𝑠𝑝𝑒 𝑎𝑛𝑜𝑑𝑒 and 𝑄 𝑣𝑜𝑙 𝑐𝑒𝑙𝑙 = 𝑄 𝑣𝑜𝑙 𝑐𝑎𝑡ℎ𝑜𝑑𝑒 . 𝑄 𝑣𝑜𝑙 𝑎𝑛𝑜𝑑𝑒 𝑄 𝑣𝑜𝑙 𝑐𝑎𝑡ℎ𝑜𝑑𝑒 + 𝑄 𝑣𝑜𝑙 𝑎𝑛𝑜𝑑𝑒 Equation I. 4 and Equation I. 5

  Figure I. 2c).

Figure I. 2 .

 2 Figure I. 2. Illustration of some intrinsic advantages of Magnesium. a) Magnesium is a solid metal in ambient conditions. b) Standard redox potentials and c) gravimetric and volumetric capacities of different anode materials: Li, Na and Mg.

  Figure I. 3b). So the dendrite free characteristic of Mg electrodeposition has to be considered with caution because it depends on the studied system.

Figure I. 3

 3 Figure I. 3. a) Smooth electrodeposited Mg film using an electrolyte made of 0.25 M of EtMgCl-2Me2AlCl (Et = ethyl and Me = methyl) in THF (THF = tetrahydrofuran) applying 1.0 mA cm -2 in a cell made of a polished Pt working electrode and a Mg foil counter electrode. 8 b) Magnesium dendrites formed in a cell using 0.5 M of MeMgCl in THF and operating at a constant current of 0.921 mA.cm -2 in a Mg-Mg symmetrical cell. 9

  25 M of Mg[B(Bu2Ph2)]2 (Bu = butyl and Ph = phenyl) in THF/DME (DME = dimethoxyethane) solvent (cf. Figure I. 4). It was a preliminary study on materials that could be used for rechargeable Mg batteries. Only few cycles were performed and the electrolyte might have decomposed because of the high potential of cobalt oxide (2.28 V vs. Mg 2+ /Mg).

Figure I. 4

 4 Figure I. 4. a) Cycling behavior of a cell Mg/0.25M Mg[B(Bu2Ph2)]2,THF/DME/Co304.10

Figure I. 5

 5 Figure I. 5. a) Typical electrochemical behavior and the basic structure of the MgxMo6S8 cathodes, 0 < x < 2, corresponding to a maximal charge capacity of 122 mAh.g -1 . b) Performance of rechargeable Mg-MgxMo6S8 coin-cell batteries during cycling at constant currents (0.2-0.3 mA.cm -2 ). The electrolyte was 0.25 M Mg(AlCl2BuEt)2 in THF. The specific capacity of the active mass is plotted against cycle number for three cells, at different optimization levels, together with voltage profiles of representative charge-discharge cycles. 11

Figure I. 6

 6 Figure I. 6. a) Discharge and charge profile of TiS2 vs. Mg at 60°C and b) its capacity and coulombic efficiency upon cycling. 14

Figure I. 7

 7 Figure I. 7. a) The discharge profile of a Mg-S cell.b) Cycling performance of S/CMK400PEG(mesoporous carbon impregnated with S) cathode in different electrolytes: diglyme/PP14TFSI (denoted as DEGIL in blue), in tetraglyme/PP14TFSI (denoted as TEGIL ingreen), in tetraglyme (denoted as TEG in red) using CMC or PVDF as binder.16

Figure I. 9 .

 9 Figure I. 9. Illustration of Mg batteries with different uses of solid electrolytes. a) Bulk solid electrolyte. b) Solid electrolyte used as protective layer between the electrode active materials and another solid electrolyte. c) Solid electrolyte used as protective layer between the electrode active materials and a liquid electrolyte.

  Figure I. 10a). The best ionic conductor was MgZr4(PO4)6 with a conductivity of 6.1x10 -3 S.cm -1 at 800°C and an activation energy Ea of 84.18 kJ.mol -1 (0.88 eV). They described its crystallographic structure as the same as the NASICON material NaZr2(PO4)3. Later, Nomura et al. demonstrated that the structure is actually monoclinic of the β-Fe2(SO4)3-type with cell parameters a = 8.863 Å, b = 8.888 Å, c = 12.408 Å, β = 90.64 Å, Z = 4 and d = 3.265 g.cm -3 . 33 Magnesium ions randomly occupy 50% of the tetrahedral interstitial sites within the structure creating vacancies. These vacancies enable Mg 2+ conduction. Ikeda et al. performed the polarization of MgZr4(PO4)6 between two platinum electrodes at 800°C and could electro-deposit Mg through a solid for the first time. To prove this, they performed Electron Probe Microanalysis (EPMA) along the electrolyte and revealed that Mg metal was deposited on the cathodic part which also confirmed the mobility of Mg 2+ in the material (cf. Figure I. 10b). A more recent study was published by Adamu et al. In their work,

Figure I. 10

 10 Figure I. 10. a) Thermal evolution of the ionic conductivity for Mg1.1(Zr0.85Nb0.15)4(PO4)6.4 + 0.4 Zr2O(PO4)2, Mg0.7(Zr0.85Nb0.15)4(PO4)6, Mg1.4Zr4P6O1.4 + 0.4 Zr2O(PO4)2 and MgZr4(PO4)6. 38 b) EPMA elementary analysis patterns along a MgZr4(PO4)6 pellet after polarization at 800°C.[START_REF] Ikeda | Solid electrolytes with multivalent cation conduction. 1. Conducting species in Mg-Zr-PO4 system[END_REF]

Figure I. 11

 11 Figure I. 11. a) XRD patterns of the (100-x)(0.6MgS-0.4P2S5)-xMgI2 samples after 1 h of ball-milling and b) after annealing. c) Composition dependence of the conductivities at 200 °C for(100-x)(0.6MgS-0.4P2S5)-xMgI2 solid electrolytes.[START_REF] Yamanaka | Preparation of magnesium ion conducting MgS-P2S5-MgI2 glasses by a mechanochemical technique[END_REF]

Figure I. 12

 12 Figure I. 12. a) Temperature and pressure changes as a function of microwave irradiation time for MBH4 M=Li, Na, or K.b) Temperature dependence of the dielectric constants εr' and εr" ofLiBH4 at a frequency of 2.16 GHz. 48 c) Crystal structures of low temperature (LT) and high temperature (HT) phases of LiBH4. d) Thermal evolution of the ionic conductivity of LiBH4. d)7 Li solid state NMR spectra of LiBH4 at different temperatures.[START_REF] Matsuo | Lithium superionic conduction in lithium borohydride accompanied by structural transition[END_REF]

  Figure I. 13f). Their study also demonstrated that borohydrides can have good physical interfaces with the electrodes even by simple cold-pressing assembly and reach very high compactness (close to 100 % of the theoretical density, cf. Figure I. 13e

  Figure I. 14b).[START_REF] Blanchard | Nanoconfined LiBH4 as a fast lithium ion conductor[END_REF]

Figure I. 13

 13 Figure I. 13. a) Cyclic voltammogram of HT-LiBH4 sandwiched between lithium and molybdenum electrodes at 120°C with a voltage scan rate of 5 mV.s -1 . 50 b) A cross-sectional SEM image of a thin-film LiCoO2 cathode coated with a thin layer of amorphous Li3PO4. 31 c) Charge-discharge curves at 120°C of a cell using LiBH4 solid electrolyte and LiCoO2 cathode coated with Li3PO4. d) Photograph of a bulk-type all-solid-state cell Li/LiBH4/TiS2 and e) cross-sectional FIB-SEM image of the TiS2+LiBH4 composite cathode. 55 f) Charge-discharge curves at 120°C of a bulk-typeall-solid-state cell Li/LiBH4/TiS2.[START_REF] Unemoto | Stable interface formation between TiS2 and LiBH4 in bulk-type all-solidstate lithium batteries[END_REF]

Figure I. 14

 14 Figure I. 14. a) Thermal evolution of the ionic conductivity of LiBH4-LiX composites (X = Cl, Br and I) compared to LiBH4 and LiI. 57 b) Ionic conductivities of bulk (red) and nanoconfined LiBH4 (black).59

Figure I. 15

 15 Figure I. 15. a) Binary LiBH4-LiNH2 phases diagram. 61 b) Thermal evolution of the ionic conductivity of Li2(BH4)(NH2) and Li4(BH4)(NH2)3 compared to LiBH4 and LiNH2. 62

  Figure I. 16b). They concluded that the electrochemical stability of the material was up to 6 V vs Na + /Na but as it was mentioned previously, this technique might highly overestimate the electrochemical stability.

Figure I. 16

 16 Figure I. 16. a) Thermal evolution of the ionic conductivity of different materials in the NaBH4-NaNH2-NaI system. b) Cyclic voltammogram measured at 333 K for Na(BH4)0.5(NH2)0.5 sandwichedbetween sodium and molybdenum electrodes.[START_REF] Matsuo | Sodium ionic conduction in complex hydrides with [BH4]and [NH2]anions[END_REF]

64

 64 Figure I. 16. a) Thermal evolution of the ionic conductivity of different materials in the NaBH4-NaNH2-NaI system. b) Cyclic voltammogram measured at 333 K for Na(BH4)0.5(NH2)0.5 sandwichedbetween sodium and molybdenum electrodes.[START_REF] Matsuo | Sodium ionic conduction in complex hydrides with [BH4]and [NH2]anions[END_REF]

  between the nearest Mg atoms in Mg(BH4)(NH2) (3.59 Å) than in Mg(BH4)2 (4.32 Å), indicating that Mg hopping diffusion may be favored in this structure. The electrochemical stability window was also estimated up to 3 V vs Mg 2+ /Mg using cyclic voltammetry with a cell made of Mg/Mg(BH4)(NH2)/Pt (cf. Figure I. 17b). This stability is interesting for a solid electrolyte application in batteries but needs to be confirmed with proper measurements. Finally, Higashi et al. tried for the first time to assemble an all-solid-state Mg battery using Mg(BH4)(NH2) as solid electrolyte, a Mg anode and different sulfur-based cathodes (S, FeS and Ag2S, cf. Figure I. 17c).

  Figure I. 18b). Despite its good ionic conductivity, it suffers from a low electrochemical stability of 1.2 V vs Mg 2+ /Mg.

Figure I. 18

 18 Figure I. 18.a) Thermal evolution of the ionic conductivity of Mg(BH4)2, Mg(en)3(BH4)2 and Mg(en)1(BH4)2, compared with MgZr4(PO4)6, Mg1.4Zr4P6O24.4 + 0.4Zr2O(PO4)2 and Mg(BH4)(NH2).b) Cyclic voltammogram of Mg(en)1(BH4)2 sandwiched between magnesium and platinum electrodes measured at 60°C with a scan rate of 10mV.s -1 .Inset: Mg stripping charge vs. cycle number.[START_REF] Roedern | Magnesium ethylenediamine borohydride as solid-state electrolyte for magnesium batteries[END_REF]

  Figure I. 19 compares the different magnesium solid electrolytes in terms of ionic conductivity.

  , we extended the study to the Mg(BH4)2-Mg(NH2)2 binary phase diagram to identify new ionicconducting phases as what has been done with LiBH4 and LiNH2 for instance. Finally, we characterized the ionic conductors as solid electrolytes using proper techniques.

Figure I. 19 . 2 - 1 .II- 2 .II- 3 .II- 4 .

 1921234 Figure I. 19. Thermal evolution of the ionic conductivity of the different Mg 2+ inorganic conductors reported in the literature. The different families of reported materials are highlighted.The range of ionic conductivity targeted for a suitable solid electrolyte in a magnesium all-solid-state battery is also displayed.

Chapter II -

  The known parts of the Mg(BH4)2-Mg(NH2)2 binary phases diagram: Mg(BH4)2, Mg(NH2)2 and

Figure II. 1 .

 1 Figure II. 1. Crystal structures of α-, β-, γ-and δ-Mg(BH4)2 polymorphs drawn with the software VESTA.13

  Figure II. 2a) crystallizing in a cubic unit cell with space group Ia-3d with Z = 24 and the cell parameter a = 15.843(1) Å. Note that all the diffractograms measured with the Cobalt radiation present an important background due to the set-up of the diffractometer and to the beryllium window of the air-tight cell used for the measurements. As γ-Mg(BH4)2 stopped to be commercially available during this thesis, another precursor had to be used for the synthesis of Mg(BH4)(NH2). This new precursor sold as "Mg(BH4)2" by Sigma-Aldrich (> 95 % pure) was, as a matter of fact, a mixture of γ and α-Mg(BH4)2. The γ and α phases are clearly detected by X-ray diffraction (cf. Figure II. 2b). Owing to a Rietveld refinement, it was determined that about 79% of the product was γ-Mg(BH4)2 and 21 % was α-Mg(BH4)2.

Figure II. 2 . 2 ,

 22 Figure II. 2. XRD patterns of commercial a) γ-Mg(BH4)2 and its profile matching and b) of γ and α-Mg(BH4)2 and its Rietveld refinement.

Figure II. 3 .

 3 Figure II. 3. 11 B MAS NMR spectra of commercial a) γ-Mg(BH4)2 and b) γ and α-Mg(BH4)2.

Figure II. 4 .

 4 Figure II. 4. Infrared spectra of commercial a) γ-Mg(BH4)2 and b) γ and α-Mg(BH4)2.

Figure II. 5 .

 5 Figure II. 5. DSC recorded at 5°C/min under argon flow of commercial a) γ-Mg(BH4)2 and b) γ and α-Mg(BH4)2.

Figure II. 6 .

 6 Figure II. 6. Diagram summarizing the synthesis for Mg(NH2)2.

Figure II. 7 .

 7 Figure II. 7. XRD pattern after different steps of the Mg(NH2)2 synthesis. a) As-received commercial MgH2. b) MgH2 after ball-milling. c) Amorphous product after ball-milling of MgH2 under NH3 pressure. d) Crystalline Mg(NH2)2 after annealing under NH3 atmosphere.

step 1 :

 1 vacuum at 1x10 -5 bar step 2: filling with 4 bar of NH3 step 3: milling in a Spex miller for 1 h

  Figure II. 7d). Figure II. 8 shows the diffractogram of as-synthesized Mg(NH2)2 crystallizing in the space group I41/acd with Z = 32 and the cell parameters a=10.393(1) Å, c = 20.089(3) Å and V = 2170.1(5)Å 3 which are close to the one of the literature (a = 10.37 Å, c = 20.15 Å and V = 2166.87 Å 3 ).

Figure

  Figure II. 8. a) XRD pattern of synthesized Mg(NH2)2 and its profile matching. b) The crystalline structure of Mg(NH2)2 drawn with VESTA.

Chapter II -

  The known parts of the Mg(BH4)2-Mg(NH2)2 binary phases diagram: Mg(BH4)2, Mg(NH2)2 and Mg(BH4)(NH2)[60]

Figure II. 9 .

 9 Figure II. 9. Infrared spectrum of as-synthesized Mg(NH2)2.

Figure

  Figure II. 10. 15 N solid state MAS NMR spectrum of as synthesized Mg(NH2)2.

Figure II. 11 .II- 4 .

 114 Figure II. 11. DSC of as-synthesized Mg(NH2)2 recorded at 5°C/min under Ar flow.

. 1 and

 1 Figure II. 12. Diagram summarizing the synthesis route for Mg(BH4)(NH2).

Chapter II -

  The known parts of the Mg(BH4)2-Mg(NH2)2 binary phases diagram: Mg(BH4)2, Mg(NH2)2 and

Figure

  Figure II. 13. XRD patterns smoothed and normalized of the products obtained with different ball-milling and annealing parameters as described in Table II. 1.

Figure II. 14 .

 14 Figure II. 14. XRD patterns smoothed and normalized of the products after ball-milling of Mg(BH4)2 and Mg(NH2)2 at different milling speeds.

  Figure II. 15). An annealing at 120°C has been performed for 48 h under 1 bar of Ar. It allowed the crystallization of a product that was still impure and mainly made of the unknown additional phase encountered before (sample 5 on Figure II. 13). A longer annealing of 96 h was performed. It leads up to a product with more Mg(BH4)(NH2) than the additional phase and no more reactant (sample 6 on Figure II. 13). Later, the ball milling was increased to 200 rpm and the sample was annealed in the same conditions as previously (i.e. at 120°C for 96h under 1 bar of Ar). The purity of this sample was similar to sample 5 with Mg(BH4)(NH2) more present than the additional phase (sample B on Figure II. 13). Another annealing was tried at 130°C for 48h after the same ball-milling of 200 rpm. This time, the sample contained mainly the additional phase (sample A on Figure II.

(sample 7

 7 Figure II. 13). Finally, after the ball-milling at 400 rpm and subsequent annealing, the product consists only in crystalline Mg(BH4)(NH2) as revealed by X-ray diffraction (sample 9 on Figure II. 13). So the ball-milling has to be strong enough for the synthesis of pure crystalline Mg(BH4)(NH2).To know if the ball milling is strong enough, one can look at the diffraction patterns of the powders right after ball-milling. On Figure II. 14, by increasing the ball-milling speed, the amorphisation of the reactants is increased and it is once the borohydride Mg(BH4)2 is completely amorphous that the purest samples are obtained after annealing.

Figure II. 15 .

 15 Figure II. 15. Infrared spectra of the equimolar mixture of Mg(BH4)2 and Mg(NH2)2 before and after a high energy ball-milling and a low energy ball-milling.

13

 13

Figure II. 13

 13 Figure II. 13). A new crystalline additional phase was present. It was different from the additional phase observed in samples A, B and C. From the previous syntheses, we know that one of the important parameters for improving the purity of the sample was the ball-milling speed. So a synthesis with an increased ball-milling speed (500 rpm) and the same annealing conditions was performed. As expected, the product appears as purely crystalline Mg(BH4)(NH2) (sample D on Figure II. 13). The diffraction pattern of this sample and its profile matching are shown on Figure II. 16. The compound crystallizes with the space group I41 with Z = 8 and the cell parameters a =

Figure II. 16

 16 Figure II. 16. a) XRD pattern of sample D and its profile matching. b) Crystal structure of Mg(BH4)(NH2) drawn with VESTA.

11 B

 11 MAS-NMR analyses were performed on samples A, B, C and D (cf. Figure II. 17a). No peaks are detected around -20 ppm which indicate that no BxHy borane groups such as (B3H8) -, (B12H12) 2-, (B10H10) 2-and BH3 are present. 22-24 In Figure II. 17, the four 11 B NMR spectra display strong peaks Chapter II -The known parts of the Mg(BH4)2-Mg(NH2)2 binary phases diagram: Mg(BH4)2, Mg(NH2)2 and Mg(BH4)(NH2)

  II. 17d) was extracted and compared to the relative proportion of the two crystalline phases observed by XRD analyses (cf. Figure II. 13). The percentage of boron atoms corresponding to the peak at -37.6 ppm increases from sample A to sample B, so with increasing proportion of Mg(BH4)(NH2). In the same time, the percentage of boron atoms contributing to the peak at -39.6

II. 13 )

 13 , only one peak was expected on the NMR spectrum as, in the crystal structure reported for Mg(BH4)(NH2), the boron atoms occupy a single Wyckoff site. Two NMR peaks are however detected in both samples. For sample C, the first peak (52 molar % of the boron in the sample) at -37.6 ppm is similar to samples A and B and confirms the assignment to Mg(BH4)(NH2) (cf. Figure II. 17b). The second peak (48 molar %) differs from that in the other samples: it is slightly less shielded at -38.6 ppm. For sample D, the contribution at -37.7 ppm due to Mg(BH4)(NH2) (79 molar %) is also present as well as the other one at -38.8 ppm (21 mol %). For sample D, Mg(BH4)(NH2) is in larger quantity than in sample C (cf. Figure II. 17d). Additional information as crystallinity degree can be extracted from the Full Width at Half Maximum (FWHM) of the NMR peaks. The FWHM is in between 1.1-1.7 ppm for all peaks except for the additional phase in sample C and D which has a value of 3.8 and 3.0 ppm respectively (cf. Figure II. 17c). This peak broadening and the pure Gaussian shape reveal the amorphous character of the new additional phase. The very amorphous character of this additional phase explains its non-observation by XRD; even if it represents 48 molar % of the total boron atoms in sample C.

Figure II. 17

 17 Figure II.17. a) 11 B MAS-NMR spectra for the samples A, B, C and D. b) Chemical shifts and c) FWHM of the peaks attributed to Mg(BH4)(NH2) and the additional phase. d) Molar percentages of boron contributing to each peak.

Figure II. 18

 18 Figure II. 18. 15 N solid state MAS-NMR spectrum for sample D.

Figure II. 19 .

 19 Figure II. 19. Infrared spectrum of sample D.

Figure 1 Figure

 1 Figure II. 20 presents the differential scanning calorimetry (DSC) curves recorded for the samples A, B, C and D. As reported in the literature1 , an endothermic phenomenon related to an amorphization occurs between 120°C and 170°C and then the product is exothermically decomposed above 200°C. A decrease in the onset temperature of the amorphization is observed, decreasing from 152°C, 142°C and finally 120°C for samples A, B and C, respectively. Then, for sample D, the onset temperature of amorphization is at 172°C with a maximum at 183°C which is close to the one reported in the literature at 181°C.1

Figure II. 21

 21 Figure II. 21. (a) TGA coupled with DTA measurement of sample D recorded at 5°C/min. (b) Gas mass spectrometry of sample D recorded at 5°C/min showing both NH3 and H2 releases.

  FigureII. 22c. The equivalent circuit is composed of an initial resistor R0 for the device and current collectors resistance, in series with a first R1//CPE1 (R1 for the bulk solid electrolyte resistance in parallel with a constant phase element CPE1 for the non-ideal capacitance) and a second R2//CPE2 for the grain boundaries conduction (resistance and capacitance) and, finally, another constant phase element (CPE3) for the impedance of both the top and bottom electrode-electrolyte junctions to consider the roughness of the electrical contact. As described previously, R1 could be associated to the conduction of Mg(BH4)(NH2) and R2 to the grain boundaries constituted by the additional phase.

Figure II. 22 . 2 ,

 222 Figure II. 22. Nyquist plots and Bode plots at 100°C from EIS measurements of a) sample D (with a 6 mm diameter pellet) and of b) samples A, B and C (with 13 mm diameter pellets). c) Equivalent circuit model used to fit the experimental data.

Figure II. 23 .

 23 Figure II. 23. a) ln(σ.T) = f(1/T) used to calculate the activation energies of the two ionic conduction phenomena (at high and middle frequency, HF and MF. b)Real capacitance values for the two conduction phenomena as a function of the temperature.

Figure II. 24 .

 24 Figure II. 24. EIS measurements of sample D fitted with one R//CPE. a) Nyquist plot at 100°C. b) Equivalent circuit model used to fit the experimental data. c) Real capacitance as a function of the temperature. d) ln(σ.T) = f(1/T) used to calculate the activation energy.

5 .

 5 The results are plotted on Figure II. 25 as a function of 1/T. The ionic conductivity increases gradually from sample A to D. This increase can be correlated to the increase of the amount of Mg(BH4)(NH2) present in the samples. The more Mg(BH4)(NH2) is present, the highest is the conductivity. Very interestingly, the ionic conductivity reaches 1x10 -5 S.cm -1 at 100°C for sample D. This conductivity is four orders of magnitude higher than the one reported by Higashi et al.1 for the same compound (cf. black dots on Figure II. 25).

  Nevertheless, it is hard to observe this kind of glass-ceramic structure by TEM on our samples without exposing them to air during the transfer into the microscope and without damaging them with the electronic beam. The enhancement of the ionic conductivity can also be related to the study of Yan et al. in which they demonstrate that adding an excess of BH4 in Li4(BH4)(NH2)3 would create a solid solution with disorder in the structure which enables fast Li + diffusion30 . In our case, the Mg(BH4)(NH2)-based composite material could have a similar behavior but the nature of the additional phase needs to be clearly characterized to support this hypothesis. Another explanation would be that the compound studied by Higashi et al. was not pure Mg(BH4)(NH2) but contained an important amount of amorphous additional phase (> 50 molar %). Again, further information as the synthesis parameters or 11 B MAS solid sate NMR measurements on their sample would have been very useful.

Figure II. 25 . 2 ,

 252 Figure II. 25. Ionic conductivity of samples A, B, C and D as a function of 1/T. The results are compared to the measurement of Higashi et al.. 1

  other compounds with different Mg(BH4)2x(NH2)2(1-x) stoichiometries was made. All the known parts of the Mg(BH4)2-Mg(NH2)2 binary phases diagram have been described in the previous chapter, namely: Mg(BH4)2, Mg(NH2)2 and Mg(BH4)(NH2). Nevertheless, other borohydride-amide binary systems have been studied in other works. The most studied one is the LiBH4-LiNH2 binary phases diagram. Besides LiBH4, LiNH2 and Li2(BH4)(NH2), other crystalline compounds were

Figure III. 2

 2 Figure III. 2. a) 11 B solid state MAS-NMR spectra of the products obtained for different Mg(BH4)2-Mg(NH2)2 ratios. b) Chemical shifts as a function of the Mg(BH4)2 amount present in the compound. The area of the circles is proportional to the amount of each component.

9 Figure

 9 Figure III. 3. XRD pattern of α-Mg(BH4)0.40(NH2)1.60. The inset is a zoom on the (200) and (004) reflections for the space group I41/acd.

Chapter III -

  Figure III. 4. XRD patterns and Rietveld refinements of α-Mg(BH4)x(NH2)2-xfor compositions x = 0, 0.17, 0.20 and 0.25.

Figure III. 5 .

 5 Figure III. 5. Cell parameters of α-Mg(BH4)2x(NH2)2(1-x) as a function of x.

Figure III. 6 11 B

 611 Figure III. 6. a) FTIR spectra of α-Mg(BH4)x(NH2)2-x for compositions x = 0, 0.17, 0.20 and 0.25. b) Zoom on the stretching vibration bands of N-H bonds.

Figure III. 7

 7 Figure III. 7. a) 11 B MAS-NMR spectra of α-Mg(BH4)x(NH2)2-x x = 0.17, 0.20 and 0.25. b) 11 B chemical shift and c) Full Width at Half Maximum (FWHM) as a function of x.

  it is interesting to see that adding Mg(BH4)2 into the structure of Mg(NH2)2 improves its ionic conductivity. Even more, the ionic conductivity continuously increases with x as presented on Figure III. 8d.

Figure III. 8

 8 Figure III. 8. a) Nyquist plots at 100°C of α-Mg(BH4)x(NH2)2-x for x = 0.17, 0.20 and 0.25. b) Ionic conductivity as a function of 1/T. c) ln(σ.T) as a function of 1/T. d) Ionic conductivity at 100°C as a function of x.

  gradually shifts towards lower temperature as x increases. This phenomenon might be due to the organization of the (BH4) -anions into the Mg(NH2)2 structure to form the α-phase. After 200°C, an exothermic phenomenon happens. This phenomenon is related to the degradation of the products and is similar to what was observed previously with Mg(BH4)(NH2) at 200°C with the release of gases (H2 and NH3) and the formation of a very foamy solid residue (cf. Chapter II).

Figure III. 9 . 4 .

 94 Figure III. 9. DSC recorded at 5°C/min under Ar flow of α-Mg(BH4)2x(NH2)2(1-x)for x = 0.17, 0.20 and 0.25 before annealing.

  Figure III. 1)

Figure III. 11 III- 4 .

 114 Figure III. 11. a) DSC recorded at 1°C.min -1 under Ar flow upon heating up to 170°C and cooling of Mg(BH4)x(NH2)2-x for x = 0.25. b) XRD patterns of the sample before and after annealing at 170°C for 12h under Ar.

  Figure III. 12. XRD patterns of β-Mg(BH4)0.5(NH2)1.5 compared with β-Mg(BH4)0.67(NH2)1.33 and sample A.

Figure III. 13

 13 Figure III. 13. a) Nyquist plots at 100°C of β-Mg(BH4)0.5(NH2)1.5. b) Ionic conductivity as a function of 1/T. c) ln(σ.T) as a function of 1/T.

Chapter III -

  The exploration of the unknown parts of the Mg(BH4)2-Mg(NH2)2 binary phases diagram[108]

Figure III. 14

 14 Figure III. 14. a) Nyquist plots at 100°C of β-Mg(BH4)0.67(NH2)1.33. b) Ionic conductivity as a function of 1/T. c) ln(σ.T) as a function of 1/T.

  Figure III. 9 and Figure III. 15). An endothermic phenomenon occurs at low temperature, 47°C. Its enthalpy of 14 J.g -1 (0.8 kJ.mol-1, cf. Table III. 6) is closed to the ones described previously. The DSC curve shows an important noise after 200°C due to the decomposition of the sample.

Figure III. 15 III- 5 .

 155 Figure III. 15. a) DSC recorded at 5°C/min under Ar flow of β-Mg(BH4)0.67(NH2)1.33 before annealing. b) TGA coupled with DSC recorded at 1°C/min under Ar flow of β-Mg(BH4)0.67(NH2)1.33 after annealing.

Figure III. 17

 17 Figure III. 17. a) XRD patterns and profile matching of γ-Mg(BH4)2x(NH2)2(1-x) for x = 0.67, 0.71 and 0.75. b) Cell parameters and unit cell volume as a function of x.

Figure III. 18 .

 18 Figure III. 18. Cell parameters of γ-Mg(BH4)x(NH2)2-x as a function of x.

Figure III. 19

 19 Figure III. 19. a) Rietveld refinement of γ-Mg3(BH4)4(NH2)2 and b) its structure drawn with the software VESTA.

Figure III. 20 .

 20 Figure III. 20. IR spectra of γ-Mg(BH4)2x(NH2)2(1-x) for x = 0.66, 0.71 and 0.75.

  Figure III. 21d).These values are the highest ever recorded for an inorganic Mg 2+ solid ionic conductor at such low Chapter III -The exploration of the unknown parts of the Mg(BH4)2-Mg(NH2)2 binary phases diagram[116]

Figure III. 21

 21 Figure III. 21. a) Nyquist plots at 100°C of γ-Mg(BH4)x(NH2)2-x for x = 0.67, 0.71 and 0.75. b) Ionic conductivity as a function of 1/T. c) ln(σ.T) as a function of 1/T. d) Ionic conductivity at 100°C as a function of x.

Figure III. 22

 22 Figure III. 22. a) 25 Mg and b) 1 H solid state MAS NMR spectra of γ-Mg(BH4)1.42(NH2)0.58 at different temperatures.

Figure III. 23 .

 23 Figure III. 23. DSC recorded at 5°C/min under Ar flow of γ-Mg(BH4)2x(NH2)2(1-x) for x = 0.67, 0.71 and 0.75 before annealing.

Figure III. 24 .

 24 Figure III. 24. TGA coupled with DTA recorded at 1°C/min under Ar flow of γ-Mg(BH4)2x(NH2)2(1-x) for x = 0.67, 0.71 and 0.75.

Figure III. 25 )

 25 FigureIII. 25). The new Mg(BH4)2x(NH2)2(1-x) phases found and described above are included in this diagram. The content of the unexplored regions are assumed from our experimental observations. On both extreme parts are the magnesium amide and borohydride, and in the center Mg(BH4)(NH2) which decompose into an amorphous compound around 180°C. Between, Mg(NH2)2 and Mg(BH4)(NH2), two new phases have been observed: the α and the β phase. The α-phase is a solid solution crystallizing in the Mg(NH2)2 tetragonal structure (s.g.: I41/acd) which expands as (NH2) -anions are substituted by larger (BH4) -. It undergoes a phase transition at about 150°C and turns into the β-phase. The β-phase was observed up to x = 0.33 but seems to be less stable as x is getting close to 0.5. On the other side, between Mg(BH4)2 and Mg(BH4)(NH2), only one new phase was observed: the γ-phase. This phase is also a solid solution crystallizing in a tetragonal structure with the space group I-4. Its structure also expands as x increases.

Figure III. 25 .

 25 Figure III. 25. Pseudo binary Mg(BH4)2-Mg(NH2)2 phases diagram. The full areas are the known ones and the shaded ones are assumed.

IV- 3 .IV- 4 .

 34 IV -Characterization of ionic conductors as solid electrolytes: towards all-solid-state Mg batteries Table of contents IV-1. Introduction ....................................................................................... IV-2. Chemical stability ............................................................................... IV-2a. Selection of electrode materials ............................................................... IV-2b. Chemical stabilities of Mg(BH4)(NH2) and γ-Mg3(BH4)4(NH2)2 ..................... Electrochemical stability ..................................................................... III-3a. Anode/electrolyte compatibility............................................................... III-3a. Electrochemical stability window ............................................................. Conclusions ........................................................................................ References ..................................................................................................

  1. The XRD pattern of the Mg powder obtained after ball milling is presented on Figure IV. 1a. The diffractogram indicates that no crystalline impurities are present: only Mg metal is visible, crystallizing in a hexagonal structure with the space group P63/mmc and the cell parameters a = 3.210(1) Å and c = 5.211(2) Å.

  additive. Carbon black C45 (Imerys) was used. It has a specific surface area of 45 m².g -1 . Before use, C45 is placed under dynamic vacuum at 750°C overnight to remove all molecules adsorbed onto the carbon surface. It is then stored in an Ar filled glovebox. Carbon black consists in a poorly graphitized carbon and this low crystallinity causes a broad diffraction peak around 26° as observed on the XRD pattern on Figure IV. 1d.

Figure IV. 1 .

 1 Figure IV. 1. XRD patterns of the different electrode materials: a) Ball-milled Mg powder, b) Bi powder, c) Mo6S8 and d) carbon C45.

2

  and Figure IV. 3), and with the carbon additive C45 (cf. Figure IV. 4), only the reflections of the solid electrolytes and the electrode materials are visible. The refined cell parameters of the Mg(BH4)(NH2) and γ-Mg3(BH4)4(NH2)2 change slightly compare to the ones obtained in the previous chapters.

Figure IV. 2 .

 2 Figure IV. 2. XRD patterns of the 1:1 mixture of Mg with a) Mg(BH4)(NH2) and b) γ-Mg3(BH4)4(NH2)2 after respective annealing at 100°C and 150°C for 12h under Ar.

Figure IV. 3 .

 3 Figure IV. 3. XRD patterns of the 1:1 mixture of Bi with Mg(BH4)(NH2) after annealing at 100°C for 12h under Ar.

Figure IV. 4 .

 4 Figure IV. 4. XRD patterns of the 1:1 mixture of C45 with a) Mg(BH4)(NH2) and b) γ-Mg3(BH4)4(NH2)2 after respective annealing at 100°C and 150°C for 12h under Ar.

Figure IV. 5 .

 5 Figure IV. 5. XRD patterns of the 1:1 mixture of Mo6S8 with a) Mg(BH4)(NH2) and b) γ-Mg3(BH4)4(NH2)2 after respective annealing at 100°C and 150°C for 12h under Ar.

Figure IV. 6 .

 6 Figure IV. 6. Galvanostatic measurement at 100°C of a symmetrical cell Mg/Mg(BH4)(NH2)/Mg cycled with a current of 5 µA.cm -2 for 2h with 5 min rest between each polarization.

Figure IV. 7 .

 7 A Rietveld refinement confirmed that a mixture of Mg3Bi2 + Bi was obtained with 46 mass % of Mg3Bi2 and 54 mass % of Bi. With this mixture, only the biphasic redox reaction (cf. Equation IV.1) would occur at both electrodes.

Figure IV. 7 .

 7 Figure IV. 7. XRD pattern of the as-prepared Mg3Bi2 + Bi electrode material and its Rietveld refinement.

  Figure IV. 9a and b show the results for Mg(BH4)(NH2) and γ-Mg3(BH4)4(NH2)2, respectively. For both phases, the Mg deposition/stripping happens around 0 V vs. Mg 2+ /Mg. No reduction other than the Mg deposition occurred so these solid electrolytes are stable at low potentials. This was confirmed by the chemical stability experiments described in the previous section. On the other side, the oxidation of the solid electrolytes was observed. It started in both cases around 1 V vs Mg 2+ /Mg and had an onset potential of about 1.49 V vs Mg 2+ /Mg for Mg(BH4)(NH2) and 1.48 V vs Mg 2+ /Mg for γ-Mg3(BH4)4(NH2)2 (cf. Figure IV. 9c and d). The composition does not seem to have a strong influence on the ESW. In the case of Mg(BH4)(NH2),

  discrepancy between the ESW measured for Mg(BH4)(NH2) and the results of Higashi et al.. With an oxidation starting at about 1 V vs Mg 2+ /Mg, it is obvious that these solid electrolytes are unstable versus cathode electrodes as Mo6S8 that has a working potential around 1.2 V vs Chapter IV [141] Mg 2+ /Mg. This explains why the solid electrolytes are oxidized by Mo6S8 to form Mg2Mo6S8 and amorphous decomposition products.

Figure IV. 9 .

 9 Figure IV. 9. Cyclic voltammetry measurements of cells made of Mg/solid electrolyte/solid electrolyte + C45 (90:10) recorded at 0.05 mV.s -1 at 100°C with a) Mg(BH4)(NH2) and b) γ-Mg3(BH4)4(NH2)2. The insets are the coulombic efficiencies of the Mg deposition/stripping. c) and d) are the zoom on the first oxidation observed on a) and b) respectively.

11 B

 11 Mg(BH4)(NH2) (cf. Chapter II). It has been synthesized by ball-milling of Mg(BH4)2 and Mg(NH2)2 and a subsequent annealing. The synthesis parameters were not described in the literature so they had to be determined. Different ball-milling speeds and annealing times, temperatures and atmospheres were tested. The purest sample was obtained with a ball-milling at 500 rpm for 1h and a subsequent annealing at 120°C for 72h under 10 bar of H2. The X-ray diffraction pattern of this sample showed only Mg(BH4)(NH2), crystallizing in the space group I41 with Z = 8 and the cell parameters a = 5.768(5) Å, c = 20.486(2) Å, and V = 681.6(1) Å 3 which are in agreement with the literature.1 During the synthesis optimization, composite samples made of Mg(BH4)(NH2) and an unknown crystalline additional phase were obtained. 11 B MAS-NMR spectroscopy experiments, that were performed for the first time on these compounds, revealed that even in samples presenting only crystalline Mg(BH4)(NH2) by X-ray diffraction, the additional phase was nevertheless present as an amorphous phase. The additional phase content determined by MAS-NMR spectroscopy strongly influences the ionic conductivity. The sample containing

  The γ-phase has also a good thermal stability as no thermal event nor mass loss was detected below 190°C. A pseudo-binary Mg(BH4)2-Mg(NH2)2 phases diagram was proposed. It helps us to understand the complexity of this Mg(BH4)2-Mg(NH2)2 system. A trend was observed for the ionic[149] conductivity of the compounds as a maximum is reached for γ-Mg3(BH4)4(NH2)2. This trend continues even if the crystal structure changes. So, good Mg 2+ ionic conductors were found in this binary phases diagram but they still needed to be characterized as solid electrolyte in order to know if they could be used in Mg batteries. Two of our best ionic conductors found in the Mg(BH4)2-Mg(NH2)2 phases diagram have been evaluated as possible solid electrolytes: Mg(BH4)(NH2) and γ-Mg3(BH4)4(NH2)2 (cf. Chapter IV). Interestingly, they do not react with the carbon additive and the anode materials: Mg and Bi. However, they reacted with the Mo6S8 cathode material. In contact with Mo6S8, they were oxidized to form Mg2Mo6S8 and amorphous decomposition products. The compatibility of Mg(BH4)(NH2) with the Mg anode was also tested by a polarization measurement in a symmetrical cell using Mg electrodes. A large interfacial resistance was observed. It could be explained by a poor contact between the electrolyte and Mg, especially because Mg is a hard metal and sticks very poorly on the solid electrolyte pellet. The interfacial resistance increased upon cycling indicating its instability upon Mg deposition/stripping. The same experiment was performed with "Mg3Bi2 + Bi" electrodes, but unfortunately the interfacial resistance was even larger than that with Mg. The interfacial resistance increased upon cycling, which could be caused by the important volume change between Mg3Bi2 and Bi. The electrochemical stability windows of the two solid electrolytes were measured by cyclic voltammetry. Magnesium deposition/stripping is observed around 0 V vs. Mg 2+ /Mg, which represents a very positive result as it has been very rarely reported for a solid Mg 2+ ionic conductor. The oxidation of the solid electrolytes starts at about + 1.0 V vs. Mg 2+ /Mg and had an onset potential around + 1.5 V vs. Mg 2+ /Mg. The oxidation potential

(

  Mg2P2S6: C2/m, Z = 2, a = 6.085(1) Å, b = 10.560(2) Å, c = 6.835 (1) Å, β = 106.97(3)° and Zn2P2S6: C2/m, Z = 2, a = 5.9717(4) Å, b = 10.3424(8) Å, c = 6.7565(5), β = 107.139(4)°). So, inorganic solid electrolytes for Mg batteries are still worth of investigation and there are plenty of rooms to improve their ionic conductivity/mobility and stability. The implementation of solid electrolytes in Mg batteries would help to develop systems with high energy densities such as the Mg-S batteries and, thus, to bring more options to the world of energy storage.

  Figure AI. 1).

Figure AI. 1 .

 1 Figure AI. 1. Picture taken while pouring deionized water on the remaining powder in a jar used to ball-mill a mixture of Mg(BH4)2 and Mg(NH2)2. It illustrates the reactivity of these compoundsand that they can lead to explosions if not manipulate with care.

Figure AI. 2 .

 2 Figure AI. 2. Picture of the Spex jar equipped with a gas inlet.

Figure AI. 3 .

 3 Figure AI. 3. Picture of the stainless steel reactor used for our syntheses.

Figure AI. 4 )

 4 photons wavelength so it has to be calculated for each radiation (cf. Figure AI. 4). The National Institute of Standards and Technology (NIST) provides a table corresponding to the mass absorption coefficient of Be for different photon energies.5 It is from this table that can be calculated the mass absorption coefficient corresponding to the photons from Co (λ = 1.78897 Å, i.e. E = 6935 eV) and Cu (λ = 1.54060 Å, i.e. E = 8053 eV) sources:

Figure AI. 4 .Figure 3 .

 43 Figure AI. 4. Mass absorption coefficient of Be for different values of photon energy.5

11 B

 11 NMR spectroscopy, the measured chemical shift can be compared with the literature to know what is the environment around the B atoms. The full width at half maximum (FWHM) of the peaks

Figure 4 .-

 4 Figure 4. Illustration of a DSC peak with its characteristic temperatures.

  5 mm is placed between two 10 mm diameter discs of cold pressed Mg powder. In this configuration, the open circuit voltage (OCV) should be close to 0 V vs Mg 2+ /Mg. Nevertheless, for the first tries, the OCV was very unstable as shown on Figure AI. 6b and the measure was very noisy (cf. Figure AI. 6c).It was observed that the spring used previously was actually too stiff so when the cell was tight, the strength of the spring was large enough to push one of the swagelock parts that slid along the glass tube. It leads to the decompression of the cell making the contacts between the different parts of the stack very bad and that is the reason why the signal was so noisy. To avoid this, the initial spring was replaced by a less stiff one and it results in a perfect conservation of the mechanical pressure applied to the stack. After this change, the OCV of the cells was much more stable and the measurement had almost no noise (cf. Figure AI. 6b and d).

Figure AI. 6

 6 Figure AI. 6. a) Design of the high temperature electrochemical cell. b) Typical open circuit voltage (OCV) before and after optimization by changing the spring of the electrochemical cell. Polarization measurement c) before and d) after optimization.

Figure AI. 7 .

 7 Figure AI. 7. Scheme of the assembling used to measure the electrochemical stability window by cyclic voltammetry.

-3 S.cm -1 at 725°C. They determined an activation energy of 0.84 eV which is in agreement with the value previously

  reported by Ikeda et al. (0.88 eV).

	In the late 90's, another Japanese team from the Osaka University continued the work
	started by Ikeda et al. and tried to improve the ionic conductivity of MgZr4(PO4)6. First, Imanaka
	et al. added excess of Zr2O(PO4)2. 35,36 They obtained composite materials made of Mg1+xZr4P6O1+x +
	x Zr2O(PO4)2. They found an optimum composition in terms of ionic conductivity for x

= 0.4 reaching 6.92x10 -3 S.cm -1 at 800°C and

	Chapter I
	from 93.1 % of the theoretical compactness for x = 0 to 99.3 % for x = 0.4. Nevertheless, the
	presence of a too large excess of Zr2O(PO4)2 would be detrimental for the ionic conductivity as this
	material is an insulator. The Zr2O(PO4)2 presence also explains the increase in term of activation
	energy as compared to pure MgZr4(PO4)6. As Ikeda et al., they demonstrated the deposition of Mg
	metal by EPMA after polarization of their composite material. Imanaka et al. worked also on the
	substitution of Zr by Nb in MgZr4(PO4)6. 37 They could synthesize a Mg1-2x(Zr1-xNbx)4(PO4)6 solid
	solution and observed the shrinkage of the cell as the radius of Nb 5+ (0.78 Å) is smaller than the
	radius of Zr 4+ (0.86 Å). The optimum composition in terms of conductivity was for x = 0.15 with a
	conductivity of 3
	Ea = 135.8 kJ.mol -1 (1.41 eV). This value is the same than
	the one obtained by Ikeda et al. for pure MgZr4(PO4)6 but it is one order of magnitude higher than
	the one they obtained in the same study for MgZr4(PO4)6 (cf. Figure I. 10a). They explained this
	result by the improvement of the compactness due to the presence of Zr2O(PO4)2. It increased
	[29]

.7x10 -3 S.cm -1 at 750°C and

	Ea ~ 80 kJ.mol -1 (0.83 eV). Whithout any information
	about the compactness, it remains difficult to know if the slight improvement of conductivity
	compared to MgZr4(PO4)6 is due to a lower porosity or to the formation of Mg vacancies. Later,
	they synthesized composite materials Mg0

.7+x(Zr0.85Nb0.15)4(PO4)6+x + xZr2O(PO4)2 and got a maximum conductivity of 9.53x10 -3 S.cm -1 at 800°C for x = 0.4. 38

  To better understand the Mg 2+ conduction process in the materials described above, Kawamura et al. performed 31 P solid state

  1 H solid-state MAS-NMR spectroscopy measurements were performed using a 20 T Avance III Bruker spectrometer operating at 272.7 MHz, 86.16 MHz, 52.05 MHz and 850 MHz, respectively. All samples were handled and packed in rotors in an Ar filled glovebox. MAS probe. The spectra were obtained at a magic angle spinning rate of 30 kHz under pure nitrogen at 278 K. The one pulse11 B experiments were recorded with a RF power of 50 kHz and a short excitation pulse (<π/18). Spinal-64 1 H heteronuclear decoupling was applied at 100 kHz during acquisition (8 ms). All spectra are referenced to a 1 mol.L -1 solution of B2O3 set at 20.5 ppm and processed with a 3 Hz exponential broadening.

	11 B solid-state MAS-NMR spectroscopy measurements were performed using a 2.5 mm
	CP-

Table of contents

Résumé étendu de la thèse (french summary)

Acknowledgements

PhD thesis and for their guidance. Thanks to the RS2E for the funding of this work through the Labex project STORE-EX.

Then, I would like to thank the members of my PhD committee Maximilian Fichtner (Helmholtz Institute Ulm), Arndt Remhof (EMPA, Zurich), Montse Casas-Cabanas (CIC energigune, Vitoria-Gasteiz) and Christian Masquelier

[83]

The nature of the additional phase remained unclear but its strong influence on the properties of the as-synthesized Mg(BH4)(NH2) is interesting. The first assumption was that this additional phase was an Mg(BH4)2-x(NH2)x substoichiometric compound. Indeed, it is not a decomposition product and it was encountered in large proportion in samples for which the reaction was not fully completed. From 11 B MAS-NMR spectroscopy, we observed that it had slightly different environments for the BH4 groups as compared to Mg(BH4)(NH2). Unfortunately, no compound of such composition has ever been reported so far. Therefore, we decided to explore the Mg(BH4)2-Mg(NH2)2 binary phases diagram to try to identify this additional phase.

III-2. First exploratory syntheses with different Mg(BH4)2:Mg(NH2)2 ratios

The unknown phases observed during the synthesis of Mg(BH4)(NH2) were actually assumed as never reported Mg(BH4)2x(NH2)2(1-x) compounds. The last sample for which one of these phases appeared was sample 12 (cf. Chapter II). So the same annealing and ball-milling parameters were used but the Mg(BH4)2:Mg(NH2)2 molar ratio was changed. As a reminder, the synthesis parameters were a ball-milling at 400 rpm for 1h and a subsequent annealing at 120°C for 72h under 10 bar of H2. The XRD patterns and the 11 The α-phase is the richest in Mg(NH2)2. By simple observation of the XRD pattern of the sample obtained for the ratio 1:4 (x = 0.20), it appears that the reflections of this new phase are similar to the ones of Mg(NH2)2 but with an important shift towards lower angles, i.e. larger reticular distances. It is the typical behavior of a solid solution. The (BH4) -are substituting the (NH2) -in the structure of Mg(NH2)2 and the shift towards lower angles is due to an expansion of the cell as (BH4) -anions are larger than (NH2) -(bond distances: B-H ~1.2 Å and N-H ~0.9 Å, see chapter II). The 11 B NMR spectrum indicates that there is only one type of BH4 in this compound with a chemical shift at -40.2 ppm. This chemical shift is specific to this new phase.

The β-phase crystallizes for ratios 1:3 (x = 0.25) and 1:2 (x = 0.33). For the ratio 1:2 (x = 0.33), important X-ray diffraction peaks of Mg(BH4)(NH2) are also observed but by 11 B NMR spectroscopy, it was probed that only 5 % of the product is made of crystalline Mg(BH4)(NH2) (chemical shift at -37.5 ppm, see chapter II). Another broad component can be fitted at -38.7 ppm, representing 51 % of the signal. This broad component is similar to the amorphous additional phase observed in sample C and D (see chapter II). The remaining 44 % is attributed to the crystalline β-phase with a shift at -39.5 ppm. For the ratio 1:3 (x = 0.25), only few crystalline Mg(NH2)2 reflections are observed by XRD and the other diffraction peaks can be attributed to the β-phase. It is probed by 11 B NMR spectroscopy that a tiny amount of Mg(BH4)(NH2) remains with a contribution at -37.3 ppm (2% of the signal). There is still a broad contribution that represents 60 % of the signal but it is shifted to -39.2 ppm. The peak corresponding to the β-phase (-39.3 ppm) is representing 38 % of the signal so its amount decreased even if by XRD, it seems well crystallized. This last observation may suggest that the broad contribution could be not only due to an amorphous phase but also to a disorder in the crystalline structure. The γ-phase is rich in borohydride. From XRD measurements, it is identical to the unknown phase appearing in addition to Mg(BH4)(NH2) in sample 12 (see chapter II). It is confirmed by 11 B NMR spectroscopy as the spectrum of this sample can be fitted with a component at -37.8 ppm which is due to Mg(BH4)(NH2) and one at -39.3 ppm which is the chemical shift observed for the γ-phase. Note that in this sample, these two components represent respectively 27 and 21 % of the signal and that the remaining 52 % are due to an amorphous phase with a chemical shift at -38.4 ppm. For the ratio 2:1 (x = 0.66), some weak diffraction peaks due to Mg(BH4)(NH2) can be observed. Then, crystalline α-Mg(BH4)2 starts also to be observed for the ratio 3:1 (x = 0.75) and even more clearly for the ratio 4:1 (x = 0.8). It was assumed that the γ-phase was a defined compound with a composition between the 2:1 ratio (x = 0.66) and the 3:1 ratio (x = 0.75). So a synthesis with a 2.5:1 Mg(BH4)2-Mg(NH2)2 ratio (x = 0.71) was done and a sample with solely the γ-phase crystallizing was obtained. Nevertheless, the diffraction pattern is evolving with the composition as the α-phase, so the γ-phase behaves like a solid solution. By looking at the 11 B NMR spectra of the composition 2:1 (x = 0.66), 2.5:1 (x = 0.71), 3:1 (x = 0.75) and 4:1 (x = 0.80), it can be observed that all of them exhibit mainly one BH4 environment with a chemical shift at about -39.3 ppm which is close to the chemical shift of pure Mg(BH4)2 (-39.6 ppm, see chapter II). Another shift around -41.4 ppm is also slightly detectable (< 7 % of the signal). It might be due to the presence of a Mg(NH2)2-rich phase as the α-phase, as a consequence of an incomplete reaction.

So, three new crystalline phases have been identified by XRD, each of them has a characteristic chemical shift for the (BH4) -anion as determined by 11 B MAS-NMR spectroscopy.

From the NMR spectra, it has been possible to determine the range of compositions where each new phase can be formed. In the next sections, the syntheses will be optimized in order to try to obtain pure phases and to perform further characterizations especially for solving their crystal structures and for measuring their ionic conduction properties.

III-7. Conclusions

In this chapter, the discovery of three new crystalline phases in the Mg(BH4)2-Mg(NH2)2 binary phases diagram has been reported. These phases behave as solid solutions with different

Two phases, named the α-phase and the β-phase, have a composition rich in amide. The αphase crystallizes with the same structure as Mg(NH2)2 in which (BH4) -anions are substituting the (NH2) -. This substitution leads to an expansion of the unit cell and an improvement of the ionic conductivity at 100°C, increasing from 4x10 -9 S.cm -1 for x = 0.17 to 1x10 -8 S.cm -1 for x = 0.25. Upon heating, the α-phase undergoes an irreversible endothermic structural transition, forming the βphase. The β-phase crystallizes in a new unknown crystal structure and might be the one crystallizing in sample A in addition of Mg(BH4)(NH2) (cf. Chapter II). This phase transition has a little impact on the ionic conductivity of the material, σ100°C = 2x10 -8 S.cm -1 , and the activation energy of the ionic conductivity increases from 1.7 eV to 2 eV. For x = 0.33, with a ball-milling at 600 rpm, poorly crystallized β-phase was obtained with an amorphous background visible by XRD.

So, it might be less stable around this stoichiometry. The ionic conductivity at 100°C of this last compound is improved, as it reaches 2x10 -7 S.cm -1 and its activation energy is 1.6 eV.

The phase crystallizing in the borohydride-rich region was named the γ-phase. It crystallizes in a tetragonal structure with the space group I-4. Similarly to the α-phase, adding more BH4 into this compound increases the unit cell volume but decreases the ionic conductivity. The γ-phase exhibits high ionic conductivity values that decrease with x, while the activation energy increases.

The best ionic conductor is γ-Mg(BH4)1.33(NH2)0.67 (which can be also written γ-Mg3(BH4)4(NH2)2)

with an ionic conductivity at 100°C of 4x10 -5 S.cm -1 and an activation energy of 0.8 eV. These values mean that it features one of the best Mg 2+ mobility in an inorganic solid ionic conductor reported so far. The good mobility of Mg 2+ was confirmed by 25 Mg solid state MAS-NMR spectroscopy with a significant narrowing of the NMR peak upon heating. The γ-phase has also a good thermal stability as no thermal event nor mass loss was detected below 190°C. (calculated from the crystallographic density: 0.9966 g.cm -3 ). The pellets were simply sandwiched between two Papyex foils. Later, the pellets were prepared with a 6 mm diameter die in which they were cold pressed directly between two Papyex foils. With this method, the Papyex foils are stuck to the solid electrolyte. It improves a lot the contacts and thus the quality of the measurements. It also needs less powder as only 30 mg of powder was pressed under 2 tons for a final thickness of about 1.1 mm, so also a density close to 95% in the case of Mg(BH4)(NH2).

The following equation was used to determine the ionic conductivity of our materials from EIS measurements:

with σ: is the conductivity (S.cm While the impedance of a CPE is expressed as:

with Y0: a prefactor n: a factor comprised between 0 and 1, depending on the surface roughness or a nonuniform current distribution Before to start fitting, the model needs to be fed with approximate values for R1 and CPE1. If R1 is difficult to approximate from the Nyquist plot, it can be extracted from the Bode plot. Indeed, the maximum on the Bode plot corresponds to the end of the semicircle (cf. Figure AI. 5b).

Knowing the frequency at which is the maximum on Bode plot allows to find the point corresponding to this frequency on the Nyquist plot and thus to find R1. For the capacitance value, the point at the top of the semi-circle has to be used. At this specific point, the cut-off frequency is reached. This frequency is linked to the resistance R1 and the capacitance C1 by the relation: From Equation AI. 6, C1 can be calculated knowing R1 and fcut-off. This value of C1 can be used as starting parameter for the CPE1, putting Y0 = C1 and n = 1. Annex I

[167]

From Equation AI. 7 and Equation AI. 8 the relationship between the parameters of the CPE (Y0 and n) and the real Capacitance (C) is obtained:

Equation AI. 9

The value of the real capacitance will change depending on the conduction phenomenon happening. The ionic conduction in the bulk and in the grain boundaries can be differentiated for instance as it is described in Chapter II.

The ionic conductivity should follow an Arrhenius law. Plotting the logarithm of the ionic conductivity values as a function of the reverse of the temperature gives us access to the activation energy of the ionic motion in our solid electrolyte. Indeed, the ionic conductivity in a solid electrolyte is given by the equation:

with σ0: a constant (K.S.cm -1 )

Ea: the activation energy of the ionic conductivity (eV)

A linear fit of ln(σ.T) as a function of 1/T give us access to the activation energy Ea from the slope which is equal to Ea/kB.

AI.5b -Polarization

Galvanostatic cycling with potential limit (GCPL) was performed using a VSP potentiostat (Biologic company) equipped with a low current channel allowing us to work with current as low as the nA. The current was applied with a positive and negative alternatively for 2h and each polarization was separated by a 5 min rest. This resting time is used to remove the accumulate charges at the interfaces. The assembly of the stacks was hardly reproducible so the current density was set at a value for which the polarization would be at reasonable potential that is the reason why it was changed from a cell to another. As our electrolytes have interesting ionic conductivities only at elevated temperature, it was necessary to perform electrochemical measurements at 100°C with a specific cell. A high temperature electrochemical cell has been developped in our lab. 7 This cell is based on the design of swagelock cells. It is made of a glass tube that is closed on both side with swagelocks cells using teflon rings for sealing (cf. Figure AI. 6a).

The sample is placed in the middle of the glass tube where it is squeezed between two long AII.1 -Neutron diffraction of the α-phase for x = 0.25

Annex II -Extra data

Neutron diffraction was used to try to differentiate the position of (BH4) -to the one of (NH2) - anions in the structure of the α-phase obtained for x = 0.25 (cf. Chapter III). Indeed, from powder X-ray diffraction, it is difficult to differentiate the positions of these two anions, especially because they contain the same number of electrons (9 electrons). By neutron diffraction, B and N should be clearly identified as they have a large difference of neutron absorption, 679.977 and 1.684 barn respectively.

The samples were sent to the ILL (Institut Laue-Langevin) in Grenoble and the diffraction measurements were performed by Emmanuel Suard with the beam line D2B (high resolution neutron powder diffractometer). The wavelength was 1.594 Å with a high flux resolution. The sample was placed in a 5.5 mm diameter vanadium tube sealed in an Ar filled glovebox.

Unfortunately, the measurement resulted in a quasi-total absorption of the neutrons and no diffraction was obtained (cf. The β-phase obtained after annealing at 170°C of the α-phase for x = 0.25, crystallizes in an unreported crystalline structure. The only symmetry found for the β-phase was monoclinic. On