Alessandro Mauro

Nono, Youcef Krishna Taki

Abdel Tom

Shu Chen

I would like to thank to Parviz Hajıyev for his valuable

This thesis study was often exciting, sometimes difficult and rarely stressful for me. I have always felt that I have been so lucky because I have had a great thesis director and talented supervisors. I would like to thank my thesis director, Sebastien Martinet, for his valuable insights and fruitful discussions during my project meetings. I always found him as a very professional and gentleman person. His vision and attitude also guided me how to become a charismatic leader. I would also like to thank to my project supervisors, Jean François Colin and David Peralta for their valuable supports during my thesis. Jean François's deep knowledge in crystalline materials and David's practical synthesis ideas shaped this project as it is today and produced valuable outcomes. With their supervision, I learned how to continue, where to stop and when I should push forward. I always admired their respect to individuals, thinking style and way of criticism. Though I have not seen them before my interview, the harmony between us was always great. I thank again for their leadership in my project and I will never forget the time I spent with them.

As like other PhD theses, we found interesting materials, but then struggled in the sea of unknowns. At that time, Prof. Claude Delmas, who is a precious visitor of CEA LITEN, shared his rich knowledge with us and showed different aspects of the problems. I would like to thank him for our fruitful discussions during our team meetings.

When I needed a practical information or opinion about any method or knowledge, I have got some hints from my lovely office mate, Carole Bourbon. I would like to thank her for her joyful attitude and the great time that we spent together. I wish to her happy retirement in 2020! I would also like to thank Anass Benayad, for his significant help for XPS measurements and discussions on the results, Laurent Gonon for his help for Raman Spectroscopy, Celine Barchasz and Alice Robba for the discussions about lithium sulfur system. I would also like to extend my thanks to all my colleagues in LITEN/STB, Sebastien Launois, Vasily Tarnopolskiy, Ognjen Hajndl, Camille Guerin, Laurent Bernard, Guillaume Lefevre and Guillaume Tonin, Gaelle Besnard, Damien Guegan, Benjamin Mercier and Oumaima Chaouachi for their friendship and personal support during my thesis.

General Introduction General Introduction

In recent years, energy requirements have increased in-line with population growth in the world. To preserve our resources and to create a sustainable world that is less dependent on fossil fuels, application of renewable energy technologies has been expanded. However, sustainability and efficiency of these technologies highly depend on the success of energy storage systems [START_REF] Armand | Building Better Batteries[END_REF][START_REF] Placke | Lithium Metal, and Alternative Rechargeable Battery Technologies: The Odyssey for High Energy Density[END_REF] ). Moreover, we are now faced with global warming and its result, climate change.

Changing our lifestyle, such as building smart homes and clean transportation are seen as possible solutions to reduce our carbon footprint. Many automakers have already started new initiatives, some of them even became actual battery producers ( 3 ). Such attempts are linked to vast investments into battery research and manufacturing, therefore batteries are becoming a hot topic again, after their first implementation in cell phones in 1990s. However, electric vehicles are still expensive, since the battery pack is the single major cost component in an electric vehicle. To reduce electric vehicle costs, US Department of Energy has set a mission for EV producers that the cost of battery packs should be reduced to 125 $.kWh -1 until 2022 ( 3 ). In order to realize this mission, new battery chemistries should be discovered, volume of production should be augmented, cell design should be optimized and manufacturing costs should be decreased. In fact, the most significant contributor to battery costs are raw material costs (50-70%). Among them, positive electrode is reported to have the highest percentage of costs (27%) ( 4 ) (Figure 0 Thus, research on new cathode material focus on Co-poor or even Co-less materials. However, it is not so easy since Co is usually needed to stabilize the layered structure, therefore new chemistries should be discovered. is the limiting factor of the battery, so its specific energy has to be increased ( 5 ).

Currently there are only a few materials that allow battery systems reaching high specific energies (250 Wh.kg -1 ). Among them, we can cite Li-rich NMC, Ni-rich NMC and high-voltage spinel materials. Despite their high capacity, low cycling stability and safety problems associated to structural deformations are also reported for these materials. This highlights the importance of discovering new high energy materials.

Recently, disordered rocksalt type positive electrodes appear to be promising because disordered rocksalt oxides offer capacities exceeding 250 mAh.g -1 at high average potential (3.6 V vs Li + /Li), while disordered rocksalt sulfides offer excellent capacities (>400 mAh.g -1 ) at relatively low average potential (2.3 V vs Li + /Li). Consequently, their energy density is much higher than today's commercial materials. This attracts research interest; thus, many groups are now working on this subject to bring new materials into the market.

In this work, we study two different disordered rocksalt families: disordered rocksalt oxides and sulfides materials. We apply different methods to synthesize each family and optimize our processes to have better electrochemical performances. Then, we perform various characterizations to understand their electrochemical and structural properties.

This thesis work is divided in five chapters:

The first chapter introduces Li-ion technology and current high-energy positive materials briefly. Thereafter, it describes disordered rocksalt positive materials and their properties in detail. It also explains why stoichiometric compositions failed before and what Li-rich compositions offer today. Finally, it shows promising electrochemical properties of Li-rich disordered rocksalts.

The second chapter contains the detail of the synthesis processes, sample preparations and procedure of electrochemical, morphological and structural characterizations.

The third chapter contains synthesis optimizations to improve electrochemical performance of disordered rocksalt Li2.2NiTi0.2Nb0.6O4. As the model material is prepared, ex situ and in situ X-ray diffraction (XRD) studies are performed to understand the structural evolution upon cycling. Finally, this chapter explains why electrochemical performances of such materials are limited at the end of the first cycle.

The fourth chapter describes preparation of disordered rocksalt Li2TiS3 powders with mechanochemical synthesis. To solve cycling stability problem of Li2TiS3, selenium substituted materials are prepared with new synthesis process and their electrochemical performances are compared with bare Li2TiS3.

The fifth chapter contains structural characterizations, such as operando neutron diffraction and ex situ XRD to investigate lithium insertion/extraction mechanism in Li2TiS3 during cycling. To elucidate redox processes in both Li2TiS3 and selenium substituted materials, powders and electrodes are analyzed with ex situ X-ray photoelectron spectroscopy (XPS).

In the last part of the thesis, we summarize our findings on disordered rocksalt oxide and sulfide materials. By referring our experience in those materials, we present a perspective and share our suggestions for future studies.

Chapter I. State of the art I.1. Rechargeable Lithium-Ion Batteries

Rechargeable lithium-ion batteries (LIBs) have become widely used power sources not only for portable devices but also for electrical vehicles. Differently from primary batteries, the electrochemical reaction is reversible, therefore the cells can be recharged several times and the delivered energy can be stored [START_REF] Daniel | Handbook of Battery Materials[END_REF].

Traditional LIBs contain four main components: positive electrode, negative electrode, non-aqueous electrolyte and separator. Transition metal oxides or phosphate based materials are generally used as positive electrodes (LiCoO2, LiNi1/3Co1/3Mn1/3O2, LiMn2O4 and LiFePO4), graphite is commonly used as negative electrode, non-aqueous electrolytes contain lithium salts (LiPF6, LiClO4, LiTFSI, etc.) and carbonate based solvents (ethylene carbonate, propylene carbonate, etc.) and separators are made of polymers (polyethylene, polypropylene). The electrolyte is a medium that can only allow ion transport, whereas it is an insulator for electron transport ( 6 ).

I.1.a. Li-ion Battery Working Principle and Characteristics

A typical LIB works following the principle that upon charging of the battery, Li ions are extracted from the positive electrode, move inside the electrolyte and insert into the negative electrode (Figure I -1). During lithium extraction, oxidation processes take place at the positive electrode and electrons flow through the external circuit from positive to negative electrode, where reduction processes happen. Thus, electricity is produced by the electrochemical reaction inside the cell. The inverse process happens during the discharge of the battery. While, reduction process at negative electrode (i.e., graphite) can be represented by

C6 + xLi + + xe -→ LixC6 eq. I -2
The overall reaction is the sum of two half-cell reactions,

LiMO2 + C6 ⇄ Li1-xMO2 + LixC6
eq. I -3

The free energy of the system, which is the driving force of the redox processes, can be expressed based on the assumption that cell is at equilibrium, ∆𝐺 = 𝑛. 𝐹. 𝑈 𝑡ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 eq. I -4

where n is number of electrons exchanged, F is Faraday constant (96485 C.mol -1 ) and 𝑈 𝑡ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 is the theoretical voltage between electrodes. At equilibrium state, there is no current flow, therefore open circuit voltage (𝑈 𝑜𝑐𝑣 ) is approximately equal to theoretical voltage. During charge and discharge, current evolves; therefore, working voltage, which is the actual operating voltage of the battery, differs from the open circuit voltage.

Among rechargeable storage systems, Li-ion battery has the highest specific energy, which is the product of the capacity and average discharge voltage. To maximize the specific energy of the cell, active materials should provide large discharge capacity and/or can be cycled at high discharge potentials. The theoretical capacity can be defined as the charge (or lithium ions) storage ability of active materials, which is calculated by 𝑄 𝑡ℎ = 𝑛.𝐹 3600.𝑀𝑤 eq. I -5

where Qth is theoretical capacity (mAh.g -1 ), n is the number of electrons (or lithium ions) exchanged during charge-discharge, F is Faraday constant (96485 C.mol -1 ) and Mw is the molecular mass of active material (g.mol -1 ). The capacity of the active material is strongly affected by the C-rate (1/h), which can be defined as a measure of the current rate at which a battery is discharging and can be represented by where i is current in A, theoretical capacity is expressed in Ah. For example, a cell charging at a C/10 rate needs 10h to reach its full capacity.

For industrial applications, Li-ion cells should have long cycle life, which is usually defined by the number of charge-discharge cycles that the cell can offer before its capacity falls below 80% of its nominal capacity.

To develop new high performance and affordable Li-ion batteries, some important parameters must be considered, such as 1. Selection of widely available and low-cost precursors for the active materials

I.1.b. Positive Materials for Li-ion Batteries

In 1980s, TiS2 was reported as the first intercalating compound by Whittingham ( 7 ) et al. after it demonstrated the ability to intercalate and de-intercalate lithium ions reversibly. Though TiS2 shows excellent cyclability, its average potential is rather low (2.1 V vs Li + /Li), consequently the energy density is not sufficient enough for high energy applications [START_REF] Whittingham | Electrical Energy Storage and Intercalation Chemistry[END_REF]. Thereafter, study of oxide type compounds gains interest because they could operate at relatively high potentials (> 4 V). These efforts results in the discovery of LiCoO2, which was reported by Goodenough ( 9 ) et al. and later commercialized by Sony in lithium-ion cells in 1991. However, a maximum of 0.5 Li ions could be deinserted from the layered LiCoO2 without destabilizing the structure; therefore the practical capacity is only half of its theoretical capacity (274 mAh.g -1 ) ( 9 ).

Increasing cobalt prices in the market is also another disadvantage. Having high theoretical capacity and cheaper price of nickel, LiNiO2 is seen as an alternative material ( [START_REF] Ohzuku | Electrochemistry and Structural Chemistry of LiNiO2 (R3̅ m) for 4 Volt Secondary Lithium Cells[END_REF][START_REF] Broussely | Li/LixNiO2 and Li/LixCoO2 Rechargeable Systems: Comparative Study and Performance of Practical Cells[END_REF] ). The studies show that the reversibility of LiNiO2 is lower than the one of LiCoO2 ( 12), thus manganese substitution is suggested. Thereafter, the best composition, which is found to be LiNi1/3Co1/3Mn1/3O2 (and called NMC111), is discovered as a solid solution of LiCoO2, LiNiO2 and LiMnO2 compositions ( [START_REF] Macneil | Structure and Electrochemistry of Li[NixCo1 -2xMnx]O2 ( 0 ⩽ x ⩽ 1 / 2 )[END_REF][START_REF] Patoux | Direct Synthesis of LiNi1/3Co1/3Mn1/3O2 from Nitrate Precursors[END_REF][START_REF] Yabuuchi | Novel Lithium Insertion Material of LiCo1/3Ni1/3Mn1/3O2 for Advanced Lithium-Ion Batteries[END_REF][START_REF] Hwang | A Combined Computational/Experimental Study on LiNi1/3Co1/3Mn1/3O2[END_REF] ).

LiNi1/3Co1/3Mn1/3O2 delivers a capacity up to 160 mAh.g -1 at 0.1 C rate in a potential range of 2.5 V -4.4 V vs Li + /Li [START_REF] Shaju | Performance of Layered Li(Ni1/3Co1/3Mn1/3)O2 as Cathode for Li-Ion Batteries[END_REF]. Here, the nickel redox provides high capacity, cobalt improves cycling stability and manganese enhances thermal stability. In parallel to LiNi1/3Co1/3Mn1/3O2, layered LiNi0.80Co0.15Al0.05O2 is also studied ( [START_REF] Zhuang | 2CO3 in LiNi0.8Co0.15Al0.05O2 Cathodes and Its Effects on Capacity and Power[END_REF][START_REF] Lee | Characterization of LiNi0.85Co0.10M0.05O2 (M = Al, Fe) as a Cathode Material for Lithium Secondary Batteries[END_REF][START_REF] Cho | Significant Improvement of LiNi0.8Co0.15Al0.05O2 Cathodes at 60 ° C by SiO2 Dry Coating for Li-Ion Batteries[END_REF] ). This material delivers 200 mAh.g -1 at 0.1 C rate in a potential range of 3 V -4.5 V ( 21) and Al substitution increases the thermal stability [START_REF] Guilmard | Thermal Stability of Lithium Nickel Oxide Derivatives[END_REF].

Moreover, polyanionic LiFePO4 is also studied in the literature ( [START_REF] Yamada | Optimized LiFePO4 for Lithium Battery Cathodes[END_REF][START_REF] Huang | Approaching Theoretical Capacity of LiFePO4 at Room Temperature at High Rates[END_REF][START_REF] Prosini | Determination of the Chemical Diffusion Coefficient of Lithium in LiFePO4[END_REF] ). Though LiFePO4

shows excellent cycling stability, its practical energy density is quite limited.

Having lower cost (<10 $.kg -1 ) and relatively stable structure compared to layer oxides, spinel LiMn2O4 attracts research interests and was also commercialized by Moli Energy. However, its practical capacity is rather small (~140 mAh.g -1 ). To utilize the advantages of different chemistries, blend of LiMn2O4 with LiNi1/3Co1/3Mn1/3O2 or LiNi0.80Co0.15Al0.05O2 have also been suggested ( [START_REF] Kim | Spinel LiMn2O4 Nanorods as Lithium Ion Battery Cathodes[END_REF][START_REF] Hosono | Synthesis of Single Crystalline Spinel LiMn2O4 Nanowires for a Lithium Ion Battery with High Power Density[END_REF] ). Addition of small amount of layer oxides suppresses the Mn dissolution as well as improves the capacity retention of LiMn2O4 cells ( 28 ).

I.1.c. High Energy Positive Materials

To extend the driving ranges of the electrical vehicles, specific energy of current lithium ion batteries should be further increased. The specific energy is related to both the capacity and the operating potential. Common approaches are focused on the use of high capacity materials, for instance Ni-rich NMC (LiNixCoyMnzO2, with x≥0.6) and Li-rich (or Mn rich) NMC [xLi2MnO3. (1-x)LiMO2] or high potential positive materials, such as LNMO (LiNi0.5Mn1.5O4).

• Ni-rich NMC (LiNi0.80Co0.10Mn0.10O2)

The capacity of the typical NMC111 (LiNi1/3Co1/3Mn1/3O2) is reported as 160 mAh.g -1 , which is smaller than the theoretical value (275 mAh.g -1 ) upon cycling at 0.1 C rate in a potential range between 2.5 V -4.4 V vs Li + /Li ( 3 ). Whereas, its Ni-rich version, LiNi0.80Co0.10Mn0.10O2 (NMC811) delivers much higher specific capacity (200 mAh.g -1 )

upon cycling at 0.1 C rate in a potential range between 3 V-4.3 V ( .

• High Voltage Spinel LNMO (LiNi0.5Mn1.5O4)

Spinel type materials are also considered as high energy materials ( [START_REF] Xiao | High-Performance LiNi0.5Mn1.5O4 Spinel Controlled by Mn 3+ Concentration and Site Disorder[END_REF][START_REF] Kim | Comparative Study of LiNi0.5Mn1.5O4-δ and LiNi0.5Mn1.5O4 Cathodes Having Two Crystallographic Structures: Fd3m and P4332[END_REF][START_REF] Zhou | LiNi0.5Mn1.5O4 Hollow Structures as High-Performance Cathodes for Lithium-Ion Batteries[END_REF] ) due to their high operating potentials (4.7 V) even though their practical capacity is rather small (140 mAh.g -1 ). Moreover, these materials could be easily produced, they show high rate capability and superior safety behavior. Those properties have already attracted many research interests. Despite such advantages, cycling at high potential causes electrolyte decomposition, that later impedes cycling stability ( 40 ). Therefore, new compatible electrolyte must be discovered for further application of these materials. In addition to these materials, recently, Li-rich disordered rocksalt oxides and sulfides have been considered as promising high energy positive materials for lithium-ion batteries ( [START_REF] Yabuuchi | Material Design Concept of Lithium-Excess Electrode Materials with Rocksalt-Related Structures for Rechargeable Non-Aqueous Batteries[END_REF][START_REF] Pralong | Lithium-Rich Rock-Salt-Type Vanadate as Energy Storage Cathode: Li2-XVO3[END_REF] ). Their cation disordered rocksalt structure is so robust that less structural problems are expected. Moreover, rocksalt oxides offer capacities exceeding 250 mAh.g -1 at high average potentials (3.6 V), while rocksalt sulfides offer excellent capacities (>400 mAh.g -1 ) at relatively low average potentials (2.3 V). Their specific energy is considerably higher than current materials. Therefore, many researches have recently been started to develop new Li-rich cation disordered rocksalt compositions. Thereafter, Li2MTiO4 (M: Ni, Co, Fe, Mn, V) type disordered rocksalts are reported at the beginning of 2004 ( [START_REF] Yang | Cation Disordered Rock Salt Phase Li2CoTiO4 as a Potential Cathode Material for Li-Ion Batteries[END_REF][START_REF] Fattakhova | Solvothermal Synthesis and Electrochemical Behavior of Nanocrystalline Cubic Li-Ti-O Oxides with Cationic Disorder[END_REF][START_REF] Prabaharan | Li2NiTiO4-a New Positive Electrode for Lithium Batteries: Soft-Chemistry Synthesis and Electrochemical Characterization[END_REF][START_REF] Küzma | Electrochemical Activity of Li2FeTiO4 and Li2MnTiO4 as Potential Active Materials for Li Ion Batteries: A Comparison with Li2NiTiO4[END_REF][START_REF] Zhang | Facile Synthesis and Electrochemistry of a New Cubic Rocksalt LixVyO2 (x = 0.78, y = 0.75) Electrode Material[END_REF][START_REF] Jiang | Electrochemical Synthesis of LiTiO2 and LiTi2O4 in Molten LiCl[END_REF] ). To eliminate diffusion limitations in disordered rocksalts, new synthesis approaches concentrate on low temperature synthesis methods, such as citric acid assisted sol-gel method ( [START_REF] Yang | Cation Disordered Rock Salt Phase Li2CoTiO4 as a Potential Cathode Material for Li-Ion Batteries[END_REF][START_REF] Fattakhova | Solvothermal Synthesis and Electrochemical Behavior of Nanocrystalline Cubic Li-Ti-O Oxides with Cationic Disorder[END_REF][START_REF] Prabaharan | Li2NiTiO4-a New Positive Electrode for Lithium Batteries: Soft-Chemistry Synthesis and Electrochemical Characterization[END_REF] ), acid base reactions ( 47 ), hydrothermal synthesis ( 48 ) and molten salt synthesis ( [START_REF] Zhang | Facile Synthesis and Electrochemistry of a New Cubic Rocksalt LixVyO2 (x = 0.78, y = 0.75) Electrode Material[END_REF][START_REF] Jiang | Electrochemical Synthesis of LiTiO2 and LiTi2O4 in Molten LiCl[END_REF] ). With these soft chemistry methods coupled with carbon coating strategies, average particle size becomes less than 100 nm and electronic conductivity increases. Thus, much higher practical capacities are reached. For In brief, the practical capacities of stoichiometric rocksalts are still much lower than the theoretical capacities. The reason is attributed to poor lithium diffusion through cation disordered structure. To overcome lithium conduction challenges in cation disordered rocksalts and to understand lithium ion conduction mechanism in these materials, theoretical calculations are required.

I.2. Cation Disordered

I.2.b. Lithium diffusion mechanism through cation disordered rocksalts

Cation disordered rocksalt crystallises in a face centered cubic (fcc) structure with a space group of Fm3 ̅ 𝑚. In this fcc structure, cations (lithium or transition metals) Previous studies describes the lithium conduction mechanism into fcc crystal lattice with a theory of Oh-Th-Oh transition in which lithium ions migrate from one octahedral site (Oh) to another octahedral site with a relaxation at tetrahedral site (Th) [START_REF] Van Der Ven | Lithium Diffusion Mechanisms in Layered Intercalation Compounds[END_REF][START_REF] Van Der Ven | Understanding Li Diffusion in Li-Intercalation Compounds[END_REF] ). Moreover, height of the tetrahedron is reported as an important parameter.

Tetrahedron height should be tall enough that the energy barrier of transition between Oh-Th sites becomes smaller, thus lithium ions can migrate easily between the sites.

This energy barrier changes based on cation distribution around tetrahedral sites ( 54 ). In 2014, A. Urban et al. perform density functional theory calculations to describe lithium conduction mechanism in disordered rocksalts ( [START_REF] Urban | The Configurational Space of Rocksalt-Type Oxides for High-Capacity Lithium Battery Electrodes[END_REF][START_REF] Lee | Mitigating Oxygen Loss to Improve the Cycling Performance of High Capacity Cation-Disordered Cathode Materials[END_REF] ). They describe that both 1-TM and 2-TM channels are inactive for lithium conduction due to strong repulsion between the cations, whereas 0-TM channels are active for lithium conduction due to low repulsion between lithium ions (or lithium-vacancy). Thus, energy barrier of transition between Oh-Th sites is smaller. Besides, they report that increasing lithium to metal (Li/Me) ratio might create a percolation network through 0-TM channels. This result perfectly illustrates that:

1. Only reduced capacities are obtained in stoichiometric rocksalts (Li/M=1.0) because the amount of lithium is not sufficient enough to create percolating network through lithium diffusion channels.

2. To achieve higher capacities, cation disordered rocksalts should contain excess lithium amount, hence sufficient lithium conduction paths can form. or Me 3+ /Me 4+ ). To increase electronic conductivity, samples are coated with carbon by mechanical milling, particle size is reduced to submicronic size. Thanks to these improvements, capacities have been increased to 250-300 mAh.g -1 , two times higher than before ( Cation Disordered Rocksalt Oxides The same authors report another disordered rocksalt Li1.3Nb0.3V0.4O2, which is prepared by substitution of Li3NbO4 with vanadium metal ( In summary, Li1.2Ti0.4X0.4O2 (X: Fe 3+ , Mn 3+ ) compositions can deliver large reversible capacities at high operating potentials. However, it should be noted that those results are obtained at high temperatures (50 °C), thus electrode kinetics is still slow.

I.3. Li-rich

I.3.c. Li4MoO5 based cation disordered rocksalt oxides

Another group prepares cation disordered in the Li4MoO5-LiMoO2 system that has the potential to deliver high capacity based on multi electron (Mo 3+ / Mo 6+ ) redox reactions.

As a comparison, Li3NbO4-LiMoO2 and Li2TiO3-LiMoO2 solid solutions are also prepared ( 61 ).

I.4.b. Li3NbS4 based cation disordered rock salt sulfides

Same author also reports a disordered rocksalt Li3NbS4 prepared with the same synthesis recipe ( [START_REF] Sakuda | Rock-Salt-Type Lithium Metal Sulphides as Novel Positive-Electrode Materials[END_REF][START_REF] Sakuda | Preparation of Novel Electrode Materials Based on Lithium Niobium Sulfides[END_REF] ). Disordered rocksalt Li3NbS4 has slightly bigger lattice parameters (5.13 Å) than Li2TiS3 (5.06 Å).

When the cells are cycled at a rate of C/10 between 3 V -1 V, a reversible capacity of 486 mAh.g -1 is achieved corresponding to 4.5 exchanged electrons (Figure I -15).

Cycling curve is reversible in subsequent cycles and its energy density is reported as high as 923 Wh.kg -1 . Li3NbS4 cells show better cycling stability than Li2TiS3 cells, 97%

capacity is reached at the end of the 50th cycle ( 66). However, average potential of Li3NbS4 is smaller (2.20 V) than average potential of Li2TiS3 (2.30 V). At high current rates (1000 mA.g -1 ), superior cycling stability is observed.

Figure I -15: Specific capacity vs cell voltage of Li3NbS4 electrodes and capacity retention vs number of lithium ions exchanged during cycling ( 66 )

Different sulfides, Li2TiS3-Li3NbS4 and Li2SnS3-Li3NbS4 are also reported by the same author and all of them are synthesized in cation-disordered rocksalt structure ( [START_REF] Sakuda | Development of Li2TiS3-Li3NbS4 by a Mechanochemical Process[END_REF][START_REF] Sakuda | Cubic Rocksalt Li2SnS3 and a Solid Solution with Li3NbS4 Prepared by Mechanochemical Synthesis[END_REF] ).

Li(8+x)/3Ti(4-4x)/3NbxS4 present reversible capacities higher than 350 mAh.g -1 corresponding

to three electrons process between 3.0 V and 1.5 V ( 67 ). On the contrary, tin-substituted composition, Li3-x/3Nb1-xSn4x/3S delivers smaller capacities with increasing amount of tin in the composition ( 68 ). The reason is explained with low electronic conductivity of tin (7.4x10 -7 S.cm -1 ). The author concludes that sufficient electronic conductivity is necessary to activate anionic (sulfur) redox process.

Overall, disordered rocksalt sulfides can provide large reversible capacities, however sufficient electronic conductivity is necessary to activate anionic (sulfur) redox process.

Moreover, capacity fading in the cells containing liquid electrolytes is a major problem for all disordered rocksalt sulfides. The reason of such a degradation in liquid cells is still unknown and needs to be elucidated.

I.5. Redox Activity in Disordered Rocksalts

Previously, large capacities delivered by disordered rocksalts were attributed to activity of both cationic and anionic redox processes. This implies that such processes should be understood well to design new high-energy chemistries ( [START_REF] Yabuuchi | Material Design Concept of Lithium-Excess Electrode Materials with Rocksalt-Related Structures for Rechargeable Non-Aqueous Batteries[END_REF][START_REF] Li | Thermodynamic Activation of Charge Transfer in Anionic Redox Process for Li-Ion Batteries[END_REF][START_REF] Grimaud | Anionic Redox Processes for Electrochemical Devices[END_REF] ).

Redox activity in a composition can be outlined with the electronic states of the materials. This can be explained with the metal-ligand band structure (Figure I -16).

Density of states of a material contains M (metal) at conduction band and L (ligand) at valence band below energy of fermi level (EF). When purely cationic redox is active, metal-ligand bands are well separated from each other, therefore there is no overlapping with these two bands ( 69 ). High ionicity between these two bands allows the metal acting as a redox center and providing necessary electrons during Li insertion. This result suggests that charge capacity is produced by both cationic (Mn 3+ / Mn 4+ ) and anionic (O 2-/ O n-) redox activities. whereas iron species are partially oxidized (Fe 3+ / Fe 4+ ). However, such reduced oxygen species disappear at the end of the charge, O K-edge spectra shows similar results for pristine and charge states. This result indicates that anionic redox is activated at the beginning of the charge, and responsible for the charge transfer. However, the reverse process is not observed during discharge, indicating anionic redox is not fully reversible in Li1.2Nb0.3Fe0.4O2 electrodes. Besides, electrode morphology changes at the end of the discharge (Figure I -18). This also indicates that such irreversible anionic redox leads to oxygen release that causes side reactions on the electrode surface as well as reorganization of particle morphology. Overall, it seems that anionic redox activity in cation-disordered rocksalt oxides is not always reversible for all the compositions and irreversible anionic redox can result in smaller reversible capacities. The reversibility of anionic redox is still unknown, therefore, further studies are required to elucidate such phenomenon.

• Anionic Redox in Cation Disordered Rocksalt Sulfides

To utilize anionic redox, it is necessary to increase covalency between metal and ligand, this requires using high electronegative (or less electropositive) metal atom and less electronegative ligand. For instance, production of persulfide (S2 2-) is much easier compared to oxide ions (O2 2-) because sulfur p orbitals are relatively higher than oxide p orbitals ( 69). This creates high covalency between transition metal-sulfur due to overlapping metal d band and ligand p band, thus anionic redox process is easily triggered in rocksalt sulfides ( 69 ). As a summary, anionic redox process is highly reversible in disordered rocksalt sulfides, flexible structural transition allows such systems to reversibly exchange the lithium ions.

I.6. Conclusion

This chapter describes lithium ion technology and current positive electrode materials.

To further increase the specific energy of lithium ion batteries, new high energy Li

Chapter II: Experimental Procedures

This chapter describes the synthesis methods, structural and electrochemical characterization techniques that are used throughout the thesis. The principles of each technique, sample preparations and experimental parameters are described.

II.1. Synthesis Methods

Molten salt Synthesis

In the previous chapter, electrochemical performances of disordered rocksalts showed size dependent properties, thus low temperature synthesis routes were suggested to prepare these materials. To prepare rocksalt oxides, we select molten salt synthesis because it allows us preparing ceramic powders at low temperatures. According to the synthesis procedure, precursors are mixed with an eutectic salt mixture and the resulting powder is calcined at a temperature above the melting point (Figure II -1).

During the calcination, the salt melts and acts as the reaction medium, thus reaction takes place at much lower firing temperatures due to high solubility and diffusivity of the active species inside the salt ( 75 ). After the synthesis, the salt is removed by washing of the powder with a solvent, such as water. Here are the important parameters for low temperature molten salt synthesis ( 75 ):

• Molten salt should be stable, shouldn't react with precursors during the synthesis • Solubility of species in molten salt should be high enough to obtain targeted phases.

• For low temperature synthesis, molten salt should be selected as it has low melting point

Mechanochemical Synthesis

Mechanochemical synthesis allows preparing powders in crystalline, meta-stable and amorphous phases. Among different millers, SPEX and planetary millers are commonly used in laboratory scale production. In this thesis study, planetary ball milling is used to prepare disordered rocksalt sulfur-selenium compositions (Li2TiSexS3-x) that are prepared with a 62:1 (wt%) ball to powder ratio and milled at 510

Figure II -3: Mechanochemical synthesis experimental set-up: stainless steel jar, zirconia balls and planetary miller (PM 100), scheme of grinding medium in PM100

The parameters that has influence on phase and the morphology of the final product are described below ( 76):

Grinding medium: Type and size of the balls affect the energy inside the jar, thus phase of the powder can change. In general, bigger balls create stronger shocks and the impact energy on the particles becomes greater. Besides, hard materials, such as stainless steel, hardened steels, tempered steel are frequently used to increase weight of the grinding medium as they create enough impact energy. In our experiments, we use zirconium jar with zirconium balls to avoid any reaction between sulfur and steel.

Ball to powder mass ratio:

Depending on the ball to powder ratio, energy input can change. High ball to powder ratio induces high energy inside the jar.

Milling speed: High milling speeds creates high energy input into the powders, consequently, temperature inside the cell increases. High temperature can be beneficial to allow the ingredients diffusing enough to trigger chemical reactions. Whereas it might also cause undesired reactions as well, such as the formation of impurities in the final product if milling takes longer durations. Therefore, powders should be milled at optimum milling speed.

Milling duration:

Milling duration often depends on the milling speed and powder load (ball to powder ratio). In general, reactions are completed at shorter durations when the powders are milled at high speeds or at high loads (high ball to powder ratios) as the input energy is sufficient. It is also noted that milling powders for longer durations increase the contamination in the final product as the balls often wear out.

II.2. Structural Characterizations Powder X-ray Diffraction (XRD)

The crystal structure of the synthesized powders is determined by X-ray diffraction using a BRUKER AXS D8 diffractometer using Cu Kα (λ=1.540 Å) anticathode, the sample is scanned in the range of 8-80° using a step size of 0.02°. As the sulfides are air sensitive, all the samples are prepared in the glovebox and the sample surface is protected with Kapton ® , which is a polyimide film, shown in Figure II -4. Kapton ® has high transmittance to X-rays, however its signal still can be detected at small 2𝜃 angles (0-28°) in the diffraction patterns of the powders. Besides, its tightness to air, and more specifically moisture, is limited to couple of hours. Therefore, duration of the X-ray diffraction is limited to a maximum of three hours.

Figure II -4: X-ray diffraction analysis: sample holder for powder and electrode

X-ray diffraction is a useful technique to determine the crystal structure of the crystalline powders. Each powder has a unique crystal structure that contains parallel crystalline planes designated by Miller indexes (hkl) that are separated by "dhkl" distance called interreticular distance.

Figure II -5: Scheme of X-ray diffraction technique and Bragg's law

As the incident X-ray wave interacts with the powder, each plane reflects X-rays differently, hence this creates a constructive interference. For instance, two different Xray waves with a wavelength of λ hit a crystal structure at incident angle θ, first wave will reflect from top atomic plane, the second wave will reflect from bottom plane (Figure II -5). Since the second wave must cross deeper into the bottom plane, it travels further than the second wave, its distance is much bigger than the first wave. If this distance is equal to nλ then these two waves are in phase, therefore they constructively interfere. This verifies Bragg's law (Eq. II -1).

Bragg's law: nλ = 2 dhkl sinθ (Eq. II -1)

where n: number of X-ray waves, λ is X-ray wavelength, dhkl is distance between the crystalline planes and θ is the reflected angle.

The most important use of X-ray diffraction method is to identify the materials. Each material has a characteristic XRD pattern that can be expressed as a fingerprint for phase identification. There are two important parameters for the phase identification: Structural factor provides information about the type of the atoms (fj) and the atomic positions (u, v and w) within the unit cell. Therefore, the target of the structural analysis is to compare observed and calculated structural factors.

The size of the crystalline domains (or crystallite size) can be calculated by using Scherrer equation (Eq. II -4). In the equation, β(θ) corresponds to peak widths obtained from the full-width at half maximum (FWHM) corrected from the instrumental broadening at corresponding angle (2θ), K is Scherrer constant that varies between 0.62 to 2.08 (0.94 for cubic systems) and λ is X-ray wavelength. The crystallite sizes (L) can be calculated at each angle, but its mean value is frequently used for a convenience.

𝛽(𝜃) =

𝐾𝜆 𝐿𝑐𝑜𝑠𝜃 (Eq. II -4)

In this thesis project, Fullprof program has been mainly used ( 77). The peak positions and peak profiles depend on the crystal structure and the space group correspondingly. At first, XRD patterns are fitted using Le Bail method in which peak shape and positions are refined iteratively until minimum Chi-square value (< 5) is reached. In this way, peak positions are fixed and the unit cell parameters are calculated with high accuracy. To perform Rietveld refinement and simulated XRD pattern, structure models (crystal information file -cif) are taken from the databases (Crystallography Open Database). Rietveld refinement is performed to minimize the difference between the observed and the calculated data by refining the positions and occupancies of the atoms in the crystal.

In the case of the in situ analysis, the data are refined using Jana2006 to enable the correction of the effect of the variable divergence slit ( 78). The refinement has been carried using the Fundamental Parameter Approach, in which the instrumental resolution is calculated from the description of the main parameter of the optical setup. This allows a direct correction of the peak profile.

• Ex situ XRD analysis

Ex situ XRD analysis is another way to detect structural evolution in disordered rocksalts. Coin cells are cycled in the standard procedure that is described in Section 2.3 and stopped at different states of charge. After cycling, cells are dismounted in the glovebox, electrodes are recuperated from coin cells, they are washed with dimethyl carbonate (DMC) and dried inside the glovebox. After drying, electrodes are placed on the surface of a glass plate and covered with Kapton ® to protect from air atmosphere (Figure II -4).

• In situ XRD analysis

To observe structural evolution in disordered Li2.2NiTi0.2Nb0.6O4 during cycling, in situ XRD analysis is performed using Biologic Instrument and BRUKER diffractometer.

The cell is cycled in GITT (Galvanostatic Intermittent Titration Technique) mode. Before the experiment in ILL, we studied and optimized the electrochemistry of Li2TiS3.

At the beginning, we observed a large polarization and poor reversibility when the cells were cycled in a standard schedule (3 V -1.5 V). Thereafter, we optimized the amount of the powder and found an adapted cycling schedule (3.5 V -1 V).

Scanning Electron Microscopy (SEM)

Morphology of synthesized powders are investigated by Scanning Electron

Microscopy equipped with a field electron gun (MEB LEO 1530 Gemini). The surface of the sample is scanned through focused beam of electrons. As electrons interact with the atoms, electrons are removed from the inner shell and empty place is filled by electrons at higher shell energy level, resulting in an energy release. Electrons emitted by the atoms are collected by detectors and the detected signal is transferred into information about surface topography. Two different types of electrons are emitted from the surface: secondary and backscattered electrons. Secondary electrons are collected close to the specimen surface; hence they create high resolution images.

Backscattered electrons emerge from deeper locations within the specimen; therefore, they produce low resolution images compared to secondary electrons. As the intensity of backscattered electron signal depends on the atomic number (Z), they provide information about the element distribution in the sample.

Energy-Dispersive X-ray Spectroscopy (EDS)

To perform the chemical mapping and determine elements in the compositions, we used Energy-Dispersive X-ray Spectroscopy using 15 kV in MEB LEO 1530 Gemini.

The surface of the sample is bombarded with focused beam of electrons and emitted X-rays are collected by the detectors. This provides the information about the element distribution in a specific region of the sample. Although this technique is easy and practical to determine the elements in a composition, most of the instruments cannot detect light elements (Z<10) precisely, thus quantitative analysis of powders is often difficult.

• Sample preparation for SEM and EDS Analyses

A specific sealed sample holder is used to transfer air sensitive samples from glovebox to the SEM vacuum chamber. The samples are placed on the surface of adhesive carbon tape that has been already sticked onto SEM stub (Figure II -10). Data is recorded for various magnifications (2000, 5000, 10000 and 30000) by using 5 kV accelerating voltage.

Figure II -10: Sealed sample holder for SEM of sulfide electrodes

Raman Spectroscopy

To observe chemical bonds in different disordered rocksalt sulfur-selenium compositions, Raman measurements are carried out on a LabRam HR instrument (Horiba Jobin-Yvon) in backscattering geometry. 632 nm laser excitation wavelength on a laser power of 10% filter (≈ 0.17 mW) is used. To prevent heating of the materials by the laser beam and to obtain a sufficient signal to noise ratio, short acquisition times (10 s) are used with multiple accumulations (30).

Working mechanism of Raman Spectroscopy is based on inelastic scattering, so called "Raman Effect". As the sample is irradiated by photons with a single frequency, the atoms vibrate and rotate as a response of electromagnetic field that is formed upon photon-sample interaction. As a result, photons are scattered at different frequencies.

In Raman spectroscopy, one can know the characteristics of each material from

• Raman frequencies, which allows distinguishing the species from each other.

• Width of Raman peaks, which gives information about the quality of crystals and amount of plastic deformations

• Polarization of Raman peak, which is more related to crystal symmetry and orientation.

Different laser can be used for Raman light source. As the wavelength of the laser gets smaller, signal quality increases, however it also causes heating of the sample, thus the sample is more easily damaged.

• Sample preparation for Raman Spectroscopy

Powders are prepared in the glovebox and a sealed sample holder is used for Raman measurements. The powder is put in the center of the sample holder that is closed by a cover containing a transparent glass that is invisible to the laser. Sample holder is transferred from glovebox to the LabRam HR instrument coupled with Atomic Force Microscopy (AFM) machine. AFM probe is used to select the particle region, and then laser beam is focused on selected particles.

X-ray Photoelectron Spectroscopy (XPS)

To determine redox active species and to elucidate the redox mechanism in disordered rocksalt sulfur-selenium compositions (Li2TiSexS3-x), X-ray Photoelectron Spectroscopy Binding energy of core electron (Eb) is expressed by Einstein relationship (Eq. II -5):

E b = hѵ -E k -Ф sp
Eq. II -5

where Ek: kinetic energy of ejected electron, which is measured by spectrometer, hυ: energy of photons, Фsp: work function induced by the analyzer.

Figure II -12: Sample holder for XPS analysis and the XPS machine

The first aim is to identify the elements present in the samples, therefore survey spectra is recorded over a spectral range of 0-1200 eV using a pass energy of 117 eV, corresponding to an energy resolution of 1.6 eV and an acquisition time of 0.5 s.eV -1 .

The high-resolution spectral analysis is carried out using a pass energy of 23 eV allowing an energy resolution of 0.5 eV.

II.3. Electrochemical Characterizations Electrochemical Cell Descriptions

• Preparation of oxide electrodes

The electrodes are prepared according to 80:10: to cycle the cells. To detect if the materials have diffusion limitations, high temperature cycling schedule is carried out. In that case, coin cells are placed into an oven that operates at 55 °C and is connected to the ARBIN instrumentation.

• Galvanostatic Intermittent Titration Technique (GITT)

To study the thermodynamics and kinetics of disordered Li2.2NiTi0.2Nb0.6O4,

Galvanostatic Intermittent Titration Technique (GITT) is used. This technique allows

obtaining galvanostatic profile at equilibrium potential.

In GITT technique, the cell is charged (or discharge) with a small constant current for a certain time, then the current is interrupted and the cell is given a rest until the potential is stabilized (Figure II -15). This procedure is repeated until the cell reached fully charged or discharged states. With imposed current, lithium is deinserted (or inserted) into the active material, thus potential changes based on concentration gradient. Without no imposed current, potential reaches its thermodynamically equilibrium condition with ion transport that occurs from the bulk to the electrodes.

Figure II -15: GITT mode of cycling

The equilibrium potential is key for various investigations in electrodes. Most of the active materials undergoes structural modifications upon cycling that adversely affect the electrochemical performances. With galvanostatic mode of cycling, it is quite difficult to detect the actual start of such modifications, one can only detect the modifications in the electrodes at the end of the charge and discharge states. On the contrary, GITT allows performing X-ray diffraction analysis at the equilibrium potentials, thus it makes possible to detect step by step the structural modifications based on lithium (de)insertion process. It also allows determining diffusion and kinetic properties of electrodes at the equilibrium potentials.

• Cycling Voltammetry (CV)

To determine redox potentials, electrode kinetics and reversibility of sulfur-selenium compositions, cycling voltammetry technique is applied. According to this technique, cell is submitted to a potential evolving linearly at a specific scan rate (V.s -1 ) and the current evolves as a response of potential change. This gives information about the different redox processes that take place at the electrodes. where n is the number of electrons exchanged.

Figure II -17:

Voltammogram of the electrodes scanned from slow scan rates (5 µV.s -1 ) to higher scan rates (100 µV.s -1 ) with gradual increments, calculation of potential difference for each scan rates to determine electrode kinetics at slower and higher scan rates

In our experiments, two different cells are cycled with a slow scan rate (5µV.s -1 ) to observe lithium (de)insertion process in the sulfur-selenium compositions. To investigate the electrode kinetics at faster scan rates and reversibility of electrode kinetics, the cells are scanned from slow scan rates (5 µV.s -1 ) to higher scan rates (100 µV.s -1 ) with gradual increments (Figure II -17). The amplitude of the current increases with scan rates, indicating the relation between electrode kinetics of the redox reactions with applied scan rates that is expressed in Randles Sevcik equation for a reversible process:

Ip= (2.687x10 5 )n 3/2 v 1/2 D 1/2 AC (Eq. II -8)
Ip is peak current, n is number of transferred electrons, v is scan rate (V.s -1 ), D is diffusion coefficient (cm 2 .s -1 ), A is electrode area (cm 2 ) and C is concentration (mol.L -1 )

By calculating potential difference (∆E) for each scan rate, diffusion coefficient of the species can be calculated using Randles-Sevcik equation (Eq. II -8). Thus, the electrode kinetics of different species can be distinguished from each other.

Introduction

In After two hours, eutectic NaCl-KCl salt in a salt ratio to precursor ratio of 4:1 is added into the mixture and then mixing is repeated for two hours. As the mixing is completed, mixture is heated to 300 °C for 2 hours to remove acetate from the precursors and later it is calcined at 670 °C for 3 hours. Between two heating steps, the powder is taken from the oven and grinded a few minutes in mortar to have a homogeneous mixture. The calcination temperature is slightly above the melting point of eutectic salt (667 °C). After the synthesis, powder is washed with hot water (60 °C)

to remove the remaining salt, filtrated and dried in the oven. The electrochemical redox process of Li2.2NiTi0.2Nb0.6O4 can be described as below:

Li2.2NiTi0.2Nb0.6O4 ⇄ Li2.2-xNiTi0.2Nb0.6O4 + xLi + + xe - (Eq. III -1)
Considering only Ni 2+ /Ni 4+ redox is active during cycling, the theoretical capacity based on 2 moles of Li + exchanged per formula is 264 mAh.g -1 .

In • Quenching: Disordered rocksalts have shown electrochemical performances dependent on particle size; especially diffusion is more problematic in powders that contain big particles. With fast cooling approach, particle size growth can be controlled, thus lithium diffusion length becomes shorter, ionic limitations might be reduced. Consequently, much greater capacities could be achieved.

• Effect of washing (type of solvent): Rocksalt Li2.2NiTi0.2Nb0.6O4 is prepared by molten salt synthesis, thus washing step is mandatory to remove remaining salts in the composition. Water is a stable solvent and solubility of various salts in water is quite high, thus it can dissolve many different salts. However, washing can cause loss of lithium in the composition, for instance lithium ions in the powders can exchange with protons from the water. This might reduce the specific capacity; therefore, new solvents should be searched to eliminate that problem.

• Effect of molten salt to precursor ratio: Molten salts have strong solving capability; thus, they act as a medium in which diffusion rate of the precursors is much higher. Ratio of salt to precursor can affect the precursors solubility;

therefore, the effective ratio should be determined.

• Type of molten salt: Each eutectic salt has a different melting point and solubility. Use of eutectic salt that has a low melting point can reduce the synthesis temperature and allows preparing powders in smaller particles size.

However, salt should be stable against reactants, should not react with the precursors throughout the synthesis and it should have enough solubility power to dissolve the reactants in order to prepare the desired phases.

• Effect of calcination temperature and duration: Low temperature synthesis methods are useful to synthesize powders in smaller particles size. Finding optimum calcination temperature and duration is important to optimize the capacity of Rocksalt Li2.2NiTi0.2Nb0.6O4.

III.1.a. Simple tempering (long firing) and quenching strategies

To improve our synthesis recipe, we first consider quench strategy. For a comparison, we also prepare a reference material prepared by classical approach. Three samples are prepared with the same recipe, however different firing and cooling steps are applied.

• The first sample (classical/reference) is heated at 670 °C for 3 hours and let to cool down in the oven

• The second sample (quench) is heated at 670 °C for 3 hours and it is taken from the oven and let to cool down outside the oven (no cooling agent is used)

• The third sample (long calcination and quench) is heated at 670 °C for 6 hours and it is taken from the oven and let to cool down outside the oven. Cycling stability is rather good at room temperature, but discharge capacity is quite small (96 mAh.g -1 ). Whereas, cycling stability at 55 °C is reduced and capacity fades quickly. Poor cycling stability at high temperatures has been already reported in different positive materials in the literature and it was attributed to the degradation of active materials ( [START_REF] Wang | Facile Molten Salt Synthesis of Li2NiTiO4 Cathode Material for Li-Ion Batteries[END_REF][START_REF] Lee | A New Class of High Capacity Cation-Disordered Oxides for Rechargeable Lithium Batteries: Li-Ni-Ti-Mo Oxides[END_REF] ). In Figure III -11, we observe that charge capacity of powder washed with ethylene glycol is slightly greater, but its discharge capacity is smaller than the powder washed with water. Although discharge capacity reduces with ethylene glycol washing, we are now sure that there is no proton-lithium exchange in our system. Therefore, we decide to wash the powders with ethylene glycol in further synthesis. charge and discharge capacities are obtained from the powder prepared with the higher salt to precursor ratio (10:1). Those results support our assumption that eutectic NaCl-KCl salt reacts with the precursors and produce an impurity phase that blocks the lithium diffusion and leads to smaller capacities. Therefore, as a next step for the synthesis improvement, we decide to change the eutectic NaCl-KCl with another eutectic salt that is more stable and does not react with the reactants during the synthesis.

III.1.d. Molten salt synthesis with a different eutectic salt (LiCl-KCl)

In the literature, different eutectic salts can be found (Table III -1). Among them, we select 0.59LiCl-0.41KCl eutectic salt that has a low melting point (354 °C.) We apply the same synthesis procedure that was described before, but we just change the type of molten salt. Instead of using eutectic NaCl-KCl salt, 0.59LiCl-0.41KCl eutectic salt in a salt to precursor ratio of 2:1 is used to synthesize Li2.2NiTi0.2Nb0.6O4 powders. The small peaks detected in diffraction patterns belongs to TiO2 and Nb2O5 precursors.

Such a result indicates that diffusion of titanium and niobium oxide precursors are still low in the 0.59LiCl-0.41KCl salt, even though melting point of 0.59LiCl-0.41KCl is lower (354 °C). On the contrary, disordered rocksalt phase is observed in the powder prepared at 670 °C. That result concludes that much higher temperatures are required to achieve disordered rocksalt phase.

In Figure III -15, new sample delivers similar charge capacity (145 mAh.g -1 ) as classical material, however the discharge capacity is slightly smaller (95 mAh.g -1 ). ion diffusion rate of reactants could be higher in LiCl-KCl salt melt.

In overall, we successfully synthesize Li2.2NiTi0.2Nb0.6O4 in disordered rocksalt phase using a different eutectic salt (0.59LiCl-0.41KCl) and the electrochemical performances are almost similar. Therefore, we decide to continue synthesizing Li2.2NiTi0.2Nb0.6O4 powders with 0.59LiCl-0.41KCl salt.

III.1.e. Effect of calcination temperature and duration

Another approach to improve the capacity of disordered rocksalt Li2.2NiTi0.2Nb0.6O4 is to synthesize the powders at low temperatures. In classical method, the mixtures are

III.2. Understanding Studies

After optimizing the electrochemical performance of Li2.2NiTi0.2Nb0.6O4, structural studies are conducted to understand why the capacity of this material is limited. almost no charge and discharge capacities are obtained upon cycling between 3.9 V -2 V. At potentials greater than 4.2 V, lithium exchange process becomes more reversible upon cycling. However, following our hypothesis, we would expect to see a good reversible behavior at low state of charge and an increase in the irreversible capacities with increasing SOC. After these findings, we conclude that the result does not align with our hypothesis. Therefore, the irreversibility might arise from the structural modifications instead of lithium diffusion limitations. Now, we have another hypothesis in which there might be a structural change at the end of the charge. Thus, new structure allows only few lithium ion exchange. This can be further investigated with ex situ and in situ X-ray diffraction techniques.

III.2.a. Ex Situ X-ray Diffraction Analysis

To perform ex situ X-ray diffraction, three electrodes are prepared and one electrode is charged to 4.8 V while another is discharged to 2 V. Detailed procedure can be found in Chapter II.2. Thereafter we decide to investigate the change of lattice parameters step by step with in situ X-ray diffraction. The change in crystalline sizes is irreversible as the final crystalline sizes are much smaller than those in pristine states. This apparition of defects (either dislocation or long-range disordering) may also be a cause of the capacity fading of this material.

III.2.b. In Situ X-ray Diffraction Analysis

III.3. Conclusion

The objective of this chapter was to develop overlithiated disordered rocksalt Li2.2NiTi0.2Nb0.6O4 with high capacity. Our initial synthesis has resulted in smaller charge (120 mAh.g -1 ) and discharge (74 mAh.g -1 ) capacities than the theoretical capacity (264 mAh.g -1 ). To reach desirable capacities, we optimize our synthesis procedure considering various synthesis parameters.

Table III -2: Comparison of initial and final synthesis parameters

After optimization of the synthesis parameters (Table III -2), we finally succeed to obtain charge and discharge capacities of 166 mAh.g -1 and 108 mAh.g -1 at room temperature respectively. However, we observe that there is always a high irreversibility at the end of the first cycle. At first, we suppose that irreversibility arises from lithium diffusion problem. After testing our hypothesis, we obtain a result that does not align with our first hypothesis. Thereafter, we test our second hypothesis that we consider a possible structural change during charge-discharge.

We observe from ex situ XRD result that there should be a second disordered phase forming during charge process. To validate this, we also perform in situ XRD analysis of electrodes at different SOC. Refinement of the initial powder indicates that second disordered rocksalt phase has already been present in the pristine state. In addition, we find out that the final phases at the end of the discharge are different from those in the pristine states. From the study of size in the crystalline domains, we conclude that disordering happening during cycling causes irreversible structural modifications.

In general, Li-rich disordered rocksalt oxides discussed in the literature offer large reversible capacities due to their high amount of lithium content. Our findings show that this could be possible if the material has pure disordered structure that can allow lithium ions exchange reversibly during cycling. It seems that diffusion limitations in disordered rocksalt oxides have been partly solved with overlithiation, nevertheless preparing those materials in pure disordered structures might gain importance for future studies.

Introduction

Metal sulfide family is extensively studied in the literature, especially after layered TiS2

showed unique lithium ion intercalating property ( [START_REF] Whittingham | Electrical Energy Storage and Intercalation Chemistry[END_REF][START_REF] Whittingham | Chemistry of Intercalation Compounds: Metal Guests in Chalcogenide Hosts[END_REF][START_REF] Whittingham | The Role of Ternary Phases in Cathode Reactions[END_REF] ). In parallel to TiS2, layered TiS3 was also studied. Upon discharging the cells, three lithium ions are intercalated into TiS3 (Eq. IV -1 and Eq. IV -2), the higher capacity of 80mAh was reached without precision on the mass of active material ( [START_REF] Whittingham | Chemistry of Intercalation Compounds: Metal Guests in Chalcogenide Hosts[END_REF][START_REF] Holleck | Transition Metal Sulfides as Cathodes for Secondary Lithium Batteries-II. Titanium Sulfides[END_REF][START_REF] Dines | Lithium Intercalation via N-Butyllithium of the Layered Transition Metal Dichalcogenides[END_REF] ). Unfortunately, this capacity was not recovered upon charging the cells because two Li ions were not reversible, thus rapid capacity fading was observed. This was attributed to the structural instability of layered TiS3 during cycling ( 62 ).

TiS3 + 2Li + +2e-→ Li2TiS3 (Eq. IV -1)

Li2TiS3 + Li + + e-→Li3TiS3 (Eq. IV -2)
In 2014, A.

IV.1. Disordered Li2TiS3 as a Positive Electrode

IV.1.a. Mechanochemical Synthesis of Li2TiS3 Powders

Li2TiS3 samples are prepared by using the same synthesis recipe as in the literature ( 62 ).

Sulfides are known as air sensitive materials; therefore, all samples are prepared in argon filled glovebox.

In the literature, Li2TiS3 powders are synthesized with planetary miller at milling speed of 510 rpm for a range of milling durations (from 20 h to 100 h) ( 62 ). Pure disordered rocksalt phase is obtained for milling durations longer than 40 hours though considerable amount of zirconium contamination is seen in the X-ray diffractogram.

At the beginning of the thesis, our target is to synthesize Li2TiS3 with the same synthesis recipe, but we would also like to improve the purity of the powders. As already mentioned in Chapter II.1, powder milling at longer durations causes undesired reactions and contaminations in the final product as the balls often wear out. Thus, we consider synthesizing the powders at shorter durations. This can be done using slightly bigger zirconium balls (5 mm) and jar (50 ml) than those in the literature (4 mm and 40 ml) with a fix ball-to jar ratio (62:1, in weight). Our synthesis recipe to prepare Li2TiS3 powders is the following one: 0.4363 g lithium sulfide (Li2S, Sigma Aldrich, 98%) and 1.0635 g titanium sulfide (TiS2, Sigma Aldrich, 99.8%) are weighed and placed into 50 ml stainless steel jar that contains 285 zirconia balls (5 mm in diameter). Then, the jar is closed tightly and transferred from an argon-filled glovebox to the planetary ball-milling machine. To find optimum milling conditions, Li2TiS3 powders are prepared using various milling durations: 15h, 20h, 25h and 40h at 510 rpm. After the milling, the jar is transferred from planetary ball-milling machine to the glovebox to avoid the air contamination during powder recovery. Thereafter, samples are analyzed with XRD to identify the phase of the powders. Sample preparation for XRD is done in the glovebox and Kapton film is used to protect the sample from air contaminations. The procedure is described in Chapter II.2.

Figure IV -1: X-ray diffraction patterns of disordered rocksalt Li2TiS3 powders prepared with planetary miller at 510 rpm for 15h, 20h, 25h and 40h durations and SEM picture of powder milled for 20 hours

In Overall, we select the 20 hours milling condition as optimum milling duration for our future synthesis though 25 hours milling in Table IV -2 seems more effective. This discrepancy arises from the fact that the first synthesis is made with fresh balls at 20 hours milling duration. We later find out that there is not a big difference in terms of cell performance and capacity retention between milling 20h and 25h when the milling is operated with used balls. To investigate the real stoichiometry and the purity of the final products, Li2TiS3 powders are analyzed with ICP-AES technique and results are shown in Table IV

-3.
After data is normalized based on fixed Ti amount (Ti=1.00 mol), quite diverse atomic ratios are observed in the powders. We find that powders contain a small trace of zirconium contamination (0.4 -4%). The atomic ratios of lithium and sulfur in the powders increase with milling durations (from 15 hours to 25 hours). However, we observe that atomic ratios obtained from the sample milled at 40 hours are rather out Here, pressed electrodes deliver much higher charge capacity (300 mAh.g -1 ) compared to the unpressed electrodes (250 mAh.g -1 ). In addition, reversible 340 mAh.g -1 charge and discharge capacities are obtained at the end of the second and third cycles. Reduction of electrode porosity by pressing should increase electronic conductivity, thus reversible capacities improve.

Besides, we observe that pressing the electrodes also have a positive impact on cycle life, which is improved considerably. As the pressed electrodes have improved electronic conductivity, we expected that their rate capability should improve as well.

To investigate this, two different cells are charged to 3 V at C/10 rate. (Figure IV -6).

To determine cell behavior at high rates, a signature test is used: the cells are discharged directly at a rate of 10D, when the discharge cut-off voltage is reached, the

Figure IV -6: Cycling curve of signature test: Current (mA) vs Test time (h) and rate capability of pressed Li2TiS3 electrodes that are charged at a rate of C/10 to 3 V and discharged by different discharge (D) rates to 1.5 V (duration of rest time is two hours between two discharge phases)

The reason of poor cycle life Li2TiS3 cells may be due to sulfur reaction with the electrolyte. In the literature, it is reported that lithium sulfur batteries also suffer from poor cycle life ( [START_REF] Pang | Advances in Lithium-Sulfur Batteries Based on Multifunctional Cathodes and Electrolytes[END_REF][START_REF] Safa | Capacity Fading Mechanism in Lithium-Sulfur Battery Using Poly(Ionic Liquid) Gel Electrolyte[END_REF][START_REF] Yan | Capacity Fade Analysis of Sulfur Cathodes in Lithium-Sulfur Batteries[END_REF] ). That is attributed to reaction of sulfur with carbonate-based electrolytes, therefore tetraglyme-dioxolane-based electrolytes are suggested ( 97 ). In our experiments, we decide to test Li2TiS3 electrodes in tetraglyme-dioxolane based electrolytes as well.

IV.1.c. Air Sensitive Property of Li2TiS3

Li2TiS3 is an air-sensitive material, thus this property slows down the rate of process from material synthesis to coin-cell preparations because it requires always working under inert atmosphere, often in an Ar-filled glovebox. For industrial perspective, material stability is important, especially for upscaling process.

In this part, the air-sensitivity of Li2TiS3 is examined. Five different electrodes are prepared for XRD analysis in the glovebox, later they are left in a fume hood and a dry room for one day and one week durations. Here, working under fume hood represents working under air condition. As the dry room has less humidity (dew point: -40 o C) compared to air atmosphere, working in dry room conditions will show oxygen sensitivity rather than humidity. Such results indicate that Li2TiS3 material is more sensitive to humidity and it is rather stable under oxygen environment. This also indicates the possibility of production and slurry coating of this material in dry room.

Figure IV -8: X-ray diffraction patterns of Li2TiS3 electrodes at different environments: pristine electrode, after 1 day in dry room, after 1 week in dry room, after 1 day in fume hood, after 1 week in fume hood

Figure IV -9: (a) Coin cells tests results of standard Li2TiS3 electrodes (black) in a comparison of electrodes after 1 day (green) and 1 week (red) in dry room

Figure IV -10: Coin cells tests results of Li2TiS3 electrodes after 1 day (blue) and 1 week (black) in fume hood

• Lithiation of Li2TiS3 with Butyl-lithium Reaction

As previously mentioned regarding Li2TiS3 electrodes, there is always extra capacity at the first cycle, thus the final composition is different from the initial composition.

Such compositional deviation disappears at the end of the second cycle, fully reversible lithium intercalation is observed. To eliminate the irreversibility at the first cycle and increase charge-discharge capacities, we aim at preparing Li2TiS3 in its more lithiated version, (Li3TiS3). In early studies on TiS2 and TiS3 systems, Dines et al. reported that butyl-lithium reaction was an effective way for lithiation of the positive materials ( 6).

Therefore, lithiation of Li2TiS3 is made by butyl-lithium experiment using the same procedure as in the literature.

• Experimental procedure of Butyl-lithium reaction

Based on the reaction in Eq. IV -3, 0.025 mol Li2TiS3 powder, 14.60 ml cyclohexane and magnetic stirrer are put into a small reactor (20 ml). 0.0375 mol butyl lithium (50% excess) is added slowly into the reactor, the reactor is closed with its cover and left to stir at 450 rpm in the glovebox (Figure IV -11a).

Butyl-lithium reaction:

nC4H9Li + Li2TiS3 → Li3TiS3 + ½ C8H18 (Eq. IV -3)
Mechanochemical reaction:

3/2Li2S + Ti + 3/2S → Li3TiS3 (Eq. IV -4)
After stirring for a week, the mixture is transferred into small flasks ( 2 This may be caused by the unsuccessful coating, because slurry gelified during preparation. It seems that unreacted Li2S stays on the particle surface and reacts with PVdF binder.

Figure IV -12: XRD pattern of Li2+xTiS3 comparison of butyl lithium reaction vs ball milling

IV.2. Selenium Substituted Li2TiS3

As already mentioned, the cycling stability of Li2TiS3 is rather poor and therefore it should be improved. We consider that substitution can be a promising solution. For instance, sulfur anions can be replaced with an alternative anion. Looking at the periodic table, selenium comes into prominence as a novel prospected material since it belongs to the same periodic group (6A), hence it is expected to show similar elemental activity as sulfur. In the literature, it is also well-known that selenium substitution improves cycling stability of Li2S cells ( [START_REF] Eftekhari | The Rise of Lithium-Selenium Batteries[END_REF][START_REF] Hu | Confining Selenium Disulfide in 3D Sulfur-Doped Mesoporous Carbon for Rechargeable Lithium Batteries[END_REF] ). However, before the synthesis, one must consider the theoretical capacity of new compositions. Theoretical capacities based on two lithium ions exchanged during cycling are reduced by selenium substitution since atomic mass of selenium (78.97 g.mol -1 ) is much bigger than atomic mass of the sulfur (32.06 g.mol -1 ) (Table IV -4). Therefore, we first consider the following compositions Li2TiSeS2 and Li2TiSe2S.

Table IV -4: Theoretical capacities of Se substituted compositions based on two Li + ions exchange process

To synthesize selenium substituted Li2TiS3 samples, necessary precursors are searched from the producers, however neither Li2Se nor TiSe2 are commercially available.

Therefore, alternative synthesis routes are searched. In the literature, the same precursor problem is mentioned in the development of Li2Se batteries and powders are synthesized by liquid synthesis routes that are quite challenging and require more than one step ( 98 ). To make one pot and much faster synthesis, we aim at preparing new compositions with mechanochemical synthesis using metallic precursors, such as metallic Ti and Se precursors rather than molecular TiSe2 precursor.

IV.2.a. Synthesis Improvements

Before preparing new materials, it is essential to adapt our previous synthesis of Li2TiS3

to the use of metallic Ti and S precursors instead of TiS2. In this way, we can compare this new synthesis route with the standard route where there are already enough structural and electrochemical characterizations. In fact, TiS2 precursor is quite expensive (22 €.g -1 ) compared to metallic Ti and S precursors (Ti: 6 €.g -1 , S: 2 €.g -1 ), therefore using metallic precursors would also bring benefits in terms of manufacturing costs. However, it should be noted that these prices are obtained from the Sigma Aldrich for lab-scale powders, therefore they may not be representative for industrial powders.

• Wet Mechanochemical Synthesis of Li2TiS3

At the beginning, we aimed at synthesizing Li2TiS3 powder with Li2S powder, metallic Ti and S powders. Our four different synthesis trials show that metallic precursors do not react with Li2S precursor, they are rather deposited on the jar surface. To eliminate this, we consider milling of the precursors with anhydrous solvent. In our first trial, we try to synthesize Li2TiS3 through a wet mechanochemical synthesis for 20 h and 25 h milling durations. However, we always detect traces of remaining Li2S precursor with rocksalt phase in XRD patterns. Pure rocksalt Li2TiS3 phase is finally obtained when the powder is wet milled for 32 hours with already used balls ( 100 ).

Synthesis Procedure in detail: 0.436 g Li2S (98%, Sigma Aldrich), 0.435 g Ti (99.98%, Alfa Aesar) and 0.609 g S (99.99%, Alfa Aesar) precursors are placed into 50 ml stainless steel jar that contains 285 zirconia balls (5 mm in diameter). Thereafter, anhydrous hexane (99.9%, Sigma Aldrich) is poured into the jar until covering the entire ball surface. Milling is operated at 510 rpm for 32 hours. The milling run time is set as 1 min, but pause time is set as 4 min. After milling, jar is transferred from planetary ball milling machine to glovebox. To evaporate hexane, jar remains open for 3 days in the glovebox. Then, powder is recovered in the glovebox. With this new process, different metal sulfides, Li2MS3 or selenides, Li2MSe3 (M: Fe, Ni,

• Wet Mechanochemical Synthesis of Li2TiSeS2 0.336 g Li2S (98%, Sigma Aldrich), 0.410 g TiS2, 0.175 g Ti (99.98%, Alfa Aesar) and 0.578 g Se (99.9%, Alfa Aesar) precursors are placed into 50 ml stainless steel jar that contains 285 zirconia balls (5 mm in diameter). Then, appropriate amount of anhydrous hexane (99.9 %, Sigma Aldrich) is added until covering the entire ball surface. Milling is operated at 510 rpm for 20 hours.

• Wet Mechanochemical Synthesis of Li2TiSe2S 0.274 g Li2S, 0.285 g Ti (99.8%, Alfa Aesar) and 0.941 g Se (99.9%, Alfa Aesar) precursors are placed into 50 ml stainless steel jar that contains 285 zirconia balls (5 mm in diameter). Then, appropriate amount of anhydrous hexane (99.9 %, Sigma Aldrich) is added until covering the entire ball surface. Milling is operated at 510 rpm for 20 hours.

After milling, jar is transferred from planetary ball milling machine to glovebox. To evaporate hexane, jar remains open for 3 days in the glovebox. Then, powder is recovered in the glovebox. During discharge, more than 2.0 Li + ions are inserted into the structure, the composition changed into Li3.08TiSeS2 and 310 mAh.g -1 is delivered. Discharge capacity should be provided by either anionic (S 2-/S2 2-or Se 2-/Se2 2-) or both anionic and cationic (Ti 4+ /Ti 3+ ) redox process. The second and third cycles are fully reversible similar to Li2TiS3 and cycling life is relatively better than Li2TiS3, 80% of capacity is reached at the end of 10 th cycle.

Coin cells tests results at the first three cycles of Li2TiSe2S upon cycling at a rate of C/10 between 3 V -1.5 V vs Li + /Li show the same trend that is observed for Li2TiSeS2 electrodes. During charge, 1.32 Li + ions are extracted from the structure and capacity is recorded as 149 mAh.g -1 . In contrary to Li2TiSeS2, much more lithium ions, 3. 44 Li + ions, are inserted into the cubic rocksalt structure and a much higher discharge capacity (379 mAh.g -1 ) is obtained. The final composition becomes as Li(4.17)TiSe2S, which can be regarded as the average (theoretical) composition. More than 3 Li + uptake cannot be explained with the same hypothesis, Ti 3+ /Ti 4+ redox process. Therefore, this should be investigated with further characterizations. Besides, we observe rapid capacity fading in Li2TiSe2S cells, 80% of capacity is reached at the end of 2 nd cycle.

Table IV -7: Summary of structural and coin cell test results of Li2TiS3, Li2TiSeS2 and Li2TiSe2S compositions: lattice parameters and cells cycled in galvanostatic mode at a rate of C/10 between 3 V -1.5 V vs Li + / Li

Table IV -7 shows a summary of structural and coin cell test results of for Li2TiS3, Li2TiSeS2 and Li2TiSe2S compositions. For Li2TiS3, charge capacity is smaller, whereas discharge capacity is equal to theoretical capacity (339 mAh.g -1 ) since rocksalt structure allows two Li + reversible exchange. For Se substituted samples, smaller charge capacities are obtained at the first cycle. On the contrary, discharge capacities are greater than the theoretical capacities (261 mAh.g -1 for Li2TiSeS2, 213 mAh.g -1 for

Li2TiSe2S). This indicates the presence of extra capacity at the end of the discharge.

Besides, there is a correlation between degree of Se substitution and the formation of extra capacity at the end of the discharge. Li2TiSe2S shows the highest discharge capacity. Given average potentials represent an average of charge and discharge potentials during 10 cycles. Average potentials decrease with selenium substitution, for instance Li2TiSe2S showed the lowest charge and discharge average potentials (2.24V and 1.98 V respectively).

• Cyclic Voltammetry Tests of Li2TiSexS3-x

To describe electrochemical processes and identify the electron transfer kinetics, cyclic voltammetry tests of Li2TiS3, Li2TiSeS2 and Li2TiSe2S cells are conducted at various scan rates (5 µV.s -1 , 10 µV.s -1 , 20 µV.s -1 , 50 µV.s -1 and 100 µV.s -1 ) between 3 V -1.5 V vs Li + / Li. As we compare the theoretical and experimental peak potential separations, we find that experimental peak separation values of both sulfur (∆Esulfur) and selenium (∆Eselenium) redox processes are much higher than theoretical values (28.2 mV for Li2TiS3, 18.8 mV for Li2TiSeS2 and 14.1 mV for Li2TiSe2S), indicating that sulfur and selenium redox processes are not completely reversible. In Figure IV -24e, we observe that peak potential separations (∆Eselenium) corresponding to selenium redox process in Li2TiSexS3-x cells increases with increasing scan rate. On the contrary, peak potential separations (∆Esulfur) corresponding to sulfur redox process is almost stable at higher scan rates (from 50 µV.s -1 to 100 µV.s -1 ) in Figure IV -24d. By comparing the two different redox processes in (Li2TiSexS3-x) cells, we can conclude that electron transfer kinetics in selenium process is much more sensitive to the scan rate.

Cyclic voltammetry can also allow determining the diffusion property of Li2TiSexS3-x cells. The peak potentials at each scan rate can be used to calculate the lithium ion diffusion coefficients using the relationship between measured current and scan rate, between peak current and square root of scan rate (slope= 0.5, R² ≥ 0.99) among

Li2TiSexS3-x cells. The relation between peak current and scan rate in Li2TiSeS2 and Li2TiSe2S cells goes in a nonlinear direction (slope> 0.5). For a diffusion controlled process, peak current is related to the square root of scan rate by i=v 1/2 , thus it allows applying Randles-Sevcik equation 102,103 (Eq. IV -8)

𝐼 𝑝 = (2.69 𝑥 10 5 )𝑛 3/2 𝐴𝐷 0 1/2 𝐶 0 𝑉 1/2 (Eq. IV -8)
where Ip is peak current in A, n the number of exchanged electrons, A the active surface in cm 2 , Co the lithium concentration in mol.cm 3-and V the scan rate in V.s -1 .

Figure IV -25b also shows the linear relation between peak current and scan rate with the corresponding regression line (R²) for Li2TiS3. The concentration (Co) is calculated as 0.06 Li + mol.cm -3 from the unit cell volume and we consider two Li + ions are exchanged during charge-discharge. Using these variables in Eq. IV -8, anodic diffusion coefficient (Do) of Li2TiS3 is calculated as 1.99x10 -9 cm 2 .s -1 . Lithium diffusion kinetics of different sulfurous compounds are studied in the literature and the diffusion coefficients are reported between 10 -9 -10 -8 ( 106-108 ). Therefore, DLi+ that we found is close to the reported values.

We further investigate capacity retention of Li2TiS3, Li2TiSeS2 and Li2TiSe2S cells cycled at different scan rates (Table IV - The goal of the selenium substitution was to improve the cycling stability of Li2TiS3, therefore, cycling stabilities of Li2TiSeS2 and Li2TiSe2S should be compared with Li2TiS3.

IV.3. Conclusion

In this chapter, Li2TiS3 powder is synthesized mechanochemically in rocksalt phase and it shows reversible charge and discharge capacities of 339 mAh.g -1 based on two lithium ions exchange process. To understand degradation process of Li2TiS3 in liquid cells, various electrochemical and structural characterizations are performed. To improve cycling stability of Li2TiS3 cells, selenium substitution is considered. Applying the same synthesis recipe, selenium-substituted samples could not be synthesized, thus the synthesis recipe had to be improved. After various synthesis strategies, a new synthesis process is successfully developed and patented ( 100,101 ).

Selenium-substituted samples (Li2TiSeS2 and Li2TiSe2S) are synthesized in disordered rocksalt phase with new synthesis process. To understand their cell performance, cycling stability and electron transfer kinetics, galvanostatic and voltammetry tests are conducted. Results show that Li2TiSeS2 and Li2TiSe2S provide larger discharge capacities than Li2TiS3 cells. Cycling voltammetry test of selenium-substituted samples at various scan rates show that oxidation and reduction potentials decrease by selenium substitution. At high scan rates in voltammetry, selenium-substituted samples show better capacity retention than Li2TiS3 cells. We also find that electron transfer kinetics of selenium-substituted samples are more sensitive to the scan rate.

Afterwards, electrochemical performances of Li2TiSexS3-x electrodes are investigated at different cycling schedules. Capacity retention of Li2TiSe2S cells is poor upon cycling between 3 V -1.5 V; however, we observe that it improves upon cycling between 3 V -1.8 V and 3 V -2.0 V. By examining the cycled electrodes with ex situ XRD and SEM analyses, we find that Li2TiSe2S electrodes degrade during discharging the cells.

Nevertheless, increasing discharge cut-off potential might limit the parasitic reactions that forms upon discharge.

Finally, Li2TiSexS3-x electrodes are also tested between 1.5 V -3 V. Only limited discharge capacities are detected in Li2TiS3 cells, whereas large discharge capacities are

Chapter V: Characterization of Li2TiSexS3-x Compositions

In previous chapter, our new rocksalt sulfide-selenide compositions show unusual electrochemical performances. We observe large discharge capacities in Se substituted samples at the first cycle. This could be due to extra Li uptake at the end of the discharge. Cubic rocksalt structure is dense and all the octahedral sites are already filled with both Ti 4+ and Li + cations, however it is still questionable where the extra lithium ions go in the structure. We also observe that cycling stability of Se substituted samples is rather poor during charge-discharge. Therefore, our next target is to find what impedes the capacity during cycling.

This chapter describes electrochemical and structural characterization of sulfurselenium compositions (Li2TiSexS3-x). Operando neutron diffraction and ex situ X-ray diffraction are described to detect Li ions in the rocksalt structure. Afterwards, results

of XPS analysis are presented to elucidate the electrochemical activity in those materials.

V.1. Operando Neutron Diffraction of Li2TiS3

Operando neutron diffraction is performed at ILL, Grenoble. We use a special cell, "neutron cell" made of Ti2.08Zr alloy, which is transparent to neutron diffraction. The experimental procedure is described in Chapter 2. At the beginning, we prepare neutron cell as it is described in Chapter 2.b., for instance the cell is tested at a rate of C/10 between 3 V and 1.5 V. However, it cannot be charged to 3 V and a wide potential plateau at 2.6 V with a large polarization (0.68 V) is observed (Figure V -1a). We suspect that the parasitic plateau at 2.6 V might be due to the moisture and a parafilm that is added around the nut to improve the cell tightness.

As our main interest in this experiment is to detect extra capacity and to find out where the extra lithium locates in the structure, we are more interested in discharge state. To produce similar results, the neutron cell is first discharged to 1. After those results, we have two different hypotheses:

1. Is there Ostwald ripening? We can imagine that some unreacted Se remains at the surface of the material and it has a very small size, therefore it is not seen by XRD. Then, when put in the electrolyte, it would dissolve and precipitate into bigger particles following an Ostwald ripening process ( 109 ).

Therefore, diffraction peak of unreacted Se would become visible due to the bigger size of the particles.

2. Is there a selenium redox activity? If selenium is redox active, the possible redox process should be Se 2-/ Se 0 that is different from sulfur redox process (S 2-/S2 2-). Therefore, Se 0 could be produced at the end of the charge and converted into soluble polyselenides containing Se 2-at the end of the discharge.

To test these hypotheses, we decide to perform a high-resolution XPS study on rocksalt sulfide-selenide pristine powders and electrodes at different states of charge.

V.3. XPS Analysis of Li2TiSexS3-x Compositions

Elucidating the electrochemical activity in those materials requires an understanding of: what species are present in the electrodes, which species are responsible for redox reactions and how those species change upon charge-discharge. Moreover, electrochemical and structural characterizations of Li2TiSexS3-x show that not only sulfur but also selenium should be redox active. To test this hypothesis, we use a complementary technique, a high-resolution X-ray Photoelectron Spectroscopy (XPS, introduced in Chapter 2.2.e), which is a surface sensitive technique that gives information about both chemical states in the compositions and does semi-quantitative analysis of the compositions. In this chapter, we present the results of a high-resolution XPS study on disordered rocksalt Li2TiSexS3-x pristine powders and electrodes at different states of charge. We also prepare tables that contain optimized parameters, such as peak positions and widths at different core levels. We believe that such results can help scientists to understand the redox processes in metal sulfides or selenides since only few publications are present in the literature.

V.3.a. Methodology: Analysis of XPS Spectra

Peak identification/indexation: In general, the peaks obtained from the XPS analysis are indexed based on the information in XPS Handbook 110 and the online XPS database 111 that contains binding energy of species in various compositions. As we are the first ones who designed sulfur-selenium compositions (Li2TiSexS3-x) and the binding energy of elements differs based on the electronic environments, we not only use the data from XPS handbook and database but we also consider similar studies in the literature to prepare fitting models.

Strategy for peak deconvolution:

To extract the hidden information from the peaks, firstly, we take the peak fittings parameters, such as peak positions, peak ratios and peak widths from the XPS handbook and literature. The latter, we fix these parameters in the Multitap program that is used to deconvolute the peaks obtained from the High-Resolution XPS spectra. Afterwards, we build a fitting model that is iteratively improved until fitting parameters do not change within a specified precision (Chi<5).

The curves are fitted using 80% Gaussian, 20% Lorenzian conditions in Multitap

Program.

• Single peaks: deconvoluting Li 1s, Se 3d and Ti 3s: The binding energy of these three core peaks are quite close (Li 1s:54-55 eV, Se 3d: 52-53 eV Ti:3s 60-61 eV)

and therefore they can be detected together at the same core level 110 . For instance, Ti 3s and Se 3d can be detected at Li 1s core. Here, each core peak is fitted and evaluated individually for quantitative elemental analysis.

• Overlapping peaks: deconvoluting Ti 2p and S 2p orbitals: For overlapping peaks, fitting is employed based on the contributions (two peaks) originated from spin orbit splitting of 2p cores (3/2p and 1/2p). Each peak is fitted within the same model considering fixed peak separations 110 and the ratio between two peaks are calculated considering these two contributions at the same core level.

We follow the same procedure for fitting of S 2p core peaks.

Quantitative elemental analysis is performed by accounting peak intensities (peak area after background removal) and the sensivity factors (Chapter 2.2.f). It gives atomic concentrations in the analyzed samples. For elemental analysis, we account of Li + in Li 1s core, Ti 4+ in Ti 2p core, S 2-(could be S2 2-) in S 2p core for Li2TiS3. It should be noted that Ti 4+ species can be seen at both Ti 2p and Ti 3s core levels, but the amount of titanium species at Ti 2p core are different from those in Ti 3s core. The difference arises from the cross-section of each elements. The number of photons obtained from Ti 2p core is much higher than those coming from Ti 3s core; thus, we only consider one contribution (Ti 4+ in Ti in 2p core) for elemental analysis. Similarly, we account Se 2- (could be Se2 2-) at Se 3d core instead of Se 3p core for selenium compositions.

V.3.b. Survey analysis of Li2TiSexS3-x samples

A survey analysis is performed to identify the elements contained in the composition with a survey spectrum that is recorded over a spectral range of 0-1200eV. This quick analysis allows us determining which core of the elements should be analyzed through XPS and what parameters should be used. Thereafter, we can optimize the quality of signal/noise ratio for High-Resolution XPS analysis. Nevertheless, it should be noted that such survey analysis gives general information about species in the compositions.

For detailed analysis, the samples are analyzed further with High-Resolution XPS (penetration length of X-rays up to 5 nm).

As illustrated in • The first target is to determine active species at specific binding energies and distinguish them in their different electronic environments.

• Thereafter, fitting models are developed for powders and then same models are used to analyze the electrodes of the same compositions.

At the beginning, one of our hypotheses was that there might have been a slight trace of unreacted selenium present in the sulfur-selenium solid solutions, therefore Se metal would have become visible in X-ray diffractograms after peak intensities reduced at the end of the charge. To test our hypothesis, we analyze three pristine powders and Se metal precursor through XPS. peak is much higher compared to intensity of Li + peak. This could be due to the elemental sensitivity to XPS analysis, for instance sensitivity factor of Li (0.02) is much smaller than sensitivity factor of Se (0.67). Besides, binding energy of Se 2-is rather smaller than Se 0 , core peak of Se 0 is detected at 56 eV. Similar binding energy of Se 0 is reported in the literature as well ( 112 ). Overall, no signature of selenium precursors is detected in XPS spectra. Such result is not aligned with our first hypothesis, consequently we conclude that there is no metallic selenium remaining in the powders after the synthesis. In Li2TiS3, we observe Li 1s core peak at 54.6 eV with a peak width of 2.1 eV. Previous studies reported similar binding energy of Li 1s core peak in Li-S systems ( 113,114 ). In the same energy window, we also observe Ti 3s at 60.5 eV with almost the same intensity as Li 1s core peak (52.4%), but a much wider peak width (3.8). We use the same peak positions, peak widths and relative intensity of Li 1s and Ti 3s core peaks for fitting selenium-substituted compositions at Li 1s and Se 3d core levels. In conclusion, we do not observe unreacted selenium in XPS analysis on powders.

V.3.c. High Resolution XPS Analysis of Li2TiSexS3-x Powders

Instead, we observe one type of Li + , but two different selenide types (Se 2-and Se2 2-) in this energy region. Besides, Ti-S and Ti-Se peaks are not easily distinguished because their binding energy are very close ( 9 ). Therefore, they are considered as one single Ti peak. In both Li2TiSeS2 and Li2TiSe2S compositions, we observed similar binding energies and peak widths of Ti 2p core levels. Here, Ti 4+ is present in sulfur-selenium, [Ti(S-Se)2] and oxygen (TiO2) environments. In addition, correlation between Se substitution and Ti 4+ in oxygen (TiO2) environment may exist because the more Se is present in the composition, the greater intensity of Ti 4+ in oxygen (TiO2) environment is detected.

Selenium compounds thus seem more air-sensitive.

• 171eV-150eV: S 2p and Se 3p cores

In Li2TiS3, we observe two different sulfur doublets (S1 and S2) at 159.9 eV and 160.6 eV with widths of 1.2 eV. According to semi-quantitative analysis, S1 presents 49.8%, while S2 presents 50.2% in Li2TiS3 composition. In the literature, many metal sulfides contains two or three sulfur species at S 2p core and they are regarded as terminal and bridge sulfur species ( 122-124 ). Since the energy difference between two different sulfur species is quite small (< 1 eV), these sulfides should be regarded as S 2-(S1) and S2 We observe LiF peak at Li 1s core. That may come from reaction of sulfur with binder.

It has stronger peak intensity in non-scratched electrode, indicating that LiF is mostly present on the electrode surface rather than in the bulk. In S 2p core, additional sulfur species (S-C) are detected in Li2TiS3 electrodes, unlike Li2TiS3 powder. That is attributed to sulfur-carbon bond in the literature since electrodes contain important amount of carbon (10%) to increase their electronic conductivity ( 123,124 ). This explains why the S-C species was not detected before in sulfur-selenium powders. Unlike powders, we also observe that the amount of S 2-reduces in electrodes (10%), whereas the amount of S2

V.3.e. XPS Analysis on Li2TiSexS3-x Electrodes at Different States of Charge

In this part, our target is to find the appropriate answer to our second hypothesis. As we observed Se metal at the end of charge, our second hypothesis was that there should have been a selenium redox activity and it occured from the change of Se 2-into Se 0 , that is why we observed Se metal at the end of the charge. According to our XPS analysis on powders, we found out that there is no Ostwald ripening since no signature of selenium precursor is found in XPS pattern of Li2TiSexS3-x compositions. However we still need to verify the origin of extra capacity and elucidate redox process of Li2TiS3 composition. In this purpose, electrodes are cycled using the standard procedure described in Chapter 2.3, dismounted and prepared for High resolution XPS as described in Chapter 2.2.

Li2TiS3 electrodes at different states of charge

At S 2p core, three different sulfur species are observed in pristine electrode, shown in electrode, all the sulfur species are found at the same binding energies as those of the pristine state. This indicates that sulfur redox process is reversible. species are detected at S 2p core level.

Li2TiSeS2 electrodes at different states of charge

At the end of charge, lithium ions are extracted from Li2TiSeS2 electrode, sulfur oxidizes, S 2-shifts to higher binding energy and peak intensity reduces down to 22%.

On the contrary, the changes in binding energies of Se 2-and Se2 2-are quite small and we observe that peak intensity of Se2 2-increases (41%) compared to Se 2-(12%). These results seems to confirm an oxidation of the anionic network, however right now, we are not able to confirm the exact selenium redox process because the deconvolution at S 2p is quite difficult due to the high number of contributions with very close energies (See Appendix).

At the end of the discharge, all the sulfur (S 2-) and selenium species (Se 2-and Se2 2-) are detected at the similar binding energies as those of the pristine state of Li2TiSeS2 electrode. It seems that selenium redox is also reversible.

V.4. Discussion on the XPS Results

In pristine Li2TiS3 electrodes, we consider S 2-as terminal disulfide species, whereas S2 showed that sulfur atoms that lose at least one lithium counter made S-S bonds (S 2-) ( 64). With formation of disulfides, the structural volume is increased. In this way, rocksalt Li2TiS3 has undergone smooth structural transition from crystalline to amorphous to crystalline structure, without direct crystalline to crystalline structural transition. At the end of the discharge process, we observe the same species at similar binding energy, this result is consistent with what we observed in ex situ XRD results previously.

In pristine Se-substituted electrodes (Li2TiSeS2 and LiTiSe2S), we detect central disulfide species (S2 2-) and selenium species (Se2 2-) with diselenides (Se 2-) without any terminal disulfides ( 115,122,125 ). Like sulfide electrode (Li2TiS3), central disulfide species (S2 2-) and selenium species (Se2 2-) lose some lithium ions and Ti-S and Ti-Se bonded sulfur and selenium species (S2 2-, Se 2-) remain at the end of the charge (3 V). Differently from sulfide electrode (Li2TiS3), diselenides species tend to make bonds with central disulfide species and produce Se-S bonds, instead of producing diselenides at the end of the charge. That is why we observe the increase of selenium species (Se2 2-) instead of diselenides (Se 2-) in the XPS spectra. This also indicates that the possible redox mechanisms between sulfur and selenium is quite different.

In previous chapter, we show that selenium redox activity appears during first discharge at 1.88 V vs Li + /Li for both Li2TiSeS2 and Li2TiSe2S electrodes. Therefore, here we only consider sulfur redox process, the shift of S2 2-species from low binding to high binding energy creates available charge capacity. At the end of the discharge, we detect similar species in selenium-subsituted samples (Li2TiSexS3-x), but it seems that more diselenides are produced. Here, we consider that disulfides and diselenides do not contribute to charge-discharge capacity, thus they are redox inactive species. If we assume continuous production of diselenides in Li2TiSe2S cells at the end of the discharge, this might also explain why we observe huge capacity fading in the Li2TiSe2S cells previously.

According to our second hypothesis, the possible redox mechanisms of selenium should be undergone from Se 2-to Se 0 . With the current study, it is quite difficult to determine this point as the actual redox process of selenium starts at the first discharge process. Therefore, further ex situ XPS analysis on Li2TiSexS3-x electrodes at different states of charges (2 nd and 3 rd cycles) are required. In sulfur 2p core level, there are high number of contributions coming from both sulfur and selenium species. To make clear observations on selenium redox process, it might be better to work with only selenium species. For instance, a new composition Li2TiSe3 might be developed and electrodes at different states of charges might be analyzed. For a comparison, high performance pure Se electrodes might also be developed and analyzed with the same approach.

V.5. Conclusion

In this chapter, we focused on characterization of Li2TiSexS3-x through electrochemical, structural and morphological studies. We observed that these materials provide extra capacity during discharge. To understand the origin of extra capacity and the reason of degradation process upon cycling, various structural characterizations were done.

Ex situ XRD showed that cubic rocksalt structure reversibly changed upon cycling. In addition, metallic selenium was observed in diffraction pattern and a columnar phase was observed in SEM images of selenium-substituted samples. After such results, we had hypotheses, such as either unreacted selenium remains in the electrodes or selenium redox process is present (Se 2-/ Se 0 ).

To investigate our hypotheses, we performed ex situ XPS study. Our first hypothesis failed because we did not observe unreacted selenium in powders. XPS analysis on Li2TiS3 showed that sulfur redox process is reversible. In selenium-substituted samples, it seems that selenium redox is also reversible, however we cannot explain and confirm the exact Selenium redox process at this point. The deconvolution at S 2p is quite difficult due to the high number of contributions with very close energies, thus further analysis is necessary.

Finally, the presence of LiF at Li 1s core and more oxidized sulfur species at S 2p core in pristine electrodes suggest that sulfur might react with PVdF binder and produce more oxidized sulfur species. This hypothesis should be investigated by further XPS study on Li2TiSexS3-x electrodes prepared with different binders. It may also help to clarify the degradation process upon cycling.

The purpose of this thesis study is to develop overlithiated disordered rocksalt materials that provide high specific energies (> 900Wh.kg -1 at material level) and present a cobalt-free alternative to conventional materials used in Li-ion batteries. Lirich cation disordered rocksalt oxides can offer large capacities (≥ 250 mAh.g -1 ), however there is always irreversibility in the first cycle and origin of this irreversibility is still unknown. Meanwhile, disordered rocksalt Li2TiS3 has attracted research interest thanks to its excellent reversible capacities (425 mAh.g -1 at 2.3 V vs Li + /Li). However, detailed electrochemical and structural studies of such a promising material are still missing, especially redox properties have not been reported yet. In addition, the question of why Li2TiS3 cells failed in traditional batteries with liquid electrolytes has not been answered yet. Therefore, we were interested to investigate those properties and examine the reason of degradation during charge-discharge.

In this thesis project, we have studied overlithiated disordered rocksalt family in two different axes: disordered rocksalt oxide (Li2.2NiTi0.2Nb0.6O4) and sulfide (Li2TiSexS3-x).

In the first part of the thesis, our aim was to optimize the performance of our patented composition, disordered rocksalt Li2.2NiTi0.2Nb0.6O4. The disordered rocksalt (cubic) structures are generally stabilized at high temperatures (1000 o C), however the performances are limited due to particle growth. Thus, an optimization of the synthesis parameters of our molten salt synthesis route was carried out. The nature of the eutectic salt, the washing solvent, as well as the duration and the temperature of the synthesis process were investigated. An optimum was found using 0.59LiCl-0.41KCl eutectic salt at 670°C for 1h followed by a quench, and ethylene glycol as washing media. The material delivered 166 mAh.g -1 of charge capacity with 108 mAh.g -1 of reversible capacity. In situ XRD was performed to investigate the irreversibility (35%)

at the first cycle. The result showed that Li2.2NiTi0.2Nb0.6O4 has undergone structural evolutions during charge-discharge. In addition, the particle size analysis showed that dislocations or defects might have arisen in Li2.2NiTi0.2Nb0.6O4 powder, this phenomenon may have been the cause of the irreversibility. This result indicated that purity of the disordered rocksalts systems is important to design high energy materials.

In the second part of the thesis, disordered rocksalt Li2TiS3 powders were prepared via mechanochemical synthesis for shorter milling duration (15h -25h) than those in the literature (40h -100h). As the Li2TiS3 electrodes were tested in carbonate-based electrolyte, they delivered a reversible of 250 mAh.g -1 , however a capacity decay starting form the second cycle was observed. Thereafter, electrode pressing was considered to improve its electronic conductivity and we observed that both reversible capacities (339 mAh.g -1 ) and cycling stability were improved. To further improve cycling stability of Li2TiS3 cells, the partial substitution of Li2TiS3 with selenium was considered. Our efforts to prepare Li2TiSexS3-x compositions have resulted in the discovery of a new synthesis method and high capacity compositions (Li2TiSeS2 and Li2TiSe2S). X-ray diffraction patterns showed that all Li2TiSexS3-x powders that were prepared in cation disordered rocksalt phase follow a solid solution behavior.

Thereafter, electrochemical properties of new materials have been investigated in comparison with Li2TiS3. Li2TiSeS2 and Li2TiSe2S cells showed better cycling stability than Li2TiS3 cells at higher cut-off discharge potentials (1.8V and 2 V) while cycling stability of Li2TiS3 cells did not improve even at higher cut-off potentials. Furthermore, Li2TiSe2S cells failed upon cycling between 3 V and 1.5 V. To find out the reason of degradation during cycling, we decided to investigate Li2TiSexS3-x with deep structural characterizations.

In the last part of the thesis, we examined the lithium insertion-extraction mechanism through disordered rocksalt Li2TiS3 with structural studies (operando neutron diffraction, FTIR and Raman). Unfortunately, poor crystallinity of Li2TiS3 powder made definitive conclusions difficult even though various characterization tools were used. Afterwards, Li2TiSeS2 and Li2TiSe2S electrodes were characterized by combining both ex situ XRD and SEM techniques. Selenium metal was observed in the charged electrodes, but it could not be detected in both pristine and discharged electrodes.

Thereafter, we had a hypothesis that there might have been Oswald ripening or Selenium redox activity during charge. To tests these hypotheses, Li2TiSexS3-x powders and Se precursors were further analyzed with X-Ray Photoelectron Spectroscopy (XPS)

analysis. The result showed that there were no remaining selenium precursors in the Li2TiSexS3-x powders, thus our first hypothesis related to Oswald ripening was not true.

Thereafter, we decided to discover redox active species and elucidate redox processes in Li2TiSexS3-x by XPS. At first, we identified species (elements) present in Li2TiSexS3-x at their core levels and developed a fitting model for both powders and electrodes.

Surprisingly, high amounts of oxidized sulfur species at 2p sulfur core and LiF species at Li 1s core were detected in pristine electrodes, showing the possible reaction of sulfur in the Li2TiSexS3-x with the PVdF binder during electrode preparations. Ex situ XPS analyses of Li2TiSexS3-x electrodes at different states of charge have confirmed our previous observations that sulfur redox process was reversible. However, it is still quite difficult to identify exact redox process in Li2TiSexS3-x due to the presence of a high number of contributions of selenium species. We suggest that further XPS analyses on Se metal or Li2TiSe3 can be useful to elucidate redox process in Li2TiSexS3-x samples.

Overall, the degradation process of Li2TiS3 in liquid cells is still a question even though various structural characterization tools were used to understand the underlying reason. Further XPS analysis of Li2TiS3 electrodes at different states of charge and during cycling (after 1 st , 5 th , 10 th and 20 th cycles for instance) might be useful to clarify this issue. The degradation might result from the Li metal anode as the lithium metal is very reactive. Thus, testing the Li2TiS3 electrodes against graphite or silicon type negative electrodes can be interesting. In this thesis study, we could only investigate reversible sulfur (anionic) redox process but did not focus on cationic process as the titanium redox is generally active at lower potentials. To determine titanium redox activity, various ex situ XPS experiments can be performed at Li2TiS3 electrodes cycled at lower discharge cut-off potentials. XPS could also be useful to examine the true selenium redox process and explain the main differences between sulfur and selenium redox processes. To this purpose, a new Li2TiSe3 electrode can be characterized at different states of charge. Consequently, different high energy compounds, such as Liexcess disordered sulfides (Li2MS3 M: Ni, Co, Fe, Mn) or chalcogenides (Li2TiX3 X: S, Se, Te) can be prepared with our patented synthesis process and characterized with the techniques that were reported in this thesis.

Figure 0 - 1 :

 01 Figure 0 -1: Cost breakdown of Li-ion cell ( 4 )

  -1). NMC [Lithium Nickel Manganese Cobalt Oxide (LiNixMnyCozO2)], NCA [Lithium Nickel Cobalt Aluminium Oxide (LiNixCoyAlzO2)] are common positive materials for electrical vehicles, while LCO [Lithium Cobalt Oxide (LiCoO2)] is common positive electrode for portable applications. Such materials contain considerable amount of cobalt and the price of cobalt is now an important concern because of its scarcity, cost and mining conditions in Democratic Republic of Congo.

Figure 0 - 2 :

 02 Figure 0 -2: Specific energy of materials from material to pack level ( 5 ) As electro-mobility applications have volume, weight and size constraints, a high energy density battery cell is mandatory (Figure 0 -2). The positive material (cathode)

Figure I - 1 :

 1 Figure I -1: Rechargeable Li ion battery: transfer of lithium ions from positive electrode (LixMO2) to negative electrode (LixC6) during charge process When the electrochemical cell is charging, oxidation process at positive electrode (i.e., LiMO2) can be represented by

2 . 3 .

 23 Development of new high-energy materials with cost effective and scalable production methods Improvement on battery cycle (shelf) life 4. Understanding of battery safety and failure mechanisms 5. Discovery of easy recycling methods

  Rocksalts as Positive Materials I.2.a. The first cation disordered rocksalts: stoichiometric compositions In 1998, Obrovac et al. discovered cation disordered rocksalts during the preparation of layer oxides compositions, LiMO2 (where M= Ni, Co, Mn and Fe) with high energy ball milling ( 43 ). Milling of LiCoO2 longer than 16 hours, results in a α-LiFeO2 type (cation-disordered rocksalt) structure rather than a α-NaFeO2 type (cation-ordered rocksalt) structure (Figure I -2). The same trend is observed for all other compositions (LiNiO2, LiMnO2, LiFeO2 and LiTiO2); longer milling always results in formation of disordered rocksalt phases. They also observe that disordered structures leads to smaller discharge capacities (< 70 mAh.g -1 ) and large irreversibility at the first cycle (Figure I -2). The reason of such low capacities is attributed to poor lithium ion conduction in disordered lattice.

Figure I - 2 :

 2 Figure I -2: X-Ray diffractogram of LiCoO2 and electrochemical performance of LiCoO2 that is prepared by ball milling ( 43 ) at different milling durations

  instance, among Li2MTiO4 compositions, Li2NiTiO4 delivers a charge capacity of 181 mAh.g -1 , upon cycling between 2.5 V-4.8 V (FigureI -3). However, 50% of charge capacity is lost at the end of the discharge and only 25% of discharge capacity is obtained at the end of the 3 rd cycle ( 50 ). That is explained by irreversible structural modifications upon cycling.

Figure I - 3 :

 3 Figure I -3: Electrochemical performance of Li2NiTiO4 electrodes that are cycled at a rate of C/50 between 4.8 V and 2.5 V vs Li + /Li ( 50 ) at the first three cycles Besides, Li2CoTiO4 shows high reversible capacity of 219 mAh.g -1 at relatively low cycling rate (0.05C), but the large polarization at the first cycle could not be reduced despite the carbon coating strategies ( 51 ).

  statistically occupy 4a positions (½ ½ ½) of octahedral interstitials, whereas anions occupy 4b positions (0 0 0) as shown in the Figure I -4.

Figure I - 4 :

 4 Figure I -4: Cation disordered rocksalt structure: face centered cubic structure that contains lithium (Li), transition metal (TM) and oxygen (O) ions

Figure I - 5 :

 5 Figure I -5: Cation arrangements around tetrahedral sites: 0-TM, 1-TM and 2-TM channels Figure I -5 shows different cation arrangements around tetrahedral sites, for instance 0-TM represents the tetrahedral sites that are coordinated by only lithium ions without any transition metals (TM), whereas 1-TM and 2-TM represent the tetrahedral sites that are coordinated by one and two transition metals respectively.

Figure I - 6

 6 Figure I -6: 0-TM percolation probability (a) and accessible Li content based on Li/Me ratio (b) in Li-rich disordered rocksalt Li1.211Mo0.467Cr0.3O2 ( 56 ) In another work, they describe the Li/Me threshold as 1.10 to form a percolating network along 0-TM channels ( 56 ). As the material becomes richer in lithium, the probability for Li + ions finding an available space in the lattice increases due to the percolating network formed through the lithium diffusion channels (Figure I -6). They also indicate that a new Li-rich rocksalt Li1.211Mo0.467Cr0.3O2 can deliver a reversible capacity of 300 mAh.g -1 associated to its Li-rich content.

  Cation Disordered Rocksalt Oxides I.3.a. Li3NbO4 based cation disordered rocksalt oxides After theoretical calculations reveal the critical Li/Me ratio as 1.10, recent research interests focus on new Li-rich cation disordered rocksalts that are prepared from electrochemical inactive ordered rocksalts such as Li2TiO3, Li3NbO4, Li4MoO5 with transition metal substitution. These materials are not electrochemically active as cathode materials because they do not have redox active centers, all transition metals are stabilized at their highest valence state, thus they cannot be oxidized to upper valence states. For instance, in the case of cation ordered Li3NbO4, which is white powder and electrochemically inactive, almost no capacity is observed upon cycling (Figure I -7).

Figure I - 7 :

 7 Figure I -7: Electrochemical performance of cation ordered Li3NbO4 ( 57 ) Yabuuchi et al. report different overlithiated rocksalts, Li1.3NbxMO2 (where, M: Ni, Fe, Mn, Co; x=0.3 or 0.43) prepared in disordered rocksalt structure by transition metal substitution ( 57 ). Such materials are synthesized in micronic size and they deliver smaller capacities (~140 mAh.g -1 ) based on single electron redox reactions (Me 2+ /Me 3+

25 Figure I - 8 :

 258 Figure I -8: Electrochemical performances of cation disordered Li1.3Nb0.3Mn0.4O2 cells cycled at a rate of 10 mA.g -1 between 4.8 V and 1.5 V vs Li + /Li, at 60 °C ( 57 ) with SEM image of Li1.3Nb0.3Mn0.4O2 powder Despite its high energy density (950 Wh.kg -1 at positive active material level), poor cyclability is observed at room temperature. When the cells are cycled at 60 °C, inital charge capacity reaches 380 mAh.g -1 (Figure I -8). Origin of such a high capacity in Li1.3Nb0.3Mn0.4O2 electrodes is investigated by X-ray absorption spectroscopy (XAS) ( 57 ).Results indicate that both transition metal and oxide ions play an active role during cycling, therefore large reversible capacities result from both cationic and anionic redox processes.

Figure I - 9 :

 9 Figure I -9: Electrochemical performances of cation disordered Li1.3Nb0.3V0.4O2 cells cycled at a rate of 10 mA.g -1 between 4.8 V and 1.5 V vs Li + /Li ( 59 )

Fe

  and Mn is suggested. Thereafter, Fe-substituted Li2TiO3 (Li1.2Ti0.4Fe0.4O2) and Mnsubstituted Li2TiO3 (Li1.2Ti0.4Mn0.4O2) cation disordered rocksalt compositions are prepared, as illustrated in Figure I -10 ( 41,60 ).

Figure I - 10 :

 10 Figure I -10: Electrochemical performances of Li1.2Ti0.4Fe0.4O2 and Li1.2Ti0.4Mn0.4O2 cells that are cycled at a rate of 5mA.g -1 between 4.8 V-1.5 V vs Li + /Li at 50 °C ( 41 )

Figure I - 11 :

 11 Figure I -11: Electrochemical performances and capacity retention of Li6/5Ti2/5Mo2/5O2, Li9/7Nb2/7Mo3/7O2 and Li4/3Mo6 + 2/9Mo 3+ 4/9O2 cells cycled at a rate of 10mA.g -1 between 4.8 V-1.0 V at 55 °C ( 61 )

I. 4 .

 4 Li-rich Cation Disordered Rocksalt Sulfides I.4.a. Li2TiS3 based cation disordered rocksalt sulfides In the literature, cation disordered rocksalt sulfides are also studied. For instance, A. Sakuda et al. mechanochemically synthesize Li2TiS3 powders at different milling durations (Figure I -12). Cation disordered rocksalt phase is obtained after 40 h milling without any remaining precursors ( 62 ).

Figure I - 12 :Figure I - 13 :

 1213 Figure I -12: (a) X-Ray diffractograms of Li2TiS3 powders milled for different durations, b) comparison of experimental and simulated diffraction pattern of Li2TiS3 ( 62 )

Figure I - 14 :

 14 Figure I -14: Capacity retention of Li2TiS3 electrodes (a) in carbonate-based electrolyte at a rate of C/10 between 3.0 V and 1.5 V vs Li + /Li and (b) in solid state electrolyte at 50°C ( 62 ) at different rates To discover structural evolution upon cycling, ex situ X-ray diffraction analysis of the electrodes cycled at different states of charge is performed. Peak intensities of Li2TiS3 reduced and structure became amorphous at the end of the charge, but it returned to cubic rocksalt structure at the end of discharge ( 64 ). Similar observation is reported for cation disordered Li8FeS5 electrodes as well ( 65 ). Cation disordered structure allows reversible structural change from disordered rocksalt to amorphous phase.

Figure I - 16 :

 16 Figure I -16: Cationic redox (Metal (M n+ d) ligand (L n-p) band structure) and anionic redox (Metal (M (n-1)+ d) ligand (L (n+1)-p) band structure)

Figure I - 17 :

 17 Figure I -17: XAS spectra of O K-edge and Mn L-edge of Li1.2Ti0.4Mn0.4O2 electrodes during charge-discharge process ( 60 ) Same study is conducted to examine redox activities in Li1.2Nb0.3Fe0.4O2 electrodes (Figure I -18). At the beginning of charge, oxygen species are partially reduced,

Figure I - 18 :

 18 Figure I -18: Li1.2Nb0.3Fe0.4O2 electrodes at the end of the charge and discharge: Fe L-edge, O K-edge spectra ( 41 ) and TEM images of Li1.2Nb0.3Fe0.4O2 before and after cycling at 50 o C

Figure I - 19 :

 19 Figure I -19: Coin cell test (a), ex situ X-ray diffraction analysis (b), pair distribution function analysis of Li2TiS3 electrodes (c) and Ti-S and S-S bonds obtained from ab initio molecular dynamics calculations (d) ( 64 )

Figure II - 1 :

 1 Figure II -1: Molten salt synthesis of Li2.2NiTi0.2Nb0.6O4 powder

  rpm for 15h, 20h, 25h and 40 h (Figure II -2). All the samples are prepared in a glovebox under inert argon atmosphere to avoid air contamination. Appropriate amount of precursors is weighed and placed into the jar that is then transferred from glovebox to planetary miller (PM 100) operating under room temperature (Figure II -2). After the synthesis, powder is recovered in the glovebox.

Figure II - 2 :

 2 Figure II -2: Mechanochemical synthesis of Li2TiSexS3-x compositionsRotating disc of the planetary miller induces planetary-earth like movement of the jar (FigureII -3). In such a configuration, the jar rotates at its own axes with an inverse direction of the one of the rotating disc. This creates a centrifugal force that allows the

  peak positions (d-spacing) and peak intensity. Peak positions allow determining the symmetry and cell parameters, whereas peak integrated intensity provides information about the atomic positions and the occupancies. X-ray diffractogram provides various information about powders, such as peak positions, peak profile, peak height and peak width, as illustrated in Figure II -6.

Figure II - 6 :

 6 Figure II -6: X-ray diffractogram that provides various information about the crystalline materialsPeak intensities (Ihkl) are proportional to structural factor (Fhkl), it is expressed below

Figure II - 7 :

 7 Figure II -7: Scheme of in situ XRD analysis with GITT mode of cyclingThe cell is charged at a rate of C/50 for 30 min, then the current is stopped and it follows a relaxation for 210 min without current pulses. During cell relaxation, XRD analysis is performed (FigureII -7). Thereafter, same procedure is applied for discharging the cell. The loop is arranged until the cell fully charge-discharged.

Figure II - 8 : 1 M

 81 Figure II -8: In situ cell assembly The in situ cell contains 30-40 mg of active material (70wt% Li2.2NiTi0.2Nb0.6O4 and 30% Carbon black) as positive electrode, glass fiber (Whatman ® ) disk as separator, LP100 (1 M LiPF6 in EC: PC: DMC= 1:1:3) as non-aqueous electrolyte, Li foil as negative

Figure II - 9 :•

 9 FigureII -9: Operando neutron diffraction cell assembly[START_REF] Bianchini | A New Null Matrix Electrochemical Cell for Rietveld Refinements of In-Situ or Operando Neutron Powder Diffraction Data[END_REF] and experiment in ILLThe cell used for neutron diffraction is designed by M. Bianchini[START_REF] Bianchini | A New Null Matrix Electrochemical Cell for Rietveld Refinements of In-Situ or Operando Neutron Powder Diffraction Data[END_REF] (FigureII -9). The bottom part of the cell is made of Ti-Zr alloy. The neutron scattering lengths of titanium (bTi= -3.37x10-12 ) and zirconium (bZr= 7.16x10-12 ) are opposite in sign and almost related to 2:1. Such alloy is prepared as Ti2.08Zr that becomes transparent to neutrons, hence it does not contribute to the background of diffraction pattern.

(

  photons with an energy of hυ and electrons are ejected from the core levels with a kinetic energy of Ek. The kinetic energies of the ejected electrons are measured, binding energies are calculated by the spectrometer (FigureII -11).

Figure II - 11 :

 11 Figure II -11: Principle of X-ray photoelectron spectroscopy

  10 formulation: 80% active mass of material (Li2.2NiTi0.2Nb0.6O4) is mixed with 10% Carbon SP (5321) in a mortar. To have a bulk electrode and to adjust the viscosity of the slurry, 10% of PVDF (Polyvinylidene fluoride) binder is added with few drops of NMP (N-methyl-2-pyrrolidone) solvent.Then, slurry is coated on an Al current collector (20 μm) by using 100 μm doctor blade and left to dry in the oven at 55 °C for 12 hours. Later, electrodes are cut (14 mm disk diameter), pressed (10 tons), weighed (loading > 4 mg) and dried in the Buchi ® oven at 80 °C for 48 hours under vacuum.• Preparation of sulfide-selenide based electrodesUnlike oxide-based electrodes, sulfide and selenide-based electrodes are prepared in a glovebox under inert argon atmosphere to avoid air contamination. Electrodes are prepared based on the same formulation (80:10:10) as described previously for preparation of oxide electrodes. Slurry is coated manually on the Al foil by using 100 µm doctor blade and left to dry in the glovebox for three days. Later, electrodes are cut (14 mm), pressed (10 tons), weighed (loaded > 2 mg) in the glovebox and dried in the Buchi ® oven at 60 °C for 12h under vacuum.• Coin cell preparationThe electrochemical characterizations of synthesized powders are performed in coin cell, which is shown in Figure II -13. The cells are prepared by using 14 mm diameter positive electrodes, 16 mm Li foil as a negative electrode and 150 μl electrolyte of LP100 containing 1 M LiPF6 in EC:PC:DMC, where EC: ethylene carbonate; PC: propylene carbonate; DMC: dimethyl carbonate, in volumetric ratio of 1:1:3 respectively. Two types of separators (Φ = 16.5 mm) are used; polypropylene separator (Celgard ® 2400) and polyolefin separator (Viledon ® ).

Figure II - 13 :

 13 Figure II -13: Coin cell assembly

Figure II - 14 :

 14 Figure II -14: Galvanostatic cycling • Galvanostatic cycling schedule Disordered Li2TiS3 is cycled at a rate of C/10 in a potential range from 3 V to 1.5 V vs Li + /Li, whereas disordered Li2.2NiTi0.2Nb0.6O4 is cycled at a rate of C/50 in a potential range of 4.8 V and 2 V. ARBIN cycling instrumentation that operates at 22 °C is used

Figure II - 16 :

 16 Figure II -16: Cyclic Voltammetry test, parameters obtained in a voltammogramWhen an electrochemical cell is scanned with a scan rate (µV.s -1 ), positive current (iox) is measured for an oxidation reaction, whereas a negative current (ired) is measured for reduction reaction in current-potential curve (cyclic voltammogram). Besides, oxidation (Eox) and reduction (Ered) redox potentials can be identified in cyclic voltammogram (FigureII -16). The difference between these two potentials (∆E) gives indications on the reversibility of the redox reactions.

Figure III - 1 :

 1 Figure III -1: Powder XRD patterns of disordered rocksalt Li2.2NiTi0.2Nb0.6O4 powder: before (1) and after (2) washing with water Figure III -1 shows XRD patterns of Li2.2NiTi0.2Nb0.6O4 powder before and after washing. In unwashed powder, diffraction peaks of disordered rocksalt Li2.2NiTi0.2Nb0.6O4, NaCl and KCl are detected. By washing the powder with water, remaining salt is totally removed and only rocksalt cubic phase is observed. Besides, sharp diffraction peaks indicate good crystallinity of Li2.2NiTi0.2Nb0.6O4 powder.

Figure III - 2

 2 Figure III -2 shows XRD pattern of Li2.2NiTi0.2Nb0.6O4 powder in Fullprof program profile matching. Li2.2NiTi0.2Nb0.6O4 powder is indexed in a rocksalt (cubic) phase (space group, Fm3 ̅ m) with lattice parameter, a= 4.1842 Å.

Figure III - 2 :

 2 Figure III -2: XRD pattern of disordered rocksalt Li2.2NiTi0.2Nb0.6O4 powder in Fullprof program profile matching To investigate the electrochemical properties of rocksalt Li2.2NiTi0.2Nb0.6O4, electrodes are prepared with the procedure explained in Chapter II.3. Three coin cells are cycled at a rate of C/50 between 4.8 V and 2.0 V vs Li + /Li at room temperature (22 °C).

Figure III - 3 :

 3 Figure III -3: Cell performances of disordered rocksalt Li2.2NiTi0.2Nb0.6O4 electrodes cycled at a rate of C/50 between 4.8 V-2.0 V vs Li + /Li at 22 °C

Figure III - 5 :

 5 Figure III -5: Summary of the parameters considered for synthesis improvement • Simple tempering (long firing): Rocksalt Li2.2NiTi0.2Nb0.6O4 is prepared using four different metal oxides reactants. To prepare homogeneous and wellcrystalline ceramic powder, simple tempering, in another words long firing, can be applied. By letting the ingredients well diffuse during the synthesis, pure rocksalt phase can be obtained.

Figure III - 6 :

 6 Figure III -6: XRD patterns of Li2.2NiTi0.2Nb0.6O4 powders with different strategies: classical calcination (3h), calcination (3h) with quench and long calcination (6 h) with quench After calcination, the three samples are washed with water and X-ray diffraction patterns are compared in Figure III -6. All samples present the targeted phase with small amount of impurities. NaNbO3 is present in every sample and is not affected by the different thermal treatments. Small amount of Li2CO3 is only found in the slowly

Figure III - 7 :

 7 Figure III -7: Cell performances of Li2.2NiTi0.2Nb0.6O4 electrodes cycled at a rate of C/50 between 4.8 V-2.0 V vs Li + /Li at 22 °C Figure III -7 shows cell performances of the samples prepared by different strategies at the first cycle. The "classical" samples show the worst behavior with a very high polarization that may rise from the formation of Li2CO3 during the cooling. On the contrary, quenched samples shows lower polarization. Between the two quenchedsamples, the samples that was submitted to the shorter thermal treatment exhibits much higher capacities with a charge capacity of 144 mAh.g -1 and a discharge capacity of 105 mAh.g -1 upon cycling at a rate of C/50 between 4.8 V-2.0 V vs Li + /Li. Charge capacity is almost 15% higher than the charge capacity of classical material, whereas discharge capacity is 30% higher than discharge capacity of classical material.Moreover, we observe irreversibility at the first cycle for quenched sample is much smaller (27%) (FigureIII -7). The lower performances of the sample quenched after 6h may be attributed to the presence of ordered phase that is known to present much poorer ionic conductivity.

Figure III - 8 :

 8 Figure III -8: Comparison of cell performances of Li2.2NiTi0.2Nb0.6O4 cells (powder fired 3h and quenched) at 22 °C and 55°C, electrodes cycled at a rate of C/50 between 4.8 V -2.0 V vs Li + /Li Figure III -9 shows cycling stability at room (22 °C) and high (55 °C) temperatures.

Figure III - 9 :

 9 Figure III -9: Evolution of discharge capacities of Li2.2NiTi0.2Nb0.6O4 (quenched) at 22 °C, electrodes cycled at a rate of C/50 between 4.8 V-2.0 V vs Li + /Li III.1.b. Effect of washing (type of solvent)

Figure III - 10 :

 10 Figure III -10: XRD patterns of Li2.2NiTi0.2Nb0.6O4 powders prepared with the same recipe separately and washed with different solvents: water and ethylene glycol. Kα2 contributions have been stripped with EVA software for clarity reason. Arrows indicates possible contributions from H + exchanged phase.

Figure III - 10

 10 Figure III -10 shows the XRD patterns of two different batches after washing. No remaining salt is detected in diffraction patterns of washed powders, both ethylene glycol and water effectively remove eutectic salt in powders. Reflections from the EGwashed samples are broader and may cover small impurities, but they present a symmetrical profile.

Figure III - 11 :

 11 Figure III -11: Half-cell performances of Li2.2NiTi0.2Nb0.6O4 washed with different solvents: water (1) vs ethylene glycol (2), electrodes cycled at a rate of C/50 between 4.8 V and 2.0 V vs Li + /Li at 22 °C, the first two cycles III.1.c. Effect of molten salt to precursor ratio

Figure III - 12

 12 FigureIII -12 shows the XRD patterns of the samples that are prepared in different salt to precursor ratios. In each case, disordered rocksalt Li2.2NiTi0.2Nb0.6O4 phases are successfully obtained. Such result indicates that even smaller salt to precursor ratio (2:1) is enough to prepare those phases. However, an impurity phase indexed as NaNbO3 is observed in diffraction patterns of salt to powder ratios (4:1) and (10:1). In contrast, we do not observe NaNbO3 impurity in powder synthesized in low salt to

Figure III - 12 :

 12 Figure III -12: XRD patterns of Li2.2NiTi0.2Nb0.6O4 powders prepared by different salt (NaCl:KCl) to precursor ratios: 2:1 (1), 4:1 (2) and 10:1 (3)

Figure III - 14 :

 14 Figure III -14: XRD patterns of Li2.2NiTi0.2Nb0.6O4 powders prepared with different salts: eutectic NaCl-KCl and 0.59LiCl-0.41KCl at 450 °C and 670 °C Figure III -14 shows XRD patterns of three different powders prepared by eutecticNaCl-KCl and 0.59LiCl-0.41KCl salts. When the powder is prepared by 0.59LiCl-0.41KCl salt and heated at 450 °C for 12 hours, no disordered rocksalt phase is formed.

Figure III - 15 :

 15 Figure III -15: Half-cell performance of Li2.2NiTi0.2Nb0.6O4 prepared by different eutectic salts: NaCl-KCl and 0.59LiCl-0.41KCl, electrodes cycled at a rate of C/50 between 4.8 V and 2.0 V vs Li + /Li at 22 °C, at the first cycle

  heated at 300 °C to remove the acetate and decomposed at 670 °C. To investigate the effect of temperature to the synthesis, two samples are prepared and decomposed at 570 °C and 670 °C separately. To discover the effect of firing duration, three samples are heated at 670 °C for different durations: standard 3 hours (standard), 1 hour and 0.5 h (30 mins).

Figure III - 17

 17 Figure III -17 shows samples calcined at different temperatures and durations. In XRD pattern of the sample calcined at 570 °C, dissymmetrical diffraction peaks are observed, indicating that new sample contains another phase. The reason might be slow diffusion of the reactants at 570 °C. For three powders calcined at 670 °C at different durations, disordered rocksalt phases are successfully obtained. Such result indicates that diffusion of reactants is high enough at 670 °C, disordered rocksalt phase can be achieved even at short firing durations (< 3 hours).

Figure III - 17 :

 17 Figure III -17: XRD patterns of Li2.2NiTi0.2Nb0.6O4 powders synthesized different synthesis temperatures and durations: powder synthesized at 670 °C for 3 hours and at 670 °C for 1 hour, 570 °C 3 hours and 670 °C 30 mins (0.5 hour)

Figure III - 19 :

 19 Figure III -19: Cycling life of Li2.2NiTi0.2Nb0.6O4 synthesized at 670 °C for different durations: 3 h, 1 h and 0.5 h In Figure III -19, electrodes show similar cycling life, except sample synthesized for 0.5 h. Here, we conclude that cell performance of Li2.2NiTi0.2Nb0.6O4 is improved by synthesizing the powders at shorter firing duration (1 h).

  Li2.2NiTi0.2Nb0.6O4 shows high irreversibility at the first cycle, but a good reversibility for the following cycles. The reason of the irreversibility at the first cycle is unknown. As previously shown, lithium diffusion in Li2.2NiTi0.2Nb0.6O4 is better at high temperatures (55 °C) compared to room temperatures (22 °C), we draw the hypothesis that Li2.2NiTi0.2Nb0.6O4 might have lithium diffusion limitations. At the end of the charge, Li ions are extracted from the shell of the particles and material's core remains richer in lithium (Figure III -20). During discharge, lithium ions insert into the structure and fill empty spaces in the material surface. As the Li diffusion is limited, Li accumulates faster at the surface than in the bulk and may create a diffusion barrier.

Figure III - 20 :

 20 Figure III -20: Hypothesis on lithium diffusion mechanism through Li2.2NiTi0.2Nb0.6O4To test this hypothesis, the cell is cycled at variable charge potentials beginning from a low potential of 3.9 V up to 4.8 V with a fixed discharge potential (2 V), as illustrated in FigureIII -21.

Figure III - 21 :

 21 Figure III -21: Cell performance of Li2.2NiTi0.2Nb0.6O4 that was cycled at variable charge potentials beginning from a low potential of 3.9 V up to 4.8 V with a fixed discharge potential (2 V) at a rate of C/50 vs Li + /Li at 22 °C Interestingly, we observe that the irreversibility arises at low state of charge (SOC) and

Figure III - 22 :

 22 Figure III -22: Ex situ XRD analysis of electrodes at different states of charge: pristine electrode, electrodes at end of charge (4.8 V) and discharge (2 V) Figure III -22 shows ex situ XRD analysis of electrodes at different states of charge. After aligning aluminum peaks, we observe that both peak intensities and lattice parameters reduce at the end of charge (4.147 Å). Besides, shape of the diffraction peak at 63.5° becomes much broader, indicating that charged electrode might contain another phase. On the contrary, lattice parameters of discharged electrode are slightly bigger than the pristine electrode (4.187 Å). Additionally, the diffraction pattern of discharged electrode is quite different from the pristine electrode. This result indicates that structural change occurring during cycling is not completely reversible. It seems that final composition might not have the same structure as the original composition.

Figure III - 23

 23 Figure III -23 shows cycling curve of in situ cell containing Li2.2NiTi0.2Nb0.6O4 powder that is cycled with GITT mode of cycling (for cell preparation, see Chapter II.2). The cell is cycled for 0.5 hour, and then current is interrupted. Afterwards, the cell is rested at 3.5 hours and X-ray diffraction is conducted during relaxation.

Figure III - 23 :

 23 Figure III -23: Cycling curve of in situ cell containing Li2.2NiTi0.2Nb0.6O4 powder, GITT mode of cycling Figure III -24 shows global view in situ XRD diffraction pattern of Li2.2NiTi0.2Nb0.6O4 powder. In diffraction patterns, aluminum and beryllium peaks are identified as well as the disordered rocksalt phase. For clarity sake, some diffraction patterns (pristine, end of charge and discharge) are selected and shown in Figure III -25.

Figure III - 24 :

 24 Figure III -24: Global view of in situ XRD diffraction pattern of Li2.2NiTi0.2Nb0.6O4 powder Focusing on the high 2θ region (62°-64°) in Figure III -25, we have similar observations previously described in ex situ X-ray diffraction section. The peak intensities reduce and the diffraction pattern of charged powder seems to contain another disordered rocksalt besides disordered rocksalt Li2.2NiTi0.2Nb0.6O4. To be sure about the second phase, we decide to refine lattice parameters using Fullprof program profile matching.

Figure III - 25 :

 25 Figure III -25: In situ X-ray diffraction pattern of Li2.2NiTi0.2Nb0.6O4 powder at different states of charge: pristine, end of charge and discharge At first, we start to do fitting the pattern according to our previous assumption that second disordered rocksalt phase is progressively forming during charging. However, we detected a discrepancy in the fitting results, thus we were not able to find a good fit after few steps of lithiation. Afterwards, we assume that the second phase should have been at the beginning, in other words initial powder (pristine state) might have already contained two phases before cycling. The refinement of the initial powder is then carried with both hypotheses, showing a real improvement when introducing a second phase (Figure III -26).

Figure III - 26 :

 26 Figure III -26: Results of the refinements of the initial powder. Top: with 1 phase, bottom: with 2 phases We find that the two phases have very close lattice parameters (Phase 1: 4.1824 Å, Phase 2: 4.1848 Å), explaining why this phase has not been detected before. During charge, the lattice parameters are decreasing in a different way and we observe different slopes (Figure III -27a).

Figure III - 27 :

 27 Figure III -27: Change of (a) lattice parameters and (b) size of the crystalline domains of Li2.2NiTi0.2Nb0.6O4 powder during in situ XRD analysis Lattice parameters increase during discharge, but they are slightly bigger than those in the pristine states (Figure III -27a). This result indicates that the final phases are different from the pristine states. The irreversible structural change may thus be one of the causes of the irreversible capacity observed during cycling.

Figure IV - 1 ,

 1 XRD patterns of the four different Li2TiS3 powders are quite similar, almost the same peak positions and intensities are observed. SEM images of Li2TiS3 powder (milled for 20h) are also displayed in Figure IV -1. As seen in the images, the powder contains agglomerates that consist of nanometer sized spherical particles.

Figure IV - 2 :

 2 Figure IV -2: X-ray diffraction pattern of disordered rocksalt Li2TiS3 powder (milled for 20 h) profile matching in Fullprof program

Figure IV - 3 :

 3 Figure IV -3: Half-cell performances of Li2TiS3 cells (powders synthesized at 510 rpm for different durations) at the first cycle, coin cells cycled at a rate of C/10 between 3.0 V -1.5 V Li + /Li at 22 °C Figure IV -3 shows the electrochemical performances of the different Li2TiS3 electrodes at the first cycle and the results are summarized inTable IV -2. For each material, discharge capacities are found greater than charge capacities, thus there is always an extra capacity in Li2TiS3 samples. Discharge average potentials are quite similar.

Li 2

 2 TiS 3the error bar. This may arise from the analysis. ICP samples are dissolved in acidic aqueous solution, so a part of sulfur may have evolve as H2S and leave the sample before the analysis.In summary, four different Li2TiS3 powders are mechanochemically synthesized as disordered rocksalt phase without any remaining precursors. With our optimized recipe, we can achieve desired phases at shorter milling durations (< 40 hours) than those in the literature (62). This can be explained by different size of balls, as slightly bigger balls are used for the mechanochemical synthesis. As bigger balls create stronger shocks, the impact energy on the particles increases. This allows reaching desired phases at shorter milling durations. When the electrochemical performances of powders are investigated, no direct relation between milling durations and zirconium contamination as well as the performances is observed. Among all milling durations, we decide to further synthesize our powders milling for 20 hours. IV.1.b. Electrochemical Performance of Li2TiS3 Now, we investigate the electrochemical performance of Li2TiS3 electrodes in detail (Figure IV -4). At the end of the first charge, 1.50 Li + ions are extracted from Li2TiS3 and the capacity reaches 250 mAh.g -1 .

Figure IV - 4 :

 4 Figure IV -4: Half-cell performances and cycling life of Li2TiS3 (unpressed) electrodes cycled at a rate of C/10 between 3 V -1.5 V vs Li + /Li: the first (black color), second (red color) and third (blue color) cycles

Figure IV - 5 :

 5 Figure IV -5: Effect of pressing to the cycling life: Half-cell performances of Li2TiS3 electrodes pressed electrodes and cycling life of pressed and unpressed Li2TiS3 electrodes cycled at a rate of C/10 between 3 V -1.5 V vs Li + /Li: the first (black color), second (red color) and third (blue color) cyclesFigure IV -5 shows electrochemical performances of pressed Li2TiS3 electrodes under 10 tons in a cold-press machine in the glovebox. Here, pressed electrodes deliver much

  current is then reduced to 5D, and this protocol is then continued with 2D, D, D/2, D/3, D/5 and finally D/10 consecutive rates to reach a full discharge of the material. Between two discharge phases, electrodes are rested for two hours. In FigureIV -6, we observe that Li2TiS3 electrodes retain 74% of discharge capacity at D rate, 64% capacity at 5D rate and 51% capacity at 10D rate.

Figure IV - 7 :

 7 Figure IV -7: Capacity and average potential of Li2TiS3 electrodes cycled at a rate of C/10 between 3 V -1.5 V in two different electrolytes: carbonate-based electrolyte (1 M LiPF6 in EC:PC:DMC 1:1:3 in volume) and tetraglyme-dioxolane based electrolyte (1M LiTFSI + 0.1M LiNO3 in TEGDME:DIOX 1:1 in volume) Figure IV -7 shows Li2TiS3 electrodes cycled in two different electrolytes, such as carbonate-based electrolyte (1 M LiPF6 in EC:PC:DMC 1:1:3 in volume) and tetraglyme-dioxolane based electrolyte (1M LiTFSI + 0.1M LiNO3 in TEGDME:DIOX 1:1 in volume). Here we observe that cycle life of Li2TiS3 electrodes in carbonate-based electrolyte is better than the one in tetraglyme-dioxolane-based electrolyte. Besides, same charge average (2.46 V) and discharge average (2.23 V) potentials are observed in the two systems. Thus, we decided to keep testing the electrodes with carbonatebased electrolyte.

Figure IV - 8

 8 Figure IV -8 shows XRD pattern of pristine Li2TiS3 electrodes exposed under differentdurations and atmospheres. When the electrodes remain for 1 day in the dry room, peak intensity at 2θ = 36° decreases, but the crystallinity is conserved. The same is observed at the electrodes remained for 1 week in the dry room. In contrast, almost no peaks are detected in electrodes remained under fume hood. We detect various TixOy compounds in electrodes that remained for 1 week under fume hood, indicating that sulfur and lithium atoms react with air and they may produce H2S gas and amorphous Li2CO3 compounds that are not observed in the X-ray diffraction patterns.

Figure IV - 9

 9 Figure IV -9 and Figure IV -10 show coin cell test results of Li2TiS3 electrodes exposed under different durations and atmospheres. Electrodes remaining in dry room deliver almost similar charge and discharge capacities for the first 8 cycles. Thereafter, small capacity fading is observed in electrodes remained in the dry room. In contrary,

  ml) and put into a centrifuge (Figure IV -11b and Figure IV -11c). They are stirred under 14000 rpm for 20 min, then the powders settle in the bottom, while excess cyclohexane solution is removed from the flask. To remove the organic impurities, the powder is washed with anhydrous cyclohexane three times (Figure IV -11d). Thereafter, it is left to dry for two days in the glovebox. As a comparison, we also prepare Li3TiS3 samples with mechanochemical synthesis. Based on reaction in Eq. IV -4, 1.5 mol of Li2S, 1 mol of Ti and 1.5 mol of S are weighed and placed into the jar. The jar is transferred into planetary ball milling machine and milling conditions are set to 510 rpm for 20 hours.

Figure IV - 11 :

 11 Figure IV -11: Butyl lithium experimental set-up: (a) stirring for a week, (b) taking the mixture into the small flask, (c) centrifuge of the samples, (d) washing the powder with cyclohexane Figure IV -12 shows the XRD patterns of Li2+xTiS3 powders prepared by butyl lithium experiment and ball milling. We observe Li2S precursor in X-ray diffraction patterns of powder prepared by mechanochemistry. It seems that powder prepared by butyllithium also contain small amount of Li2S, which is a side product. The presence of Li2S impurity indicates that it is more difficult to stabilize Ti 3+ in the disordered rocksalt structure.

Figure IV - 13 :

 13 Figure IV -13: Half-cell performances of LixTiS3 that is lithiated by butyl lithium experiment, electrodes cycled at a rate of C/10 between 3 V -1.5 V vs Li + /Li; the first cycle (black color), second cycle (red color) and third cycle (blue color) In Figure IV -13, we observe that Li2+xTiS3 cells show almost fully reversible chargedischarge curve though capacities are rather smaller (~110 mAh.g -1 ) than expected.

Figure IV - 14

 14 Figure IV -14 shows the XRD pattern of Li2TiS3 prepared by wet and dry mechanochemical syntheses. In the two different synthesis processes, Li2TiS3 powders are successfully achieved in disordered rocksalt phase. Both powders have similar cell parameters (awet: 5.0853 Å and adry: 5.0855 Å). Moreover, coin cell test results in Figure IV -15 show similar charge and discharge capacities of Li2TiS3 prepared by wet and dry mechanochemical syntheses. Those results confirm that disordered rocksalt Li2TiS3 can be prepared by a new synthesis process that is more flexible and less expensive.

Figure IV - 16 :

 16 Figure IV -16: XRD patterns of Li2TiSexS3-x powders prepared with wet mechanochemical synthesis

Figure IV - 17 :

 17 Figure IV -17: XRD pattern of Li2TiSeS2 profile matching in Fullprof program

Figure IV - 19 :

 19 Figure IV -19: Simulated XRD patterns of Li2TiSexS3-x powders

Figure IV - 21 :

 21 Figure IV -21: SEM images of Li2TiSeS2 (left) and Li2TiSe2S (right) powders

Figure IV - 22 :

 22 Figure IV -22: Half-cell performances and cycle life of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes cycled at a rate of C/10 between 3 V -1.5 V vs Li + /Li at 22 o C

Figure IV - 23 : 5 𝑛𝑚𝑉

 235 Figure IV -23: Redox potentials comparison of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes at the first (5µV.s -1 ) and the second (10 µV.s -1 ) cycles between 3 V -1.5 V vs Li + /Li

  which is expressed as i=av b ( 102,103 ). Here a and b parameters are variables and b parameter is expressed as a sum of faradaic (diffusion controlled) and non-faradaic (capacitive/surface controlled) currents ( 103 ). By plotting logi vs logv, b parameter can be obtained from the corresponding slope. According to b parameter, the reaction kinetics are either surface controlled (b=1) or diffusion-controlled process (b=0.5) ( 103,105 ).

Figure IV - 25

 25 Figure IV -25: a) Logarithmic and b) linear relation between peak current and scan rate of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes In Figure IV -25a, only Li2TiS3 shows a linear relation (diffusion controlled process)

  8). By ignoring the fact that the scan rates are quite slow (µV.s -1 ), Li2TiSeS2 and Li2TiSe2S electrodes retain the capacities in a better way than the Li2TiS3 cells at high rates. Although such a result is quite different from what is observed in galvanostatic tests, it also supports our previous observation in which electron transfer kinetics in Li2TiSeS2 and Li2TiSe2S compositions are much more sensitive to scan rate.

Figure IV - 26

 26 Figure IV -26 shows capacity retention of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes cycled in galvanostatic mode at different discharge cut-off potentials. Cycling the cells between 3 V -1.5 V vs Li + /Li, Li2TiSeS2 cells show better cycling stability than Li2TiS3 cells, whereas Li2TiSe2S cells show rapid capacity fading. At higher cut-off discharge potentials (3 V -1.8 V) and (3 V -2 V), all selenium substituted cells (Li2TiSeS2 and Li2TiSe2S) show better cycling stability than Li2TiS3. To discover why Li2TiSe2S cells fail upon cycling between (3 V -1.5 V), we perform ex situ XRD and SEM analyses of electrodes at different states of charge.

Figure IV - 26 :

 26 Figure IV -26: Capacity retention (%) of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes cycled in galvanostatic mode at different cycling discharge cut-off potentials: 3 V -1.5 V, 3 V -1.8 V and 3 V -2.0 V vs Li + /Li.

Figure IV - 27 :

 27 Figure IV -27: SEM images of Li2TiSe2S electrodes cycled at a rate of C/10 between (3 V -1.5 V) and (3 V -1.8 V)

  2.b.Before performing operando diffraction experiment, Li2TiS3 powder (≈ 120 mg) is analyzed with ex situ neutron diffraction in ILL. The target is to collect the diffraction patterns and to determine the quality of data before the real experiment. In ex situ neutron diffraction pattern, we observe a large background with very small diffraction peaks. We conclude that the powder amount is too low to obtain sufficient diffraction peaks. Afterwards, we decide to increase powder amount (200 mg) to get sufficient 123 intense diffraction peaks during the real experiment. However, we also know that working with the cells containing large amount of powders is quite challenging, it often results in large cell polarization with smaller capacities. Therefore, our first goal is to optimize the electrochemical performances of the neutron cell by applying different cell testing schedules (Figure V -1).

Figure V - 1 :

 1 Figure V -1: Cycling of operando neutron cell cycled at a rate of C/10: a) Cell is charged to 3 V and discharged to 1.5 V b) Cell is discharged to 1.5 V and charged to 3 V, c) cell is discharged to 1.5 V and charged to 3.5 V d) cell is charged to 3.5 V and discharged to 1.0 V

  5 V and charged to 3 V (Figure V -1b). Thereafter, the high cut-off potential is increased to 3.5 V and we finally obtain a reversible cycling curve, as shown in Figure V -1c. Nevertheless, we do not observe the extra capacity during discharge. We further decrease the low cut-off potential and the cell is charged first to 3.5 V and discharged to 1.0 V. After such optimizations, we succeed to obtain reversible cycling curve and observe the extra capacity during discharge.

Figure V - 2 :

 2 Figure V -2: Operando neutron diffraction pattern of Li2TiS3 powder: pristine powder, discharged powder and empty neutron diffraction cell (neutron wavelength: 1.594 Å)

Figure V - 2 V. 2 .

 22 Figure V -2 shows experimental and Figure V -3 shows simulated operando neutron diffraction patterns of Li2TiS3 powders. Based on simulated diffraction pattern, we should see a strong peak at 31°. However, in Figure V -2, we observe that the first peak at 31° is lost in the background coming from the cell itself. We could only see a small peak at 62°, but it is difficult to make a conclusion based on this result. Differently from the simulated diffraction pattern, two additional peaks at 47 Å and 56 Å in both pristine and discharge states are seen in Figure V -2. Those peaks should arise from the separator within the cell.

Figure V - 4 :

 4 Figure V -4: Cell performances of Li2TiS3 electrodes at different states of charge: electrode at the end of charge (3 V) and electrode at the end of discharge (1.6 V)

Figure V - 6

 6 Figure V -6 shows SEM images of Li2TiS3 electrodes cycled at different states of charge.In Li2TiS3 electrode at the end of the charge, expansion is obvious and some degree of porosities are observed. However, the morphology of Li2TiS3 electrode at the end of the discharge is quite similar to Li2TiS3 pristine electrode. Overall, we observe reversible structural change and detect similar electrode morphologies after charge and discharge of Li2TiS3 electrodes.

Figure V - 6 :

 6 Figure V -6: SEM images of Li2TiS3 electrodes at different states of charge: pristine electrode, charged (3 V) and discharged (1.6 V) electrode

Figure V - 7 :Figure V - 8 :

 78 Figure V -7: Cell performances of Li2TiSeS2 electrodes at different states of charge: electrode at the end of charge (3 V) and electrode at the end of discharge (1.6 V)

Figure V - 9 :

 9 Figure V -9: Cell performances of Li2TiSe2S electrodes at different states of charge: electrode at the end of charge (3 V) and electrode at the end of discharge (1.6 V)

Figure V - 11 :

 11 Figure V -11: SEM images of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes at different states of charge: pristine, at the end of charge (3 V) and at the end of discharge (1.6 V)

Figure V - 12

 12 Figure V -12: a) SEM images and b) EDX images of Li2TiSe2S electrodes cycled to 3V and kept for 24 h

  Figure V -14, similar elements are detected in electrodes as well, but differently from the powders, F 1s core peak (687.7 eV) is detected in electrodes. This originates from PVdF binder used during electrode preparation.

Figure V - 13 :

 13 Figure V -13: XPS analysis of powders in survey mode

Figure V - 15 :

 15 Figure V -15: XPS spectra of powders: Li2TiS3, Li2TiSeS2, Li2TiSe2S and selenium precursor at Se 3d core Figure V -15 shows the Se 3d core peak of Li2TiS3, Li2TiSeS2, Li2TiSe2S and selenium precursor. The four spectra were calibrated based on C 1s core peak at 284.8 eV. In the figure, Li 1s and Se 3d binding energies are quite close to each other, Li 1s and Se 3d core peaks are detected between 53 eV and 54 eV respectively. The intensity of Se 2-

Figure V - 16

 16 Figure V -16 shows the XPS spectra and peak fittings of Li2TiS3, Li2TiSeS2 and Li2TiSe2S powders at Li 1s, Ti 2p and S 2p core levels. The curves are fitted using 80% Gaussian, 20% Lorenzian conditions in Multitap program.

Figure V - 16 :

 16 Figure V -16: XPS spectra of Li2TiS3, Li2TiSeS2 and Li2TiSe2S powders at Ti 2p3/2-1/2 (Ti1_orange: Ti 2p3/2, Ti2_yellow: Ti 2p1/2), S 2p3/2-1/2 (S1_light blue: S 2p3/2 and S2_dark blue: S 2p1/2), Li 1s core (navy blue), Ti 3s (purple) and Se 3d5/2-3/2 (Se1_green:Se 3d5/2, Se2_pink:Se 3d3/2) • 72-44eV: Li 1s, Ti 3s and Se 3d cores

•

  473eV-449eV: Ti 2p core In Li2TiS3 powder, two different Ti 2p doublets, such as Ti1 and Ti2 are observed at 455.1 eV and 457.4 eV, indicating Ti 4+ present in sulfur (TiS2) and oxygen (TiO2) environments. Presence of TiO2 peak at Ti 2p core indicates that there should be air exposure during sample transfer from transfer box to XPS vacuum chamber. Similar phenomenon was observed in the literature as well ( 119-121 ).

  Figure V -17 (Li2TiS3), Figure V -18 (Li2TiSeS2) and Figure V -19 (Li2TiSe2S).

Figure V - 17 :

 17 Figure V -17: XPS spectra of Li2TiS3 at Ti 2p3/2-1/2, S 2p3/2-1/2 and Li 1s core: Li2TiS3 powder vs non-scratched Li2TiS3 electrodes and scratched Li2TiS3 electrodes In Figure V -17, similar active species are observed in Li2TiS3 powder and electrodes.

Figure V - 18 :

 18 Figure V -18: XPS spectra of Li2TiSeS2 at Ti 2p3/2-1/2, S 2p3/2-1/2 and Li 1s core: powder vs scratched electrodes and non-scratched electrodes

Figure V - 20 .

 20 Figure V -20. S1 and S2 represents S 2-(23.7%) and S2 2-(70.7) respectively, whereas S3 (S-C) (5.6%) corresponds to sulfur-carbon bond.

Figure V - 20 :

 20 Figure V -20: XPS spectra of Li2TiS3 scratched electrodes at different states of charge: pristine, charged electrode (3 V) and discharged electrode (1.6 V)

Figure V - 21

 21 Figure V -21 shows XPS spectra of Li2TiSeS2 scratched electrodes at different states of charge. In pristine Li2TiSeS2 electrode Se 2-(22%), Se2 2-(28%), S 2-(45%) and S-C (5%)

Figure V - 21 :

 21 Figure V -21: XPS spectra of Li2TiSeS2 and Li2TiSe2S scratched electrodes at different states of charge: pristine, charged electrode (3 V) and discharged electrode (1.6 V) In Figure V -22, reversible selenium redox in Li2TiSeS2 is also seen at Se 3d core. Similarly to Figure V -21, shifts in binding energies are very small, however change in the intensities is evident. It seems that Se 2-converts to Se2 2-during charge, but it returns to the original position during discharge. It suggests a reversible redox activity.

Figure V - 22 :

 22 Figure V -22 : XPS spectra of Li2TiSeS2 electrodes at different states of charge: Comparison of Se 3d and S 2p core levels

  [START_REF] Zhang | Understanding of Performance Degradation of LiNi0.80Co0.10Mn0.10O2 Cathode Material Operating at High Potentials[END_REF]

  ).To improve structural stability and specific capacity of NMC, Thackeray et al.[START_REF] Thackeray | Advances in Manganese-Oxide 'Composite' Electrodes for Lithium-Ion Batteries[END_REF] suggest to prepare the solid solution of LiMO2 with electrochemically inactive Li2MnO3 composition. The new lithium rich composition delivers unusual charge capacity exceeding 250 mAh.g -1 between 2 V and 4.8 V vs Li + /Li. ([START_REF] Zhang | Hierarchically Porous Micro-Rod Lithium-Rich Cathode Material Li1.2Ni0.13Mn0.54Co0.13O2 for High Performance Lithium-Ion Batteries[END_REF][START_REF] He | Enhanced Electrochemical Performance in Lithium Ion Batteries of a Hollow Spherical Lithium-Rich Cathode Material Synthesized by a Molten Salt Method[END_REF]

	That increases cell resistance and lowers the cycling stability. Surface coating
	techniques ( 31 ), doping ( 32 ) and core-shell structural design ( 33 ) have been proposed to
	solve these problems.
	• Li-rich (or Mn-rich) NMC [xLi2MnO3. (1-x)LiMO2]

  [START_REF] Yabuuchi | High-Capacity Electrode Materials for Rechargeable Lithium Batteries: Li3NbO4-Based System with Cation-Disordered Rocksalt Structure[END_REF]

  ). The compositions are shown in TableI-1. Among all, Mn doped

	Li3NbO4 (Li1.3Nb0.3Mn0.4O2) shows the lowest polarization with a reversible capacity of
	260 mAh.g -1 at room temperature (Figure I -8). Similar composition
	(Li1.25Nb0.25Mn0.50O2) is studied by another group as well ( 58

Table I -

 I 1: Summary of cation disordered rocksalt oxides reported in the literature

						Initial	Reversible	Potential
	Type	Composition	T (°C)	rate		charge	Capacity	window	Reference
						capacity	(mAh.g -1 )	(V)
		Li 1.30 Nb 0.30 Mn 0.40 O 2				260	260		[57]
		Li 1.30 Nb 0.30 Fe 0.40 O 2				300	250		[57]
		Li 1.30 Nb 0.43 Ni 0.27 O 2				280	280		[57]
		Li 1.30 Nb 0.43 Co 0.27 O 2				310	310		[57]
		Li 1.30 Nb 0.30 V 0.40 O 2				180	180	4.8 -1.5	[59]
		Li 1.30 Nb 0.30 Mn 0.40 O 2	60	10 mA.g -1	380	270	4.8 -1.5	[57]
		Li 1.25 Nb 0.25 Mn 0.50 O 2	55	10 mA.g	-1	280	290	4.8 -1.5	[58]

Li 2 TiO 3 -Li 2 MnO 3

	Li 1.20 Ti 0.40 Mn 0.40 O 2	5 mA.g	-1	360	310	4.8 -1.5	[41]
	Li 1.20 Ti 0.40 Fe 0.40 O 2	5 mA.g -1	315	210	4.8 -1.5	[41]
	Li 1.20 Ti 0.40 Fe 0.40 O 2	10 mA.g -1	350	250	4.8 -2.0	[60]
	Li 6/5 Ti 2/5 Mo 2/5 O 2	10 mA.g -1	130	290	4.0 -1.0	[61]
	Li 1.20 Ni 0.333 Ti 0.333 Mo 0.133 O 2	20 mA.g	-1	250	230	4.8 -1.0	[55]

Li 3 NbO 4 -Li 3 MoO 4

Li 3 MoO 4

	Li 9/7 Nb 2/7 Mo 3/7 O 2	25	10 mA.g -1	150	280	4.0 -1.0	[61]

Li 3 MoO 4 -LiCrO 2

	Li 4/3 Mo 6+	2/9 Mo 3+	4/9 O 2	25	10 mA.g -1	180	330	4.0 -1.0	[61]
	Li 1.211 Mo 0.467 Cr 0.3 O 2	25	C/20	280	280	4.3 -1.5	[56]

  [START_REF] Nakajima | Lithium-Excess Cation-Disordered Rocksalt-Type Oxide with Nanoscale Phase Segregation: Li1.25Nb0.25V0.5O2[END_REF]

  [START_REF] Yabuuchi | Material Design Concept of Lithium-Excess Electrode Materials with Rocksalt-Related Structures for Rechargeable Non-Aqueous Batteries[END_REF]

)

  In Li1.2Ti0.4Fe0.4O2 electrodes, initial charge capacity reaches 350 mAh.g -1 and a potential plateau is observed at 4 V. During discharge, polarization increases, thus discharge capacity reduces to 250 mAh.g -1 . Similar observation is made on Li1.2Ti0.4Mn0.4O2

	electrodes: discharge capacity (320 mAh.g -1 ) is slightly smaller than the charge capacity
	(360 mAh.g -1 ) at the first cycle though subsequent cycles are fully reversible. The
	difference between first and subsequent cycles in Li1.2Ti0.4Mn0.4O2 is much smaller
	compared to the one in Li1.2Ti0.4Fe0.4O2. The author explains the reason with the
	reversibility of anionic redox process. Anionic redox process is almost reversible for
	Li1.2Ti0.4Mn0.4O2, whereas it is irreversible in Li1.2Ti0.4Fe0.4O2 electrodes.

Redox in Cation Disordered Rocksalt Oxides

	When anionic redox is active, high covalency of metal-ligand band is present. P orbital
	of ligand is overlapping with d orbital of metal atom, so ligand transfers electrons, fills
	d orbital of metal atom and leaves behind oxidized species (or holes). That means part
	of transition metal is reduced to lower valence state, this is called Reductive Coupling
	Effect in the literature ( 70 ). Such effect is reported for both Li-rich layered ( 69-74 ) and
	cation-disordered rocksalts ( 62,73 ) when anionic redox is activated.
	• Anionic As previously mentioned, Li1.2Ti0.4Mn0.4O2 electrodes delivered large reversible
	capacities. To examine such large reversible capacities in Li1.2Ti0.4Mn0.4O2 electrodes,
	XAS study is conducted by Li et al. ( 60 ). XAS spectra of Ti shows that valence state of

titanium does not change, therefore Ti does not contribute to charge transfer. Whereas, XAS spectra of Mn and oxygen progressively changed upon charging (Figure I -17).

  Figure III -3, disordered rocksalt Li2.2NiTi0.2Nb0.6O4 delivers 120 mAh.g -1 charge and 74 mAh.g -1 discharge capacities, corresponding respectively to 0.90 Li + and 0.56 Li +

	exchanged upon cycling. The irreversible capacity at the first cycle is 38%. On the
	contrary, second cycle is fully reversible and less polarized, but a much smaller
	capacity of 69 mAh.g -1 is delivered. Moreover, the cycle life at room temperature is

rather poor (Figure

III -4

). Those results indicate that synthesis improvement is required to achieve better specific capacities.

Table III -

 III

	Salt	Eutectic molar ratio	melting point/°C
	LiNO 3		255
	LiOH		477
	LiCl		610
	LiBr		552
	LiOH-LiCl	0.65 : 0.35	277
	LiCl-KCl	0.59 : 0.41	354
	LiOH-LiNO 3 0.40 : 0.60	243
	NaCl		801
	KCl		771
	NaCl-KCl	1.0 : 1.0	667

1: Melting point of salts and eutectic salts

At the beginning, firing at low temperature is considered because the melting point of the new salt is 354 °C. Two different powders are prepared based on the same recipe.

One sample is heated at 450 °C for 12 hours, thereafter it is quenched. Another is synthesized with the classical method; it is heated at 300 °C for 2 hours and calcined at 670 °C for 3 hours and quenched. The powders are washed with ethylene glycol, filtered and dried.

  Rapid capacity fading was observed in the liquid electrolytes, whereas 92% capacity was maintained after 80 cycles in solid-state electrolytes.

	Therefore, this material was suggested for solid-state batteries rather than for
	conventional batteries. In this chapter, our first aim is to develop disordered rocksalt
	Li2TiS3 with mechanochemical synthesis. Various structural, morphological and
	electrochemical characterizations are performed to understand the properties of
	Li2TiS3 compound. To solve cycle life problem, we propose substitution of Li2TiS3 with
	selenium. Therefore, this chapter also describes the development of a new synthesis
	process and new selenium substituted compositions (Li2TiSexS3-x).

Sakuda et al. prepared

Li2TiS3 as a disordered rocksalt phase by mechanochemical synthesis and they observed that this material could provide 425 mAh.g -1 capacity based on 2.5 Li + ion reversible exchange process. Beyond its interesting electrochemical performances, authors reported that the cycling life of this material was poor.

Table IV -

 IV

	Table IV -2: Coin cell results of Li2TiS3 compositions synthesized by ball milling for 15h, 20h,
	25h and 40h , where CQ1: first charge capacity, DQ1:first discharge capacity, Vav: average
	discharge potential at the first cycle

3: ICP-AES (Inductively coupled plasma atomic emission spectroscopy) of Li2TiS3 powders mechanochemically synthesized at different milling durations. Data is normalized

based on fixed Ti amount (Ti=1.00 mol) in the structure.

Table IV - 8 :

 IV8 Cyclic voltammetry test results of Li2TiS3, Li2TiSeS2 and Li2TiSe2S cells cycled at different cycling rates between 3 V -1.5 V vs Li + / Li•

Cycle life improvement with different cycling schedules

Table V

 V Se2 is found at 53.8 eV. The parameters can be found in Table V -1. The difference between these two energies and the presence of two Se doublets indicate that their valence states are not the same, Se1 could be regarded as Se 2-and the binding energy of Se2 is quite close to the Se2 2-(115,116 ). In the literature, Se 3d core peaks were reported between 53 and 54 eVs for different metal selenide systems ( 117,118 ). By semi-quantitative analysis, we also observe that Se1 is largely present (72.4%) compared to the amount of Se2 (27.7%) in Li2TiSeS2. Normally, we expect to detect only one type of selenide at Se 3d, but it seems that nearly one third of the Se 2-is already oxidized to Se2 2-at least at the surface of the material. This means that the amount of Li in Li2TiSeS2 is less than

	-1 shows all
	optimized fittings parameters and atomic percentages of Li2TiS3, Li2TiSeS2, Li2TiSe2S
	and selenium precursor at Li 1s (or Se 3d), Ti 2p3/2-1/2, and S 2p3/2-1/2.
	In Li2TiSeS2, we observe two different selenides (Se1 and Se2) at Se 3d core level and
	each selenide contains a doublet, the Se 3d5/2 and Se 3d3/2 peaks, with widths of 0.86 eV.
	Binding energy at Se 3d5/2 core of Se1 is detected at 52.7 eV, whereas binding energy of
	what we calculated (Li2-xTiSeS2).
	In Li2TiSe2S, similar peak positions are detected. Besides, we observed that ratio of Se1
	to Se2 increased since 81.6% of Se1 is present compared to Se2 (18.4%). The amount of
	Se 2-is much higher in Li2TiSe2S, as expected. Nevertheless, we observe again that some
	part of Se is oxidized, thus the amount of Li in Li2TiSe2S, is less than what we calculated
	(Li2-xTiSe2S).

  2-(S2).

	Contrary to Li2TiS3, only one S 2p doublet is observed in both Li2TiSeS2 and Li2TiSe2S
	compositions; however, two different selenium doublets, Se1 and Se2, are observed in
	Li2TiSeS2 at 158.8 eV and 159.4 eV in percentages of 21.2% and 7.2% respectively
	(Sulfur: 71.6%). As already mentioned before, Se2 might be regarded as Se2 2-. Similar
	species are detected in Li2TiSe2S, the amount of selenium species increase in Li2TiSe2S,
	for instance Se2 now reaches 20%, much higher than the one in Li2TiSeS2. On the

contrary, the percentage of Se1 found is 38.5% (Sulfur: 41.5%). In conclusion, all the species in the powders are detected at Li 1s, Ti 2p and S 2p cores and curve fittings models are developed. Such models are used for curve fitting of XPS analysis on electrodes at different states of charge.

Table V - 1 :

 V1 Optimized fitting parameters of Li2TiS3, Li2TiSeS2, Li2TiSe2S and selenium precursor at Li 1s/Se 3d, Ti 2p3/2-1/2, and S 2p3/2-1/2

	V.3.d.

XPS Analysis of Li2TiSexS3-x Pristine Electrodes

  After finding the best fitting models through powder analysis, we decide to continue XPS analysis on various electrodes. Before analyzing the electrodes at different states of charge, our next target is to investigate the surface and the bulk properties of the

  2-increases (80%) with S-C (10%). We explain this in a new hypothesis that sulfur may react with PVdF binder (production of LiF) and produce more oxidized sulfur species. That also indicates that less lithium is present in the composition, the actual formula might be Li2-xTiS3. Additionally, it might be the reason of why S-C bond is only seen in electrodes, not in powders. Sulfur may react with carbon in the PVdF and created S-C compound. In Ti 2p core, signal of Ti 4+ in TiS2 environment becomes much lower in electrodes, whereas signal of Ti 4+ in TiO2 environment becomes stronger in electrodes.

	Similar phenomenon is observed for Li2TiSeS2 and Li2TiSe2S, LiF and S-C species are
	detected in Li 1s core and S 2P core levels respectively. That also strengthens our
	hypothesis that sulfur might react with PVdF binder.
	Overall, peak intensity of lithium and sulfur species are much stronger in scratched
	electrodes, therefore we decide to perform post mortem analysis on scratched
	electrodes. Since our next target is to analyze sulfur-selenium redox process, S 2p core
	levels of scratched electrodes are investigated through high-resolution XPS.

Same analyses are conducted for both Li2TiSeS2 and Li2TiSe2S compositions. It should be noted that all the peak positions and peak separations are fixed at the beginning of the XPS analysis of scratched and non-scratched electrodes, thus the results obtained in both electrodes and powders are similar (Figure V -18 and Figure V -19). The main differences between scratched and non-scratched electrodes appear in Li 1s core and S 2p core levels. Much higher selenium peak intensities at Li 1s core level are observed in non-scratched electrodes, whereas intensities become lower in S 2p core peaks.

  2- as central disulfide species that are bonded with Li + and Ti 4+ ions ( 122 ). Upon charging, lithium ions are extracted from structure, central disulfide species (S2 2-) lose some lithium ions and only Ti-S bonded sulfur species (S2 2-) remain. Meanwhile, central disulfide species that lose neighbor lithium ions prefer creating disulfides (S 2-), therefore we observe an increase of S 2-species at the end of the charge. The similar observation is reported by A. Sakuda et al., their Monte Carlo simulation results

Acknowledgments

Co, Nb, Mo, etc) and different chalcogenides, Li2TiX3, (X: S, Se, Te, etc.) can be prepared ( 100,101 ). second cycle, indicating that the second redox process is reversible. This new peak becomes much more defined in Li2TiSe2S cell, it is observed at 2.12 V and 1.88 V. V. The same porous morphology is also seen in the electrodes cycled at 3 V -1.8 V, but it is slightly eliminated by increasing the cut-off discharge potential.

Figure IV -14: Comparison of wet and dry mechanochemical syntheses: XRD pattern of Li2TiS3 powders prepared by wet and dry mechanochemical syntheses

In Despite detailed characterization of Li2TiSexS3-x compounds, we still do not know the redox active species and possible redox mechanisms. To discover such properties, we further characterize these compounds with X-ray Photoelectron Spectroscopy (XPS),

Raman Spectroscopy and neutron diffraction in the following chapter.

General Conclusion and Perspectives Appendix

Abstract « Synthesis and characterization of new positive electrode materials for Li-ion high energy applications »

This thesis focuses on the development of overlithiated disordered rocksalts for high-energy

Li-ion systems. Firstly, synthesis parameters have been optimized to improve the performances of the disordered rocksalt Li2.2NiTi0.2Nb0.6O4. To examine its high irreversibility (35%) at the first cycle, in situ advanced structural and electrochemical analyses have been performed. Results show that a structural change and disordering happen during the first charge. In a second part, the disordered rocksalt Li2TiS3 has been prepared with our patented process. To improve cycling stability of the cells, Li2TiS3 has been partially substituted with selenium and new Li2TiSexS3-x compositions have been prepared. Li2TiSexS3-x cells have large discharge capacities at slightly lower potentials. Reversible sulfur redox activity is confirmed by electrochemistry and ex situ surface analyses, however further characterizations are required to elucidate the relatively complex selenium redox process.