First and foremost, I want to convey my deep appreciation to Jean-Marie Tarascon for dedicating significant time and effort to assist me in constructing my articles and presentations. His guidance has been a crucial factor throughout these three years. I also want to express my gratitude for creating a remarkable laboratory and for consistently being accessible and supportive of my thesis work while allowing me a significant degree of autonomy.

My gratitude also goes to Romain Dugas for his great help. His scientific knowledge and the discussions we had were extremely helpful for this work. I also highly benefited from the three-electrode cell that he designed. I wish to express my appreciation for Jürgen Janek and Laurence Croguennec for willing to undertake the task of refereeing this thesis, and to Sylvain Franger and Olivier Guillon for agreeing to be part of the jury. I am grateful for the time spent on reviewing my PhD work.

Many thanks to the collaborators who have contributed significantly to my PhD work.

Special mentions go to Anshuman Chaupatnaik for his work on Li4Ti5O12 synthesis, Ronan

Chometon and Juan Forero-Saboya for their contributions to SEM, Gwenaelle Rousse for her expertise in crystallography and XRD interpretation, and Benjamin Hennequart for his valuable input on Halide solid electrolyte.

My warmest thanks goes to my wonderful co-workers Ronan Chometon and Benjamin

Hennequart for the friendly and supportive environment I benefited from working alongside you. It has been a true pleasure to work in the solid-state team with you. I am also grateful to many of my past and present lab colleagues for creating a joyful environment despite the ups and downs of the PhD. Special mentions go to Ivette, Ronan, Benjamin, Clémence, Aleksandra, Juan, Zhenying, Alexia, Laura and others for the memorable times we spent together inside and outside the lab. I wish you all the best. Finally, my deepest appreciation goes to my family, for their unwavering support throughout this journey.

Broader context and thesis outline

Broader context

The development of the steam engine in 1780 marked the beginning of the industrial revolution, leading to an era of significant anthropogenic influence on Earth's geology and ecosystems, known as the Anthropocene era [START_REF] Lewis | Defining the Anthropocene[END_REF] . This period is characterized by the unregulated consumption of finite energy sources such as coal and petroleum, resulting in greenhouse gas emissions and climate change. To limit the environmental and health impacts of these products, it is necessary to reduce energy consumption and to transition to renewable energy sources, such as wind, geothermal, and solar power. Energy-storage technologies are essential in this transition, as they provide flexibility to fluctuating renewable energies.

Li-ion batteries have undergone significant advancements and are now the dominant energy storage system for portable devices, providing high energy density and reliability [START_REF] Larcher | Towards Greener and More Sustainable Batteries for Electrical Energy Storage[END_REF][START_REF] Goodenough | The Li-Ion Rechargeable Battery: A Perspective[END_REF] .

Batteries also play a crucial role in the development of electric vehicles, which are expected to exponentially increase in demand in the coming years due to EU policies (no more gasolinefuelled car sold in EU after 2035) 4 . To meet the demand for higher safety and energy density in EV batteries, research is focusing on "beyond Li-ion technology", particularly All-Solid-State Batteries (ASSB), which are proposed as a superior alternative to Li-ion batteries [START_REF] Masias | Opportunities and Challenges of Lithium Ion Batteries in Automotive Applications[END_REF][START_REF] Placke | Lithium Metal, and Alternative Rechargeable Battery Technologies: The Odyssey for High Energy Density[END_REF][START_REF] Scrosati | Lithium Batteries: Status, Prospects and Future[END_REF] . This thesis aims to contribute to the development of ASSB, as elaborated in the subsequent section.

Thesis outline

The aim of this thesis is to contribute to ASSB development, and more specifically to explore the interfacial detrimental phenomena occurring in cathode composite, based on Li6PS5Cl solid electrolyte (SE). Indeed, mastering solid-solid interfaces is at the core of ASSB optimization.

This manuscript is structured in the following chapters.

Chapter 1: Background and motivation

Chapter 1 gives a brief overview of the lithium-ion battery technology history. It is followed by the core of this thesis subject, namely the ASSB. We explain why ASSB are seen as the next generation of battery, especially to power the EV. Then, we will go deeper by studying the cathode composites interfacial challenges. Finally, an outline of the best ASSB performances released in recent years is given.

Chapter 2: Electrochemical/chemical stability of Li6PS5Cl based cathode composites

Chapter 2 serves as a base to understand what drives the performances of a cathode composite based on Li6PS5Cl (Argyrodite) solid electrolyte. This solid electrolyte small stability window hinders the stability and kinetic performances. It is therefore crucial to understand its mechanistic in contact or not with other materials. The influence of electronically conductive additive is discussed, so as the active material nature. We decoupled two interfacial reactivities affecting battery performances and tackled their properties and parameters of importance. This study uses impedance spectroscopy as tool to decorrelate interfaces.

Chapter 3: In situ electronic conductivity of Li6PS5Cl based composites

The electronic conductive additive is proven to enhance degradation in Chapter 2. Thus, Chapter 3 explores the possibility to suppress it. Focusing on the electronic conductivity of carbon-free cathode composites during cycling, we looked at the effect of particle size, volumetric changes of active material and active material nature on composites electronic conductivity. To proceed, we use a homemade unique device to track in situ the electronic conductivity changes of cathode composites.

Chapter 4: Exploring high potential degradation of Li6PS5Cl based composites

To suppress interfacial reactivities and enhance the stability of Li6PS5Cl composites, we explore the impact of plural parameters, namely the assembly pressure, the cathode loading, the impact of carbonate coating on SE and active material (AM), and formation cycle protocols. In the end, we combined the best parameters and propose an optimized composites/cycling protocol. negative electrode, which is typically made of graphite, and a positive electrode, which is usually composed of a layered oxide material or LiFePO4. These two electrodes are separated by a porous separator, which is filled with an organic electrolyte to prevent short circuit and enable Li + ions conduction. The electrolyte contains LiPF6 salt dissolved in a mixture of alkyl carbonate solvents, such as ethylene carbonate (EC) and dimethyl carbonate (DMC).

From liquid electrolyte-based

Commercial electrolytes may also contain other additives to suit specific operating conditions.

The electrodes are fabricated using a slurry process on a metal foil, which involves active material particles (typically of micrometre-scale), an electrically conductive additive (such as carbon-based ones: Super P, C65, or VGCF), and a polymeric binder. All electrodes alike the separator are porous, and filled by the electrolyte to ensure ionic motion near active surfaces. Several parameters are required for optimal performances of rechargeable batteries.

These parameters encompass high energy density, high specific energy, low cost, extended cycle life, high power density, high safety, and elevated temperature tolerance. Among these parameters, energy density and specific energy are of paramount importance for electric vehicle (EV) development. These parameters are typically calculated at a particular current (i.e., a specific C-rate) and are the product of the cell potential by the capacity. It can be expressed in per unit of weight (Wh/kg) (following equation below), or per unit of volume (Wh/L) : 𝐸 (𝑊ℎ. 𝑘𝑔 -1 ) = 𝑄(𝐴. ℎ. 𝑘𝑔 -1 ). 𝑈(𝑉) Specific Energy (or massic energy, in Wh/kg) relates to the amount of energy that can be stored per mass. As example, lead-acid battery can store 30-70 Wh/kg, while Li-ion battery reach 250-370 Wh/kg. Energy density (or volumetric energy, in Wh/L) relates to the amount of energy that can be stored per volume. It is an important measure for application needing small volume storage, such as smartphones or EV. As example, lead-acid battery can store around 80 Wh/L, while a Li-ion battery pack reached 600 Wh/L.

The energy density and specific energy of a cell mainly depend on the specific capacity (mA.h/g) and potential (V) of the two electrodes. Thus, tuning the properties of the active material unlocked the performances of Li-ion batteries [START_REF] Abakumov | Solid State Chemistry for Developing Better Metal-Ion Batteries[END_REF] . To tune those properties, one needs to understand the specific capacity and the cell voltage.

The specific capacity of an active material is defined by the theoretical amount of charge species, which can be insert/disinsert in the structure and is calculating with the following equation: Owing to the high theoretical capacity of lithium (3860 mA.h/g), its low potential (-3.04

V vs. SHE) and its light molar weight (6.94 g/mol), lithium-insertion compounds were early on explored in primary and secondary batteries. The road to material development for Li-ion battery is a rich history dating back to the 1970's. In 1975, chalcogenide TiS2 was first reported by Whittingham as an interesting lithium intercalation material for Li metal batteries, following the reactions described below [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 | Lithium Batteries and Cathode Materials[END_REF] .

(𝑎𝑡 𝑐𝑎𝑡ℎ𝑜𝑑𝑒) 𝑇𝑖𝑆 2 + 𝐿𝑖 + + 𝑒 -⇆ 𝐿𝑖𝑇𝑖𝑆 2 (𝑎𝑡 𝑎𝑛𝑜𝑑𝑒) 𝐿𝑖 ⇆ 𝐿𝑖 + + 𝑒 -

The Li metal/TiS2 and Li metal/MoS2 systems were initially considered as promising and commercialization with Moli Energy was launched, but high safety risks were revealed. Such failure originates from the lithium dendrite growth issues, leading to internal short-circuits and fire accidents [START_REF] Blomgren | The Development and Future of Lithium Ion Batteries[END_REF] . To overcome this issue, the "rocking-chair battery" concept was proposed in 1980 [START_REF] Lazzari | A Cyclable Lithium Organic Electrolyte Cell Based on Two Intercalation Electrodes[END_REF] . This idea supposes the replacement of Li metal by an insertion material at the negative electrode, especially LiC6 graphite. Graphite is a competitive anode material, due to its high surface area and theoretical capacity of 372 mA.h/g [START_REF] Yazami | A Reversible Graphite-Lithium Negative Electrode for Electrochemical Generators[END_REF] . The cathode is replaced by a lithiated oxide intercalation material. This choice is driven by the need for a new Li source in the battery, as Li anode is removed. Additionally, the increased in cell potential of oxides as compared to chalcogenide, and the higher electronegativity of Oxygen as compared with Sulphur, are beneficial for cell performances. This new technology led to the commercialization in 1991 of a Sony cell, using graphite as anode and LiCoO2 as cathode, an oxide layered compound developed by John Goodenough [START_REF] Blomgren | The Development and Future of Lithium Ion Batteries[END_REF]15,[START_REF] Ozawa | Lithium-Ion Rechargeable Batteries with LiCoO2 and Carbon Electrodes: The LiCoO2/C System[END_REF] . Despite the good reversibility of the LiCoO2 reactions depicted below, a limited capacity is obtained, and cobalt mining poses serious ethical issues [START_REF]This Is What We Die for: Human Rights Abuses in the Democratic Republic of Congo Power the Global Trade in Cobalt[END_REF] .

(𝑎𝑡 𝑐𝑎𝑡ℎ𝑜𝑑𝑒) 𝐿𝑖𝐶𝑜𝑂 2 ⇆ 𝑥𝐿𝑖 + + 𝑥𝑒 -+ 𝐿𝑖 1-𝑥 𝐶𝑜𝑂 2 (𝑎𝑡 𝑎𝑛𝑜𝑑𝑒) 𝐶 𝑛 + 𝑥𝐿𝑖 + + 𝑥𝑒 -⇆ 𝐶 𝑛 𝐿𝑖 𝑥 Thus, other compounds were discovered, such as the "spinel" LiMn2O4 and the "polyanionic"

LiFePO4 [START_REF] Tarascon | The Li-Ion Battery: 25 Years of Exciting and Enriching Experiences[END_REF][START_REF] Padhi | Phospho-olivines as Positive-Electrode Materials for Rechargeable Lithium Batteries[END_REF][START_REF] Thackeray | Lithium Insertion into Manganese Spinels[END_REF] . Moreover, the layered oxide materials family largely expanded, with the exploration of transition metals that are cheaper and more abundant than cobalt [START_REF] Lee | Can Cobalt Be Eliminated from Lithium-Ion Batteries?[END_REF] . As such, in 2000's the NMC and NCA family (LixNi1-yMnyCoyO2 and LixNiwCoyAlzO2) started to be explored [START_REF] Rozier | Review-Li-Rich Layered Oxide Cathodes for Next-Generation Li-Ion Batteries: Chances and Challenges[END_REF][START_REF] Yoon | High-Energy Ni-Rich Li[Ni x Co y Mn 1 -x-y ]O 2 Cathodes via Compositional Partitioning for Next-Generation Electric Vehicles[END_REF][START_REF] Noh | Comparison of the Structural and Electrochemical Properties of Layered Li[END_REF] . In those compounds, the cobalt is partially substituted by other transition metals, favouring greater stability and energy density. For instance, transitioning from LiNi0.33Mn0.33Co0.33O2 (NMC 111) to LiNi0.6Mn0.2Co0.2O2 (NMC 622) enabled a capacity enhancement from 150 mA.h/g to 200 mA.h/g [START_REF] Noh | Comparison of the Structural and Electrochemical Properties of Layered Li[END_REF] . Today, NMC/NCA active materials are widely commercialized for industrial application.

To even push further the energy density of the batteries, research is heading towards new generations of storage systems, especially All-Solid-State Batteries (ASSB).

Expected benefits of All-Solid-State Batteries

After a decade of oblivion, the discovery of LGPS (Li10GeP2S12), a highly conductive solid-state electrolyte with ionic conductivity comparable to that of liquid electrolytes (10 -2 S/cm), sparked renewed interest in the development of all-solid-state batteries in 2011 [START_REF] Kamaya | A Lithium Superionic Conductor[END_REF] . The ideal ASSB cell design includes a negative lithium metal electrode, a solid electrolyte separator, and a positive lithium intercalation electrode, as shown in Figure 1.2 [START_REF] Janek | A Solid Future for Battery Development[END_REF] .

These batteries are viewed as the next major generation of energy storage systems for applications such as electric vehicles, offering higher energy density, specific energy, and individual safety. Indeed, ASSB use solid ionic conductors, either oxide, halide or sulphidebased, rather than the organic liquid electrolytes used in conventional Li-ion batteries. As a result, ASSB are expected to be safer, with no risk of flammable electrolyte leakage, and to have a higher cycling kinetics, as the Li + transference number in solid electrolytes is unity, eliminating concentration polarization observed in liquid systems [START_REF] Bachman | Inorganic Solid-State Electrolytes for Lithium Batteries: Mechanisms and Properties Governing Ion Conduction[END_REF][START_REF] Zhang | New Horizons for Inorganic Solid State Ion Conductors[END_REF] .

Moreover, great hopes were placed on the possibility to use a lithium metal anode, thanks to greater mechanical strength of SE as compared to liquids. However, in recent years numerous studies have demonstrated that SE cannot prevent dendrite growth [START_REF] Kasemchainan | Critical Stripping Current Leads to Dendrite Formation on Plating in Lithium Anode Solid Electrolyte Cells[END_REF][START_REF] Schnell | All-Solid-State Lithium-Ion and Lithium Metal Batteries -Paving the Way to Large-Scale Production[END_REF][START_REF] Lee | High-Energy Long-Cycling All-Solid-State Lithium Metal Batteries Enabled by Silver-Carbon Composite Anodes[END_REF] . Thus, despite the promising performances of ASSB, remaining issues need to be tackled by the industrial and academic community, as explained in next section.

Context and challenges regarding ASSB development

Despite the promising potential of all-solid-state batteries, their performances remain poor due to various micro to macro level issues. On the road to full cell battery packs, numerous problematics remain.

Firstly, interfacial instability issues at both cathode and anode sides, stemming from the small SE stability window, remain significant obstacles [START_REF] Albertus | Publisher Correction: Status and Challenges in Enabling the Lithium Metal Electrode for High-Energy and Low-Cost Rechargeable Batteries[END_REF] . Poor stability is often reported, as compared with commercial liquid Li-ion cells. Especially, developing a stable lithium metal/solid electrolyte interface would unlock the ASSB performances.

Secondly, the cathode composite pressure needs to be reduced to atmospheric or low pressure to allow for integration of the lithium metal anode and to achieve a competitive volumetric energy. The issue of maintaining intimate contact at low external pressure, while AM particles undergo internal volume changes upon cycling, requires efficient processing and wise choice of materials [START_REF] Hlushkou | The Influence of Void Space on Ion Transport in a Composite Cathode for All-Solid-State Batteries[END_REF][START_REF] Zhang | Chemical Expansion during Cycling: Monitoring the Pressure Changes in Operating Solid-State Lithium Batteries[END_REF] .

Thirdly, a higher loading and higher AM wt. % (> 80 %) need to be reached in the cathode, without compromising the performances [START_REF] Zaman | Processing and Manufacturing of next Generation Lithium-Based All Solid-State Batteries[END_REF] . The current literature proposes good performances with very low cathode composite loading (< 5 mg/cm²) and 70% weight percentage of AM. A factor of 10 increase in loading is needed to achieve acceptable volumetric and mass energy density [START_REF] Schnell | All-Solid-State Lithium-Ion and Lithium Metal Batteries -Paving the Way to Large-Scale Production[END_REF][START_REF] Kato | All-Solid-State Batteries with Thick Electrode Configurations[END_REF] .

Lastly, decreasing and mastering solid electrolyte layer thickness is crucial for cost and energy density. Indeed, one kilogram of Li7La3Zr2O12 (LLZO), LGPS and Li6PS5Cl costs respectively $2000, $69,500 and $36,000 [START_REF] Zaman | Processing and Manufacturing of next Generation Lithium-Based All Solid-State Batteries[END_REF][START_REF] Schnell | Prospects of Production Technologies and Manufacturing Costs of Oxide-Based All-Solid-State Lithium Batteries[END_REF][START_REF] Mauler | Battery Cost Forecasting: A Review of Methods and Results with an Outlook to 2050[END_REF] . In fact, as reported in Figure 1.3, the typically used solid electrolyte layer of 50 µm can bring 75% of dead weight to the battery. Note that a commercial Li-ion battery tend to have a separator thickness between 20 -25 µm [START_REF] Horváth | Quantifying the Effect of Separator Thickness on Rate Performance in Lithium-Ion Batteries[END_REF] . The high cost of SE materials underscores the importance of thin, defect-free solid electrolyte processing for large-scale applications, which is challenging due to the rigid and brittle nature of most SE [START_REF] Balaish | Processing Thin but Robust Electrolytes for Solid-State Batteries[END_REF] . Reproduced from reference [START_REF] Zaman | Processing and Manufacturing of next Generation Lithium-Based All Solid-State Batteries[END_REF] .

Bottlenecks regarding cathode composites and current state of the art

One of the main obstacles lies in attaining dense and highly energetic composites, which necessitates efficient formulation and processing techniques. Distinct methods and strategies were employed to reach this objective, what this section is all about.

Processing and formulation of solid-state composites

A solid-state composite is typically composed of a lithium-ion conducting phase (SE), and an active electronically conductive phase (AM and possibly carbon additives). For effective performances, a composite cathode design should minimize voids, maximize the AM/SE exchange surface, and include the minimum amount of SE necessary to ensure sufficient Li diffusion from the bulk SE to the AM/SE interfaces. Theoretically, this can be achieved by using infinitely small interconnected ionic/electronic conductive materials, to create an infinite exchange surface. Obviously, reality deviates from this model as the composites are composed of particles with nm to µm scale sizes (AM: 100 nm to 20 µm, SE: 100 nm to 100 µm), and their architecture being akin to an isotropic system (see Figure 1.4). This is translated in the formation of isolated or disconnected regions of lithium-ion/electron conducting phase, which lead to poor ionic/electronic composites conductivity, reduced battery efficiency and forces the use of low composites loading [START_REF] Bielefeld | Microstructural Modeling of Composite Cathodes for All-Solid-State Batteries[END_REF][START_REF] Bielefeld | Modeling Effective Ionic Conductivity and Binder Influence in Composite Cathodes for All-Solid-State Batteries[END_REF] . One of the key metrics to assess the quality of a composite is the porosity [START_REF] Hlushkou | The Influence of Void Space on Ion Transport in a Composite Cathode for All-Solid-State Batteries[END_REF][START_REF] Minnmann | Designing Cathodes and Cathode Active Materials for Solid-State Batteries[END_REF] . The level of porosity varies depending on the type of SE and their processing, with oxide SE having the lowest porosity due to high-temperature sintering, typically reaching less than 5 % [START_REF] Raju | Crystal Structure and Preparation of Li7La3Zr2O12 (LLZO) Solid-State Electrolyte and Doping Impacts on the Conductivity: An Overview[END_REF] . However, this manuscript will focus on sulphide-based SE, generally presenting between 20 and 30 % porosity. Altogether, these above limitations need to be addressed. To do so, previous studies have focused on various parameters, including assembly pressure, the AM/SE ratio, AM and SE particle sizes, and mixing technique efficiency, which will be discussed here. Reproduced from reference [START_REF] Tsai | A Garnet Structure-Based All-Solid-State Li Battery without Interface Modification: Resolving Incompatibility Issues on Positive Electrodes[END_REF] .

Of utmost importance are the mixing technique and protocol used. Indeed, the production of ASSB composites encompasses multiple processes that offer varying degrees of homogeneity and repeatability. This variability in processes finds analogy in the domains of asphalt and concrete design, wherein numerous protocols where investigated. Their efficiency necessitates the consideration of several factors, including the distribution of voids and the quantity and spatial arrangement of coarse aggregate particles (coarse aggregates are a construction component made of rock quarried from ground deposits) [START_REF] Behnood | Morphology, Rheology, and Physical Properties of Polymer-Modified Asphalt Binders[END_REF][START_REF] Rubio | Warm Mix Asphalt: An Overview[END_REF][START_REF] Ferraris | Concrete Mixing Methods and Concrete Mixers: State of the Art[END_REF] . Our research aligns with these concerns and can draw insights from these researches, while distinguishing between dry and wet processes.

Dry processing commonly involves hand grinding of the components, using an agate mortar and a pestle, or ball-milling techniques, with the latter being preferred for higher homogeneity. Both approaches offer the advantages of rapid implementation and ease of use [START_REF] Kerman | Review-Practical Challenges Hindering the Development of Solid State Li Ion Batteries[END_REF][START_REF] Park | Design Strategies, Practical Considerations, and New Solution Processes of Sulfide Solid Electrolytes for All-Solid-State Batteries[END_REF] . In this perspective, current research is exploring two paths, using innovative methods such as mechano-fusion, as well as the optimization of various ball-milling parameters [START_REF] Kim | High-Performance All-Solid-State Batteries Enabled by Intimate Interfacial Contact Between the Cathode and Sulfide-Based Solid Electrolytes[END_REF][START_REF] Cronau | Ionic Conductivity versus Particle Size of Ball-Milled Sulfide-Based Solid Electrolytes: Strategy Towards Optimized Composite Cathode Performance in All-Solid-State Batteries[END_REF] .

Considering ball-milling parameters, it typically include the sequence of material addition, rotation speed, composite mass, number of balls, jar volume, and duration of rotation. Given the complexity of these factors, it is not surprising that ball-milling techniques often lack repeatability.

In order to address these limitations, researchers have turned their attention to wet processes, which typically offer improved homogeneity and repeatability. These wet processes involve the use of slurry-based techniques and ball milling with the aid of a solvent, with the latter still submitted to the aforementioned considerations. Slurry based techniques are industrially adaptable for ASSB, and their processing is comprised of conventional Li-ion batteries coating strategies: wet/dry mixing of composites, tape formation, stacking and lamination [START_REF] Inada | Fabrications and Properties of Composite Solid-State Electrolytes[END_REF][START_REF] Nam | Bendable and Thin Sulfide Solid Electrolyte Film: A New Electrolyte Opportunity for Free-Standing and Stackable High-Energy All-Solid-State Lithium-Ion Batteries[END_REF] . In this perspective, current research considers the premixing of some components to achieve higher electronic/ionic conductivity, the reactivities of the distinct battery compounds with the solvents and proper ink formulation [START_REF] Laue | Modeling the Influence of Mixing Strategies onMicrostructural Properties of All-Solid-State Electrodes Vincent Laue[END_REF][START_REF] Nam | Toward Practical All-Solid-State Lithium-Ion Batteries with High Energy Density and Safety: Comparative Study for Electrodes Fabricated by Dry-and Slurry-Mixing Processes[END_REF] . Thus, transitioning to a slurry-based technique necessitates a revaluation of the composite formulation as compared to dry processes based ASSB composites, where all the difficulty lies.

Turning to battery assembly pressure, it is of utmost importance as it can affect the relative density of the composite and the electronic/ionic percolation, by enhancing interparticle contact [START_REF] Zhang | A Review of the Effect of External Pressure on All-Solid-State Batteries[END_REF] . This is exemplified by switching formation pressure from 55 MPa to 370 MPa, the relative density of Li6PS5Cl pellets increases from 68 to 77 %, as shown in Figure 1.5 [START_REF] Doux | Stack Pressure Considerations for Room-Temperature All-Solid-State Lithium Metal Batteries[END_REF] . It results in a threefold increase in ionic conductivity, as seen in Figure 1.6, when going to 370 MPa. Additionnaly to these processing aspects, composites formulation is critical as it significantly impact battery performances [START_REF] Bielefeld | Microstructural Modeling of Composite Cathodes for All-Solid-State Batteries[END_REF][START_REF] Minnmann | Designing Cathodes and Cathode Active Materials for Solid-State Batteries[END_REF][START_REF] Shi | High Active Material Loading in All-Solid-State Battery Electrode via Particle Size Optimization[END_REF][START_REF] Zhang | The Detrimental Effects of Carbon Additives in Li 10 GeP 2 S 12 -Based Solid-State Batteries[END_REF] . In composites typically comprising active material, solid electrolyte and a carbon additive, an ideal particle size ratio is typically achieved when the SE size is 2-3 times smaller than the AM particle size, as illustrated in Figure 1.7 [START_REF] Shi | High Active Material Loading in All-Solid-State Battery Electrode via Particle Size Optimization[END_REF] .

Alongside the particle size optimization, the AM/SE mass ratio should also be optimized, as highlighted through experimental and computational studies, with a high fraction of AM desired (> 80 wt. %) to reach high energy density [START_REF] Shi | High Active Material Loading in All-Solid-State Battery Electrode via Particle Size Optimization[END_REF][START_REF] Froboese | Effect of Microstructure on the Ionic Conductivity of an All Solid-State Battery Electrode[END_REF][START_REF] Strauss | Impact of Cathode Material Particle Size on the Capacity of Bulk-Type All-Solid-State Batteries[END_REF] . However, it is important to acknowledge certain limitations, such as the adverse effects of solid-state diffusion, particularly when the particle size of the AM exceeds a certain threshold, as illustrated in Overall, to optimize the performance of ASSB composites, it is necessary to conduct a comprehensive experimental plan, which is time-consuming and limited to a single combination of AM and SE materials and particle size distributions. Thus, computational holistic studies are valuable for addressing these challenges. Despite the difficulty, combining experimental and theoretical studies can provide a solid foundation for understanding the effects of various design factors such as material quantity, particle size, pressure, and mixing process on the performance of ASSB composites. Thus, by carefully selecting the optimal values for each of these factors, it is possible to access the full potential of composites. It is noteworthy to emphasize that the mixing process efficiency holds significant importance although the literature on this subject is limited, particularly in terms of comparative analyses and repeatability assessments, with such investigations predominantly confined to industrial secret research.

Additionally to this complex spectrum of parameters, one should also add the different chemical/electrochemical degradation processes, occurring at all composites interfaces during cycling. It will depend on the carbon and AM exchange surfaces, the nature of the compounds, and the cycling parameters. Such phenomena need to be taken into account in the optimization process, as we will see in the following sections.

Interfacial chemical/electrochemical reactivities in ASSB composites

Various parameters such as the chemical and electrochemical nature of the reaction, the conductive properties of the new phases formed, the potential, growth rate, and the location of the reaction can influence battery performances. Three types of degradation should be considered for ASSB composites, namely environmental reactivity, intrinsic electrochemical decomposition of the solid electrolyte, and reactivity towards the active materials. The type of SE used determines the distinct interfacial degradations observed, and different scenarios occur among the three main ionic conductor families (oxide, sulphide, and halide).

Environmental reactivity is a major concern for large-scale battery fabrication. Comparing the three SE types, oxide-based SE demonstrate the highest environmental stability [START_REF] Ren | Oxide Electrolytes for Lithium Batteries[END_REF] . Both halide and sulphide SE are reacting with humidity, with halide presenting a lower health hazard risk, as compared to sulphides that are well-known for releasing H2S and lithium hydroxide (LiOH) when exposed to air and humid environment [START_REF] Wang | Prospects of Halide-Based All-Solid-State Batteries: From Material Design to Practical Application[END_REF][START_REF] Nikodimos | Chemical Stability of Sulfide Solid-State Electrolytes: Stability toward Humid Air and Compatibility with Solvents and Binders[END_REF] . Despite this well-known process, two other degradation pathways have been identified, namely the intrinsic stability of the SE and the reactivity with layered oxide AM.

In relation to the intrinsic stability of a solid electrolyte, it corresponds to the range of potential values in which the SE remains stable. This stability is evaluated using both experimental techniques, especially cyclic voltammetry (CV), and computational methods such as Density Functional Theory (DFT) [START_REF] Zhu | Origin of Outstanding Stability in the Lithium Solid Electrolyte Materials: Insights from Thermodynamic Analyses Based on First-Principles Calculations[END_REF][START_REF] Dewald | Experimental Assessment of the Practical Oxidative Stability of Lithium Thiophosphate Solid Electrolytes[END_REF] . These two approaches rely on distinct foundations. Specifically, the DFT computational method employs thermodynamic considerations to determine the stability window, but it often conflicts with experimental data due to the significant influence of kinetics on battery performance. On the contrary, cyclic V versus Li-In/In [START_REF] Zhang | The Detrimental Effects of Carbon Additives in Li 10 GeP 2 S 12 -Based Solid-State Batteries[END_REF][START_REF] Zuo | A Mechanistic Investigation of the Li10GeP2S12|LiNi1-x-yCoxMnyO2 Interface Stability in All-Solid-State Lithium Batteries[END_REF][START_REF] Koerver | Capacity Fade in Solid-State Batteries: Interphase Formation and Chemomechanical Processes in Nickel-Rich Layered Oxide Cathodes and Lithium Thiophosphate Solid Electrolytes[END_REF][START_REF] Walther | Influence of Carbon Additives on the Decomposition Pathways in Cathodes of Lithium Thiophosphate-Based All-Solid-State Batteries[END_REF][START_REF] Auvergniot | Redox Activity of Argyrodite Li6PS5Cl Electrolyte in All-Solid-State Li-Ion Battery: An XPS Study[END_REF][START_REF] Tan | Elucidating Reversible Electrochemical Redox of Li 6 PS 5 Cl Solid Electrolyte[END_REF] . It leads to the formation of highly insulative products (such as LiCl and P2S5), and especially elementary sulphur [START_REF] Han | Electrochemical Stability of Li 10 GeP 2 S 12 and Li 7 La 3 Zr 2 O 12 Solid Electrolytes[END_REF][START_REF] Tan | Elucidating Reversible Electrochemical Redox of Li 6 PS 5 Cl Solid Electrolyte[END_REF] . The reaction mechanism proposed for Li6PS5Cl oxidation is the following [START_REF] Tan | Elucidating Reversible Electrochemical Redox of Li 6 PS 5 Cl Solid Electrolyte[END_REF] :

𝐿𝑖 6 𝑃𝑆 5 𝐶𝑙 ⇒ 𝐿𝑖𝐶𝑙 + 1 2 𝑃 2 𝑆 5 + 5 2 𝑆 + 5𝐿𝑖 + + 5𝑒 -
This reactivity is characterized by the additional slope observed in first charge of Li6PS5Clbased batteries (see Figure 1.10) and a fast capacity fading [START_REF] Walther | Influence of Carbon Additives on the Decomposition Pathways in Cathodes of Lithium Thiophosphate-Based All-Solid-State Batteries[END_REF][START_REF] Tan | Elucidating Reversible Electrochemical Redox of Li 6 PS 5 Cl Solid Electrolyte[END_REF][START_REF] Strauss | Influence of Electronically Conductive Additives on the Cycling Performance of Argyrodite-Based All-Solid-State Batteries[END_REF] . The degree of such reactivity is primarily reliant on the specific surface area of the electronically conductive additive, such as Vapour Grown Carbon Fibre (VGCF) or carbon nanotubes [START_REF] Tan | Elucidating Reversible Electrochemical Redox of Li 6 PS 5 Cl Solid Electrolyte[END_REF][START_REF] Tan | From Nanoscale Interface Characterization to Sustainable Energy Storage Using All-Solid-State Batteries[END_REF] . Note that this decomposition is seen in Lithium-Sulfur ASSB, resulting in capacities above theoretical [START_REF] Manthiram | Challenges and Prospects of Lithium-Sulfur Batteries[END_REF][START_REF] Pan | A Review of Solid-State Lithium-Sulfur Battery: Ion Transport and Polysulfide Chemistry[END_REF] . Lastly, stability of SE versus layered-oxide type AM is challenging, as it will depend on the AM utilization, the Ni-content of the AM, the amount of structural defects and surface species, among others. While oxide-type SE were previously considered stable in comparison to cathode active materials, recent studies have shown that they induce rapid capacity fading when cycled with NMC 811 or LiCoO2 (LCO) [START_REF] Tsai | A Garnet Structure-Based All-Solid-State Li Battery without Interface Modification: Resolving Incompatibility Issues on Positive Electrodes[END_REF][START_REF] Roitzheim | All-Solid-State Li Batteries with NCM-Garnet-Based Composite Cathodes: The Impact of NCM Composition on Material Compatibility[END_REF][START_REF] Finsterbusch | High Capacity Garnet-Based All-Solid-State Lithium Batteries: Fabrication and 3D-Microstructure Resolved Modeling[END_REF] [START_REF] Kuo | Origin of Structural Phase Transitions in Ni-Rich Li x Ni 0.8 Co 0.1 Mn 0.1 O 2 with Lithiation/Delithiation: A First-Principles Study[END_REF] . This latter possibility is well-known in liquid systems, where the release of O2 and the structural densification of the AM lead to a disastrous reactivity with the organic electrolyte [START_REF] Kong | Insights into the Structural Evolution and Li/O Loss in High-Ni Layered Oxide Cathodes[END_REF][START_REF] Ruan | Structural Evolution and Capacity Degradation Mechanism of LiNi0.6Mn0.2Co0.2O2 Cathode Materials[END_REF][START_REF] Xu | Phase Behavior during Electrochemical Cycling of Ni-Rich Cathode Materials for Li-Ion Batteries[END_REF] . As key metrics to characterize such phenomenon in ASSB, dioxygen release is monitored during cycling with solid-state OEMS/DEMS (Online/Differential Electrochemical Mass Spectroscopy), with its release approximately starting at 3.6 V vs. Li-In/In until the end of charge [START_REF] Walther | The Working Principle of a Li 2 CO 3 /LiNbO 3 Coating on NCM for Thiophosphate-Based All-Solid-State Batteries[END_REF][START_REF] Kim | Stabilizing Effect of a Hybrid Surface Coating on a Ni-Rich NCM Cathode Material in All-Solid-State Batteries[END_REF] . Interestingly, the gas release causes a porous interface around the AM particles, as exemplified in Figure 1.12. Consequently, this porous interface is likely to affect the AM/SE interfacial contacts, thereby necessitating further research in this area. Within the recent advancements and unresolved inquiries in the field, a captivating topic pertains to the greater high potential stability of composites based on nickel-rich NMC materials, as compared to liquid-based Li-ion batteries. The influence of structural densification on this stability remains unclear, raising questions regarding the possibility of an electrochemically active rock salt layer [START_REF] Mcgrogan | Compliant Yet Brittle Mechanical Behavior of Li2S-P2S5 Lithium-Ion Conducting Solid Electrolyte[END_REF] . Furthermore, a novel form of interfacial degradation has recently been identified in Li6PS5Cl-based composites, involving the interdiffusion of transition metals. Cryogenic transmission electron microscopy (cryo-TEM) was employed to visualize this phenomenon for the first time in 2023 [START_REF] Zhang | Transition-Metal Interdiffusion and Solid Electrolyte Poisoning in All-Solid-State Batteries Revealed by Cryo-TEM[END_REF] . Consequently, additional research is needed to fully comprehend the consequences of such phenomena on the durability of ASSB. Overall, this section has compiled information on various types of instabilities that manifest in different types of solid electrolytes, either of chemical or electrochemical origin.

The discussion has then been narrowed down to the specific issues encountered in composites alike Li6PS5Cl-based ones. Li6PS5Cl presents a narrow stability window and a high reactivity with NMC-like cathode materials at high potentials. Despite these reactivities, sulphide-based SE offer attractive benefits such as high ionic conductivity, ductility, abundancy and ease of fabrication [START_REF] Mcgrogan | Compliant Yet Brittle Mechanical Behavior of Li2S-P2S5 Lithium-Ion Conducting Solid Electrolyte[END_REF] . Hence, they hold potential as ideal materials for high-energy dense cathode composites if their reactivities can be effectively managed. In order to address these complex challenges while preserving the favourable properties of sulphides, protective coatings are commonly used to safeguard the interfaces between the solid electrolyte and anode materials. The subsequent section will delve further on this topic. Typical coatings usually composed of ternary or binary oxides, should exhibit stability within the operating potential window and act as a lithium-ion conducting barrier and be made of an electronic insulating material. Moreover, coatings must be conformal and thin to enable electron tunnelling and ensure electronic percolation despite their electronically insulating nature. Indeed, full coverage of the particle would hinder AM-AM electronic contact and inhibit the composite's functionality, as depicted in the Capacity wise, numerous studies have demonstrated that applying a coating to ASSB can lower interfacial resistances prior to cycling, leading to reduced polarization and increased capacity [START_REF] Walther | The Working Principle of a Li 2 CO 3 /LiNbO 3 Coating on NCM for Thiophosphate-Based All-Solid-State Batteries[END_REF][START_REF] Strauss | 2 ZrO 3 -Coated NCM622 for Application in Inorganic Solid-State Batteries: Role of Surface Carbonates in the Cycling Performance[END_REF] . This effect may arise from three potential mechanisms. First, adding a coating creates two new interfaces, the AM/coating and SE/coating interfaces, instead of the previous AM/SE interface. The sum of these two interfaces can lead to a lower interfacial resistance.

Especially, better contact can be achieved. On this, the impact of coating deposition technique is crucial (Atomic Layer Deposition, Wet chemical approaches, Spray coating, Chemical Vapour Deposition) as it can influence the wettability and uniformity of the coating. The second possibility is the reduction of space-charge layer effects. Specifically, the lithium chemical potentials of AM and SE are believed to provoke the formation of an interfacial barrier, which is a lithium deficient layer [START_REF] Haruyama | Space-Charge Layer Effect at Interface between Oxide Cathode and Sulfide Electrolyte in All-Solid-State Lithium-Ion Battery[END_REF][START_REF] De Klerk | Space-Charge Layers in All-Solid-State Batteries; Important or Negligible?[END_REF] . However, although computational studies have investigated this effect, experimental proof of detrimental space-charge layers is still lacking. Finally, it should be noted that surface species that are present on the surface of the AM particles because of storage and synthesis processes, can chemically react with the sulphide-based SE.

Coating the AM particles can prevent this chemical reaction from occurring, provided that there are no surface species or less reactive ones present at the coating surface.

Stability-wise, coatings have been shown to diminish capacity fading through multiple possible mechanisms. Firstly, they effectively reduce the reactivity between high potential active material and solid electrolyte, as evidenced by various studies which show a decrease in the release of O2 gas in coated active material [START_REF] Kim | Stabilizing Effect of a Hybrid Surface Coating on a Ni-Rich NCM Cathode Material in All-Solid-State Batteries[END_REF][START_REF] Strauss | 2 ZrO 3 -Coated NCM622 for Application in Inorganic Solid-State Batteries: Role of Surface Carbonates in the Cycling Performance[END_REF][START_REF] He | Countering Voltage Decay, Redox Sluggishness, and Calendering Incompatibility by Near-Zero-Strain Interphase in Lithium-Rich, Manganese-Based Layered Oxide Electrodes[END_REF][START_REF] Liang | A Gradient Oxy-Thiophosphate-Coated Ni-Rich Layered Oxide Cathode for Stable All-Solid-State Li-Ion Batteries[END_REF] . This suggests that coatings may confine or obstruct the release of O2 gas. Furthermore, coatings can hinder interdiffusion, as exemplified by reduced diffusion of Co in a Li2SiO3-coated LCO/Li2S-P2S5 composite, as shown in Figure 1.14 100 . Finally, coatings are hypothesized to limit and accommodate pressure changes by serving as an elastic layer surrounding active material particles, thereby enhancing mechanical cohesion [START_REF] Culver | On the Functionality of Coatings for Cathode Active Materials in Thiophosphate-Based All-Solid-State Batteries[END_REF]101 . Although research on this topic is still in its early stages, recent studies demonstrated this highly interesting impact of coating. To do so, polylithic layered oxide AM are preferentially used, as they are known for mechanically degrading rapidly due to their anisotropy which favours particle fracturing during chemo-mechanical effect of Li insertion/disintertion 102-104 . This is exemplified by the effect of infiltrating LiAlO2 inside polylithic Li1.2Mn0.26Ni0.26O2 which led to improved stability (see Figure 1.15).Through nanoindentation, HAADF-STEM (High-Angle Annular Dark-Field Scanning Electron Transmission Microscopy) and theoretical consideration, this coating is clearly proved to act as a mechanical binder [START_REF] He | Countering Voltage Decay, Redox Sluggishness, and Calendering Incompatibility by Near-Zero-Strain Interphase in Lithium-Rich, Manganese-Based Layered Oxide Electrodes[END_REF] . particles (Loading speed: 20 mN min -1 , the standard deviations of the applied crushing forces are less than 5%) (d) Cross-sectional SEM micrographs of Pristine-LRMN (blue dots) and LAO@LRMN (red dots) with calendering density of 3.3 g cm -3 . Adapted from reference [START_REF] He | Countering Voltage Decay, Redox Sluggishness, and Calendering Incompatibility by Near-Zero-Strain Interphase in Lithium-Rich, Manganese-Based Layered Oxide Electrodes[END_REF] .

Among the recent studies, the SE coating is starting to be investigated. This is Overall, significant progress has been achieved through the incorporation of coatings; however, several aspects warrant careful consideration. Firstly, the lack of consistency in coating methodologies and performances across different studies, even when utilizing coatings of the same nature. This discrepancy can be attributed to challenges associated with achieving precise uniformity and thickness of the coatings, as their synthesis involves numerous parameters regardless of the methodology employed. Furthermore, it should be noted that the quality of the active material's surface prior to coating plays a crucial role in the coating quality, despite being insufficiently explored in the literature. Secondly, there remains uncertainty regarding the differentiation between incidental doping and coating effects, as the high-temperature processing steps commonly involved in coating techniques can inadvertently introduce dopants to the surface of the active material. Especially, the introduction of dopants into the active material has demonstrated effectiveness in mitigating the occurrence of high potential O2 gas release and modifying the chemo-mechanical properties of the material, particularly in studies conducted in liquid environments [107][108][109][110] .

Thus, further studies are needed to fully comprehend the impact of coating and/or doping.

Current state of the art performances

In the latest years, remarkable performances were released from articles using innovative methodologies of cell design, new ionic conductor materials, the implementation of high energy anodes (lithium metal or silicon) and efficiently structured cathode composites.

One of the first disruptive solid-state papers was presented by Lee et al. in 2020, proposing an advanced design for both cathode and anode part. The optimized cathode composite contains small Li6PS5Cl particles (< 1 µm), a Li2O-ZrO2 coated NMC and VGCF carbon additive [START_REF] Lee | High-Energy Long-Cycling All-Solid-State Lithium Metal Batteries Enabled by Silver-Carbon Composite Anodes[END_REF] . The important breakthrough arise from the usage of an "anode less" system, were no excess lithium is used. The implementation of an Ag-C anode enables to override the persistent issue of lithium dendrite formation. This anode proved itself to enable homogeneous and efficient lithium lithiation/delithiation through silver particles alloying. This design enabled the formation of a prototype pouch cell (0.6 A.h) and long cycle life (> 1000 cycles). Noteworthy, it is the first time such performances were obtained at low pressure (20 bar). However, 60°C operating temperature is used to increase capacity and prevent dendrite formation. Nonetheless, this study is one of the first to demonstrate a long-term stability and high energetic battery design, paving the way for new high performances studies (see Figure 1.17). (1.0 C = 6.8 mA cm -2 ). Reproduced from reference [START_REF] Lee | High-Energy Long-Cycling All-Solid-State Lithium Metal Batteries Enabled by Silver-Carbon Composite Anodes[END_REF] .

In 2021, another breakthrough is reached on the anode side by Tan et al., with the successful implementation of the highly energetic silicone material (> 3500 mA.h/g), so far prevented by its volumetric change upon cycling (> 300 % expansion during lithiation) 111 . Such strategy also overcomes the hurdle of lithium metal integration. By using high pressure (1 t/cm²) during a formation cycle and a small loading of silicon (1.67 mg/cm²), 80% capacity retention after 500 cycles at 1C is reached. Noteworthy is the high silicon content of the anode (99.9 %) compared to traditional silicon anode. However, no low C-rate performances are shown, possibly hiding a long-term instability, well known for Li6PS5Cl SE using carbon additive.

Nonetheless, it is the first time silicon anode is successfully integrated in ASSB, opening the path for future optimization (see Figure 1.18). Such performances rely on a stable halide solid electrolyte, presenting high ionic conductivity, coupled with a high capacity active material. The relatively small SC-NMC811 particle size (2 µm), successfully match with 2 µm LYC particles, created a high electronic and ionic percolation made possible by tailored particle sizes. Noteworthy, no coating is used for the active material, pointing out at the oxidative stability of LYC and at the importance of solid of this SE, which can be easily synthesized by a simple solid-state synthesis route. Also, the lack of carbon additive, known to be detrimental with Argyrodite, suggests that a high C-rate in non-carbon containing system is possible. However, the extremely low loading of the cathode composite (2.5 mg/cm²) questions the impact of a higher loading on the high rate performances. In

Conclusion of the chapter

We have provided an overview of the development of Li-ion batteries, which has been driven by innovative research on electrode materials since the 1970s. However, the demand for batteries for electric vehicles requires higher energy and safety standards. Thus, we focused on advanced solid-state batteries, which have gained increasing interest as a nextgeneration battery and are nearing commercialization. Out of all SE material for ASSB, Li6PS5Cl exhibits excellent properties such as ionic conductivity, ductility, and easy processability.

Nevertheless, several scientific challenges must be addressed. Especially, we explored the current state of knowledge regarding the impact of composite's composition and processing and the impact of the distinct reactivities occurring in the ASSB composites interfaces. Both considerations hinder the ASSB performances. To overcome these challenges, an efficient and comprehensive optimization plan for composite formulation and the careful selection of appropriate coatings are being explored. Despite the continuous advancements in ASSB performances, research still lags behind the expected level required for commercialization.

This assessment is attributed to the formidable task of addressing specific issues related to solid-solid interfaces, including contact and mechanical evolutions, while simultaneously dealing with the low stability observed in current SE materials, versus lithium metal and in the cathode composites.

Introduction

As discussed in the preceding chapter, All-Solid-State Batteries have emerged as a prospective battery technology to meet the demands of the expanding electric vehicle market.

Among the various systems being investigated, sulphide-based ASSBs exhibit favourable characteristics such as ease of processing and promising performances. However, these systems encounter a significant drawback, namely long-term capacity degradation caused by increased decomposition reactions. Although previous literature has addressed these instabilities to some extent, a comprehensive analysis encompassing all aspects of these issues in our system is essential.

Thus, the current chapter presents an exhaustive investigation into the reactivities of Li6PS5Cl-based systems, using NMC 622 as reference material and considering diverse operational conditions, including changes of active material, voltage window and carbon additive amounts. Firstly, the study focuses on SE decomposition, by evaluating the influence of the type and quantity of carbon additives on the initial cycle and cell capacity retention Secondly, the reactivity at the interface between NMC and solid electrolyte is examined, with a detailed study of the effects of potential window and active material characteristics. Finally, the study delves into the effects of both active material and electronic conductive additive by employing the Potentiostatic Electrochemical Impedance Spectroscopy (PEIS) method to effectively analyse and differentiate the two interfaces and reactivities under investigation.

Intrinsic decomposition of Li 6 PS 5 Cl

Influence on first cycle performances

To investigate the intrinsic reactivity of Li6PS5Cl, we assessed its stability window and scrutinized the impact of a carbon additive on the initial cycle performance. In order to accomplish this, we employed a commercial Argyrodite solid electrolyte provided by NEI company, which exhibits a wide particle size range spanning from 100 nm to 30 µm, as depicted in occurs through oxidation within the potential window of 2.1-3.9 V vs. Li-In/In. To further investigate this phenomenon, X-ray diffraction (XRD) analysis was conducted on the VGCF/SE composite after oxidation at 55°C for a duration of one month, revealing the disappearance of Li6PS5Cl peaks and the formation of an amorphous component (for more details, refer to the Appendix 2.2). These findings align with the decomposition mechanism reported in the literature for Li6PS5Cl, where the oxidation of Argyrodite within the same potential range results in the generation of highly insulating products such as S and P2S5 [START_REF] Auvergniot | Redox Activity of Argyrodite Li6PS5Cl Electrolyte in All-Solid-State Li-Ion Battery: An XPS Study[END_REF][START_REF] Tan | Elucidating Reversible Electrochemical Redox of Li 6 PS 5 Cl Solid Electrolyte[END_REF] .

In a typical composite comprising NMC 622, the intrinsic decomposition of the solid electrolyte described above, is reflected by a characteristic first cycle early delithiation slope.

This behaviour is demonstrated by examining the first cycle performance of our Li6PS5Cl/NMC 622 composites (70/30 wt. %), with the inclusion of 1 or 5 wt. % of VGCF or without any additive (refer to Figure 2.1.c) [START_REF] Zuo | A Mechanistic Investigation of the Li10GeP2S12|LiNi1-x-yCoxMnyO2 Interface Stability in All-Solid-State Lithium Batteries[END_REF][START_REF] Walther | Influence of Carbon Additives on the Decomposition Pathways in Cathodes of Lithium Thiophosphate-Based All-Solid-State Batteries[END_REF] . As anticipated, an increasing additional slope is observed within the voltage range of 2.1 to 3.1 V vs. Li-In/In, with higher amount of VGCF.

At this point, it is imperative to explore the influence of the newly formed degradation products on the performances of the first cycle. To address this, we proceed to vary the amount of carbon additive in composites, between 0 to 10 wt. % and study its impact on the initial cycle. As demonstrated in In order to gain a deeper understanding of the system's behaviour, we employed the concept of first cycle irreversible capacity as a valuable tool. The irreversible capacity observed during the initial cycle is commonly attributed to either slow kinetics or irreversible phenomena 115 . To distinguish between these possibilities, the cells potential was kept at the end of discharge, and the amount of capacity recovered was used as an indicator of the kinetic limitations within the cells. Concurrently, the resulting final irreversible capacity provided valuable insights into the occurrence of side reactions, as depicted in Figure 2.3.a. We will further refer to the "apparent irreversible capacity" and the "true irreversible capacity" for the irreversible capacity prior and after potential imposition, respectively. It is important to note that the extent of kinetic limitations and reactivity is highly influenced by the nature of the solid electrolyte, as demonstrated in the Appendix 2.4 through a comparison of the first cycle of solvent and/or solid-state synthesized Li6PS5Cl, Li3InCl6, and Li3PS4. Furthermore, it is noteworthy that this methodology has been successfully employed in the investigation of lithium-ion cells 115,116 . However, when comparing liquid and solid-state cells, we observed a lower true irreversible capacity in the liquid system, as exemplified by our NMC 622 in Figure 2.3.b, suggesting an enhanced reactivity in the solid-state.

PEIS

We implemented this methodology on coated NMC 622 system with various carbon additives nature and content, offering different morphologies and specific surfaces: C 65 (62 m²/g), VGCF (24 m²/g) and Super P (80 m²/g). We discharged Carbon-free (C-free) and

Carbon-containing (C-containing) cells on their first cycle to 2.1 V vs. Li-In/In in a galvanostatic mode at C/10, prior switching to a constant potential mode for 30 hours when the discharge cut-off potential was reached. As seen in Figure 2.3, we could retrieve most of this irreversible capacity for the C-containing cells under this procedure, indicating that the apparent irreversible capacity is mainly arising from kinetic limitations. Furthermore, both apparent and true irreversible capacities are increasing with the electronic surface, pointing at both an enhanced side reaction with Li6PS5Cl and worse kinetic of the first cycle linked to the new highly insulative phases formed. Comparing the three carbon additives, VGCF shows a lower reactivity with 10 mA.h/g of true irreversibility, while C65 and Super P-containing systems respectively present 13 and 17 mA.h/g for 5 wt. % additive addition. This small difference is most likely due to the VGCF lower BET (Brunauer-Emmet-Teller) specific surface (24 m²/g) [START_REF] Walther | Influence of Carbon Additives on the Decomposition Pathways in Cathodes of Lithium Thiophosphate-Based All-Solid-State Batteries[END_REF]117 . Moreover, the optimized nature of the long fibers has been proven to allow efficient electronic percolation in solid-state systems 118 . If the reader is interested, we also made high temperature experiment on a variety of systems using distinct AM, as depicted in Appendix 2.5, and we found an enhanced final irreversibility alongside a lower apparent irreversibility when using 55 °C, matching a reactivity governed by kinetics. Surprisingly, a discrepancy in the apparent irreversibility trend is found for 10 wt. % VGCF in Figure 2.3.c, exhibiting poor kinetics with an apparent capacity loss of 37-45 mA.h/g. Such behaviour questions the impact of morphology on the kinetics, additionally to the side reactions effect.

Indeed, a legitimate question arise from whether the slow kinetics observed are only linked to the newly formed interfaces, or also to a carbon additives physical presence effect. % VGCF, 5 wt.% C65 and 5 wt.% Super P. The theoretical exchange surface is computed by taking the mass of carbon additive and their expected BET surfaces given by the manufacturers (VGCF: 24 m²/g, Super P: 80 m²/g and C65: 62 m²/g). We also add the surface of our NMC, which is computed by making the hypothesis of spherical particles.

Sluggish kinetics of carbon-containing systems

As evidenced through last section, a question remains regarding the kinetic limitations of the 10 wt. % VGCF composite. To explore this surprising result, we investigated the impact of carbon additive presence on the ionic and electronic resistance of ASSB composites, prior any cycling. Firstly, the composites ionic conductivity is assessed via Potentiostatic Electrochemical Impedance Spectroscopy in a three-electrode cell setup developed in house 119 , so as to decouple WE impedance evolution from the CE one. Detailed explanation on three-electrode assembly can be find in Appendix 2.3. A scheme of the cell is shown in For interpreting such impedance spectra, an equivalent model is necessary as solidstate composite systems present complex spectra. Transmission Line Models (TLM) is used as a model of choice in the literature since it present a direct transcription of the physics behind such mixed ionic/electronic systems 120-125 . The system is assumed to be 1D, and we consider uniform values of the parameters, to highly simplify the fitting process by using the mathematical equations proposed by Ioroi et al. 126 . This model features two railways representing ionic and electronic conduction paths along the mixed conductive system.

Between them lies the interfacial phenomena arising from AM/SE. Ohms, respectively, using a loading of 16 mg/cm² and an electrode thickness of approximately 70 µm. Switching to carbon containing cells, the electronic resistance is dropping below 1 Ohm by adding only 2 wt. % of VGCF. Thus, electronic resistance is supposed negligible in the TLM model when more than 2 wt. % carbon is added, in accordance with the DC conductivity measurements. More surprising results are seen for the ionic conductivity. Indeed, increasing carbon amount goes along with an unexpected rise of the ionic resistance, counterbalancing the positive effect of the electronic conductivity decrease. Hence, the carbon addition is not only promoting side reactions, but is also detrimental for the ionic motion when a large amount is used. This is evidenced by the limited kinetics observed in 5 wt. % VGCF containing systems (see the C-rate tests presented in Appendix 2.6). A plausible hypothesis lies in the increasing tortuosity of Li + pathway when carbon is added. Another one is the chemical reactivity of Li6PS5Cl with surface species existing on carbon additive surfaces. Overall, as carbon is essential for improving the overall electronic conductivity of ASSB composites, our results suggest that keeping the amount of additive between 1 to 2 wt. % is sufficient to provide a proper electronic motion while maintaining a low increase of ionic resistivity.

Effect on stability upon cycling

Turning to carbon additive impact on cell lifetime, C-free composites are compared with 5 wt. % VGCF containing ones. Here in, a large amount of carbon additive is chosen to neglect the active material electronically conductive surface, which therefore becomes 10 times lower than the VGCF one in C-containing cells, in the aim of solely focusing on VGCF/SE interfaces.

We decided to monitor the resistance evolution of C-free and C-containing cells, following the procedure described in Using NMC 622 as AM, Figure 2.7.b-d shows that the higher the imposition potential, the higher the loss of capacity and the polarization of the following cycle. This degradation is the worst for 3.9 V vs. Li-In/In imposition with more than 12 mA.h/g lost on the 4 th discharge.

In contrast, we found that a 3 V vs. Li-In/In imposition does not affect the cycling, while 3.4 V and 3.6 V vs. Li-In/In demonstrate a negligible impact with less than 5 mA.h/g losses. This is not new, as it has been proven that going above 3.6 V vs. Li-In/In induces high surficial reactivity in liquid cells, particularly in the case of Ni-rich NMC. This phenomenon arises from the oxygen release, the H2-H3 phase transition, and the highly reactive Ni 4+ appearing during charged state of NMC [START_REF] Kong | Insights into the Structural Evolution and Li/O Loss in High-Ni Layered Oxide Cathodes[END_REF]129 . Moreover, using similar methodology, we show the dependency of this effect with the nature of the layered oxide used in our composite (Figure 2.8.a). The worst behavior of NMC 811 (6 % capacity loss) as opposed to 4 and 2.5 % for NMC 111 and LiNi0.9Co0.1O2 (NC9010) respectively, suggests different degradation rates. This effect is suspected to arise from the nickel content, as the stability of cathodes in liquid electrolytes was shown to decrease with increased Ni content [START_REF] Kong | Insights into the Structural Evolution and Li/O Loss in High-Ni Layered Oxide Cathodes[END_REF][START_REF] Ruan | Structural Evolution and Capacity Degradation Mechanism of LiNi0.6Mn0.2Co0.2O2 Cathode Materials[END_REF]129 . However, as Ni content also affects the electronic conductivity of NMC, it raises a legitimate question regarding the impact that it could have on the observed degradation mechanism, especially considering the decomposition of our SE explained in the previous section 130 . Thus, we assessed the electronic conductivity of both the active materials mixed with 1 wt. % PTFE or with 30 wt. % Li6PS5Cl, which are displayed in

Effect on stability upon cycling

To tackle the impact of the AM/SE detrimental phenomena upon cycling, the cells were cycled using a cut-off potential of 3.9 V vs. Li-In/In to further enhance instabilities issues previously spotted.

Decorrelating the two interfacial reactivities using PEIS

To decouple the distinct effects of VGCF as function of cut-off potential, and of both interfacial reactivities (AM/SE and VGCF/SE), we model resistance evolutions through impedance spectroscopy measurements. To build the equivalent model used, we complexify the transmission line circuit previously employed in section 2.2.2, to account for the phenomena occurring during cycling, such as charge transfer. The resulting detailed equivalent model used in this study is depicted in Figure 2.10.e. Different circuit elements are added between the two rails to represent the charge transfer (R CT, Q CT) and the solid-state diffusion of the AM presenting a Warburg-like impedance (W). Additionally, a R//Q (R Int // Q Int) is implemented to account for a low frequency additional feature. However, considering the uncertainty of its origin, we did not further investigate its values. We would like to emphasize that its addition does not alter the estimated values of interest, (R Ion, R Elec and R CT) as those phenomena are decorrelate from higher frequencies (see Appendix 2.7 for further explanation). As a necessary first step, we will transform this model into the one depicted in Figure 2.10.f, due to the mirror effect of R Ion, composite and R Elec, composite, making them impossible to decorrelate. We will therefore refer to a global conduction resistance (R Global, conduction = f (R Ion, composite, R Elec, composite)) merging ionic and electronic evolution in a non-linear equation. For a complete explanation on our impedance modelling and on the impact of each parameter, see Appendix 2.7. For this, two methodologies were applied to extract R Ion, composite and R Elec, composite.

Equivalent model development

Experimental data results

Firstly, for cycle 0 (pristine pellet) R Elec, composite, cycle 0, is obtained from pellets DC conductivity measurements, collected using 80 mg of composites and placing the pellets in between two carbon current collectors in two-electrode cells. Turning to R Ion, composite, cycle 0, its values were obtained from PEIS measurements in three-electrode cells, using exactly the same methodology as in Figure 2.4. This procedure is schematically explained in Figure 2.12.a. For cycle 10, it is not possible to collect pellet from the cycled cell without breaking it. Thus, we developed another strategy to decorrelate R Elec, composite, cycle 10 and R Ion, composite, cycle 10. The cells undergoes different temperatures and spectra are collected between 0 °C and 30 °C. We assume that R Elec, composite is not dependent on temperature, and that R Ion, composite follows an Arrhenius law. It enables us to get a system of non-linear equations in function of temperature to optimize, as explained in Figure 2.12.b (for complete explanation on the method, see

Appendix 2.8).

The resulting evolution of both parameters before cycling and after 10 cycles is depicted in Figure 2.12.c. This method reveals R Ion, composite to be the major contributor to the C-containing system's impedance rise in low potential window, growing from 200 to 2000 Ohm. In C-free system, R Elec, composite slightly rises with a minor increase of about 20-40 Ohm, while it remains below 1 Ohm in C-containing cells, which is not a surprise considering the large amount of VGCF used (5 wt. %). In the high potential window, minor changes of R Ion, composite and R Elec, composite are obtained. This is expected, considering that R CT is the only parameter evolving in Figure 2.11.b. Those results reassert the distinct behaviours previously observed regarding the mixed effect of carbon additive and cut-off potential on Li6PS5Cl-based systems. On one hand the increasing charge transfer resistance when cycling up to 3.9 V, suggests a detrimental phenomenon originating from the active material interfaces with Li6PS5Cl or from the material degradation itself. On the other way, the increase of R Ion, composite in cells comprising VGCF and cycled up to 3.6 V, results from the VGCF/SE reactivity arising from SE decomposition.

Conclusion

Altogether, the results presented in this section demonstrate a degradation mechanism of Li6PS5Cl in ASSB that is ruled by two interfaces, namely Li6PS5Cl/VGCF and Li6PS5Cl/AM (see Figure 2.13). Note that the first one, the intrinsic decomposition of Li6PS5Cl when put in contact with an electronic conductor, should also occurs versus the active material. However, we showed that such effect is minimized when using the C-free cells, as mirrored by their high stability and small resistance increase (see Figure 2.5). When cycled below 3.6 V vs. Li-In/In, the main source of capacity fading originates from the SE/VGCF interfaces and yield to 10 % fading in 25 cycles when 5 wt. % VGCF is added. Moreover, the addition of VGCF is proven detrimental for the ionic conduction of the cathode prior cycling, when more than 2 wt. % is added. Such statements push the belief that small amount of carbon additive should be used for both high kinetics and stability. Going to higher potential window (3.9 V vs. Li-In/In cut-off potential), distinct behaviour are observed as the main degradation source shifts toward the SE/AM interfaces. This high potential reactivity yields to a building up strongly resistive interface around AM particles, resulting in 20-25 % capacity loss in 25 cycles. Surprisingly, no additional effect of VGCF is found is this potential window and this phenomenon will be further investigated in Chapter IV. Toward the goal of higher stability and capacity, the next chapters will investigate how we can alleviate those instabilities, by focusing on two possibilities. Firstly, we will investigate the kinetics of C-free cells by focusing on the electronic conductivity evolution of composites, linked with particle size effect, percolation and phase transitions. Secondly, Chapter IV will investigate the impact on distinct parameters to try to negate both AM/SE high potential and VGCF/SE reactivity.

Introduction

As discussed in the preceding chapters, important polarization and transport resistances undermine sulphide-based ASSB performances, due to the distinct reactivities occurring at composites interfaces and unoptimized ionic and electronic percolation.

Especially, the addition of carbon additives leads to a fast capacity decay when cells are cycled in the conventional potential window.

In order to investigate these issues and determine kinetic parameters, we propose a unique new device, enabling to track in situ the global variation of a composite electronic conductivity upon cycling. A composite's electronic evolution upon charge and discharge can reflect multi-scale phenomena that enlist inter-and intra-particle contacts, percolation and electronic band structure changes of the active material, among others [131][132][133][134] . Thus, we will explore these phenomena in a variety of composites based on lithium transition metal layered oxide materials (LiCoO2, LiNiO2, NMC622, NMC811, NMC111 and LiNi0.9Co0.1O2 (NC9010)) and Li4Ti5O12 (LTO) active materials. Especially, we investigate the effect of phase transitions, chemo mechanical Li driven process and particle size on the global electronic conductivity first cycle evolutions.

Validation and presentation of the device

Presentation of the cell

In our attempt to monitor electronic conductivity of ASSB composites upon cycling, we developed a new setup. This setup is based on the three-electrode cell configuration previously employed for performing PEIS measurements in last chapter. It relies on the use of two potentiostat channels connected to this cell, with one used for galvanostatic cycling and the other one for direct current (DC) conductivity measurements. To enable the measurement of direct current conductivity, a 60 µm thick aluminium mesh is positioned at the reference connection level, as illustrated in Figure 3.1. Regarding the assembly procedure, we utilize a trilayer configuration, wherein the cathode composite is initially pelletized to fit between the aluminium mesh and the upper piston. Additionally, to prevent any electrical short circuit, a large amount of cathode material (60-80 mg/cm²) is employed, the cell integrates 30 to 40 mg of Li6PS5Cl, and a large quantity (> 120 mg/cm²) of counter electrode composite is added as well. Lastly, the entire stack undergoes a compression force of 1 t/cm² during the cycling process, using similar procedure as for the two-electrode assembly.

Turning to the cycling protocol, it encompasses two channels. To ensure a good synchronization between those two, the trigger IN and trigger OUT of the channels are connected to one another. Moreover, to avoid any leakage current, the CE to GROUND connection is used for both channels. One channel enforces galvanostatic cycling, periodically halted at intervals of t1, followed by a relaxation period of t2 to ensure that the system reaches the equilibrium state. The second channel is utilized for measuring the DC electronic conductivity when the system is at equilibrium. A minor additional polarization for a duration of t3 is applied to attain a steady-state condition (refer to Figure 3.2) at which R Elec is evaluated. Note that the value of the additional polarization allows a good signal-to-noise ratio and maintains system stability throughout the experimental procedure. Lastly, the data extracted from each channels are combined to form the resulting σ Elec vs. lithiation state of the AM curves, by doing a thickness estimation to transform R Elec in σ Elec. This thickness is estimated via the average value of thickness per mass of composite, as measured on at least two pellet measurements per system. Additionally, 60 µm is subtracted to account for the aluminum mesh, as it is embedded in the cathode composite upon assembly. hour and 30 minutes on channel 1, followed by a relaxation period of 1 hour. Subsequently, channel 2 was activated, imposing a potential step of 10 mV between the aluminum mesh and the upper piston for a duration of 7 minutes. This experimental procedure was repeated on three cells, demonstrating exceptional repeatability of the electronic conductivity evolution.

During the early stages of delithiation, when approximately 0.04 Li is removed from LiCoO2, a remarkable increase in electronic conductivity is observed. This behaviour is attributed to a first-order Mott transition, which has been experimentally and theoretically validated 135,136 . The significant five orders of magnitude increase in conductivity is effectively captured by our in situ tool (Figure 3.3.e) and exhibits a similar magnitude as ex situ measurements conducted on delithiated LCO obtained from liquid cells. More precisely, we delithiated LCO powders in Swagelok cells, using the experimental procedure described in

Chemo mechanical effect in TLMO-based composites

Active material vs. composites electronic conductivity

Following the successful demonstration of our proof of concept, our research subsequently shifted towards investigating more interesting active materials for solid-state research, specifically NMC 111, NMC 622, NMC 811, LiNi0.9Co0.1O2, and LiNiO2 (see Appendix

for further details).

To initiate this exploration, we interrogate the issue of electronic percolation in ASSB composites. In order to accomplish this, we use Li6PS5Cl powder from NEI, possessing electronic conductivity in the range of approximately 10 -9 to 10 -10 S/cm 140,141 and exhibiting a significant dispersion in particle size (ranging from 1 to 50 µm, as illustrated in

First cycle evolution of five TLMO composites

To go even deeper, we monitored the first cycle electronic conductivity evolution in our five TLMO-based composites, by using our in situ setup. The electronic conductivity profiles are presented in For the majority of TLMO compounds, namely NMC 622, NMC 811, NC 9010, and LiNiO2, the electronic conductivity exhibits slight variations in the early stage of lithium removal, prior to displaying a distinct bell-shaped evolution upon further delithiation.

Interestingly, this behaviour is absent in the NMC 111 composite, which instead demonstrates a significant increase in conductivity when x < 0.5, surpassing the electronic conductivity values of NMC 622 and NMC 811 at the end of the charge cycle (Figure 3.5.f). The observed rise in electronic conductivity (σ Elec) of NMC particles can be attributed to the presence of Ni 3+ /Ni 4+ valence states, as previously reported through experimental and computational studies 130,132,142,143 . However, the subsequent decrease in conductivity at higher levels of lithium removal for high Ni-content TLMO composites represents a novel observation in our knowledge. Especially, it can even lead to values of σ Elec lower than the initial measurement at Li content x = 1, such as for NMC 811.

To understand this intriguing phenomenon, we investigated the influence of AM morphology and phase transitions. Firstly, the polylithic nature of NMC 111 cannot explain this effect, as this decline in conductivity is also occurring in polylithic NMC 622, using particles of similar sizes (see Appendix 3.4). Secondly, to explore the potential influence of structural phase transitions, the present study employed Galvanostatic Intermittent Titration Technique measurements on NMC 111 and NMC 811 samples (refer to Appendix 3.5). The GITT profile of NMC 811 at a high state of charge exhibited a notable reduction in the diffusion coefficient, indicating the occurrence of the H2-H3 transformation as reported in existing literature [144][145][146] .

This transformation is known to be accompanied by a decrease in the c lattice parameter 137,147,148 . In contrast, the observed phenomenon was less pronounced in NMC 111, suggesting that the decline in electronic evolution for phases with high nickel content at high state of charge may be attributed to this phenomenon. Indeed, the chemo-mechanical Li driven effect, which origin lies in the c lattice parameter collapse, is largely increased with higher Ni content. To test the hypothesis of a chemo-mechanical origin, the electronic conductivity of NMC 811 was examined using in situ and ex situ measurements in liquid cells (see Figure 3.6). Detailed explanation on the procedures are described in Appendix 3.6.

Neither approach indicated a decrease in electronic conductivity during the delithiation process. On the contrary, they demonstrate a continuous increase of electronic conductivity, thereby confirming that the loss of electronic contact is associated with a mechanical effect occurring in solid-state, which competes with the influence of Ni 3+ /Ni 4+ valence states. Thus, this behaviour arises from substantial inter particle contact loss when NMC undergoes volumetric shrinkage. This is significantly amplified in the solid-state configuration due to the high porosity of the composites and the limited ductility of the Li6PS5Cl solid electrolyte. This behaviour persists and is reversible throughout subsequent cycles, as demonstrated in cells based on NC 9010 and NMC 622 (see Figure 3.7). Overall, considering the critical importance of mechanical integrity in all-solid-state batteries, the ability to qualitatively monitor contact loss provides a valuable advantage for future optimization.

Particle size impact on Li 4 Ti 5 O 12 composites electronic conductivity

To delve further into the morphology and particle size impact on first cycle electronic conductivity, solid-state and solution-combustion lithium titanium oxide powders were synthesized using established experimental procedures available in the literature. This material was chosen for its original electronic property of a metal-insulator transition in early lithiation, and for its zero-strain ability, enabling us to suppress the chemo-mechanical effect studied in previous section.

AM synthesis and characterization

For the in-house LTO solid-state synthesis, TiO2 (Sigma, purity ≥ 99%) and Li2CO3 (5% excess, refer to Figure 3.8 for optimization of precursor amount) (Sigma, purity ≥ 99%) were combined in stoichiometric quantities. The mixture was manually ground for 10 minutes using an agate mortar, followed by pelletization. Subsequently, the pellets were subjected to a heat treatment at a temperature of 900 °C for a duration of 2 hours with a heating rate of 5°C/min.

In some instances, an additional round of manual grinding and heat treatment under similar conditions was performed to minimize the presence of any impurities. The LTO solution-combustion homemade samples were synthesized using a previously reported self-combustion method 149 . Anhydrous lithium nitrate (10 % excess, refer to Figure 3.9 for precursor optimization) (Alfa, purity 99 %) and titanyl nitrate (TiO(NO3)2) were employed as oxidizers (O), while glycine (NH2CH2COOH) served as the fuel (F) in a 1:1 O/F ratio.

For each synthesis, the TiO(NO3)2 oxidizer was freshly prepared by combining 3 ml of titanium isopropoxide (C12H28O4Ti, Alfa, purity 97 %+) with 9 ml of isopropyl alcohol under ice-cold conditions. Subsequently, 3 ml of a mixture containing 1: To investigate this phenomenon further, we utilized the four LTO samples synthesized through solution-combustion, which exhibit distinct primary and secondary particle sizes as discussed in the previous section. Solid-state batteries were built using these different LTO samples, mixed with 50 wt. % Li6PS5Cl as the positive electrode composite. These Collectively, these findings demonstrate the influence of particle size on the electronic conductivity of LTO-based composites during lithium incorporation, with the larger agglomerates governing the composite's electronic conductivity at high lithium contents. This observation is somewhat counterintuitive, considering that dispersed mixtures generally enable higher active material utilization and faster kinetics due to improved percolation and enhanced surface interactions between ionic and electronic conductors. Surprisingly, composites with agglomerate sizes nearly ten times larger than most NEI Li6PS5Cl particles exhibit the highest electronic conductivity, suggesting a different trend in the electrodeelectrolyte particle size ratio's effect on performances in LTO composites. These observations can be rationalized by considering the unique nature of the lithium reaction pathway in LTO and the mixing process of solid-state composites.

1 v/v
To begin with, it is important to note that lithium intercalation/deintercalation in LTO occurs via a two-phase reaction between Li4Ti5O12 (an electronic insulator) and Li7Ti5O12

(an electronic conductor) during the 1.5 V plateau, as the Ti d-states become partially filled (e.g., Ti 3+ ) during lithiation 150 . However, the lithiation pathway in LTO has been a subject of controversy, resulting in various models (core-shell, radial, etc.) for phase transformation kinetics, with their validity dependent on the size, morphology, and concentration of Ti 3+ 151- 154 . In situ electron holography studies on liquid cells have described LTO nanoparticles as following a core-shell process within the narrow solid solution range at both ends of the plateau, which aligns with the rapid increase in electronic conductivity observed during early delithiation (as shown in Figures 3.14 and 3.15) 155 . Furthermore, using conductive Atomic Force Microscopy, it has been demonstrated that the evolution of electronic conduction pathways in thin film LTO follows a biphasic reaction that does not propagate laterally through grain boundaries or uniformly form across the material. Instead, it proceeds via percolation channels within individual grains 156 . This one-by-one filling of multiple particles emphasizes the significance of good inter-particle contact for the electronic propagation of the LTO phase, which is consistent with the larger agglomerates exhibiting higher efficiency in electronic motion within our systems. In other words, as conductive Li-rich and insulative Li-poor phases coexist, a Li-poor grain is filled when it is in close proximity to a Li-rich phase, following the path of lower resistance. Consequently, the lower tortuosity of the electron pathway in larger agglomerates results in higher electronic conductivity at the end of lithiation. This behaviour can be schematically represented as a function of particle size, as illustrated in Figure 3.16.

Grain per grain propagation

Lower tortuosity Higher tortuosity

Conclusion

We have presented a novel homemade electrochemical cell that allows for in situ measurement of electronic conductivity of all-solid-state battery composites. This unique diagnostic tool is achieved by connecting two potentiostat channels to the same cell, providing convenient monitoring of electronic conductivity. The device combines two programs: one for galvanostatic cycling and the other for DC conductivity measurement. It can be easily assembled, maintains airtightness, and the presence of an aluminum mesh does not interfere with the cycling process. Furthermore, its simplicity and accuracy make it suitable for daily usage, providing valuable insights into the apparent electronic transport of composites.

To demonstrate its functionality, we successfully investigated the metal-insulator transition in LiCoO2, which was further validated by ex situ liquid measurements. Additionally, we presented direct evidence of how this new cell can differentiate between the electronic properties of the active material and the overall electronic transport in three-dimensional solid-state mixtures. In essence, this cell paves the way for studying electronic transport in a wide range of cathode-solid electrolyte composites. Furthermore, we have also chosen to demonstrate the benefit of our new cell to track changes in electronic transport associated with Li-driven chemo-mechanical changes of TLMO-based composites. Specifically, we observed a significant reduction at high SOC in electronic conductivity of composites containing high nickel content, as a result of pronounced volume changes driven by lithium insertion, notably associated with the well-known H2-H3 transition. Additionally, through an examination of the Li insertion process in zero-strain Li4Ti5O12, we were able to monitor the insulator-metal transition during the initial lithiation of this compound. By using distinct LTO primary/secondary particle sizes, this investigation allowed us to gain insights into how electronic percolation is influenced by particle size and morphology, thereby providing valuable information regarding the intricate lithium reacting pathway.

Overall, numerous parameters influence the electronic conductivity of solid-state composites, including particle size and phase transitions. Especially, the usage of carbon-free high Ni-content NMC based composites is hindered by the high SOC decrease of electronic conductivity, pushing the belief that a minim amount of carbon additive could be highly beneficial for C rate performances in such systems. The following chapter will explore this possibility, by taking the advantage of carbon additive while simultaneously mitigating the long-term capacity losses associated with carbon addition.

Introduction

As we have seen in previous chapters, significant polarization and transport resistances compromise the performance of sulphide-based ASSB performances, due to the distinct reactivities occurring at composites interfaces. Of particular note is the marked drop in capacity observed with the introduction of carbon-based additives during standard cycling.

However, considering the problems of Li-driven chemo mechanical loss of contact between particles occurring in high Ni-content NMC, and the lower AM utilization in C-free ASSB composites, the addition of carbon remains necessary provided that carbon-induced reactivities are managed. Thus, this chapter focuses on distinct ways to alleviate the long-term stability issues pertaining to C-containing composite, using 1.5 wt. % VGCF and a NMC 811 (MSE supplies) as standard AM. We also chose to cycle at up to 3.9 V vs. Li-In/In, as we found no additional detrimental effects of VGCF in this potential window (see Chapter II) while experiencing a greater capacity gain as compared to a conventional potential window cycling. Thus, this chapter aims at briefly introducing some optimization work on the Li6PS5Cl/NMC811 + 1.5 wt. % VGCF composite, with a special interest on stability. For this, we firstly assessed the repeatability of our cycling on 9 cells originating from two NMC 622-based composites, prior switching to NMC 811. Then, we assessed the cycling behaviour of the NMC 811 composites as function of C-rate, cut-off potential and presence of carbon additive. Afterward, we explored the impact of distinct parameters on stability, namely the assembly pressure, the loading, the impact of SE and AM carbonate coating and formation cycles protocols. Finally, the different approaches are combined and the best composite/cycling protocol obtained is presented.

Cycling behaviour of high Ni-content NMC composites

Repeatability of cycling and cathode composite processing

Prior conducting any study on improving the stability of ASSB composites, we initiated an assessment of the repeatability of a specific composite consisting of NMC 622 (Umicore) and Li6PS5Cl in a weight ratio of 70/30, with the addition of 1.5 wt. % VGCF. Note that a new NMC 622 from Umicore was used as compared to the one from Chapter II, due to powder shortage. It displays the same particle size and morphology but a better stability, as demonstrated below. the capacity decay after 100 cycles was found to be 15.87 % and 15.71 % for composite 1 and 2, associated with a standard deviation of 0.53% and 0.44 %, respectively. Regardless of the slightly more dispersed dataset of composite n°1, our cell preparation process, including the manual grinding step, exhibited good reproducibility in terms of stability, as mirrored by the low standard deviation values. Notably, we can now discern what constitutes a statistically significant variance in the decrease capacity, which in this case is equivalent to around 1 % of the mean value of 15.78 %. Ideally, a larger dataset will enable us to achieve a statistically higher level of confidence in our analyses. However, in view of the practical problems we are facing, we will accept these values and only take into account deviation of more than 1.6 % in the following figures. Note, that we will consider this as the discrepancy associated with the repeatability of the assembly and processing, independently of the active material. At this point, an investigation of the capacity retention characteristics of the NMC 811 composite becomes imperative, to ascertain whether this new composite exhibits comparable trends to those elucidated in Chapter II for NMC 622. Thus, we initiated a comprehensive study on the capacity retention of both NMC 622 and NMC 811 composites, considering varying C-rates (C/2 or C/10), cut-off potentials (3.6 or 3.9 V vs. Li-In/In), and the presence of VGCF (1.5 wt. % VGCF or no VGCF). Each experimental condition was replicated in two cells, leading to the cycling of 32 cells, and the resulting average capacity retention data are summarized in Table 1.

Several noteworthy conclusions can be drawn from this study. Firstly, it is evident that lower C-rates correspond to higher degradation. For instance, NMC 811 exhibits a decay of -12.5 % and -38 % in C-containing cells cycled at C/2 and C/10, respectively, and up to 3.9 V vs. another active material, but also reveal a remarkably similar behaviour when comparing NMC 622 with NMC 811. Furthermore, no additional capacity decay is observed for NMC 811, implying that we can harness the additional capacity offered by NMC 811 while maintaining degradation at levels comparable to our previously utilized reference composite.

Li-In/In (refer to

Using this new composite, we launched a preliminary investigation into different approaches to mitigate capacity degradation observed when cycling up to 3.9 V vs. Li-In/In in C-containing cells. composites, including 1.5 wt. % VGCF or no VGCF. Two cut-off potentials are used: 3.6 and 3.9 V vs. Li-In/In and two C-rate: C/2 and C/10. Each value represent the average of two cells.

V vs. Li-In

Impact of assembly pressure and loading

Assembly pressure

First of all, assembly pressure has been proposed as an essential parameter for determining the performances of All-Solid-State Batteries, as it directly affects contact phenomena and the chemo-mechanical behaviour of composites [START_REF] Zhang | A Review of the Effect of External Pressure on All-Solid-State Batteries[END_REF][START_REF] Doux | Stack Pressure Considerations for Room-Temperature All-Solid-State Lithium Metal Batteries[END_REF]157,158 . In this study, we investigate the influence of both low and high assembly pressures and analyse their effects on the initial cycle performance and capacity retention. To achieve this, we changed the assembly pressure to 2, 3, 4, 5 and 6 t/cm² (1 t/cm² = 98 MPa), which is applied for 10 minutes on the whole battery stack, and the resulting first cycle capacity is depicted in t/cm² cells present the worse performances, likely linked with large porosity resulting from the low pressure applied. Interestingly, the 3 t/cm² cells present the highest stability, with 9.5 % capacity decay after 120 cycles, while at higher pressures, there is no difference between the capacity retention of the systems assembled at 4 to 6 t/cm². Such results contradict our expectations that higher pressures should lead to enhanced stability and capacity. One possible hypothesis is the impact of pressure on the properties of the SE and AM used. This topic has begun to be explored for SE, but its counterpart for AM is still lacking [159][160][161] .

Especially, an interesting investigation would be to assess the impact of pressure on the AM properties (diffusion, morphology…) and on its mechanical evolution upon cycling. Note that it has already been reported that high stack pressure provokes fractures and disintegration of AM particles 162 .

Cathode loading influence

A typical constraint in achieving high energy density in present-day ASSBs is the low active material loading of the composite, as indicated in previous studies [START_REF] Bielefeld | Microstructural Modeling of Composite Cathodes for All-Solid-State Batteries[END_REF][START_REF] Minnmann | Designing Cathodes and Cathode Active Materials for Solid-State Batteries[END_REF][START_REF] Shi | High Active Material Loading in All-Solid-State Battery Electrode via Particle Size Optimization[END_REF] . Enhancing the composite loading is frequently accompanied by challenges related to limited ionic and electronic percolations. Thus, we examined the influence of the loading on both the initial cycle capacity at C/10 and the capacity retention and results are depicted in mg/cm²) and the highest loading (24-28 mg/cm²) show slightly higher polarization than the rest of the loadings. The behaviour upon cycling is presented in Figure 4.9, revealing a consistent capacity decay trend across all loading conditions (with an average capacity loss of in 12.2 % 100 cycles, with a standard deviation of 1.7 %), thereby implying the negligible impact of this parameter on the overall stability of the system. Interestingly, the capacity at C/2 is similar between 12 and 26 mg/cm² (150-160 mA.h/g), meaning that we can increase our cathode loading up to 26 mg/cm² without penalty on the capacity. Overall, no significant changes are observed with the loading, in the range studied. However, we must bear in mind that this behaviour is significantly contingent on the specific surface of the active material used, as demonstrated in the Appendix 4.4. In this case, we used an NMC 622 from UMICORE, which had been stored outside an inert environment and whose capacity had initially deteriorated. The heat treatment under air was shown to largely improve both the initial capacity and stability in this specific case of a degraded NMC surface. As a result, the applicability of this methodology for active materials presenting a clean and fresh surface is questionable.

Exploring formation cycles

We next conducted a comprehensive investigation into the effects of various formation cycle conditions, encompassing temperature, potential window, and the number of formation cycles. As reference composite, we still employed NMC 811 (MSE) + Li6PS5Cl (70/30 wt. %) + 1.5 wt. % VGCF.

Conclusion

This chapter is dedicated to the investigation of distinct approaches for mitigating the capacity degradation observed in Li6PS5Cl-NMC composites during cycling up to 3.9 V vs. Li-In/In. Because of the large number of parameters considered, we realized that we risked losing the reader. We therefore decided to summarize all the parameters studied with their advantages and disadvantages, in a consumer-style report table in Table 4.2.

General conclusions and perspectives

This thesis elucidated the fundamental parameters and criteria governing the performance of Li6PS5Cl/NMC cathode composites in the context of All-Solid-State Batteries.

We have conducted a comprehensive assessment of the pivotal role played by interfacial evolutions in solid electrolytes and have introduced novel tools and methodologies for their characterization.

Through this approach, we have conducted an in-depth exploration of the degradation phenomena occurring within NMC 622 + Li6PS5Cl composites, yielding fresh insights into the dynamics governing this system. By employing techniques such as global resistance monitoring, Potentiostatic Electrochemical Impedance Spectroscopy in three-electrode cells, and classical galvanostatic charge-discharge cycling, we have identified two critical parameters: the quantity of carbon additive and the cut-off potential.

To start with the carbon additive, our investigation revealed a dual negative impact of its presence on the composite's performance. Firstly, we demonstrated that the presence of more than 2 wt. % of VGCF adversely affects the ionic conductivity of the composite, a phenomenon substantiated by DC and PEIS measurements conducted prior cycling. Secondly, we established that the presence of VGCF exacerbates capacity losses during cycling in the voltage range of 2.1 to 3.6 V vs. Li-In/In. This effect arises from the high intrinsic reactivity of Li6PS5Cl, which undergoes oxidation throughout the entire potential window. This behaviour is closely associated with a decline in the ionic conductivity of the composites during cycling due to the formation of highly insulating by-products, such as elementary sulphur.

Adding the impact of potential on cycling stability, we have elucidated the distinct behaviours that depend on the chosen cut-off potential and the presence of VGCF. To this end we designed and successfully used a transmission line model for Potentiostatic Electrochemical Impedance Spectroscopy measurements during cycling in a three-electrode cell configuration. When subjected to cycling up to 3.6 V vs. Li-In/In, we have identified that the primary source of capacity degradation stems from the interfaces between the solid electrolyte and VGCF, resulting in a 10% reduction in capacity over 100 cycles when utilizing 5 wt. % VGCF and a C-rate of C/10. This deterioration significantly affects the ionic resistance of the composite. After a further increase in potential to 3.9 V against Li-In/In, we observed a shift in the primary source of degradation towards the interfaces between the active material and solid electrolyte, manifested by an increase in charge transfer resistance. This transition leads to the formation of a passivating, highly resistive interface surrounding the AM particles, resulting in a substantial capacity loss of 20-25% in 25 cycles for our specific protocol.

Intriguingly, we noted that similar decreases in capacity are observed, regardless of the presence or absence of VGCF in this potential window. This observation underlines the need for further research to better understand the underlying mechanisms.

It should be noted that these interactions between carbon and cut-off potential were subsequently reaffirmed by the evaluation of another analogous NMC 622 and a NMC 811 active material, as detailed in Chapter IV. This corroboration strengthens the belief that we have here a general mechanism with high Ni-content NMC and Li6PS5Cl here.

In light of the uncertainty surrounding the utility of VGCF incorporation, we directed our attention towards investigating the evolution of electronic conductivity in carbon-free composites. To address this inquiry, we developed an innovative apparatus designed to monitor the in-situ electronic conductivity dynamics of ASSB composites. Our initial efforts involved assessing the reliability and advantages of this experimental setup by examining composites based on LCO, known to exhibit a distinctive metal-insulator transition during the early delithiation stages. Following the successful validation of our method, we proceeded to investigate the influence of primary and secondary particle sizes in the context of zero strain anode material LTO, shedding light on how particle size and morphology impact electronic percolation-a question of paramount importance. Subsequently, we transitioned we focused on NMC-based composites and carried-out a comparison between LNO, NC 9010, NMC 111, 811, and 622. This examination allowed us to qualitatively assess the influence of Li-induced chemical-mechanical effects on the electronic conductivity of the composites, particularly when operating at high states of charge.

The accumulation of these new and original results strongly indicates that, if aspects reactivity are judiciously controlled, a minimal quantity of carbon additive could bring substantial benefits in such systems.

Advices for ASSB research

This section presents a list of pragmatic guidelines pertaining to the construction of ASSBs. It is important to note that the majority of these recommendations are founded upon my personal experimental insights and intuitive understanding cultivated throughout the course of my doctoral research, rather than being substantiated by empirical evidences.

Glovebox and degradation over time:

An essential preliminary measure in conducting research on ASSB entails the establishment of a dedicated glovebox exclusively designated for ASSB experimentation, and free of any solvents (with the exception of DMC (Dimethyl Carbonate) for washing purposes in our case), as this aligns with the inherent reactivity of sulphur-based SE. Note that, using such glovebox for our experiments we did not see an increase of stability upon cycling comparing inside and outside cycling of our cells, thus assessing the airtightness capability of our cycling cell.

In the context of composites storage, the Li-In composites displayed excellent resistance to degradation over time, sustaining their integrity for at least six months in the glovebox while maintaining favourable cycling performance in two-electrode cells.

Conversely, numerous cathode composites experienced significant capacity loss after 4 to 5 months. Thus, it is crucial to assess their time degradation in your own glovebox when initiating ASSB research. In our case, we prepared limited quantities of composites (approximately 300 mg per batch) for 1 to 2 months of use.

Composites hand grinding:

During the preparation of Li-In composites through hand grinding, it is imperative to diligently mix the components until a state of homogeneity is achieved. In my case, this process typically demands approximately 10 minutes. Failure to attain homogeneity may result in inadequate adhesion of the composite to the piston after pressing, attributed to incomplete mixing of certain segments of lithium or indium foils with Li6PS5Cl. To address this concern, we also recommend the addition of an aluminum foil layer onto the Li-In composite prior to the application of assembly pressure (note that the addition of Al foil does not affect potential nor capacity retention, despite minim alloying of Al and Li might occur). This precautionary measure ensures the subsequent removal of pistons, if necessary during the continuation of the assembly process, as well as a facilitate recovery of clean pistons for future utilization.

The investigation pertaining to cathode composites grinding procedure highlights the paramount importance of hand-grinding duration and applied force. Notably, an extended grinding period (for my case > 12 minutes) was found to render the composite brittle and resulting in diminished performances. Additionally, excessive force during grinding led to the fracturing of the active material. Consequently, we propose employing a delicate handgrinding approach, applying minimal force, for a duration ranging from 6 to 10 minutes.

Importantly, our experimental results demonstrate no disparity between the 6 and 10-minute intervals, though shorter durations were not explored.

Cell assembly:

Multiple practical considerations arise during cell assembly to optimize performance.

Firstly, caution must be exercised regarding the speed of piston insertion, as excessively rapid insertion can lead to composite powder deposition on the cell's side walls. Moreover, to ensure adequate coverage of the cell surface, the operator should rotate the piston in contact with the added powder before pressing.

During assembly pressure, meticulous alignment of the pistons is of paramount importance to maximize the contact surface between the cathode composites and the pistons.

Failure to achieve proper alignment has been shown to result in significant performance discrepancies of C-free cathode composites, with deviations of approximately 50 mA.h./g observed when alignment is not executed. Another potential remedy for this issue involves incorporating carbon paper on the cathode composite side following assembly pressure.

Lastly, when tightening the screws to apply pressure during cycling, a strategic approach should be used. Employing the "triangles" method with the six screws and incrementally increasing pressure steps (consecutively 0.5, 0.75, 1, 1.5, 2, and 2.3 N.m torque on screws in our case) enables a well-distributed pressure throughout the cell. Such meticulous attention to assembly details can contribute to optimizing cell performance and most importantly, to the repeatability.

AM surface & synthesis:

An exceedingly critical aspect in the investigation of interfaces and composite performances lies in acknowledging the significance of active material synthesis and state of the material's surface. Even when employing NMC 622 active materials with comparable particle sizes, pronounced discrepancies in behaviours can manifest. These disparities may originate from distinct synthesis conditions (leading to varying quantities of transition metals in lithium sites for example), or may be influenced by the storage environment of the material.

Consequently, it becomes impractical to generalize interfacial behaviour without considering the specificities of the AM synthesis and surface state. To lend greater credibility to your interpretations, conducting multiple experiments utilizing diverse AM or, ideally, incorporating AM synthesis studies into the research approach would be advised. This more rigorous approach will ensure a comprehensive understanding interface behaviour and composite performance, however being time-consuming and/or unpractical with most laboratory facilities.

Repeatability:

In the fast-paced and dynamic world of battery research, there is a prevailing temptation to conduct numerous experiments in haste to swiftly arrive at conclusions, often deferring repeatability measurements for later consideration. However, such a rush to conclusions without ensuring repeatability is highly dangerous, particularly in the context of ASSBs where the assembly process plays a critical role. The consequences of premature conclusions can be severe, as researchers may find themselves confronted with contradictions and dilemmas after months of study, necessitating the choice between disregarding certain data or embarking on an extensive redo of the entire study. Ethically, the intelligent course of action would be to embark on a comprehensive redo, but this entails a substantial loss of time and effort. Thus, the most important counsel I can offer is as follows:

-Firstly, prioritize understanding the repeatability of your reference system by launching 5 to 10 cells and meticulously assessing their consistency in terms of

2.7: Detailed explanations regarding impedance fitting and model:

Transmission line models are generally used when the impedance response of a system is distributed in multiple directions regarding the electronic current, such as in porous electrode. The base model consists of an upper and a bottom line, representing ionic and electronic currents. In this configuration, the system's response to the alternative excitation imposed while taking EIS is not uniform and both ionic and electronic contribution are dependent on x, y and z positions. Therefore, a discretized and position dependent model is theoretically needed to represent each specific active material/electrolyte/carbon additives distinct interfaces sites. However, as such approach makes it difficult or impossible to obtain an analytical solution, a convenient way to model data is to consider that the values of the cathode elementary components does not depend on position. This macroscopically uniform scenario enables to obtain the model transfer function of the system via an infinite number of infinitesimal component elements. The solution are solved analytically with a differential equation problem, and this approach has been widely used and has proved its efficiency to model complexed systems.

Modelisation:

Overall, plural phenomena needs to be represented in the equivalent circuit of the cathode composites under study in this article:

-The ionic and electronic transport phenomena, which are represented by the two railways of the TLM, already shown before.

-The charge transfer phenomena, which is represented by a CPE in parallel with a charge transfer resistance.

-The solid-state diffusion occurring inside our active material particles (here NMC 622).

Fitting experimental data solely with a TLM just having charge transfer in between the two railways does not able to fit well the low frequency part of the spectra, as seen in Figure S2.11.a. Indeed this one is attributed to the solid-state diffusion due to the 45° slope observed, characteristic of transport diffusion phenomena. However, implementing solid-state diffusion

MATLAB programming:

The optimization program created for the fitting procedure is based on a Broyden-Fletcher-Goldfarb-Shanno code to diminish our experimental-model distance. The choice of this method has been done after comparing three other optimization methods namely, quasi-Newton method (David-Fletcher-Powell, DFP), Conjugated gradient method and a trust-region method. Both DFP and BGFS methods were presenting smaller distances, with a little advantage for BGFS.

The experimental-model distance (d) is calculated via the following formula: Note that the prefactor takes into account the usual σ0 of the Arrhenius law, but also the cathode composite thickness. With the small amount of cathode composite (~ 16 mg/cm²), we can expect a variation of thickness with cell repetition, surface roughness, flatness of the current collector… We can assume a similar thickness in between cells. However, the different values of Rion and Relec cannot be compared in between cells if they are close to each other.

To evaluate the optimization efficiency, we plotted the evolution of Rion (T) in logarithm plot and the difference of Z (RGlobal) and Z (Relec, Rion)) compared with Z (Rglobal) in Bode plots.

The optimization shows remarkable results as the temperature evolution of the fitted Rion fits perfectly with an Arrhenius law (Figure S2.17 On cells without VGCF, the residues are in the same order as Z (Rglobal) at high frequencies, certainly linked with the additional feature observed at high frequencies in those systems. This does not bring concern on the optimization results validity, as the residues become more than 10 times lower than Z (Rglobal) starting from 3 kHz, below which frequency the spectra features of interest are largely observed.

3.9: Accusizer optical sizing protocol

Particle size distribution of LTO materials is made with an Accusizer using single-particle optical sizing to count one particle at a time. The device range is 0.5 -500 µm in term of particle diameter. LTO powders are diluted in MilliQ water and at least 100 000 particles are measured for one experiment. Each measurement is repeated twice. showed a higher capacity than Li6PS5Cl, with an additional 15 mA.h/g (resulting in a capacity of around 195 mA.h/g at C/10). However, despite the capacity advantage, the capacity retention was found to be quite similar to Li6PS5Cl, with an average capacity loss of 13 % over 100 cycles. Note that decreasing the cut-off potential to 3. To conclude, despite the preliminary nature of our results, we demonstrated that LYBC halide SE can bring additional capacity. However, the conventional belief that halide solid electrolytes exhibit superior stability compared to the conventional argyrodite is questionable, as similar or worse stability is found using halide. Instead, a more interesting approach to deal with cathode composite stability could be to introduce a well-design coating, as we demonstrated that it can significantly improve the stability. Such methodology necessitates the design of a conformal and thin coating with a repeatable process, something industrials are better equipped to do than academics.

Halide solid electrolytes materials and methods

Solid electrolytes:

LYBC,LYC and LYB were kindly provided by Saint Gobain.

The Li3InCl6 SE was synthesized by mechanochemical milling of a stoichiometric mixture of InCl3

and LiCl using planetary ball-mill followed by annealing step, as previously reported 140 . The ball-

  Li-ion to all-solid-state battery 1.2.1 Operating principles of liquid electrolyte-based Li-ion battery

Figure 1 . 1

 11 Figure 1.1 illustrates a standard schematic of a Li-ion battery. The battery comprises a

Figure 1 . 1 :

 11 Figure 1.1: Scheme of a common lithium ion battery. Reproduced from reference.7

  number of electron involved, F is the faraday constant (96485 C/mol), M is the molar mass (g/mol). For instance, LiNi0.6Mn0.2Co0.2O2 (NMC 622) theoretical capacity is around 275 mA.h/g, while the one of LiFePO4 (LFP) is around 170 mA.h/g. The total voltage of a cell (U) can be determined by the difference between the potentials of the two electrodes, which are determined by the active materials used and the chemical potential of lithium associated with them. This establishes the open-circuit voltage (OCV) of the battery before use. During cycling, the cell voltage changes due to the thermodynamics of the electrochemical reaction. The Nernst equation provides a way to relate this voltage to the activity of the reactants in the system, particularly for a single reaction.

Figure 1 . 2 :

 12 Figure 1.2: Typical battery architectures for the conventional lithium-ion and solid-state batteries. The volumetric and gravimetric energy densities are represented by wvol and wgrav, respectively. Reproduced from reference 26 .

Figure 1 . 3 :

 13 Figure 1.3: Schematic of (a) battery-driven applications, namely ear pods, computer and electric vehicle (b) battery architecture from low to high energy density, corresponding weights of battery components (c) manufacturing processes of ASSB with respect to scalability and (d) cell type with respect to production cost.

Figure 1 . 4 :

 14 Figure 1.4: SEM (Scanning Electron Microscopy) cross section images of a ASSB that underwent 100 galvanostatic charge-discharge cycles at 50 °C, using Tantalum-doped Li7La3Zr2O12 as SE, LCO as AM and indium foil as anode.

Figure 1 . 5 :

 15 Figure 1.5: SEM images of FIB (Focused Ion Beam) cross-sections of Li6PS5Cl electrolyte pellets made with a fabrication pressure of (a) 50 MPa and (b) 370 MPa, with their corresponding relative density estimated by physical measurements. The resulting FIB reconstructions show the porosity in blue for the (c) 50 MPa pellet and (d) 370 MPa pellet. Reproduced from reference 58 .

Figure 1 . 6 :

 16 Figure 1.6: (a) Conductivity and relative density of Li6PS5Cl electrolyte as a function of the fabrication pressure, and (b) Nyquist diagrams of the electrochemical impedance spectra at the same fabrication pressures. All measurements were performed with a stack pressure of 25 MPa. Error bars represent the standard deviation on four samples. Reproduced from reference 58 .

Figure 1

 1 Figure 1.8.

Figure 1 . 7 :

 17 Figure 1.7: (a) First-cycle using different-sized LPS particles with fixed NMC size (5 µm) and 60 wt. % NMC. (b) Comparison of experimental capacities with model-predicted capacities. (c),(d) First-cycle voltage curves using different-sized NMC particles (5 and 12 µm) with fixed LPS size (3 µm for (c) and 1.5 µm for (d)) and 80 wt. % NMC. (e) Comparison of experimental capacities in (c) and (d) with model-predicted capacities. Reproduced from reference 59 .

Figure 1 . 8 :

 18 Figure 1.8: Influence of cathode active material and SE particle size on the different aspects of requirements for ASSB cathodes. Reproduced from reference 43 .

  voltammetry provides a practical assessment of the stability window and combines thermodynamic and kinetic considerations. This technique utilizes various parameters, including sweep voltage and electronic exchange surface, which must be carefully selected to reveal the stability window. Considering these statements, CV is typically favoured for comparing the stability window of different SE. Thus, Figure1.9 illustrates the stability window through CV experiments, for representative halide, argyrodite, and oxide SE. Oxide SE like Li7La3Zr2O12 (LLZO) present a large stability window and are consequently ideal for lithium metal integration, as illustrated in Figure1.9.a[START_REF] Han | Electrochemical Stability of Li 10 GeP 2 S 12 and Li 7 La 3 Zr 2 O 12 Solid Electrolytes[END_REF] . Turning to Halide, their composites often demonstrate favourable cycling stability, but a common halide such as Li3InCl6 typically experience decomposition starting from ~ 3.3 V vs. Li-In/In (where 0 V versus Li-In/In is equivalent to 0.62 V versus Li + /Li), as seen in Figure1.9.c[START_REF] Koç | Toward Optimization of the Chemical/Electrochemical Compatibility of Halide Solid Electrolytes in All-Solid-State Batteries[END_REF] . This intriguing observation arises from two considerations. Firstly, Li3InCl6 can not only oxidize, but can also be reduced.Secondly, the resulting decomposition products are likely to be less resistive and/or passivating faster than those observed for sulphide-based SE. This high cycling stability of Li3InCl6-based composites contrasts with the well-known stability issues associated with sulphide-based composites. For instance, sulphides such as Li3PS4, Li6PS5Cl (Figure1.9.b), or Li10GeP2S12 are known to undergo electrochemical oxidation starting from approximately 2.7

Figure 1 . 9 :Figure 1 . 10 :

 19110 Figure 1.9: (a) Cyclic voltammetry of Li/LLZO/LLZO+Carbon/Pt cell within the voltage range of 2.6-10.0 V. Adapted from reference 68 . (b) Cyclic voltammograms of Li6PS5Cl-Carbon composites for the first two cycles;between 0 and 4.2 V (vs Li/Li + ) at 100 μV s -1 . Adapted from reference[START_REF] Tan | Elucidating Reversible Electrochemical Redox of Li 6 PS 5 Cl Solid Electrolyte[END_REF] . (c) Cyclic voltammetry measurement at 0.01 mV/s upon 10 cycles between 2 and 3.9 V vs Li-In/In. Li3InCl6 + VGCF is used as cathode. Adapted from reference[START_REF] Koç | Toward Optimization of the Chemical/Electrochemical Compatibility of Halide Solid Electrolytes in All-Solid-State Batteries[END_REF] .

Figure 1 . 11 :

 111 Figure 1.11: (a) Local fragment distribution of NiO2 -(blue) and POx -(green) fragments for the reference SE and the uncycled and cycled composite cathode. The sum of the first four scans of the depth profile is shown. The solid electrolyte shows local POx -fragment enrichments, which can also be identified in the composite cathodes (middle picture, most intensive signals). Signals from POx -fragments at the NMC622/SE interface significantly increase after the battery cycling (total counts (TC) of POx -fragments become double). (b) Three-dimensional reconstruction of the depth profile of the cycled composite cathode. Shown are local fragment distributions of NiO2 -(blue), POx -(green), and SOx -(red) on the surface and in the bulk. This shows that a SEI (Solid Electrolyte Interphase) layer is formed directly at the NMC622/SE interface due to battery cycling. Reproduced from reference 90 .

Figure 1 . 12 :

 112 Figure 1.12: SEM characterization of the Li6PS5Cl/NMC721 interfaces after 4000 cycles. A low magnification SEM image of the Li6PS5Cl /NMC721 interface 92 .

Figure 1

 1

Figure 1 . 13 :

 113 Figure 1.13: Schematic illustrating (a) the possible scenarios for achieving electrical contact between coated CAM particles during processing, as well as (b) influence of volumetric changes of the electrode materials on the coating. Reproduced from reference 94 .

Figure 1 .

 1 Figure 1.14: (a) Cross-sectional HAADF-STEM image of LiCoO2 electrode/Li2S-P2S5 solid electrolyte interface after initial charging and (b) cross-sectional EDX line profiles for Co, P, and S elements. The arrow in a presents the positions at which EDX measurements were taken (c) Similar procedure with Li2SiO3-coated LiCoO2/Li2S-P2S5 interface after initial charging and (d) cross-sectional EDX line profiles for Co, P, S, and Si elements. Adapted from reference 100 .

Figure 1 .

 1 Figure 1.15: (a) Schematic diagram showing the reactive infiltration of LiAlO2 (LAO) in the grain boundaries of LRMN (Li1.2Mn0.54Ni0.26O2) secondary particles to form a stable and interlocked coherent interface (b) cycling performance of the samples at 1 C (c) Experimental crushing test results for Pristine LRMN and LAO@LRMN

Figure 1 . 16 .

 116 evidenced through the introduction of a new coating methodology in the year 2023, capitalizing on the reactivity of lithium surface species with CO2 gas, creating a Li2CO3 layer around SE particles105 . This approach improved the atmospheric stability of Li6PS5Cl and exhibited favourable stability performance of composites, as demonstrated in the accompanying Likewise, Al2O3 deposited on Li6PS5Cl by atomic layer deposition (ALD) technique showed comparable results. Specifically, ALD alumina coating was found to improve the stability of argyrodite under both humid and oxidizing conditions, and effectively stabilized the SE against reactivity with Li metal 106 .

Figure 1 .

 1 Figure 1.16: (a) Schematic illustration of the solid-gas reaction. (b) The ionic conductivity of gas-treated Li6PS5Cl samples before and after exposure to ambient air with a relative humidity of 17% for 1 hour and its conductivity retention. (c) Long-term cycling performance at 0.5C. Adapted from reference 105 .

Figure 1 . 17 :

 117 Figure 1.17: (a) Characterization of a 0.6 Ah class prototype pouch cell and illustration of a bi-cell structure (b,c) Discharge capacities were monitored under 0.1 C/0.1 C charge/discharge conditions as the discharging temperature was varied from 60 to -10 °C. The charging temperature was fixed at 60 °C. (d) Cycling performance and Coulombic efficiency of the Ag-C|SSE|NMC prototype pouch cell (0.6 Ah) are plotted against the cycle numbers. A constant current mode with the charge/discharge rate of 0.5 C/0.5 C was applied (voltage window, 2.5-4.25 V versus Li + /Li at 60 °C). The areal capacity loading of the NMC cathode was 6.8 mAh cm -2

Figure 1 . 18 :

 118 Figure 1.18: (a) Schematic of 99.9 wt. % micro Si electrode in an ASSB full cell. During lithiation, a passivating SEI is formed between the micro Si and the SE, followed by lithiation of micro Si particles near the interface. The highly reactive Li-Si then reacts with Si particles within its vicinity. The reaction propagates throughout the electrode, forming a densified Li-Si layer. (b) Cycle life at room temperature. All cells were tested under similar charge and discharge conditions between 2.0 and 4.3 V. The first cycle voltage profile of each respective cell is plotted in black. Reproduced from reference 111 .

  the composite optimization and design (see Figure 1.19). However, the usage of Yttrium in this SE raises the question of industrial cost and resources available for commercialization.

Figure 1 . 19 :

 119 Figure 1.19: (a) Schematics of ASSB cell configuration. Charge/discharge voltage profiles of (b) Polycristalline-NMC811 (PC-NMC 811) and (c) Single crystal-NMC811 (SC-NMC811) ASSB cells. (d) Capacity retention plots for the cells cycled at 0.5 C for 200 cycles followed by 3 cycles at 0.2 C. The same sequence repeats throughout the test. Note that the performance fluctuation is due to changes in laboratory ambient temperature during the test. (e) Rate capability comparison of PC-NMC811 and SC-NMC811 ASSB cells. Reproduced from reference 112 .

Figure 1 . 20 :

 120 Figure 1.20: Rate capability and cyclability evaluation of the LiNi0.6Mn0.2Co0.2O2/ Li5.5PS4.5Cl1.5/ In-Li SSBs. The rate capability of the fabricated ASSB working at (a) RT and (b) -20°C. Cycling performance of LiNi0.6Mn0.2Co0.2O2/Li5.5PS4.5Cl1.5/In-Li ASSB working under different temperatures and charge/discharge rates: (c) cycled at 1C under -20 °C. Solid-state batteries cycled under room temperature at (d) 5C and (e) 10C. The loading amount of cathode used for assembling solid-state batteries in this section was 2.5 mg composite cathode/cm. Reproduced from reference 113 .

Figure 1 . 21 :

 121 Figure 1.21: Long-term cycling of the NCM85 ASSB (performed after rate cycling) at a 3C rate (a) and ultra-high voltage NCM85 ASSB cycled between 2.8 and 4.8 V versus Li + /Li (b) and the corresponding charge-discharge voltage profile (c). Reproduced from reference 82 .

Figure 1 . 22 :

 122 Figure 1.22: Results of S/LYC electrodes with an excessive amount of LYC (40.7 wt. %) for all-solid-state half cells tested at 30 °C. Cross-sectional SEM-BSE images of S/LYC electrodes with (a) 29.1 wt. % and (b) 40.7 wt. % LYC after first discharge to 3.0 V (vs Li/Li + ). Note the cracks formed inside the single NCA, indicated by the arrows in (a). (c) Cycling performances at 0.5C for the S/LYC electrodes with varied amounts of LYC. The cells were charged with or without CC-CV (constant current-constant voltage) mode. In the inset, the charge-discharge voltage profiles at the 2nd and 200th cycles for S/LYC with 40.7 wt. % LYC are shown. A result for S-NCA electrodes in liquid electrolyte cells is also compared. (d) Rate capabilities. Reproduced from reference 114 .

1 46 2. 2 47 2. 2 . 1 47 2. 2 . 2 51 2. 2 . 3 54 2. 3 55 2. 3 . 1 58 2. 4 59 2. 4 . 1 59 2. 4 . 2 61 2. 5

 1462472147225123543553158459415942615 Introduction.....................................................................................................................Intrinsic decomposition of Li6PS5Cl...............................................................................Influence on first cycle performances…..............................................................Sluggish kinetics of carbon-containing systems..................................................Effect on stability upon cycling............................................................................High potential reactivity of Li6PS5Cl with NMC active materials..................................Impact of cut-off potential....................................................................................55 2.3.2 Effect on stability upon cycling............................................................................Decorrelating the two interfacial reactivities using PEIS…………………....................Equivalent model development...........................................................................Experimental data results...................................................................................Conclusion ......................................................................................................................64 This chapter includes the following publication: Quemin, E.; Dugas, R.; Koç, T.; Hennequart, B.; Chometon, R.; Tarascon, J.-M. Decoupling Parasitic Reactions at the Positive Electrode Interfaces in Argyrodite-Based Systems. ACS Appl. Mater. Interfaces 2022, 14 (43), 49284-49294. https://doi.org/10.1021/acsami.2c13150.

Figure 2 . 1 .

 21 a. To start with, Li6PS5Cl particles are mixed with 5 wt. % VGCF (Vapour Grown Carbon Fibre) and tested in a two-electrode configuration, versus Li-In/Li6PS5Cl anode composite. The cell assembly procedure and composite preparation are outlined in the Appendix 2.1.Cyclic-voltammograms were recorded for this system using a scan rate of 0.05 mV/s, covering a voltage range up to 3.9 V vs.Li-In/In at different temperatures: 25°C, 55°C, and 75°C. The results presented in Figure 2.1.b demonstrate that higher temperatures are associated with an increase in current amplitude, indicating a degradation process of the solid electrolyte governed by kinetic factors. Notably, the decomposition of Li6PS5Cl exclusively

Figure 2 .

 2 2.a, the composites with VGCF exhibit lower active material utilization, and a higher level of polarization arising from an enhanced resistance with increased VGCF content, at C/10 (1C is calculated with the theoretical capacity value of the AM in the whole manuscript). Such effect is highlighted by the impedance spectra evolution, taken at end of first charge in three-electrode setup, which presents a growing resistance with increasing amount of VGCF (see Figure2.2.b). Note that the three-electrode assembly used here is built via a multistep process described in Appendix 2.3, using Li0.5In + Li6PS5Cl (60/40 wt. %) composite as reference electrode.

Figure 2 . 1 :

 21 Figure 2.1: (a) SEM pictures of Li6PS5Cl (b) Cyclic voltammetry of VGCF + Li6PS5Cl (5/95 wt. %) at different temperatures. Cyclic voltammograms at 0.05 mV/s and at three different temperatures (25°C, 50°C and 75°C).

  Early delithiation of first cycle of NMC 622 + Li6PS5Cl (70/30 wt. % ratio) with additional VGCF put in 1 or 5 wt. %. Cells are cycled at C/10 and a loading of 16 mg/cm² of composite is used.

Figure 2 . 2 :

 22 Figure 2.2: (a) First cycle at C/10 using NMC 622 + Li6PS5Cl (70/30 wt. % ratio) as cathode composites with 5 wt. % or 10 wt. % additional VGCF, or without VGCF. A loading of 16 mg/cm² of composite is used. (b) Composites of same formulation are implemented in three electrode setup, to take PEIS measurements at the end of the first charge (3.6 V vs. Li-In/In) at C/10, 10 hours of relaxation is made prior the measurements.

Figure 2 . 3 :

 23 Figure 2.3: (a) Scheme of the first cycle irreversibility measurement. The "true" irreversibility is obtained after one cycle, by clamping the potential at 2.1 V vs. Li-In/In at the end of discharge for 30 hours, ensuring an ending current below C/500. Galvanostatic cycling is made between [3.6 V-2.1 V vs. Li-In/In] at C/10. (b) Apparent and true irreversibility of NMC 622 liquid and solid-state based cells, with 5 wt. % VGCF (c) Apparent and true irreversibility of cells using 70/30 wt. % NMC 622-Zr coated/Li6PS5Cl with pristine or 2.5 wt. %, 5 wt.% and 10 wt.

Figure 2 . 4 .

 24 Figure 2.4. We used Li0.5In/In + Li6PS5Cl (60/40 wt. %) as both reference and counter electrode composites, 1 t/cm² pressure is applied upon cycling and the system is designed to be airtight. Two solid electrolyte layers of 30 mg (adding approximately 0.2 mm to the stack thickness each) are added on both sides of the reference layer, creating a five-layer configuration for PEIS measurements. Secondly, electronic resistance values are obtained by DC conductivity measurement of an 80 mg composite, pelletized at 1 t/cm² in a two electrode cell setup. To do so, 10 mV polarization is imposed until a steady-state current is reached, and the electronic resistance is computed by Ohm's law. A scheme of the experiment is shown in Figure 2.5.a.

Figure 2 . 4 :

 24 Figure 2.4: Schematic of the three-electrode setup used to take PEIS measurements.

Figure 2 .

 2 Figure 2.5.b, presents the distinct ionic and electronic resistance evolution, with our C-free reference system (70/30 wt. % NMC 622/SE) showing Relec and Rion values of 60 and 120

Figure 2 . 5 :

 25 Figure 2.5: (a) Scheme of the experiment used for both DC conductivity measurements in pellet system and ionic conductivity measurement in 3-electrode system. (b) Resulting electronic and ionic resistance of the composites using different addition of carbon additives. Pristine state is made with 70/30 wt. % ratio of NMC 622/Li6PS5Cl. Same amount of composites (80 mg) is used for all experiments.

Figure 2 .in

 2 6.a. The cells are maintained at a defined cutoff potential (3.6 V vs. Li-In/In here) and a relaxation of 10 hours is done after each 10 hours of imposition. This relaxation is followed by a 10 mV potential imposition step for 20 min, and the steady-state current obtained after 1000 seconds is used to compute the resistance R1000s of the cell, by simple ohmic law. This procedure is repeated 15 times. As a result, the C-containing system is showing a linear increase with time, while the C-free one rapidly reach a constant resistance after 20 hours, mirroring a passivating interface formation (see Figure 2.6.b). On the contrary, the C-containing cells reflect a continuous insulative phase formation building up at VGCF/SE interfaces, which does not passivate in 150 hours. Such evolutions are mirrored by the capacity retention at C/10 of C-free and C-containing cells, which are shown Figure 2.6.c. Especially, the C-containing system shows a clear decay by losing approximately 10 % of capacity in 25 cycles, contrasting with the stability of the C-free system.

Figure 2 . 6 : 2 . 3

 2623 Figure 2.6: (a) scheme of the protocol. A monitoring of resistance is done during a constant potential imposition (at 3.6 or 3.9 V vs. Li-In/In). A first charge at C/10 is directly followed by a potential imposition. Every 10 hours of constant voltage, 10 hours of relaxation is done, followed by a 10 mV step for 20 minutes. (b) Resistance evolution during potential imposition at 3.6 V vs Li-In/In. Two cells per system are shown. (c) Capacity retention evolution in C-containing and C-free systems, cycling at C/10 until 3.6 V vs. Li-In/In. Two cells per system are shown.

Figure 2 . 7 :

 27 Figure 2.7: Potential imposition effect on NMC 622 + Li6PS5Cl (70/30 wt. %) system (a) Scheme of the procedure (b) Galvanostatic charge-discharge at C/10 in NMC 622/Li6PS5Cl/Li-In-Li6PS5Cl system with a 60 hours potential imposition on the third charge at different potential (3 V, 3.4 V, 3.6 V and 3.9 V vs. Li-In/In). Additionally, the average discharge potential of the 4 th cycle is shown.

Figure 2 .

 2 8.b.Although the electronic conductivity of NC90 material, either mixed with Argyrodite or not, is greater than other NMC, it demonstrates slightly lower capacity loss on Figure2.8.a. This suggests that the degradation mechanism in presence of SE and transition layered oxide materials is not solely governed by the electronic conductivity. Note that, we are comparing similar average particle size systems (~ 4 µm) as seen on Figure2.8.c-e, but distinct morphologies and possibly distinct particle size dispersions are observed in between those three active material particles. Thus, a legitimate question regards the effect of such parameters on the electronic percolation inside composites. Nevertheless, similar trend were demonstrated between 1 wt. % PTFE and 30 wt. % Li6PS5Cl containing systems. Therefore, we can attribute the composites electronic conductivity changes as mostly pertaining to the nickel content present in the active material.

Figure 2 . 8 :

 28 Figure 2.8: Potential imposition effect on various transition metal oxide (a) Similar procedure as in Figure 2.7 is used, with a 3.6 V imposition. Capacity loss (%) on the 4 th cycle after 60 hours imposition is shown for different active material. Two cells are plotted. (b) Electronic conductivity of systems with AM +1 wt. % PTFE and AM + Li6PS5Cl (70/30 wt. %) (30-50 mg of composite is used for those last systems). SEM micrographs of (c) NMC 622 (d) NMC 811 and (e) NC 9010.

Figure 2 . 9

 29 Figure 2.9 show capacity retentions and resistance evolutions at C/10 for C-free and C-containing cells cycled up to 3.9 V. Both show a ~ 20 % capacity loss after 20/25 cycles, and possibly a slight beneficial effect of carbon. By employing a similar methodology as depicted in Figure 2.6 (voltage clamping), both cell types exhibit a substantial and rapid increase in cell resistance within the initial 25 hours of cycling, reaching nearly 13 000 Ohm after 80 hours.This phenomenon suggests the occurrence of an AM/SE interfacial passivation, as observed in

Figure 2 . 9 .

 29 Figure 2.9. Collectively, these findings imply a significant deleterious reactivity, surpassing the

Figure 2 . 9 :

 29 Figure 2.9: (a) Capacity retention of galvanostatic cycling at C/10 with 3.9 V vs. Li-In/In. Two cathode composites are used: NMC622 + Li6PS5Cl (70/30 wt. %) and NMC622 + VGCF + Li6PS5Cl (66.5/5/28.5 wt. %). (b) Resistance evolution during potential imposition at 3.6 V vs Li-In/In. Two cells per system are shown.

Potentiostatic

  Electrochemical Impedance Spectroscopy is a technique of choice to assess degradation phenomena occurring in a battery. Especially, it can efficiently decorrelate phenomena affecting distinct interfaces or/and having different kinetics, provided that they have distinct characteristic frequencies.

Figure 2 .

 2 10.a-d presents the PEIS spectra evolution at end of charge of NMC 622 composite cathodes for 10 cycles, cycled until both 3.6 V and 3.9 V and with or without 5 wt. % VGCF. PEIS is done at equilibrium, after a 10 hours relaxation, in the frequency range of 500 kHz to 3 mHz, and we have used a 20 mV perturbation to ensure a limited noise and the stability of the system along the measurement. Data were reproduced at least once for every system. Moreover, we proceeded to fit the data through a homemade Matlab program, and the fitting results are presented in open black circles alongside the experimental data on Figure 2.10.a-d. Both data set greatly match with each other.

Figure 2 . 10 :

 210 Figure 2.10: Effect of VGCF and cut-off potential via impedance spectroscopy. (a), (b) Spectra of cells cycling with 3.6 V vs. Li-In/In cut-off potential in C-free and C-containing systems (c), (d) Spectra of cells cycling with 3.9 V vs. Li-In/In cut-off potential in C-free and C-containing systems. PEIS spectra are taken at the end of charge of cycle n°1, 2, 3, 6, 7 and 10. 20 mV of amplitude is applied after 10 hours of relaxation, [500 kHz -3 mHz]. The fitted spectra obtained by a personal Matlab optimization program are plotted along with each experimental data in open black circle. (e) Original TLM model representing all phenomena occurring inside our composites (f) TLM practically used for fitting with R Global, conduction = f(R Elec, composite, R Ion, composite).

Figure 2 .

 2 Figure 2.11 presents the capacity evolution at C/30 and the R Global, conduction and R CT values, both extracted from the four cells used in Figure 2.10. Despite a lower C-rate and an enhanced capacity fading, the three-electrode cell follows the same trends as the one of twoelectrode cells (see Figure 2.8.a and Figure 2.5.c). Turning to the evolution of R Global, conduction and R CT, it presents different behaviour in respect to cut-off potential and carbon additive presence. At 3.6 V vs. Li-In/In cut-off, the C-free cell metrics remain quite stable, in line with its stable capacity, while the R Global, conduction continuously increases in C-containing cells (growing from 200 to nearly 2000 Ohm) (Figure 2.11.a). Note that the two cells show equally a small increase of R CT. Interestingly, cells reaching 3.9 V vs. Li-In/In experience an important

Figure 2 . 11 :

 211 Figure 2.11: Result of PEIS fitting. Results are presented for cells cycling (already shown in Figure 2.10) with cutoff potentials reaching 3.6 V vs. Li-In/In (a) and 3.9 V vs. Li-In/In (b) Galvanostatic cycling at C/30 of the threeelectrode cells are shown, next to R Global and R CT evolution during the first 10 cycles, computed by PEIS fitting of the spectra. R Global account for the global conduction resistance.

Figure 2 . 12 :

 212 Figure 2.12: Decorrelation of R Ion, composite and R Elec, composite from R Global, conduction (a) Scheme of the fitting procedure to extract R Elec, composite, cycle 0 and R Ion, composite, cycle 0. (b) Scheme of the fitting procedure to extract R Elec, composite, cycle 10 and R Ion, composite, cycle 10 (c) Results of the fitting procedures comparing cycle 0 and 10 in log scale, on the cells presented in Figure 2.10. R Global account for the global conduction resistance.

Figure 2 . 13 :

 213 Figure 2.13: Global schematic of the distinct behaviour of NMC 622/Li6PS5Cl cathode composite as function of both VGCF presence and cut-off potential.

68 3. 2 71 3. 3 74 3. 4 81 3. 5

 682713744815 .....................................................................................................................Validation and presentation of the device ....................................................................69 3.2.1 Presentation of the cell........................................................................................69 3.2.2 Accuracy and repeatability of the cell: a LiCoO2 study .......................................Chemo mechanical effect in TMLO-based composites................................................72 3.3.1 Active material vs. composites electronic conductivity .......................................72 3.3.2 First cycle evolution of six TLMO composites......................................................Particle size impact on Li4Ti5O12 composites electronic conductivity........................78 3.4.1 AM synthesis and characterization......................................................................78 3.4.2 First cycle electronic conductivity evolution.........................................................Conclusion ......................................................................................................................88 This chapter includes the following publication: Quemin, E.; Dugas, R.; Chaupatnaik, A.; Rousse, G.; Chometon, R.; Hennequart, B.; Tarascon, J. An Advanced Cell for Measuring In Situ Electronic Conductivity Evolutions in All-Solid-State Battery Composites. Adv. Energy Mater. 2023, 2301105. https://doi.org/10.1002/aenm.202301105.

Figure 3 . 1 .

 31 Figure 3.1. (a) Scheme of the in-situ electronic conductivity measurement upon cycling. (b) Picture of a cycling cell in CE to GROUND connection. (c) Picture of the cell body with the aluminium mesh placed inside.

Figure 3 . 2 .

 32 Figure 3.2. Presentation of the data treatment using LiNi0.9Co0.1O2 composites as example (a) GITT-like curve obtained from channel 1 (b) Current versus time obtained from channel 2 (c) Resulting electronic conductivity vs. x curve. (d), (e) Insets showing specific DC polarization points in potential and current versus time. All current versus time points present a stable behaviour, with small fluctuation restricted to a range of merely 0.01 mA (which would only result in an insignificant deviation below 0.5% for points 1 to 3).

σ 2

 2 Elec vs. State of charge Li 6 PS 5 Cl (SE) Active material + SE Li 0Accuracy and repeatability of the cell: a LiCoO2 study To illustrate the significance of this approach, our study focused on investigating the phase transitions occurring in a well-established material, specifically monolithic LiCoO2 powders (provided by UMICORE) composed of primary particles ranging in size from 2-3 µm, which exhibit partial agglomeration into 10 µm particles as observed in scanning electron microscopy images (Figure 3.3.a-b). In order to conduct this investigation, LiCoO2 was combined with Li6PS5Cl for both galvanostatic intermittent titration technique (GITT) and in situ direct current conductivity measurements, as shown in Figure 3.3.c-f, respectively performed in two-electrode and in situ conductivity cell setups. For the in situ conductivity measurements, galvanostatic cycling at a rate of C/50 was periodically interrupted every 1

Appendix 3 . 1 .

 31 The delithiated powders are further washed and dried, prior being mixed with Li6PS5Cl to take DC conductivity measurements in two-electrode cells. Following this metalinsulator transition, and in accordance with existing literature, a biphasic system (0.75 < x < 0.96) emerges, consisting of two end members, Li0.75CoO2 and Li0.96CoO2, exhibiting an O3 structure with distinct lattice parameters 137 . This gives rise to a characteristic bell-shaped variation in diffusion coefficients, as manifested by the evolution of the apparent diffusion coefficient, D App (Figure3.3.c). Such behaviour is commonly observed in biphasic insertion materials, such as LiFePO4, for instance138,139 . Subsequently, a solid solution behaviour is observed until x = 0.4, without any notable features in terms of electronic conductivity or apparent diffusion coefficient. Overall, the observed evolution aligns perfectly with the reported phase transitions in LCO. More importantly, we have established the excellent repeatability and accuracy of our device (Figure3.3.d-f).

Figure 3 . 3 (

 33 Figure 3.3 (a) (b) SEM images of LCO particles. (c) The apparent diffusion coefficient (DApp) and corresponding GITT curves of an LCO cell cycled in two-electrode setup (16-18 mg/cm²). (d) Evolution of electronic conductivity during first cycle of three cells. (e) Inset of electronic conductivity of LCO during early delithiation, with ex situ liquid points added as blue dots. (f) Galvanostatic cycling at C/50 of the cells used for monitoring electronic conductivity.

Figure 3 . 4 .

 34 Figure 3.4.a). Li6PS5Cl is further mixed with the five targeted transition metal layered oxides, each possessing distinct particle sizes and morphologies (as observed in SEM micrographs depicted in Figure 3.4.b-f), in a weight ratio of 30/70 wt.%. The electrochemical performance of all the investigated composites, cycled for 15 cycles at a rate of C/10, is presented in Figure 3.4.g. It is noteworthy that all the composites exhibit a consistently low and stable capacity ranging from 90 to 150 mA.h/g, depending on the composition of the sample. In contrast, all of these layered oxides demonstrate capacities exceeding 150 mA.h/g in liquid cells. This disparity can be attributed to the comparatively sluggish kinetics observed in the solid-state due to unoptimized percolation.

Figure 3 . 4 :

 34 Figure 3.4: SEM micrographs of (a) Li6PS5Cl, and (b) NMC111, (c) NMC622, (d) NMC811, (e) NC9010 and (f) LiNiO2.(g) Cycling performances of NMC111, NMC622 and NMC811, NC9010 and LiNiO2 at C/10, in the [2.1 V-4.5 V vs. Li + /Li] potential window, and using a 70/30 wt. % active material/Li6PS5Cl ratio (first 15 first cycles are presented). (h) Electronic conductivity measurements of active material + 2 wt. % PTFE films at 1 t/cm² in 2electrode setup and electronic conductivity measurements of active material + 30 wt. % Li6PS5Cl at 1 t/cm² prior cycling, in 2-electrode setup. Three cells are done for each measurements. The average and the standard deviation are represented.

Figure 3 . 5 , 3 .

 353 while exemplary corresponding potential versus capacity curves can be found in Appendix 3.It should be emphasized that the initial electronic conductivity values depicted in Figure 3.5 align closely with our prior findings (Figure 3.4.h).

Figure 3 . 5 :

 35 Figure 3.5: Electronic conductivity measurements during first cycle of (a) LiNiO2 (b) NC 9010 (c) NMC 811 (d) NMC 622 and (e) NMC 111, when mixed with 30 wt. % Li6PS5Cl. Galvanostatic cycling is made at C/50 (loading of ~70 mg/cm²), every 1.5 h a point of conductivity is taken after 1 h relaxation. [2.1-4.5 V vs. Li + /Li] potential window is used. (f) Overall dataset represented in logarithm scale.

Figure 3 . 6 :

 36 Figure 3.6: Comparison between the in situ and ex situ electronic monitoring in liquid or solid-state cells. For liquid in situ monitoring (black curve), It uses the three-electrode cell from EL CELL company, using a similar setup as in our advanced solid-state cell. NMC 811 with 1 wt. % PTFE film is used as cathode, and LP30 as electrolyte.For the in situ solid-state monitoring (dark green curve), we used our usual setup in the advanced cell, with NMC 811 mixed with 30 wt. % Li6PS5Cl as cathode composite. For the ex situ measurements (blue points), 80-100 mg of NMC 811 is delithiated in Swagelok liquid cells, further washed with DMC and dried. The resulting powder is mixed with 30 wt. % Li6PS5Cl and pelletized, to be embedded in a routine cell for DC conductivity points.

Figure 3 . 7 :

 37 Figure 3.7: First three cycles with electronic conductivity evolution of a) NMC 622 and b) NC 9010.

Figure 3 . 8 :

 38 Figure 3.8: XRD diffractograms of LTO made by solid-state synthesis. (a) Different excess ratio of Li2CO3 are tried to diminish impurities, the hand-grinded and further pelletized pellets are put for 2h at 900°C (5°C/min) (b) XRD diffractogram of the 5 % excess sample, re hand-grinded and pelletized for a second heat treatment for 2h at 900°C (5°C/min).

7 .

 7 HNO3 diluted with deionized water was added while stirring. The resulting combustion ash was divided into four batches, three of which underwent different heat treatment durations (t0 = 0 h, t1 = 15 hours, t2 = 48 hours, t3 = 95 hours) at a temperature of 900 °C with a heating rate of 5 °C/min to achieve distinct particle sizes. X-ray diffraction patterns of the four samples are illustrated in Figure 3.10.a-b, and Rietveld refinements can be found in the Appendix 3.Regardless of the thermal treatment, all samples exhibited a single phase, except for the 0 hour LTO sample which contained a small quantity of rutile TiO2 (2.5 (2) wt. %) (refer to the Appendix 3.7 for details).The four samples examined in this study exhibit an identical cubic lattice parameter, indicating their identical chemical composition. The peak profiles of these samples exhibit broadening, which can be attributed to microstructural effects and microstrain, as determined through Rietveld refinements. Williamson-Hall plots, presented in the Appendix 3.8, provide a clear representation of the evolution of sizes (intercept) and microstrain (slope). The analysis revealed a calculated crystallite size of approximately 96 (6) nm, along with some microstrain (ε = 0.7784 (6) × 10 -4 ), for the non-annealed LTO sample (0 hours at 900°C). On the other hand, the samples annealed for longer durations at 900°C showed crystallite sizes ranging from 145(13) to 183 (12) nm, with no broadening due to microstrain. Scanning electron microscopy images of the four LTO samples (Figure3.10.c-n) exhibit various morphologies. The nonannealed LTO sample (0 hours at 900°C) consists of primary particles with a size of approximately 300 nm (Figure3.10.c, g, h). With increasing annealing time at 900°C, a coarsening of the primary grains is observed, with the largest grains exceeding 3 µm for the sample annealed for 95 hours (Figure3.10.f, m, n). To further analyse the secondary particle size, an Accusizer employing single-particle optical sizing was utilized, revealing that large agglomerates (see Appendix 3.9 for further details) define a majority of the powder volume.Interestingly, longer heating times result in sintering of the larger agglomerates, leading to a shift in particle size from the range of 100-500 µm to the range of 10-70 µm.

Figure 3 . 9 :

 39 Figure 3.9: XRD diffractograms of LTO made by solution-combustion synthesis. Different excess ratio of Lithium additive are tried to diminish impurities.

Figure 3 . 3 . 11 )

 3311 Figure 3.10: a) XRD powder patterns (λ Cu) of the four active materials (b) Zoom of the XRD patterns around the (440) peak. (c)-(n) SEM images of the 0 h, 15 h, 48 h and 95 h solution-combustion LTO materials, depicted from left to right, at 10000, 40000 and 160000 magnification. (o) Average among three measurements of particles dispersion between 500 nm and 500 µm, represented in volume fraction (%).The volume fraction is obtained assuming spherical particles.

Figure 3 . 11 :

 311 Figure 3.11: Electronic conductivity evolution during first cycle of NEI (60/40 wt.%) and solid-state homemade LTO (70/30 wt.%). Their morphology is depicted with SEM pictures.

  batteries were initially subjected to cycling tests in standard cells at a rate of C/10 (1C corresponds to 3 moles of Li per mole of active material in 1 hour). The first cycle results revealed reversible capacities that decreased as the duration of the annealing step increased (refer to the Appendix 3.11 for detailed). Furthermore, additional charge-discharge tests were conducted on the aforementioned composites at various C-rates, ranging from C/10 to 2C. Notably, the LTO sample synthesized with a 0-hour annealing time demonstrated the highest power rate performance among the tested composites (see Figure 3.12).

Figure 3 . 12 :

 312 Figure 3.12: C rate tests with the four solution-combustion LTO (0 h, 15 h, 48 h, 95 h at 900°C) mixed with 50 wt. % argyrodite. Cycling is done from C/10 to 2C in two-electrode setup, using Li0.8In + Li6PS5Cl (60/40 wt. %) as counter electrode.

Figure 3 .

 3 Figure 3.13: (a) Electronic evolution of the 0 h, 15 h, 48 h and 95 h solution-combustion LTO mixed with 50 wt. % Li6PS5Cl. Galvanostatic cycling is done at C/30, a point of conductivity is made every 1 hour and after 1 hour relaxation, and represented in a logarithm scale. (b) Linear scale of the electronic conductivity points between 0-0.012 S/cm.

Figure 3 . 14 :

 314 Figure 3.14: (a) Electronic evolution of the sieved 0 h solution-combustion LTO mixed with 50 wt. % Li6PS5Cl. The pristine composite is presented, as well as the > 20 µm, 10 -20 µm, 5 -10 µm and < 5 µm ones. Galvanostaticcycling is done at C/30, a point of conductivity is made every 1 h and after 30 minutes relaxation, and represented in logarithm scale. (b) Linear scale of the electronic conductivity points between 0-0.012 S/cm. For ease of interpretation, the delithiation points were removed.

Figure 3 . 15 :

 315 Figure 3.15: (a) First galvanostatic cycle of the sieved LTO (pristine, > 20 µm, 10-20 µm, 5-10 µm and < 5 µm) mixed with 50 wt. % Li6PS5Cl. Cycling is done at C/10 in two-electrode setup, using Li0.8In + Li6PS5Cl (60/40 wt. %) as counter electrode. (b) C rate tests.

Figure 3 . 16 :

 316 Figure 3.16: Proposed schematic of lithiation pathway in LTO composites

Figure 4 . 1

 41 Figure 4.1 illustrates the initial cycle at a C/10 rate in the [2.1-3.9 V vs. Li-In/In] potential window, for nine cells fabricated from two distinct composites, which were prepared through manual grinding. All cells exhibited excellent repeatability, especially for composite n°2.Figure 4.2 presents the discharge capacity profiles of the aforementioned cells for a

Figure 4 . 2

 42 presents the discharge capacity profiles of the aforementioned cells for a minimum of 80 cycles at a C/2 rate. Overall, the results demonstrate remarkable similarity across all nine cells. Note that the oscillations observed are linked to temperature changes over the days since the experiments were carried-out under non-thermostatic conditions. The average initial cycle capacity at C/10 is 168.1 mA.h/g, with a standard deviation of 2.7 mA.h/g. Subsequently, the second cycle at a C/2 rate yielded a capacity of 135.42 mA.h/g, accompanied by a standard deviation of 3.6 mA.h/g. Additionally, as highlighted in Figure 4.3,

Figure 4 . 1 :

 41 Figure 4.1: (a) First cycle of NMC 622 + Li6PS5Cl + 1.5 wt. % VGCF first composite prepared by hand grinding, at C/10 and on four different cells. Potential window is 2.1-3.9 V vs. Li-In/In. (b) Same experiment with a second composite prepared by hand grinding and on five different cells.

Figure 4 . 2 :

 42 Figure 4.2: (a) Cycling performances of NMC 622 + Li6PS5Cl + 1.5 wt. % VGCF first composite prepared by hand grinding and on four different cells. The first cycle is made at C/10 while the following ones goes at C/2. Potential window is 2.1-3.9 V vs. Li-In/In. (b) Same experiment with a second composite prepared by hand grinding and on five different cells.

Figure 4 . 3 :

 43 Figure 4.3: Average capacity loss estimated by linear extrapolation after 100 cycles at C/2 for the nine cells belonging to the two composites. one and two standard deviation areas are shown for each composites, and the global values of the nine cells are also written.

Figure 4 . 4 :

 44 Figure 4.4: (a) SEM pictures of NMC 811 (MSE). (b) First cycle at C/10 of both NMC 622 (Umicore) and NMC 811 (MSE) Li6PS5Cl composites, with 1.5 wt. % VGCF addition. (c) C-rate tests of NMC 811/Li6PS5Cl composite with 1.5 wt. % VGCF or without VGCF. The cathode loading is comprised between 16 and 18 mg/cm².

Figure 4 . 5 .

 45 a). Secondly, the introduction of VGCF has a perceptible negative impact on cells cycled up to 3.6 V vs. Li-In/In (as exemplified in Figure 4.5.b), whereas its effect becomes negligible when employing a 3.9 V cut-off (as shown in Figure 4.5.c). Lastly, a substantial increase in capacity decay is observed when employing higher cut-off potentials, (as exemplified in Figure 4.5.d), with the notable exception of the NMC 811 C-containing cells cycled at C/2. All capacity retention curves are provided in the Appendix 4.1 and 4.2. The outcomes of this study does not only confirm the behaviour observed in Chapter II with

Figure 4 . 5 :

 45 Figure 4.5: Cycling behaviour of NMC 811 + Li6PS5Cl composites (70/30 wt. %). 1.5 wt. % additional VGCF can be used. The red dashed lines represent the linear extrapolation of the average of two cells. (a) Impact of C-rate, using C-containing cells and cycling up to 3.9 V vs. Li-In/In. C/2 and C/10 are compared, with two cells per systems. (b) Impact of VGCF addition, using cells cycling up to 3.6 V vs. Li-In/In and a C-rate of C/2. C-free and C-containing cells are compared, with two cells per systems. (c) Impact of VGCF addition, using cells cycling up to 3.9 V vs. Li-In/In and a C-rate of C/2. C-free and C-containing cells are compared, with two cells per systems. (d) Impact of cut-off potential, using C-containing cells and a C-rate of C/10. Cells cycled up to 3.6 and 3.9 V vs. Li-In/In are compared, with two cells per system. The capacity retention is extrapolated up to 100 cycles, and we compute the average of the two cells.

Figure 4 . 6 . 3 .

 463 The lowest pressure (2 t/cm²) shows the largest polarization, while higher pressures present similar polarization between each other. If we focus on capacity retention, Figure 4.7 shows the discharge capacity at a rate of C/2 in relation to the cycle number, along with the capacity retention derived from the moving average capacity curve. In this context, the moving average of discharge capacity serves the purpose of mitigating the fluctuations induced by temperature variations. All original curves can be found in Appendix 4.As expected, the 2

Figure 4 . 6 :

 46 Figure 4.6: First cycle as function of assembly pressure. The assembly pressure is defined between 2 t/cm² and 6 t/cm², and applied for 10 minutes on the battery stack. The first cycle is made at C/10, in the [2.1-3.9 V vs. Li-In/In] potential window. Inset shows the early delithiation polarization. Two cells per system are presented. All cells were launched on the same week.

Figure 4 . 7 :

 47 Figure 4.7: (a) Discharge capacity versus cycle number as function of assembly pressure. The assembly pressure is defined between 2 t/cm² and 6 t/cm², and applied for 10 minutes on the battery stack. The first cycle is made at C/10, in the [2.1-3.9 V vs. Li-In/In] potential window. (b) Capacity retention of the discharge capacity moving

Figure 4 . 8 ,

 48 using distinct ranges of loading, specifically 8-10, 12-14, 16-18, 20-22, 24-26 mg of composite per cm². Each loading range is represented by two individual cells. The lowest loading (8-10

Figure 4 . 8 :

 48 Figure 4.8: First cycle galvanostatic curve at C/10, using different loading of cathode composites: 8-10, 12-14, 16-18, 20-22 and 24-26 mg/cm². NMC 811 (MSE) + Li6PS5Cl (70/30 wt. %) composite is used, and an additional 1.5 wt. % VGCF is added. Cells are cycled in the 2.1 -3.9 V vs. Li-In/In potential window. Two cells per system are shown. All cells were launched on the same week.

Figure 4 . 9 :

 49 Figure 4.9: Capacity retention at C/2 after one cycle at C/10, using different loading of cathode composites: 8-10, 12-14, 16-18, 20-22 and 24-26 mg/cm². NMC 811 (MSE) + Li6PS5Cl (70/30 wt. %) composite is used, and an additional 1.5 wt. % VGCF is added. Cells are cycled in the 2.1 -3.9 V vs. Li-In/In potential window. Two cells per system are shown with at least 90 cycles for each.

4. 4

 4 Coating benefits on high potential stability4.4.1 Carbonates-coated NMC 811Previously, NMC materials heat-treated in air environment have been used to improve the capacity stability of Li6PS5Cl-based solid-state composite batteries, through the deposition of carbonate species on the NMC surface[START_REF] Walther | The Working Principle of a Li 2 CO 3 /LiNbO 3 Coating on NCM for Thiophosphate-Based All-Solid-State Batteries[END_REF][START_REF] Strauss | 2 ZrO 3 -Coated NCM622 for Application in Inorganic Solid-State Batteries: Role of Surface Carbonates in the Cycling Performance[END_REF] . This coating occurs by a reaction between CO2 and the lithium residues on the NMC surface. In this study, we have exposed our NMC 811 material to heat treatment in air at temperatures ranging from 100°C to 800°C, using a heating rate of 5°C/min and for a duration of 8 hours. The resulting active materials were subsequently combined with argyrodite in a 70/30 wt. % ratio, followed by the addition of 1.5 wt. % VGCF. After an initial cycle at C/10, a galvanostatic cycle at C/2 was performed and the resulting curves are shown in Figure4.9, with two cells per composite. Notably, all composites exhibited reduced capacity compared to the non-heated NMC 811-based composite, and the 800°C treated one presenting the worst capacity. This behaviour is linked to the increase in electronic resistance resulting from a thicker carbonate deposit, as demonstrated in Figure4.9.c. Additionally, all the NMC 811-based composites display similar capacity decay. Thus, these cumulative findings indicate no advantage of the carbonate coating in this context.

Figure 4 . 9 : 2 . 5 .

 4925 Figure 4.9 : Impact of heat treatment of NMC 811 (MSE) under air for 8h at different temperature (100°C, 200°C, 300°C, 400 °C, 600 °C and 800°C). 5°C/min heat rate is used. (a) The resulting NMC 811 composites are cycled in the [2.1-3.9 V vs. Li-In/In] potential window, at C/2 after one cycle at C/10. Two cells per composites are shown. The cells are cycled under ambient temperature. (b) First cycle at C/10. (c) Electronic conductivity of NMC 811 powders mixed with 1 wt. % PTFE, took with DC conductivity on pellet system. All cells were launched on the same week.

Figure 4 . 10 :

 410 Figure 4.10: Impact of carbonate-coated Li6PS5Cl by CO2 gas flow treatment. (a) Capacity retention of cells integrating normal argyrodite and argyrodite with 10 min, 20 min, 45 min and 60 min CO2 treatment in their cathode composites. Cycling is done between 2.1 and 3.9 V vs. Li-In/In. One cycle is done at C/10 followed by C/2 cycling. Two cells are done per system with at least 100 cycles for each. (b) First cycle at C/10 of the aforementioned cells. (c) Ionic resistance of the distinct argyrodite. Three pellets are done per system. Average values are depicted, alongside the standard deviations. (d) Arrhenius plot of the distinct argyrodite, took between -10 to 40 °C.

Figure 4 . 11 :

 411 Figure 4.11: (a) Two first cycles between 2.1 and 3.9 V at C/20 and at 55°C. (b) Two first cycles between 2.1 and 3.9 V at C/20 and at 0°C. (c) Discharge capacity evolution after formation cycles and at ambient temperature.

Figure 4 . 14 :

 414 Figure 4.14: (a) Schematic of the cycling protocol for formation cycles: 2 cycles are done at C/20 with distinct low cut-off potential, namely 2.5, 2.1, 1.8, 1.5 and 1.3 V vs. Li-In/In. The upper cut-off potential is defined at 3.9 V vs. Li-In/In. The following cycles are done at C/2 in the potential window of [2.1 -3.9 V vs. Li-In/In]. (b) The resulting discharge capacities are plotted as moving average and as function of cycle number. Two cells are plotted per system with at least 100 cycles for each. (c) Capacity retention of the discharge capacity moving average of the aforementioned cells. The moving average is done by selecting 8 data points for each average. All cells were launched on the same week.

Figure 4 . 15 :

 415 Figure 4.15: Capacity retention at C/2 between 2.1 and 3.9 V vs. Li-In/In, after (a) two cycles at C/20 in the 1.8-3.4 V vs. Li-In/In potential window or (b) 10 cycles at C/20. We use a NMC 811 (MSE) + Li6PS5Cl (20 min under CO2) in a 70/30 wt. % ratio, with the addition of 1.5 wt. % VGCF. The assembly pressure is set at 3 t/cm². All cells were launched on the same day.

Figure S2. 1 : 2 . 2 :Figure S2. 2 :

 1222 Figure S2.1: Two-electrode cell setup

Figure S2. 3 :

 3 Figure S2.3: Three-electrode cell setup with a) detailed assembly of the cell body and b) detailed assembly of the complete cell.

2. 4 :Figure S2. 4 :

 44 Figure S2.4: First cycle of NMC 622 + 30 wt. % solid electrolyte + 5 wt. % VGCF at C/20, with 3.6 V vs. Li-In/In cutoff potential. A constant potential of 2.1 V vs. Li-In/In is applied for 20 hours at the end of discharge. Two cells are done for each system. a) Galvanostatic cycling b) Extracted points of initial and final irreversibility.

2. 6 :Figure S2. 6 :

 66 Figure S2.6: C rate tests using NMC 622 + Li6PS5Cl (70/30 wt. % ratio) as cathode composites with 5 wt. % or without VGCF. The cut-off potential is 3.6 V vs. Li-In/In.

Figure S2. 8 :

 8 Figure S2.8: Theoretical spectra evolution with (a) changing Rion (b) changing Relec and (c) changing R CT.

Figure S2. 10 :

 10 Figure S2.10: Theoretical spectra evolution with (a) changing R global and (b) changing R CT.

Figure S2. 13 :

 13 Figure S2.13: Comparison of two uncoated-NMC + Li6PS5Cl (70/30 wt. %) PEIS spectra after building in threeelectrode setup, with two different loadings: 18 or 61 mg/cm².

  Four equations and two additional unknowns (Ea and σ0) are obtained. Overall, we have 7 unknowns [Rion (0°C) | Rion (10 °C) | Rion (25 °C) | Rion (30 °C)| Relec | C | Ea] and 8 nonlinear equations. Using a similar optimization program as the one presented previously, we can solve this problem. The values obtained for the cycle 10 of the 4 cells presented in Figure 2.11 are shown in

  ) and the residues are low enough (FigureS2.16).

Figure S2. 16 :

 16 Figure S2.16: Bode plots of Z (Rglobal) -Z (Rion, Relec) module and Z (Rglobal) module for each frequency between 300 kHz and 3 mHz.

Figure S2. 17 : 3 . 4 :Figure S3. 2 :

 17342 Figure S2.17: Arrhenius plot of the Rion (T) optimized values

Figure S3. 3 : 3 . 6 :

 336 Figure S3.3: Apparent diffusion coefficient upon first cycle in NMC 111 and NMC 811 took in two-electrode cell setup.

Figure S3. 5 :

 5 Figure S3.5: Zoomed Rietveld refinement of X-ray diffraction pattern of the 0 h LTO (λCu). The black continuous line, red continuous line, represent the observed and calculated patterns, respectively. Vertical green tick bars are the Bragg positions. Orange arrows highlight TiO2 rutile peaks.

3. 8 :Figure S3. 6 :

 86 Figure S3.6: Williamson-Hall plots deduced from the Rietveld refinements of the LTO samples made from solution-combustion synthesis.

3. 10 :Figure S3. 7 : 3 . 11 :Figure S3. 8 :

 1073118 Figure S3.7: Optimization of the AM/SE wt. ratio in (a) LTO from NEI and (b) LTO from solid-state synthesis.

Figure S4. 5 :

 5 Figure S4.5: Galvanostatic cycling of C-free carbonate-coated NMC 622 (UMICORE, stored) composites, at C/10 in the [2.1-3.9 V vs. Li-In/In] potential window.

Figure S4. 6 :Figure S4. 7 :

 67 Figure S4.6: Electronic conductivity with one pellet done for each air treated NMC 622-based composites

4. 8 :

 8 Figure S4.11: (a) Schematic of the cycling protocol for formation cycles: 2 cycles are done at C/20 with distinct cut-off potential, namely 3.4, 3.6 and 3.9 V vs. Li-In/In. The lower cut-off potential is defined at 2.1 V vs. Li-In/In. The following cycles are done at C/2 in the potential window of [2.1 -3.9 V vs. Li-In/In]. (b) The resulting discharge capacities are plotted as function of cycle number, and compared with the pristine curves. Two cells are plotted per system with at least 100 cycles for each. (c) Capacity retention of the aforementioned cells.

S4. 17 .

 17 7 V vs. Li-In/In did not improve capacity retention for LYBC, but made it worse, with an average capacity loss of 25 % after 100 cycles, as shown in the Figure S4.16. Finally, Li3InCl6 exhibited a capacity of 160 mA.h/g at C/2 and an average capacity loss of 9 % over 100 cycles, which is slightly higher than our reference cells with Li6PS5Cl. We additionally tried Li3YBr6 (LYB) and Li3YCl6 (LYC) solid electrolytes provided by Saint Gobain, but they both display disastrous capacity, as depicted in Figure In summary, the capacity retentions obtained in this chapter are shown in Figure S4.18. The optimized NMC 811-Li6PS5Cl cells (shown in Figure 4.15) and the NMC 622 (LiNbO3)-Li6PS5Cl cells exhibited the best stability, with respectively 8 and 2.8 % capacity loss in 100 cycles. Moreover, they present a good C/2 capacity ranging from 140 to 160 mA.h/g.

Figure S4. 14 :

 14 Figure S4.14: Capacity retention at C/2 between 2.1 and 3.9 V vs. Li-In/In, after one cycle at C/10 in the 2.1-3.9 V vs. Li-In/In potential window. We use NMC 811 (MSE) as active material, and distinct solid electrolyte in a 70/30 wt. % ratio, with the addition of 1.5 wt. % VGCF. The solid electrolyte used are: Li6PS5Cl (NEI), Li3YBr2Cl4 (LYBC) sieved < 5 µm or not from Saint Gobain and Li3InCl6 (LIC) homemade synthesized. Two cells per system are shown. All cells were launched on the same week.

First cycle at C/10 of halide compositesFigure S4. 15 :

 15 Figure S4.15: First cycle at C/10 and up to 3.9 V vs. Li-In/In of two NMC 811 (MSE) + SE cells per system assembled in a 70/30 wt. % ratio and with an additional 1.5 wt. % VGCF. (a) LYB (b) LYBC (< 5 µm) (c) LYBC (d) LYC and (e) LIC.

Figure S4. 18 :

 18 Figure S4.18: Summary of the capacity retention obtained in the systems having a first cycle at C/2 capacity above 135 mA.h/g.

  Potentiostatic Electrochemical Impedance Spectroscopy TLM: Transmission Line Model TLMO: Transition Layered Metal Oxide GIIT: Galvanostatic Intermittent Titration Technique VGCF: Vapour Grown Carbon Fiber EV: Electric Vehicle XRD: X-ray diffraction ASSB: All-Solid-State-Battery SEM: Scanning Electron Microscopy CV: Cyclic Voltammetry

  Considering the non-faradaic case of an uncycled electrode (no charge transfer or solid-state diffusion), we place a Constant Phase Element (Q) to account for the blocking electrodes. Note that, Constant Phase Elements (CPE) are used to represent capacitive behaviour in battery systems, as real devices capacitive response displays a distribution of relaxation times127,128 .

Table 4 . 1 :

 41 Resulting capacity decays estimated by linear extrapolation for 100 cycles on NMC 622 and NMC 811

	/In	3.9 V vs. Li-In/In

Table 4 . 2 :

 42 Summary of the impact of each parameter on the stability of the cycling and the capacity.

	By combining the best parameters, our study resulted in a capacity loss of
	approximately 23-24 % after 300 cycles in cells undergoing 2 formation cycles and to 18-22 %
	after 200 cycles in cells undergoing 10 formation cycles, as demonstrated in Figure 4.15. Such
	results are somewhat disappointing, and there is no synergetic beneficial effects of the distinct
	optimized parameters. On the contrary, these results suggest that there is a detrimental
	interaction between at least two parameters.

  Overall, although these results are very interesting from a methodological point of view, they do not solve the problem of the stability of Li6PS5Cl-based cathode composites. This highlights the need to explore other ways of stabilizing ASSB cathode interfaces. As a first step, we present the performances of coated NMC 622 and halide-based composites in Appendix

	4.10. Despite the preliminary nature of our results, they manifest considerable promise in
	terms of both stability and capacity. Especially, by replacing Li6PS5Cl by Li3InCl6, we can
	increase the stability retention up to 91 % after 100 cycles, with an initial capacity of 160
	mA.h/g. Additionally, the usage of LiNbO3-coated NMC 622 yields to considerable stability
	improvement with 97.2 % capacity retention after 100 cycles. Thus, a transition towards halide

solid electrolytes and a meticulous design of active materials may represent a promising direction in the context of ASSB cathode composites.

Table S1 .

 S1

							Prefactor
		R ion,0°C	R ion,10°C	R ion,25°C	R ion,30°C	R elec	(C)	Ea
	3.6 V VGCF	9638,6	5052,5	2079,8	1577,7	1,0	9088,6	0,43
	3.9 V VGCF	712,6	432,3	217,5	175,6	2,6	1969,5	0,33
	3.9 V No VGCF	429,2	272,1	145,4	119,7	50,1	938,3	0,30
	3.6 V No VGCF	201,2	120,7	59,6	47,9	52,8	9871,0	0,34

Table S1 :

 S1 Results of the optimization program to decorrelate R ion and R elec from R global.

wt. % excess reheated at 900°C (2h)

Acknowledgements

Slow cycling and temperature effects

The formation stage of liquid Li-ion batteries encompasses multiple possible protocols involving high temperature, slow cycling, pulses... Of interest for this study, literature suggests the application of slow charge and discharge rates (ranging from 3 to 5 cycles at C/20), followed by cycles at elevated temperatures as a mean to stabilize the battery interfaces 163,164 .

Despite the ongoing exploration of faster and more efficient protocols, this method is known to generate a robust solid electrolyte interphase, thereby enabling stable cycling performance.

In this context, we have adopted a similar methodology to investigate the combined influence of temperature and slow cycling on the capacity retention of all-solid-state batteries.

First, the cells undergo two cycles at C/20 at temperatures ranging from 0°C to 55°C. Subsequently, we assessed the following C/2 cycling at ambient temperature. Switching to low temperature (0°C), the capacity at C/2 is around 165 mA.h/g, but an enhanced capacity decay is found, reaching 25 % loss in 100 cycles. Overall, there was no improvement in performance in either high-or low-temperature formation cycles. inclines us to think that the process of stabilization achieved through ten formation cycles might be reproduced with prolonged cycling at C/2. Further studies are necessary to fully unravel this behaviour.

Potential window influence

We initially start by following the methodology outlined in Figure 4.13.a, to tackle the impact of upper cut-off potential. This involves applying two formation cycles at C/20, utilizing distinct cut-off potentials (3.4 V, 3.6 V, and 3.9 V). Subsequently, normal cycling was performed at C/2 within the potential window of [2.1 -3.9 V vs. Li/In/In]. The resulting moving average discharge capacity and corresponding capacity retention evolutions are presented in Note that certain data points were lost due to a software issue on the related computer. Overall, our findings demonstrate that high potential formation cycling has a detrimental impact on capacity retention. Cells subjected to 3.6 V and 3.9 V cut-off potentials exhibit high capacity decays with capacity losses reaching 15.5 % and 18.5 % in 120 cycles, respectively. In contrast, two formation cycles up to 3.4 V vs. Li-In/In demonstrate similar stability as our reference cycles, with 12.5 % capacity losses in 120 cycles. However, the origin of such stability with low cut-off potential during formation cycle is questionable. Indeed, we might just observe the detrimental impact of longer time at high potential, as mirrored by our previous experiments in Figure 4.12. Thus, further studies are necessary to comprehend this phenomenon. Interestingly, the use of a 2.5 V low cut-off potential in the formation cycles led to disastrous capacity decay during C/2 cycling, with capacity losses reaching up to 17.5% within 100 cycles. On the contrary, reducing the low cut-off potential to 1.5 V or 1.8 V vs. Li-In/In resulted in increased stability with capacity decay of 8 %. However, further decreasing the low cut-off potential to 1.3 V led to a decrease in capacity retention, reaching values of 9.5% capacity loss within 100 cycles. We can hypothesize that this positive impact of low cut-off potential might be linked to the reduction of the degradation products formed with Li6PS5Cl intrinsic decomposition. Indeed, S and P2S5 are reported to reduce below 1.8 V vs. Li-In/In [START_REF] Tan | Elucidating Reversible Electrochemical Redox of Li 6 PS 5 Cl Solid Electrolyte[END_REF] . All in all, the low and high cut-off potentials are very interesting parameters that strongly influence stability, while not having any effect on the capacity. The best performances should arise from cells that underwent formation cycles between 1.8 V and 3.4 V vs. Li-In/In. For this reason, we have embarked on a systematic exploration aimed at solving the problem of the stability within Li6PS5Cl-based composites containing a minimum carbon content (1.5 wt. %). Our research was designed to improve the stability of composites composed of NMC 811, Li6PS5Cl, and 1.5 wt. % VGCF, when cycled over a voltage range spanning from 2.1 to 3.9 V vs. Li-In/In. In this comprehensive experimental design, we examined eight key parameters of interest, encompassing assembly parameters (assembly pressure and loading), coating strategies (including the application of carbonate coatings on both the solid electrolyte and the active material), as well as specific formation cycles attributes (such as temperature, number of formation cycles, and high and low cut-off potentials). By careful optimizing each parameters, we were able to achieve optimized stability with a capacity at C/2 of 140-160 mA.h/g and a capacity retention of 92 % after 100 cycles.

Collectively, the ideas presented in this manuscript contribute to a better understanding of the interfaces and critical parameters that influence the performance of Li6PS5Cl-based cathode composites. In addition, this research introduces a novel experimental setup, expanding the methodological toolbox for the study of All-Solid-State Batteries. We expect these results to enrich the experimental landscape of ASSB research, offering new methodological approaches and avenues for further experimentation.

Ultimately, our investigations have led to improvements in the stability of systems with significant capacity for ASSBs. Nevertheless, it is essential to recognize that while these while these improvements are notable, they fall short in addressing the challenge of capacity retention over long cycling. Thus, given the complex and multifaceted nature of the interfacial reactions inherent to Li6PS5Cl, these results underscore the imperative need for a transition toward more stable solid electrolytes. More specifically, halide-type solid electrolytes appear to be promising alternatives, offering the possibility of enhanced stability. capacity and capacity retention. In cases where repeatability is not achieved, it indicates discrepancies in the assembly procedure that warrant immediate attention and rectification.

-Secondly, when initiating any new experiment, diligently undertake two cells for the given investigation concurrently, rather than deferring the second one for a later time. This approach ensures a more robust experimental dataset and minimizes the risks associated with uncertainties in data interpretation.

By adhering to these prudent guidelines, researchers can navigate the ASSB research with greater confidence and methodological rigor, ultimately enhancing the reliability and validity of their findings.

Appendix Appendix for chapter 2 2.1: Protocol to assemble ASSB cell

Li-In/In composite having the 2/1 atomic ratio was prepared via a two steps route. Firstly, the proper amount of Lithium foil was enveloped in an Indium foil. Secondly, the foil was laminated and folded several times with a glass tube until the Li-In alloy became darker and brittle. Finally, Li6PS5Cl powder is added in a mortar and hand-ground for 10 minutes to ensure optimum homogeneity to form the Li0.5In/SE composite (60:40 wt. % ratio). The composite anode has a constant potential of 0.62 V vs. Li + /Li. Lithium metal was not used since interfacial and dendrite issues prevent its usage. Besides, it cannot withstand a 1 t/cm² upon cycling without short-circuiting the cells.

Cathode composites are prepared by weighting the desired amount of active material and

Li6PS5Cl solid electrolyte in the glovebox, to reach a ratio of 70/30 wt. % and hand-ground for 10 minutes to make composites. Occasionally a defined amount of vapor-grown carbon fibers (VGCF) was added to the mixture, after weighting the 70/30 wt. % AM/SE. Two-electrode cells assembly is made via a three steps process using a hardware developed in house that consists of a PEI body and two stainless steel pistons of 0.5 cm² surface (8 mm diameter) (a scheme of the cell is shown in Figure S2.1). Herein, if not otherwise specified, 30 mg of Li6PS5Cl as ionic separator are first pressed at 0.5 t/cm².Subsequently, 25 mg of anode (Li0.5In/Li6PS5Cl 60:40 wt. %) and 8 -9 mg of cathode composites (loading of 16 mg/cm²) are placed on both sides and the stack is then pressed at 4 t/cm². Finally, a pressure of 1 t/cm² is imposed during cycling by the mean of six screws.

Instead of fitting spectra with Relec and Rion values that does not make sense, we choose to consider an Relec value equal to zero, and replace R ion by an R global (Figure S2.9). Equalizing the two transfer functions (when ZC = 0 or when ZC = Relec) and knowing that the charge transfer related values (RCT, QCT and αCT) will be equal in both TLM representations, we found the following relationship between the three resistances:

In Figure S2.9, the 45° slope originating from the TLM is having the same tendency as in the previous case, as it grows with the increase of R global. It is not surprising, as this scenario is similar to one of a cathode composite containing enough carbon additive so that the electronic conductivity becomes negligible (and therefore R global would equalize R Ion).

(which is actually another transmission line model) in parallel with RCT enables to mimic the low frequencies, but one additional feature is missing Figure S2.11.b.

To account for such difference we added another R//CPE in series with (RCT + W)//CPECT, which could represent an additional layer arising from side reactions around NMC particles.

This new feature will be called Rint//CPEint. We therefore succeed to model efficiently our cathode composite spectra

Case without carbon additive:

In system without VGCF, an additional feature is seen at high frequency and does not evolve upon cycling. For the fitting procedure, this feature is model by an R//Q prior to cycling, and the values are held constant for the following fits. This particularity has already been

2.8: Detailed explanations on RIon and RElec decorrelation from RGlobal :

To enable a decoupling of Relec and Rion contribution to Rglobal, we proceeded to resolve a system of nonlinear equations, playing on Rion characteristic dependence with temperature.

The electronic conductivity of the cathode composites is assumed constant versus temperature. This system uses two types of equation that will be described in the following sections. To fasten computation, we only fit the high/mi-frequencies part of the spectra, not taking the solid-state diffusion and the extra interfacial semi-circle features.

Four temperatures are took per spectra: 0 °C, 10 °C, 25 °C and 30 °C.

Equalizing transfer function models: Assuming temperature change has negligible impact on Relec, we have 4 equations for 5

unknowns (Relec, Rion,0°C, Rion, 10 °C, Rion, 25 °C and Rion, 30 °C). There is a need for more equations to solve this system.

Arrhenius evolution of the ionic conductivity:

Further equalities can be obtained from the dependence of Rion with temperature. We assume the ionic resistance of the composite to follow an Arrhenius behaviour:

Appendix for chapter 3

4.10: Exploring new paths to enhance the stability of ASSB composites

We decided to explore the impact of two coatings made by industrials and of halide solid electrolytes on the stability of ASSB cells. mill apparatus was used, with 6 milling process set at 510 rpm for 24 h, applying 15 min milling followed by 15 min of cooling step. After the ball-milling process, the obtained powder was pelletized and annealed at 260 °C (heating rate 5 °C/min) for 5 h in a quartz tube sealed under vacuum. After the annealing process, the sample was cooled naturally.

Starting

Capacity retention of LYB and LYC based composites

Cell assembly:

The protocol for the two-electrode cell assembly for halide-based composites follow our typical procedure, with the exception that both the solid electrolyte and anode layer is composed of LYBC. This choice is driven by the well-known reactivity of LIC and LYB at the anode side,

LYB: Li3YBr6

LYC: Li3YCl6

LYBC: Li3YBr2Cl4

LIC: Li3InCl6

Abstract

While Lithium-ion batteries dominate portable devices, growing safety and energy density demands in electric vehicle batteries have led to the exploration of "beyond Li-ion" technology. All-Solid-State Batteries (ASSBs) have emerged as a promising alternative to Li-ion batteries. Thus, this doctoral research focuses on overcoming challenges hindering the practical implementation of ASSBs, with a specific emphasis on cathode composites. The investigation revolves around a common composite comprising Li6PS5Cl solid electrolyte (SE) and NMC active material (AM). The research unveils the degradation mechanisms within ASSBs, governed by SE/Carbon additive and SE/AM interfaces. It is observed that capacity deterioration, occurring below 3.6 V vs. Li-In/In, is primarily attributed to SE/Carbon interfaces. Conversely, elevating the voltage to 3.9 V shifts the primary degradation source to SE/AM interfaces. Then, the adverse effects of carbon additives on the ionic conduction of composites are demonstrated, particularly when exceeding 2 wt. % VGCF. Moreover, the study delves into the electronic conductivity of carbon-free composites using innovative in situ monitoring. This reveals Li-induced alterations hindering electronic conductivity, especially at high charge levels, notably in high Ni-content NMC. Furthermore, the influence of particle size and morphology on electronic percolation is extensively examined, advocating for minimal VGCF to enhance kinetics and stability. Strategies for effectively incorporating carbon additives while mitigating long-term capacity loss are explored, encompassing assembly pressure, loading, formation cycles, temperature, and carbonate coating. By mixing these optimal conditions, an enhanced cathode composite is introduced, holding promising potential for the progression of All-Solid-State Battery technology.

Keywords: Solid-state, Interfaces, Impedance spectroscopy, Argyrodite, Electronic conductivity