The combustion of fossil fuels is scientifically recognized for releasing greenhouse gases, primarily carbon dioxide, initiating a complex chain reaction involving climate change, amplified global temperatures, and increased frequency of severe climatic events. 2 This combustion also leads to air pollution through emissions of sulfur dioxide, nitrogen oxides, and particulate matter. 3 The finite nature of fossil fuel reserves raises concerns about long-term energy security, potentially leading to geopolitical uncertainties and supply chain disruptions. 4 Additionally, fossil fuel dependence introduces economic volatility due to fluctuating fuel prices, impacting various sectors of society. The transition towards renewable energy sources, such as solar, wind, and hydroelectric power, emerges as a critical imperative to address these pressing challenges.

As the global transition towards sustainable energy gains momentum, increasing attention is being directed toward renewable sources for electricity generation (as illustrated in Figure 1.b). However, the inherent intermittency and variability characteristic of most renewable energy systems presents significant challenges. 5 For instance, the unpredictable nature of solar and wind resources necessitates the establishment of mechanisms to capture excess energy during periods of peak production and distribute it during periods of low generation. In response to this challenge, the integration of energy storage solutions emerges as a critical necessity.

The concept of energy storage systems relies on the conversion of energy into a storable form that can be later converted back into electricity when required. Several diverse systems exist for energy storage, including chemical storage through battery energy storage systems, 6 harnessing hydropower through pumped hydroelectric storage, 7 employing compressed air energy storage, 8 utilizing flywheel kinetic energy storage, 9 implementing thermal energy storage, 10 tapping into hydrogen storage, 11 and adopting hybrid systems that synergize various technologies. Notably, the combination of different technologies, such as the integration of batteries with flywheels or supercapacitors, makes it possible to optimize strengths and compensate for limitations in the energy storage landscape. Among these energy storage technologies, batteries are one of the pioneering energy storage technologies. More especially, lithium-ion (Li-ion) batteries are widely employed in portable electronics, electric vehicles, and grid-scale applications [12][13][14][15][16][17] . They can efficiently store and discharge electrical energy (= secondary battery), significantly contributing to modern energy systems. 18 At the heart of the Li-ion battery are three essential elements: the negative electrode, the positive electrode, and the electrolyte. The electrolyte orchestrates the migration of lithium ions between the negative and the positive electrode during the charging and discharging cycles. The negative electrode is primarily constituted of graphite. Charging triggers the extraction of lithium ions from the positive and their insertion into the graphite electrode. Diverse materials constitute the positive electrode such as lithium cobalt oxide (LiCoO2), lithium iron phosphate (LiFePO4), and lithium manganese oxide (LiMn2O4) emerging as prominent choices. 18 Discharging involves the migration of lithium ions from the negative electrode back to the positive side through the electrolyte. The electrolyte, of paramount significance, facilitates the lithium-ion flow between the negative and positive. It commonly consists of a solvent in which lithium salts are dissolved. 19 In particular, the electrolyte must enable the mobility of ions while avoiding the development of dendritic formations that could cause short circuits in the battery. Li-ion batteries extend several merits, catalyzing their widespread adoption such as high energy density, exhibiting negligible selfdischarge rates, and cycle longevity. 19 Despite their merits, Li-ion batteries face a certain number of challenges. Safety is probably one of the biggest concerns, as internal short circuits or overcharging can trigger thermal runaway incidents, leading to fires or battery explosions. 20 Additionally, these batteries depend on materials such as lithium and cobalt, resources whose availability is subject to constraints and ethical concerns linked to extraction practices. 21 Over time, Li-ion batteries lose capacity, which reduces their effectiveness as energy storage. 19 Additionally, the disposal and recycling of batteries necessitate improved practices to mitigate their environmental impacts.

A straightforward approach to solving the flammability issues encountered by Li-ion devices is to replace the organic electrolyte commonly used in these systems with a safer, nonflammable aqueous electrolyte.

The objective of the Ph.D. and manuscript outline

The main objective of this thesis is to explore and develop strategies that enhance the reactivity of titanium oxide concerning proton insertion, making it a viable candidate for use as a negative electrode in aqueous battery conditions. To achieve this objective, we aim to examine the influence of cations inserted within the confined water environment of a titanate-layered structure. Our focus is to understand the impact of these cations on proton transport properties, and then to assess their implications for proton capacity and reversible insertion. Layered materials have distinct advantages, making them very promising as electrode components.

These materials offer a high surface area, facilitating enhanced capacity and improved overall performance. 22 Their layered arrangement promotes efficient movement of charges and ions, making them ideal for high-power output. The easy diffusion paths for ions and electrons within the layers facilitates electrochemical reactions. Additionally, their inherent capacity to accommodate to volume fluctuations during cycling further enhances their relevance. The ability to tailor these materials to specific requirements and their compatibility with additives contribute to obtain specific properties. Our team has previously demonstrated the successful synthesis of nanostructured layered titanate compounds via soft chemistry methods, including the effective incorporation of chemically inserted Zn 2+ leading to an increase by a factor of 450 the proton bulk conductivity compared to the pristine layered titanate sample. 23 Another route for enhancing proton intercalation capacity involves investigating the effects of cationic vacancies within a three-dimensional (3D) framework. 3D open frameworks exhibit substantial promise as electrode materials in energy applications due to their inherent advantages. These frameworks also provide a notable surface area, translating into increased capacity and increased performance. 24 Interestingly, vacancies in their structure facilitates the rapid movement of charges and ions, making them also suitable for high-power applications. 25 Additionally, the short ion/electron diffusion paths facilitate the electrochemical reaction and the three-dimensional (3D) framework offers mechanical stability. This latter ensures longevity and extended operating cycles, making them very promising as electrodes for batteries.

Interestingly, recent research in our group has highlighted the favorable intercalation behavior of cations such as Li + , Mg 2+ , or Al 3+ within these vacancies. 26 We can therefore anticipate a similar behavior concerning proton intercalation.

The manuscript comprises five chapters:

In Chapter I, we provide an overview of aqueous batteries, their advantages, and drawbacks. Additionally, we looked into TiO2 and its use in electrochemical devices and explored its various crystalline forms for proton insertion. This allows us to introduce the mechanism of proton-electron transfer within TiO2 materials. The chapter then looks at the Hydrogen Evolution Reaction (HER) and its kinetics with TiO2. The focus is on sol-gel methods for synthesizing TiO2 nanomaterials, as well as strategies for developing titanate-based electrodes.

Chapter II focuses on the synthesis and characterization of diverse TiO2 materials serving as negative electrodes for proton intercalation. Among them are nanostructured layered titanate compounds (L-TiO2) hosting nanoconfined water, a TiO2 sample with titanium defects in a 3D structure (Vac-TiO2), and a defect-free 3D TiO2. These samples are prepared using solgel methods, and a detailed microstructural analysis is conducted using various techniques including PDF, X-ray diffraction in temperature and at ambient temperature, TG analysis, exsitu and in-situ HR-TEM.

In Chapter III, we study the impact of nanoconfined water composition on proton electrochemical properties of lepidocrocite type TiO2 (L-TiO2). To do so, we tailor a layered titanate sample by chemically incorporating and stabilizing Zn 2+ cations within its nanoconfined water. The subsequent characterization of these compounds with added zinc ions explores their influence on proton transport properties using BDS and electrochemistry tools. Further, the electrochemical performance of zinc incorporated L-TiO2 is assessed in a half-cell setup.

In Chapter IV, our investigation centers on the advantageous effects of cationic vacancies on electrochemical performance. More specifically, we address the central issue of optimizing proton capacity through defect engineering. Our objective is to analyze the process of proton intercalation and diffusion through cyclic voltammetry, while also examining the stability and rate performance using galvanostatic techniques.

Chapter V aims to establish a complete cell utilizing these electrodes, by discerning the most relevant parameters-composition of the interlayer or structural defects-for the intercalation of proton in a battery environment. To build the full cell, we employ homemade electrospun carbon nanofibers, developed by our team, as the positive electrode that serves as a current collector and support for the conversion of MnO2/Mn 2+ . This chapter is devoted to an in-depth study of the electrochemical properties of this battery configuration.

Chapter I. State of the Arts

I. Aqueous battery context

Aqueous batteries have undergone substantial development, with the lead-acid battery, invented by French physicist G. Planté in 1859, representing the first rechargeable instance. 1 This aqueous-based system comprises lead dioxide (PbO2) positive plates, lead (Pb) negative plates, and a diluted sulfuric acid (H2SO4) electrolyte. During discharge, a reaction forms lead sulfate (PbSO4) and water, while during charge, the reverse reaction reverts lead sulfate to lead dioxide and lead, accompanied by the release of oxygen and hydrogen gases. 1 Lead-acid batteries have a relatively lower energy density, typically ranging from 30-50 Wh.kg -1 . 2 They are commonly used in automotive starter and ignition systems to start the engine and for backup power applications. This seminal lead-acid battery serves as a foundation for the development of aqueous battery technologies.

Aqueous batteries have emerged as a significant contender in the landscape of energy storage technologies, presenting a compelling alternative to conventional choices such as lithium-ion batteries. Regarding water-based electrolytes, these batteries offer distinct advantages that resonate with the demands of contemporary energy systems and sustainability objectives. 3 Notably, one of their paramount strengths lies in safety. Aqueous electrolytes are not prone to flammability when compared to certain non-aqueous counterparts frequently employed in lithium-ion batteries. 4 Furthermore, aqueous batteries exhibit the potential for costeffectiveness. 5 Water-based electrolytes are abundant and economically accessible, contributing to the realization of more economically viable battery solutions. The utilization of readily available materials can yield a reduction in manufacturing expenditures. 6 Equally significant is their environmental compatibility. These applications span from compact devices like portable electronics to more extensive grid-scale energy storage systems. 3 Beyond leadacid batteries, other notable aqueous battery variants also command prominence in the energy storage landscape:

Nickel-Cadmium (Ni-Cd) Battery: Invented in 1899, consists of a nickel oxyhydroxide (NiOOH) positive electrode, a cadmium (Cd) negative electrode, and an alkaline potassium hydroxide (KOH) electrolyte. 7 During discharge, the NiOOH electrode transforms into nickel Chapter I. State of the Arts hydroxide (Ni(OH)2), while the cadmium-hydrogen adsorbed electrode transforms into Cd.

This system constitutes the first reported battery with proton (de)insertion in strong basic condition. Advantages include good energy (40-60 Wh.kg -1 ), 8 durability, wide temperature range, and long cycle life (1500 cycles). 8 However, cadmium's toxicity poses environmental, and health concerns, potential "memory effect" and self-discharge.

Nickel-Metal Hydride (NiMH) Battery: Invented in 1967, consists of an improvement over Ni-Cd batteries, NiMH batteries present higher energy density typically ranging from 40-90 Wh.kg -1 , 2 and enhanced environmental compatibility, as they eschew toxic cadmium. 7 The NiMH batteries consist of a positive electrode containing nickel oxyhydroxide (NiOOH), a negative electrode with a hydrogen-absorbing alloy, and an alkaline KOH electrolyte.

Advantages include good energy density, eco-friendliness, memory effect reduction, and low self-discharge. Drawbacks include lower energy density than lithium-ion batteries, limited charge efficiency, and lower cycle life (300-500 cycles). 8 Zinc-Manganese Dioxide (Zn-MnO2) Battery: The prototype of a rechargeable Zn-MnO2 battery was introduced by Yamamoto in 1986. 9 It consists of a Zn negative electrode and a mixture of black carbon and EMD-γ-MnO2 (Electrolytic Manganese Dioxide) as the positive electrode in a ZnSO4 electrolyte. During discharge, Zn 2+ is formed at the negative electrode and MnO2 can be electro-dissoluted into Mn 2+ , converted into Mn2O3, or MnOOH at the positive.

Advantages include lightweight design, cost-effectiveness, long cycle life (5000 cycles), 10 and wild energy density, generally ranging from 70-400 Wh.kg -1 . 11 Their energy density is influenced by the specific battery design and electrolyte composition. It's important to note that these energy densities can vary depending on factors such as electrode materials, electrolyte composition, and overall battery design. The drawbacks are lower energy density compared to lithium-ion batteries.

Aqueous Sodium-Ion Battery: These batteries have recently garnered attention and utilize sodium ions as carriers of electric charge. 12,13 A variety of positive and negative materials can be employed in aqueous sodium-ion batteries. Some promising options for positive materials are layered oxides, Prussian blue analogs (PBAs), and polyanions. 14 On the negative side, a notable material is NaTi2(PO4)3, which belongs to the Na superionic conductor category.

This material claims a high theoretical capacity, an open three-dimensional structure, and a suitably low redox potential. 13 The strengths of aqueous sodium-ion batteries encompass the abundant and cost-effective nature of sodium, their environmental friendliness as they avoid toxic or rare elements, and their safety due to the use of water-based electrolytes. Nonetheless, these batteries encounter challenges including lower energy density (50-80 Wh.kg -1 ), limited cycle life (~500), and voltage restrictions. 15 Finally, the typical power density for Li-ion batteries ranges from 1-10 kW.kg -1 which is quite important compared to aqueous batteries. 2 For instance, Ni-MH Batteries have a power density of 0.5-3 kW.kg -1 while Lead-acid batteries generally exhibit lower power density (typically 0.1-0.7 kW.kg -1 ). 2 The most competitive one is the Zn-MnO2 batteries series, they exhibit a power density comprise between 0.1 to 10 kW.kg -1 . 11 Aqueous batteries, despite their many benefits, are constrained by inherent limitations that influence their practicality and potential uses. A notable drawback of aqueous batteries lies in their comparatively lower energy density compared to non-aqueous options. This divergence largely arises from the narrower electrochemical stability window (ESW) they possess, as illustrated by the stability of pure water at 1.23 V (Figure 1.1). 3,12,16,17 6 When surpassing the potential stability range of the electrolyte, the aqueous electrolyte experiences the occurrence of dihydrogen formation through the Hydrogen Evolution Reaction (HER) at the negative electrode, and the generation of dioxygen through the Oxygen Evolution

Chapter I. State of the Arts Reaction (OER) at the positive electrode. These reactions are commonly referred to as watersplitting reactions, and their progression is contingent upon the pH level: 3 In acidic solution:

OER: 2H2O + 4H2O → O2 + 4H3O + + 4e -(eq. 1.1)

HER: 2H3O + + 2e -→ H2 + 2H2O (eq. 1.2)

In basic solution:

OER: 4OH -→ O2 + 2H2O + 4e -(eq. 1.

3)

HER: 2H2O + 2e -→ H2 + 2OH -(eq. 1.4)

Consequently, the limited electrochemical stability range of aqueous electrolytes constrains the range of suitable materials. Furthermore, the longevity of certain aqueous battery configurations can be undermined, largely attributable to material dissolution and degradation resulting from the HER and OER processes during battery operation. 16 This leads to shorter lifespans in comparison to non-aqueous battery technologies.

Nevertheless, aqueous batteries continue to be a dynamic realm of research in the field of energy storage. Considerable focus is directed toward metallic ions serving as charge carriers, including Li + , Na + , K + , Mg 2+ , and Ca 2+ . [18][19][20][21] These ions are intriguing prospects due to their potential for high volumetric capacity through involvement in multielectron transfers, and/or their abundance in Earth's crust. However, these ions suffer from a sluggish solid-state diffusion and high hydration energy of multivalent ions. 12,14,22 The process of solid-state diffusion encounters obstacles due to two specific types of electrostatic interactions involving the guest ions and the anions/cations present in the host material. 23 Multivalent ions exhibit strong interactions with the anionic framework, resulting in a reduction of the potential energy of the intercalation electrode and the formation of robust bonds between guest cations and anions (localized ionic bonds). 23 Another crucial factor is charge redistribution. 23 When guest ions are incorporated, it necessitates adjusting the oxidation state of the transition metal hosts to maintain charge balance. This adjustment becomes especially complex for multivalent ions, which require multi-electron redox reactions. However, this type of reaction is thermodynamically unfavorable and often triggers pronounced local distortions in the material's structure.

Yet, non-metallic charge carriers have received less attention, they can be categorized into anionic species (such as OH⁻, F⁻, Cl⁻) and cationic species (like H⁺, NH₄⁺). 24 A notable distinction between metallic and non-metallic charge carriers lies in their manner of interaction upon insertion into the active material. The metallic charge carriers make localized pure ionic bonds whereas cationic non-metallic charge carriers can establish hybrid delocalized covalentionic bonds. 24 Non-metallic ions such as proton-based aqueous batteries appear to have several advantages compared to metal-ion aqueous batteries. The concept of proton-ion battery is not novel and has been already experimented with success since 1859 in lead-acid batteries (Figure

1.
2.e). The use of protons as charge carriers offers clear advantages compared to other existing chemistry in batteries:

i/ Hydrogen as an element is ubiquitous on Earth, with a much greater abundance (≈1400 ppm) than, for example, Li (20 ppm) which promises the sustainability and affordability of proton batteries (Figure 1.2.a). 13 ii/ The proton has the lowest molar mass of 1.0 g mol -1 , which minimizes the mass of charge carriers in the entire battery (Figure 1.2.b).

iii/ The proton, with or without its solvation sphere, has the smallest radius, which should minimize stress or volume expansion during intercalation and de-intercalation (Figure

1.2.c

). 1,25,26 iv/ The diffusion mechanism appears to often be Grotthuss type, similar to a Newton cradle (Figure 1.2.d). 27 On the positive side, metal oxides, such as MnO2, have emerged as promising materials capable of undergoing electro-dissolution/precipitation processes. This is attributed to the presence of protons before the initiation of the OER. 28 Recently, our group conducted a study on Zn||MnO2 batteries, investigating the discharge mechanism of low-valence MnOx electrodes in aqueous electrolytes. 29 The results revealed that the mechanism involves a proton-coupled electrochemical dissolution of the active material. This mechanism, similar to the established process for high-valence MnO2 electrodes, contradicts previously proposed models involving reversible cation or monovalent ion insertion/intercalation. This further underscores that, regardless of Mn's oxidation state (2 < Mn's average oxidation state ≤ 4) or crystallinity degree, the discharge mechanism uniformly follows a combined proton-coupled electrodissolution reaction, with protons sourced from the electrolyte.

Furthermore, the OER, known for its intricate four-electron process (eq. 1.1), frequently requires a significant overpotential to facilitate its advancement. This requirement stems from the intricate kinetics associated with this reaction, which can be attributed to two fundamental mechanisms: the Adsorbate Evolution Mechanism (AEM) and the Lattice-Oxygen-Mediated Mechanism (LOM). 30 The prevalence of these mechanisms depends on whether the reaction involves a single metal ion or the interaction between two adjacent metal ions. Elevated activation energies, combined with low energy conversion efficiencies, exacerbate the required overpotentials. This need for significant overpotentials is essential for expanding the Electrochemical Stability Window (ESW). Specifically, in the domain of lead-acid batteries, the water-splitting reaction at the lead (Pb) electrodes encounters sluggish kinetics. [30][31][32] As a result, it becomes feasible to achieve an elevated operational voltage of approximately 2.1 V.

On the negative side, materials that can be used for inserting proton are organics, 33,34 Prussian blue analogs, 35 Metal-Organic Frameworks (MOFs), 36 MXenes, 37 PolyOxoMetalates (POMs), 38 and transition metal oxides. 14,[39][40][41][42][43] The exploration of the latter for proton insertion has emerged only recently. Notable examples among the diverse set of metal oxides include WO3.0.6H2O, 44 H2W2O7, 45 MoO3, [46][47][48] H3+xIrO4, 49,50 V2O5, 51 VO2, 52,53 H2Ti3O7, 54 (H3O + )1.35Ti3O6.67□0.35.1.73H2O, 55 TiO2, 41 amorphous TiO2. 39,40 Among this short list of compounds, only Titanium dioxide exhibits the lowest environmental impact. 56 TiO2 stands out as an environmentally favorable option owing to its abundant occurrence in nature and its high recyclability. Its extraction process is generally less reliant on toxic chemicals in comparison to certain other metals. This aspect contributes to its reduced environmental footprint. Conversely, metals like vanadium, 57 tungsten, 58 molybdenum, 59 and iridium tend to exhibit elevated environmental impacts (Figure S1.1). 60 This can be attributed to various factors, such as the intricacies involved in their extraction processes, habitat disruption caused by mining activities, and the potential contamination of water and soil. The processing of these metals often requires the use of chemicals, which, if not properly managed, can give rise to environmental pollution issues. Furthermore, certain elements are situated in area known for their rich biodiversity and fragile ecosystems, amplifying the potential environmental consequences and the vulnerability of these delicate ecosystems. This can further magnify their overall environmental impact.

The phenomenon of proton insertion in TiO2 raises fundamental questions. Before discussing these questions, it is important to note that some TiO2 material exhibit nanoconfined water within their layered structure, while others are identified as hydrogen titanates or cationic deficient titanates. In this PhD work, our objective is to investigate the impact of proton transport in both layered structures containing nanoconfined water and 3D open TiO2 materials with vacancies.

II. Titanium dioxide for electrochemical devices

Titanium dioxide (TiO2) is a versatile and widely studied material that has found various applications. 56 It exists in three main crystalline forms or polymorphs: anatase, rutile, and brookite depending on the arrangement of the TiO6 octahedra within the structure. TiO2 is a wide-bandgap semiconductor material with excellent optical and electronic properties. It is known for its high chemical stability in a wild range of pH, non-toxicity, and low cost. 56,61 Rutile, characterized by its tetragonal crystal structure within the P42/mnm space group, has a band gap of 3.0 eV, with the distinction of being the thermodynamically favored phase of TiO2. 56 It can be obtained at elevated temperatures, and often other phases of TiO2 undergo transitions into the rutile phase when heated. 62,[START_REF] Manzoli | a Sometimes under Evaluated TiO2polymorph[END_REF] In this compound, the interconnectedness of TiO6 octahedra is made through the sharing of two edges. A noteworthy attribute of rutile lies in its remarkable thermal stability. This stability, coupled with its comparatively heightened electronic conductivity relative to anatase positions rutile as a valuable candidate for applications such as solid oxide fuel cells and specific types of sensors. [START_REF] Reich | Niobia Based Rutile Materials as SOFC Anodes[END_REF][START_REF] Bai | Titanium Dioxide Nanomaterials for Sensor Applications[END_REF] Brookite exhibits an orthorhombic crystal structure (Space group: Pbca) and features a band gap ranging from 3.1 to 3.4 eV. [START_REF] Manzoli | a Sometimes under Evaluated TiO2polymorph[END_REF] Within its structure, the TiO6 octahedra are interconnected through the sharing of three edges. With an energy band gap that falls between those of anatase and rutile, brookite occupies an intermediate position in terms of its electronic properties. This characteristic renders it potentially well-suited for a range of applications involving photocatalysis and solar energy conversion. [START_REF] Kandiel | Brookite versus Anatase TiO2 Photocatalysts: Phase Transformations and Photocatalytic Activities[END_REF] Its bandgap energy makes it a candidate for initiating light-driven reactions, including water splitting and pollutants' degradation processes. Brookite's attributes start to extend its applicability in electrochemical devices such as batteries, [START_REF] Zhang | Brookite TiO2 Mesocrystals with Enhanced Lithium-Ion Intercalation Properties[END_REF] supercapacitors, [START_REF] Devan | Highly Stable Supercapacitive Performance of One-Dimensional (1D) Brookite TiO2 Nanoneedles[END_REF] and even nanosensors. [START_REF] Di Paola | the Least Known TiO2 Photocatalyst[END_REF] Anatase has a tetragonal crystal structure (space group: I41/amd) and has a band gap of 3.2 eV. [START_REF] Katal | A Review on the Synthesis of the Various Types of Anatase TiO2 Facets and Their Applications for Photocatalysis[END_REF] Within this compound, the TiO6 octahedra interact by sharing edges along four sides.

While anatase is the less stable thermodynamic phase compared to other polymorphs, it shows the lowest surface energy. [START_REF] Manzoli | a Sometimes under Evaluated TiO2polymorph[END_REF][START_REF] Fernández-García | Anatase-TiO2 Nanomaterials: Morphological/Size Dependence of the Crystallization and Phase Behavior Phenomena[END_REF] This unique property simplifies the synthesis process for producing nanoparticles. Anatase's remarkable attributes lie in its superior photocatalytic properties. [START_REF] Katal | A Review on the Synthesis of the Various Types of Anatase TiO2 Facets and Their Applications for Photocatalysis[END_REF] Its notably wider bandgap facilitates the absorption of a substantial portion of ultraviolet (UV) light. This photonic characteristic renders anatase highly prized for various applications. Among them, it plays a pivotal role in dye-sensitized solar cells (DSSCs) as the photoanode, harnessing its light-absorbing prowess to efficiently convert sunlight into electricity. [START_REF] Fazli | Dye-Sensitized Solar Cell Using Pure Anatase TiO2 Annealed at Different Temperatures[END_REF] Additionally, the usefulness of anatase extends to energy storage technologies, particularly lithium-ion batteries, which have been the subject of extensive research. [START_REF] Kanamura | Diffusion of Lithium in the TiO2 Cathode of a Lithium Battery[END_REF][START_REF] Cava | The Crystal Structures of the Lithium-Inserted Metal Oxides Li0.5TiO2 Anatase, LiTi2O4 Spinel, and Li2Ti2O4[END_REF][START_REF] Van De Krol | Spatial Extent of Lithium Intercalation in Anatase TiO2[END_REF][START_REF] Ma | Lithium Intercalation in Anatase Titanium Vacancies and the Role of Local Anionic Environment[END_REF][START_REF] Wagemaker | Large Impact of Particle Size on Insertion Reactions. A Case for Anatase LixTiO2[END_REF][START_REF] Spreafico | Excess Electrons and Interstitial Li Atoms in TiO2 Anatase: Properties of the (101) Interface[END_REF][START_REF] Liang | Anatase Titanium Dioxide as Rechargeable Ion Battery Electrode -A Chronological Review[END_REF][START_REF] Chen | Titanium-Based Anode Materials for Safe Lithium-Ion Batteries[END_REF] Anatase TiO2 holds particular significance as a compound of interest in the realm of proton intercalation. Its suitability stems from its extensively studied open crystal structure, primarily as a negative electrode in lithium batteries. [START_REF] Kanamura | Diffusion of Lithium in the TiO2 Cathode of a Lithium Battery[END_REF][START_REF] Cava | The Crystal Structures of the Lithium-Inserted Metal Oxides Li0.5TiO2 Anatase, LiTi2O4 Spinel, and Li2Ti2O4[END_REF][START_REF] Van De Krol | Spatial Extent of Lithium Intercalation in Anatase TiO2[END_REF][START_REF] Ma | Lithium Intercalation in Anatase Titanium Vacancies and the Role of Local Anionic Environment[END_REF][START_REF] Wagemaker | Large Impact of Particle Size on Insertion Reactions. A Case for Anatase LixTiO2[END_REF][START_REF] Spreafico | Excess Electrons and Interstitial Li Atoms in TiO2 Anatase: Properties of the (101) Interface[END_REF][START_REF] Liang | Anatase Titanium Dioxide as Rechargeable Ion Battery Electrode -A Chronological Review[END_REF][START_REF] Chen | Titanium-Based Anode Materials for Safe Lithium-Ion Batteries[END_REF] In this context, anatase has demonstrated a good cycle life and a notable resistance to forming lithium dendrites. [START_REF] Liang | Anatase Titanium Dioxide as Rechargeable Ion Battery Electrode -A Chronological Review[END_REF] Nevertheless, its capacity is comparatively lower when compared with materials like graphite electrodes. [START_REF] Liang | Anatase Titanium Dioxide as Rechargeable Ion Battery Electrode -A Chronological Review[END_REF][START_REF] Zhang | Graphite as Anode Materials: Fundamental Mechanism, Recent Progress and Advances[END_REF] Additionally, the diffusion of Li + ions within anatase is relatively sluggish. [START_REF] Chen | Titanium-Based Anode Materials for Safe Lithium-Ion Batteries[END_REF] Despite these characteristics, anatase TiO2 retains its pertinence due to its capacity to maintain stability at the nanoscale, thereby exhibiting heightened surface reactivity. It's worth noting that this reactivity can be meticulously tuned, influenced by factors such as exposed crystal facets, [START_REF] Xu | Tailored Hydrothermal Synthesis of Specific Facet-Dominated TiO2Nanocrystals from Lepidocrocite-Type Layered Titanate Nanosheets: Systematical Investigation and Enhanced Photocatalytic Performance[END_REF] defects, [START_REF] Hahn | Something from Nothing: Enhancing Electrochemical Charge Storage with Cation Vacancies[END_REF][START_REF] Dambournet | Cationic Vacancies in Anatase (TiO2): Synthesis, Defect Characterization, and Ion-Intercalation Properties[END_REF] particle size, [START_REF] Wagemaker | Large Impact of Particle Size on Insertion Reactions. A Case for Anatase LixTiO2[END_REF] and dopants. [START_REF] Davoisne | Room-Temperature Synthesis of Iron-Doped Anatase TiO 2 for Lithium-Ion Batteries and Photocatalysis[END_REF] Given the reduced size of proton ions, we can expect enhanced kinetics to facilitate more efficient charge transfer and ion diffusion. This enhancement can contribute to the optimization of electrochemical performance such as power and energy density.

II.1. Proton-coupled electron transfer

Proton-coupled electron transfer (PCET) is a fundamental process in chemistry that involves the transfer of both an electron and a proton in a concerted manner. PCET in TiO2 was first introduced by B. Lemon and J. Hupp in 1996 by performing photochemical quartz crystal microbalance. [START_REF] Lemon | Photochemical Quartz Crystal Microbalance Study of the Nanocrystalline Titanium Dioxide Semiconductor Electrode/Water Interface: Simultaneous Photoaccumulation of Electrons and Protons[END_REF] They demonstrated the evidence that illuminated nanocrystalline TiO2 electrodes undergo simultaneous photo-accumulation of electrons and protons. According to L.

Lyon and J. Hupp in 1999, the conduction band edge energy is found to vary systematically with the pH of the electrolyte solution. [START_REF] Lyon | Energetics of the Nanocrystalline Titanium Dioxide/Aqueous Solution Interface: Approximate Conduction Band Edge Variations between H0 = -10 and H-= +26[END_REF] The potential shift follows a Nernstian behavior of approximately -60 mV per pH unit between strongly acidic and basic pH conditions, demonstrating that the conduction band edge energy is highly pH-dependent. 39 Traditionally, most aqueous electrolytes employed for proton intercalation have been based on H2SO4, a strong acid. However, the application of such a strong acid electrolyte has its limitations, potentially promoting the dissolution of the active material or triggering corrosion in the current collector. 1 This realization has paved the way for exploring alternative electrolyte options. The pivotal innovation in utilizing neutral or mildly acidic buffer electrolytes for achieving proficient proton intercalation yields several advantages. Notably, it enhances the stability and cyclability of the electrode. 16 Moreover, this approach offers the potential to elevate safety levels and reduce environmental impact, rendering it particularly conducive to the development of environmentally sustainable batteries. 16 The insertion mechanism of the proton within the TiO2 structure, when a proton donor mild acidic electrolyte is used, is as follows: 39,41 𝑇𝑖𝑂 2 + 𝑥𝑒 -+ 𝑥𝐴𝐻 ↔ 𝑇𝑖𝑂𝑂𝐻 𝑥 + 𝑥𝐴 -(eq. 1.8)

Here, x represents the extent of TiO2 reduction (0 < x < 1), and AH/A -denotes the conjugate weak Bronsted acid/base pair of the buffer. Additionally, they have also demonstrated that bulk proton insertion in both amorphous and crystalline anatase can occur. 41 These advances constitute a good starting point for developing titanate material for proton aqueous batteries.

Chapter I. State of the Arts

III. HER electrochemistry and kinetics

The Hydrogen Evolution Reaction (HER) constitutes an electrochemical process entailing the reduction of protons (H⁺), leading to the generation of hydrogen gas (H₂) as expressed by (eq. 1.2). In this context, two protons and two electrons unite to yield a single molecule of hydrogen gas, predominantly occurring at negative potentials. However, the intricacies of this reaction go beyond its simple formulation and can be broken into two distinct steps. The initial step involves adsorption, a stage that greatly hinges on the surface's chemical composition of the active material. [START_REF] Lasia | Mechanism and Kinetics of the Hydrogen Evolution Reaction[END_REF] This process, recognized as the Volmer reaction (eq. 1.9), plays a pivotal role in shaping the overall HER mechanism. Subsequently, the recombination step follows, manifesting through two feasible pathways. The first involves electrochemical desorption, wherein an adsorbed proton combines with another proton originating from the electrolyte, facilitated by the introduction of a second electron. [START_REF] Lasia | Mechanism and Kinetics of the Hydrogen Evolution Reaction[END_REF] This mechanism is referred to the Heyrovsky reaction (eq. 1.10). The second pathway termed the Tafel reaction, entails a chemical desorption process wherein two adsorbed protons amalgamate to yield H₂ (eq. 1.11). [START_REF] Lasia | Mechanism and Kinetics of the Hydrogen Evolution Reaction[END_REF] Importantly, it is pertinent to note that both forms of recombination, i.e., the Heyrovsky and Tafel reactions, can occur concurrently. [START_REF] Lasia | Mechanism and Kinetics of the Hydrogen Evolution Reaction[END_REF][START_REF] Koper | Thermodynamic Theory of Multi-Electron Transfer Reactions: Implications for Electrocatalysis[END_REF][START_REF] Gennero | Hydrogen Evolution Reaction : Analysis of the Volmer-Heyrovsky-Tafel Mechanism with a Generalized Adsorption Model[END_REF][START_REF] Nørskov | Trends in the Exchange Current for Hydrogen Evolution[END_REF][START_REF] Parsons | THE RATE OF ELECTROLYTIC HYDROGEN EVOLUTION AND THE HEAT OF ADSORPTION OF HYDROGEN[END_REF] Volmer: 𝐻 + + 𝑒 -+ * ↔ 𝐻 𝑎𝑑𝑠 (eq. 1.9)

Heyrovsky: 𝐻 𝑎𝑑𝑠 + 𝐻 + + 𝑒 -↔ 𝐻 2 + * (eq. 1.10)

Tafel: 2𝐻 𝑎𝑑𝑠 ↔ 𝐻 2 + 2 * (eq. 1.11)
with * corresponds to a free adsorption site on the metal electrode surface. In both scenarios, whether it refers to proton insertion or to HER, the adsorption stage following the Volmer reaction is a requisite process (eq. 1.9). In this context, the binding adsorption energy assumes paramount importance within the electrocatalytic HER. [START_REF] Nørskov | Trends in the Exchange Current for Hydrogen Evolution[END_REF] Drawing upon the Sabatier principle, which posits that catalyst efficiency is optimal when the reactant energies closely mirror those of the transition state, the concept is that a catalyst achieves its maximum efficacy when it neither excessively binds nor weakly binds. [START_REF] Sabatier | How I Have Been Led to the Direct Hydrogenation Method by Metallic Catalysts[END_REF] This balanced interaction facilitates an efficient transition from the reactant state to the transition state where the reaction takes place.

To classify catalysts, several Volcano plots exists, depicting the current density as a function of the hydrogen binding Gibbs energy (ΔGH + ). 30,[START_REF] Nørskov | Trends in the Exchange Current for Hydrogen Evolution[END_REF][START_REF] Wang | Hydrogen Production from Water Electrolysis: Role of Catalysts[END_REF] An extremely weak ΔGH + can impede the overall reaction due to a limiting adsorption step, whereas an excessively strong ΔGH + can lead intriguing question: Does the concept of an aqueous electrolyte remain intact within the framework of WiSE? In the WiSE electrolyte, the water content is significantly reduced, thereby raising queries about the continued classification of such systems within the aqueous electrolyte category.

IV. Strategies to improve electrode stability

Another approach to mitigate water-splitting reactions involves the application of coatings to the electrode surface or the introduction of additives into the electrolyte. 17 In a study conducted by Makivic et al., they employed a Nafion coating on the surface of a TiO2 film electrode. 41 Nafion is a well-known brand of perfluorosulfonic acid-based polymer widely utilized as a proton-exchange membrane (PEM) in various electrochemical applications, particularly in fuel cells. [START_REF] Choi | Thermodynamics and Proton Transport in Nafion[END_REF] One of its notable characteristics is its exceptional proton conduction capability. In this configuration, the researchers observed an improvement in both cyclability and proton capacity through the reduction of H2 production. The Nafion coating acts as a barrier, selectively allowing the passage of ionic species such as protons while limiting the passage of water molecules. This effect diminishes the electrocatalytic activity of water at the interface between the electrode and the electrolyte, contributing to the reduction of water-splitting reactions.

Furthermore, the introduction of electrolyte additives has demonstrated the potential to effectively reduce the overpotentials associated with the HER and OER by obstructing the active sites on electrodes. 17 Consequently, these substances are classified as electrolysis inhibitors. [START_REF] Hou | Surfactant Widens the Electrochemical Window of an Aqueous Electrolyte for Better Rechargeable Aqueous Sodium/Zinc Battery[END_REF] These additives encompass various categories: (1) non-electrolyte solutes capable of diminishing free water molecules through the establishment of hydrogen bonds. [100][101][102] (2) cosolvents that displace water in primary solvation shells to engage with cations. 103 (3)

antisolvents that supplant water in primary solvation shells to facilitate cation-anion coordination. 104 In this thesis, our approach does not involve the manipulation of the electrolyte; instead, we concentrate on enhancing material properties. We believe that by improving intrinsic material properties, such as proton transport properties or through defect engineering, we can achieve favorable electrochemical performance regarding reversible proton insertion.

V. Titanate nanomaterials via sol-gel method

Different strategies to tune titanate materials have been employed recently by our group.

Among them, we can cite the generation of structural defects such as vacancies, that can enable increased capacities in non-aqueous electrolytes for Li + , Mg 2+ , or Al 3+ insertion. For instance, Fluorinated anatase TiO2 succeeded in stabilizing 22% of titanium vacancies through sol-gel methods. [START_REF] Dambournet | Cationic Vacancies in Anatase (TiO2): Synthesis, Defect Characterization, and Ion-Intercalation Properties[END_REF] The vacancies are stabilized within the TiO6 octahedra by charge compensation where divalent O 2-are substituted by monovalent OH -or F -groups. This sample in the presence of vacancies has demonstrated thermodynamically favorable insertion of Li + , Mg 2+ , or Al 3+ (Figure 1.5). [START_REF] Ma | Lithium Intercalation in Anatase Titanium Vacancies and the Role of Local Anionic Environment[END_REF]105,106 In addition, compared to defect-free anatase, the capacity reached higher values and both diffusion and cycling stability were improved, for instance, Mg 2+ reached ~160 mAh.g -1 (vs ~20 mAh.g -1 ) or Al 3+ reached 90 mAh.g -1 (vs ~30 mAh.g -1 ), and Li + diffusion was improved to 10 -13 cm 2 .s -1 . [START_REF] Ma | Lithium Intercalation in Anatase Titanium Vacancies and the Role of Local Anionic Environment[END_REF][START_REF] Dambournet | Cationic Vacancies in Anatase (TiO2): Synthesis, Defect Characterization, and Ion-Intercalation Properties[END_REF] Another pathway consisted of stabilizing chemically cations into a layered structure.

Layered material is an interesting candidate regarding their good ionic transport. 107,108 In our group, Kang et al. succeeded in chemically stabilizing zinc ions into nanoconfined water in a layered lepidocrocite-type TiO2 material (Figure 1.5). 109 This compound exhibits exceptionally high proton conductivity (10 -2 S.m -1 ) working through a Grotthus mechanism. 27 The Grotthus mechanism consists of a hydrogen-bonded network, such as water molecules, Both of the previous strategies are based on the sol-gel method. [START_REF] Dambournet | Cationic Vacancies in Anatase (TiO2): Synthesis, Defect Characterization, and Ion-Intercalation Properties[END_REF]110 This method is a prominent and versatile approach within the realm of materials science. The nomenclature "sol" designates a colloidal dispersion of solid particulates within a liquid medium, while "gel" translates the transformation of this sol into a structured, solid-like framework.

The fundamental of the sol-gel process for titanate material is explained as follows: 111

Sol preparation: Metal alkoxide such as Titanium isopropoxide (TIP) is commonly used as a precursor. It is dissolved in a suitable solvent, such as isopropanol, to create a precursor solution. This solution contains titanium ions in the form of Ti (IV) alkoxide.

Sol-Gel transition: The precursor solution undergoes a controlled hydrolysis and condensation reaction. Water is added to initiate hydrolysis by nucleophilic substitution, leading to the formation of titanium hydroxide species (eq. 1.12). This step highly depends on the nucleophilic and electrophilic character of entering groups, metal atoms, and also the coordination unsaturation (coordination number Noxidation state z). As these species condense, they create nanoscale clusters suspended in the solvent, forming a sol. Three distinct condensation processes via nucleophilic substitution are observed. These reactions include olation (eq. 1.13), which entails the creation of bridging hydroxo groups. On the other hand, oxolation (eq. 1.14) and alcoxolation (eq. 1.15) processes yield bridging oxo groups. The olation reaction occurs when the metal atom remains incompletely coordinated.

Hydrolysis: M-OR + H2O → M-OH + ROH (eq. 1.12)

Condensation:

Olation: M-OH + M-OH2 → M-OH-M + H2O (eq. 1.13)

Oxolation: M-OH + M-OH → M-O-M + H2O (eq. 1.14)

Alcoxolation: M-OR + M-OH → M-O-M + ROH (eq. 1.15)
The sol is then left to age, allowing the clusters to further aggregate and form a threedimensional network. This aging step is critical for achieving the desired crystallinity and structure. This step can be thermally activated. After aging, the gel is carefully dried to remove the solvent, resulting in a porous and solidified gel structure.

Calcination (optional): The dried gel is subjected to calcination in a furnace at temperatures that can go up to generally 500 °C for Anatase. [START_REF] Katal | A Review on the Synthesis of the Various Types of Anatase TiO2 Facets and Their Applications for Photocatalysis[END_REF] This thermal treatment induces further condensation reactions, crystallization, and removal of organic residues. The calcination temperature and duration influence the crystalline phase and the particle size of TiO2 nanoparticles.

This method confers numerous advantages such as the capability to synthesize a wide range of materials, ranging from thin films to bulky ceramics. This approach provides precise control over the composition, structure, and morphological characteristics at the nanoscale, allowing for meticulous tailoring of material properties. Sol-gel reactions often occur at moderate temperatures, making them suitable for eco-conceived electrodes. Materials produced through the sol-gel process typically exhibit notable chemical uniformity and purity, highlighting the inherent homogeneity of the resulting products. 111,112

VI. Conclusion

This chapter delved into the intricate landscape of aqueous batteries' development, elucidating a fundamental challenge stemming from their constrained energy density due to a limited electrochemical stability window. This window is marked by the occurrence of gas production, such as H2 and O2, through the hydrogen evolution reaction and oxygen evolution reaction. These parasitic reactions significantly curtail the overall battery performance, imposing constraints on the choice of suitable electrode materials. Only a few negative electrodes are compatible with proton intercalation and sustainability. In alignment with the pursuit of environmentally friendly energy storage solutions, we embarked on investigating TiO2 as a promising candidate. TiO2 has displayed remarkable performance as a negative electrode material within the context of Li-ion technology. Moreover, it has unveiled a fascinating proton-coupled electron transfer mechanism, enabling bulk proton insertion.

Building upon these intriguing attributes, our focus now shifts towards a comprehensive investigation of proton intercalation in TiO2 materials, particularly within the use of a mildly acidic buffer electrolyte. This exploration seeks to uncover new avenues for harnessing TiO2's multifaceted capabilities in the realm of energy storage and conversion. Our focus remained centered on addressing these challenges through the manipulation of material properties, avoiding direct modifications of the electrolyte.

By harnessing defect engineering within 3D TiO2 structures by sol-gel method, we sought to increase capacity and amplify proton diffusion kinetics. Furthermore, we directed our attention towards the enhancement of proton transport properties, specifically targeting the nanoconfined water within layered titanate materials-a pivotal host of proton conduction. This endeavor aimed at elevating the reversibility of proton insertion, a crucial factor for the Chapter II: Synthesis and characterization of titanate nanomaterials via soft chemistry method

I. Introduction

The sol-gel synthesis method is a versatile and widely used technique for preparing titanate compounds. 1 Among the various TiO2 polymorphs, anatase stands out as a particularly interesting one. It exhibits a metastable tetragonal crystal structure and finds extensive use in fields such as solar cells, 2 environmental purification, 3 photocatalysis, 4,5 gas sensors, 6 and lithium batteries. 7,8 As a negative electrode material for lithium-ion batteries, anatase demonstrates a relatively high operating voltage (1.4 ~ 1.8 V vs. Li/Li + ) and stable cycling performance, making it an attractive choice for battery applications. 9 Moreover, anatase is a promising candidate for proton intercalation, considering its natural abundance, chemical stability over a wide range of pH, and environmental compatibility. 10,11 Recent studies have revealed the potential of anatase for reversible proton intercalation within the bulk, exhibiting good proton diffusivity, which is vital for efficient charge and discharge cycles. 12,13 However, despite its promising properties as a negative electrode, some challenges remain, such as enhancing overall capacity and inhibiting secondary reactions. In the context of aqueous batteries, the low reduction potential of TiO2 (Ti IV /Ti III ) can favor the parasitic Hydrogen Evolution Reaction (HER), necessitating strategies to overcome this limitation and improve the performance of TiO2 electrodes for proton-ion batteries. Nanostructuring, surface modification, and enhancing proton transport are some key material-related strategies being investigated to make TiO2 electrodes viable for proton-ion batteries.

Our research group has undertaken recent studies on the synthesis of titanium dioxide compounds, including 2D lamellar TiO2 and cationic-deficient TiO2 using the sol-gel method. 14- 20 These compounds hold significant potential for investigation under proton electrochemistry.

The 2D lamellar TiO2 consists of a lepidocrocite-type titanate with an orthorhombic structure (Immm) featuring layers of edge-sharing TiO6 octahedra and interlayer regions occupied by crystal water. 21 Prior studies have reported the synthesis of potassium or cesium lepidocrocite titanates through solid-state reactions, subsequently transformed into hydrogen lepidocrocitetype titanate through ion exchange in an acidic environment. [21][22][23][24][25][26][27] Furthermore, low-temperature solvothermal synthesis has been employed, resulting in the formation of an amorphous material by reacting Ti (IV) isopropoxide with a controllable content of water at 90 °C. 14 Structural investigations using high-energy X-ray synchrotron pair distribution function (PDF) have confirmed the lepidocrocite-type structure, with electrically neutral titanate layers and absence of interlayer cations. The low-temperature synthesis pathway promotes the formation of hydroxy groups, generating a positive charge to arise due to the charge difference between O 2-and OH -groups. To balance this charge disparity cationic vacancies are formed, becoming stabilized within the structure. The lepidocrocite-type titanate exhibits promising electrochemical behavior for the insertion of monovalent ions such as Li + , Na + , and K + . 16 Considering proton behaviors, nanoconfined water within the structure is expected to display distinctive behaviors, impacting various processes such as self-dissociation, proton mobility, and proton transport.

The sol-gel method has emerged as a well-established route for the controlled synthesis of nanostructured materials, offering precise control over composition and morphology.

Structural defects, particularly vacancies, have proven crucial in achieving higher capacities in certain materials. 28,29 For example, fluorinated anatase TiO2 synthesized via the soft chemistry method stabilizes titanium vacancies through the substitution of divalent O 2-with monovalent F -and OH -. [17][18][19] This modification leads to significantly higher capacities compared to pure TiO2 anatase, as the cationic vacancies provide thermodynamically favored insertion sites. 19 The presence of vacancies also facilitates charge distribution, mitigating sluggish diffusion of guest ions, and resolving issues associated with ion insertion kinetics. 18 Moreover, vacancies help mitigate volume changes induced by the intercalation of multivalent ions, further enhancing the electrochemical properties.

In this study, we aim to synthesize three titanate materials via the sol-gel method, following well-established protocols. These titanate compounds will serve as negative electrodes for proton-ion batteries in the following chapters. Our focus will be on synthesizing lamellar lepidocrocite-type TiO2, cationic-deficient Anatase, and stoichiometric Anatase.

II. Synthesis and structure

II.1. Synthesis process

The sol-gel method allows for the controlled formation of oxide, hydroxide, or oxyhydroxide networks via the inorganic polymerization of molecular precursors, typically metal alkoxides. 1 In the sol-gel process of titanium dioxide, metal alkoxides (e.g., titanium isopropoxide) are used as the starting materials. These alkoxides are highly reactive with water due to the low electronegativity of the titanium metal and its ability to coordinate with ligands. 1 When the metal alkoxide reacts with water, hydrolysis reactions occur, leading to the replacement of alkoxy ligands with hydroxyl groups by nucleophilic substitution (eq. 2.1):

Ti(OR)4 + xH2O → Ti(OR)4-x(OH)x + xROH (eq. 2.1)
The hydrolyzed metal alkoxide precursors then undergo various condensation reactions, such as olation, oxolation, and alcoxolation, forming bridging hydroxo (eq. 2.

2) and oxo groups (eq. 2.3 & eq. 2.4) in the titanate networks. These reactions play a crucial role in tailoring the structure and properties of the final product.

Olation: Ti-OH + H2O-Ti → Ti-OH-Ti + H2O (eq. 2.2)

Oxolation: Ti-OH + OH-Ti → Ti-O-Ti + H2O (eq. 2.3)

Alcoxolation: Ti-OH + RO-Ti → Ti-O-Ti + ROH (eq. 2.4)
The sol-gel synthesis method provides several advantages, enabling precise control over the stoichiometry, composition, and morphology of titanate compounds. Key reaction conditions, such as the hydrolysis ratio (h = [H2O]/[Ti(OR)4]) and temperature, exert critical influences on the properties of the final product. Temperature plays a significant role in the process by activating both hydrolysis and condensation reactions. Research using titanium isopropoxide precursors revealed that when h < 2, condensation is primarily governed by oxolation and alcoxolation, leading to the formation of chain polymers. 1 However, increasing the hydrolysis ratio poses challenges to the formation of oxo bridge due to the rising partial charge of the OR groups, thereby favoring polymerization via olation. Excessive water (h > z = oxidation state of Ti) results in gel formation and precipitation. Furthermore, the sol-gel process allows for the incorporation of various dopants or additives, enabling the tailoring of Chapter II. Synthesis and characterization of titanate nanomaterials via soft chemistry method 40 titanate properties for specific applications. It also offers opportunities for forming nanostructured titanate materials, which can exhibit enhanced properties and functionalities compared to their bulk counterparts.

According to Ma et al. 14 the reaction system involved titanium isopropoxide dissolved in isopropanol, with water serving as the hydrolysis agent at a constant H2O/Ti molar ratio of At 90 °C, the HRTEM image displayed an amorphous state with no lattice plan. At 110 °C, the number of crystallized particles increased, but they remained aggregated, making it difficult to estimate the average particle size. At 150 °C, the particles were homogeneously distributed, and well-defined geometrical crystallites characteristic of Anatase were observed. The TEM analysis revealed an average particle size of 12.4 nm along the width and 20.9 nm along the length (see Figure S2.1), suggesting an anisotropic crystal growth, typical of anatase without a shaping agent. 30 The range of temperatures where crystallization occurs is small suggesting that the structures between the amorphous and crystalline samples are closely related.

coherence length, representing crystal periodicity, is relatively short, not exceeding 8 Å, indicating the absence of long-range order. On the other hand, at T = 110 and 150 °C, the coherence lengths exceed 30 Å, characteristic of long-range order in crystalline materials. To gain insights into the local arrangement, we attempted to refine the PDF data. In Figure 2.2.b, the sample at T = 90 °C was refined using a lepidocrocite model (Space group: Immm).

However, for T = 110 and 150 °C, a single-phase refinement using an anatase model (space group: I41/amd) was utilized to obtain a good fit. Indeed, the attempt to fit the T = 90 °C data with an anatase model was unsuccessful (see Figure S2.

2).

Nevertheless, the structural arrangement of these two structures is related. 31 The lepidocrocite structure displays a lamellar arrangement of TiO6 octahedra ( The refinement reliability of both crystalline phases and amorphous L-TiO2 is good, with Rw < 20% and equal to 33%, respectively. These refinement values are well in agreement with the values found by Ma et al. 14 Chapter II. Synthesis and characterization of titanate nanomaterials via soft chemistry method 44 the water present in the structure, and especially after the removal of OH groups, which is consistent with a topotactic dehydration. 23 Collaborating with M. Sorriaux, a PHENIX Ph.D. candidate, we estimated the chemical composition using thermogravimetric data and a general chemical formula: Ti2-x□xO4-4x(OH)4x

• nH2O. 14 Where the presence of OH groups enables the generation of titanium vacancies by

II.4. Crystallization mechanism

L-TiO2 is characterized by distinct layers separated by crystal water molecules. As reported, upon thermal decomposition of this interlayer water molecules in Lepidocrocite-type TiO2, a topotactic dehydration process occurs, leading to layer collapse and rearrangement, resulting in the formation of anatase compound. 14,23 To investigate the crystallization mechanism under thermal annealing, we employed Powder X-ray thermodiffraction, a technique well-suited for studying structural changes of materials as a function of temperature using X-ray diffraction. Therefore, it is particularly suited to studying the phase transition occurring within L-TiO2, especially when the transition is thermally controlled. is well in agreement with the appearance of Bragg peaks observed in the X-Ray thermal diffraction. The condensation mechanism is believed to proceed through an oxolation process as follows:

2(Ti-OH) → Ti-O-Ti + H2O (eq. 2.5)

Therefore, and in agreement with the reported solid-state transformation, we can assess that the solid phase transformation mechanism under thermal treatment happens as follows: 14,23 1. Precipitation of L-TiO2:

Ti2-x□xO4-4x(OH)4x • nH2O (eq. 2.6)
2. Decomposition with the departure of adsorbed and crystal water:

Ti2-x□xO4-4x(OH)4x + nH2O (eq. 2.7)
3. Condensation of the zig-zag chains to adopt the anatase structure:

y * Ti1-x□xO2-4x(OH)4x + nH2O (eq. 2.8)
Here, y =

(2-𝑥) (1-𝑥)
due to the collapse of the layer and the condensation, which induces a reduction of the unit formula, giving Ti1-x□xO2-4x(OH)4x • nH2O. A further increase in time or temperature reaction induces the condensation of the OH groups leading to (1-x) TiO2.

II.5. Morphology evolution upon heating

To study the structural and morphology changes upon heating, we conducted SEM-FEG The in situ TEM study aims to extend the investigation to the original synthesis process by directly observing the phase transition occurring in the liquid phase. However, up to date, only in situ TEM experiments of L-TiO2 powders have been achieved. The observed morphology evolution during crystallization confirms the occurrence of the two-phase process, where the amorphous phase first precipitates and then converts to anatase. This two-phase process occurs via a solid-state transformation, involving decomposition and condensation of the intermediate amorphous particles.

II.6. Anisotropic growth

To investigate the kinetics and anisotropic growth in the liquid phase, we maintained the synthesis at a specific temperature, here 110 °C, and interrupted the process at various time

III.1. Structural refinement

To characterize the crystal properties of Vac-TiO2, we conducted a microstructural analysis using refinement techniques. For the Rietveld analysis performed in this study, highquality data were essential, requiring proper sample preparation and acquisition parameters.

The obtained XRD pattern was fitted using a TiO2 space group: I41/amd, with lattice parameters a = b = 3.87200 Å, c = 9.61600 Å, and α = β = γ = 90.0°.

To achieve a good fit, a crystallographic tetragonal anisotropic broadening, specifically affecting the Lorentzian component, was incorporated using spherical harmonics (Figure 2.9.a, Annexes A.5). 35 Indeed, the 004 Bragg peak, could not be fitted without considering the anisotropic effect (Figure 2.9.b). This approach allows for the consideration of peak asymmetry and different broadening along distinct crystallographic directions. Spherical harmonics expansion enables the decomposition of the broadening function into specific harmonics, corresponding to different orders and orientations. 35 In the case of tetragonal anisotropic broadening, spherical harmonics are commonly employed to model and separate the Lorentzian component, which accounts for broadening due to the size and shape distribution of the crystallites in the sample.

In Table 2.1, we have reported the results of refinements. The obtained reliability of the fitting is good with a low Bragg R-factor. To obtain a good fit for Vac-TiO2, it was necessary to refine the Ti occupancy, where the X² was significantly improved switching from 9.48 to 4.18 and the Bragg R-factor was reduced from 7.46% to 2.927% respectively. This yields to a content of ~16.7% of Ti 4+ vacancies, which is in between the PDF and TGA analysis and in agreement with the reported TiO2 synthesized in the same condition at 110 °C, where Ma et al.

have found ~12% by PDF refinement. 14 Additionally, the refinement of lattice parameters revealed some differences between Vac-TiO2 and Pure Anatase.

III.2. Microstructural study

In order to investigate the influence of vacancies on size & strain within the Vac-TiO2 structure, we conducted an analysis using the Williamson-Hall method. This method takes advantage of the broadening observed in diffraction peaks of recorded patterns to examine the microstructural characteristics. 36 The broadening of these peaks is influenced by multiple factors, including the size of the crystallites, lattice strain, and instrumental effects. By carefully examining the shape and width of the diffraction peaks, the Williamson-Hall method allows for the separation of the different contributions to the broadening and provides quantitative insights into the underlying microstructural parameters.

To perform the Williamson-Hall analysis, we plot the integral breadth on the y-axis and the magnitude of the scattering vector (related to the Bragg angle) on the x-axis. This plotting technique enables us to determine various contributions to the broadening of the diffraction peaks, such as crystallite size and lattice strain. By carefully analyzing the slope and intercept of the Williamson-Hall plot, we can extract valuable information about the average size of the crystallites, the level of strain present, and the potential presence of other mechanisms causing broadening. This analysis allows for a deeper understanding of the microstructural characteristics of the Vac-TiO2 sample and sheds light on the effects of vacancies on the strain within its structure.

By using a Cauchy function as a sum of Lorentzian and Gaussian component, we have β that correspond to the integral breadth: 36 𝛽 = 𝛽 𝑜𝑏𝑠 -𝛽 𝑖𝑛𝑠𝑡𝑟 = 𝛽 𝐿 + 𝛽 𝜂 (eq. 2.10)

With βL due to size effect and βη due to micro-strain:

𝛽 𝐿 = 𝜆 𝐿 cos 𝜃 (eq. 2.

11)

𝛽 𝜂 = 𝜂 tan 𝜃 (eq. 2.12)

With L as the crystallite size, and 𝜂 = 4𝜀 where 𝜀 = the slopes in both directions are attenuated, suggesting a nearly strain-free behavior. By examining the ASS plot, we find that the apparent strain within Pure Anatase is relatively low at 0.06% (Figure S2.6), while Vac-TiO2 exhibits a slightly higher apparent strain of 0.08% (Figure 2.10.b). Interestingly, considering the higher vacancy content (~16.7%) in the Vac-TiO2 sample, we initially expected a more pronounced strain effect. This suggests, that vacancies are well integer and stabilized into the crystallographic structure of this defective anatase.

The calculated average apparent size of the Pure Anatase crystallites is found to be approximately 12.5 nm, which closely matches the measured value of 12.4 nm obtained from TEM imaging. This agreement highlights the robust quality of the Rietveld fit and validates its accuracy in determining the crystallite size.

However, in the case of Vac-TiO2, obtaining an accurate average apparent size through TEM characterization was challenging due to the high aggregation of particles. Consequently, we turn to the Rietveld fit to estimate the apparent size, which yields a value of approximately 9 nm for the Vac-TiO2 crystallites. This finding is consistent with the lower synthesis temperature, which typically leads to slower particle growth.

IV. Transport properties

In collaboration with J-C. Badot and O. Dubrunfaut from LGEP, Paris-Saclay, by measuring the dielectric response over a wide frequency range, typically here from 50 to 10 10

Hz, we can study the dynamics and behavior of protons and water within the material of interest.

This technique allows for the analysis of dielectric properties such as permittivity () and conductivity () as a function of frequency ( = 2 with  the frequency). () and () are linked by the following equation: () = i0() where 0 is the vacuum permittivity (0  8.854×10 -12 F.m -1 ). The complex dielectric permittivity of the material is influenced by various processes that depend on the frequency. These processes include microscopic fluctuations of molecular dipoles and the movement of charge carriers. 37 In a layered material with confined water, additional mechanisms come into play, namely the rotation of water molecules and charge hopping at high frequencies, and grain polarizations at low frequencies. 38 To study the dielectric relaxation, different model functions can be used. In an ideal condition, where dipole interactions are absent, the Debye function is often employed (eq.

2.15

). This function involves parameters 𝜀∞ and 𝜀𝑠, representing the permittivity at the lowest and highest frequencies, respectively, while 𝜏 denotes the relaxation time. The Nyquist plot (" vs ') of Debye function is represented by a semi-circle centered on real axis '.

ε * (ω) = ε ∞ + ε s -ε ∞ 1+iωτ (eq. 2.

15)

In many experimental data, non-Debye relaxations are commonly observed. To describe these relaxations, the Cole-Cole function (eq. 2.16) is utilized, which is a generalization of the Debye function. In the Cole-Cole function, the parameter α is introduced as an empirical value that quantifies the degree of deviation from the Debye mode. 39 This provides a more flexible and accurate representation of dielectric relaxation phenomena that deviate from ideal Debye behavior. The Nyquist plot (" vs ') of CC function is represented by a circular arc centered under real axis '.

ε * (𝜔) = ε ∞ + ε 𝑠 -ε 1+(𝑖𝜔𝜏) 1-𝛼 (eq. 2.

16)

The Havriliak-Negami (HN) functions have found considerable use in the analysis and interpretation of dielectric relaxation phenomena in a variety of materials, including polymers, glasses, biological systems, and complex fluids. 40 It is an empirical model from the Debye relaxation model where it encompasses a broad distribution of relaxation times, provided by incorporating both a dispersive and a dielectric relaxation component. The decomposition of the dielectric spectra revealed three relaxations (Figure 2.13) and were fitted using an HN function (eq. 2.17).

ε * (𝜔) = ε ∞ + ε 𝑠 -ε (1+(𝑖𝜔𝜏) 1-𝛼 ) 𝛽 (eq. 2.17)
where β are an empirical model related to the degree of deviation from the Debye model.

The Nyquist plot (" vs ') of HN function is represented by a skewed arc. Higher frequency part of HN function is a straight line in Nyquist representation

The transport properties of layered Lepidocrocite-type TiO2 were investigated by Kang et al. 15 They found that the mobility of protons within the material is primarily facilitated by proton hopping between water molecules located in the interlayer spaces. These water molecules form hydrogen bonds with each other, creating a network that allows for the shortrange movement of protons. While L-TiO2 exhibits low ionic conductivity, it demonstrates the ability for local proton hopping due to the presence of these hydrogen bonds. In contrast, high protonic conductivity is typically attributed to a Grotthuss mechanism involving long-range hopping. The formation of hydrogen bonds between adjacent water molecules (O-H...O) is the key factor enabling the local proton hopping behavior. The activation Energy (Ea) for the proton hopping process, measured at 0.49 eV, represents the energy barrier that protons must overcome to move within the material. These mild hydrogen interactions provide a favorable environment for proton mobility within the material. 41 This relatively low activation energy indicates the ease of proton movement facilitated by the presence of hydrogen bonds (Table 2

.2).

To investigate the influence of vacancies in Vac-TiO2, we analyzed the dielectric response using Nyquist plots. The first Nyquist plot shows a prominent relaxation, labeled as P1, in the low-frequency region. This relaxation is attributed to the presence of an ionic double layer at the silver paint/sample interface. After subtracting the contribution P1 ( with the HN and CC functions respectively, we obtained their characteristic frequencies of 7.2x10 4 Hz and 1.4x10 9 Hz at room temperature (Table 2 compared to reported L-TiO2, 15 and by of factor of 10 compared to pure Anatase (Table 2.

2).

We estimated Ea for the different processes by measuring the frequencies at various temperatures and plotting them using the Arrhenius equation (Table 2.2). However, due to the poor ionic conductivity of pure Anatase, we were unable to extract the corresponding Ea for proton hopping and water rotation in this sample. Interestingly, the Ea of bulk ionic conductivities were found to be in a similar range, approximately 0.51, 0.52, and 0.56 eV, for L-TiO2, Vac-TiO2, and pure Anatase, respectively. Interestingly, the Ea for proton hopping showed similar behavior in L-TiO2 and Vac-TiO2, with values of 0.49 eV and 0.53 eV, respectively, indicating mild hydrogen interactions. 41 The similar Ea of bulk ionic conductivities and Ea of proton hopping suggest that the proton is the charge carrier responsible for ionic conduction.

Furthermore, the water movement and proton hopping in Vac-TiO2 exhibited a unique behavior. The water rotation frequency was found to be the lowest among the different samples, while the attempt frequency (ν0), which is proportional to the charge holder, was surprisingly high. Additionally, the dielectric strength () relative to the intensity of polarization (≈ length of jump) of Vac-TiO2 was similar to L-TiO2. These findings suggest that a greater number of charge carriers participate in the conduction process in the Vac-TiO2 sample, while the water movement is less mobile. This behavior can be attributed to the OH groups occurring in the vacancies, where their mobility is strongly affected by their coordination. The participation of vacancies in proton conduction can explain the significant increase in ionic bulk conductivity observed in Vac-TiO2 compared to L-TiO2 and pure Anatase.

V. Conclusion

We We have seen that morphological evolution during heating is closely linked to the crystallization process of L-TiO2. The particle size reduction during calcination is mainly attributed to the loss of interlayer water and the occurrence of crystallization into anatase with condensation of vacancies.

The refinement study of the resulting Vac-TiO2 compound revealed an amount of 16.7% of cationic vacancies, which is closely supported by pdf and TG analysis. Then, by microstructural analysis, we have investigated the role of these vacancies on the contribution of strain and have found that they do not have a huge impact on it. For the next chapters, all compositions will be expressed regarding TG analysis, as L-TiO2 cannot be refined by XRD Rietveld refinement.

The proton transport properties were investigated via the BDS technique. The Vac-TiO2 sample exhibits a higher bulk ionic conductivity compared to L-TiO2 and pure Anatase. This gain in conductivity was attributed to a greater number of protons that participate in the conduction and especially the role of vacancies in the proton motion.

These materials will serve as negative electrode for proton intercalation. In the next chapter, we aim to tune the water interlayer with cation to enhance even more the proton transport properties. The presence of vacancies in 3D structure within Vac-TiO2 will be investigated for the proton intercalation. Pure Anatase will be our reference materials.

I. Introduction

The aqueous battery faces a significant challenge in terms of limited energy density, primarily stemming from the constrained electrochemical stability window (ESW).

Thermodynamically, the pure water's ESW is only 1.23 V, imposing substantial limitations on both energy density and the choice of materials compatible with such conditions. The negative side is typically restricted by the competitive hydrogen evolution reaction (HER), while the positive side contends with the oxygen evolution reaction (OER).

TiO2 is a promising material for negative electrode applications due to its costeffectiveness, abundance, non-toxicity, and stability across various pH levels in aqueous environments. 1,2 Its relatively low one-electron reduction potential (Ti IV /Ti III ) allows for a theoretical gravimetric capacity of up to 335 mAh.g -1 with one electron stored per Ti redoxactive center. 3 Proton insertion into bulk amorphous and anatase TiO2 has been successfully demonstrated. 4 Additionally, our research group has investigated nanostructured amorphous layered titanates as potential candidates for lithium intercalation. 5 This specific material, known as Lepidocrocite type TiO2, adopts a layered structure comprising TiO6 octahedra separated with interlayer water molecules (Figure 3.1.a). This arrangement offers high disorder, allowing for the accommodation of cationic vacancies up to 25%, stabilized with OH groups by charge compensation. 5 The focus of the recent work done by Kang et al. has been on the exploration of nanoconfined water within the layers of Lepidocrocite type TiO2 to enhance proton mobility. 6 The unique properties of confined water have drawn significant interest, particularly its role in proton conduction mechanisms. Extensive research has identified the Grotthus mechanism as a vital pathway for proton transport within confined water, involving the sequential transfer of protons through a network of hydrogen bonds among adjacent water molecules. 7 Building on this foundation, recent advancements in our research group have suggested the potential to tailor the electrochemical characteristics of Lepidocrocite type TiO2 by chemically introducing cations into its structure. Based on it, the interlayer water becomes a key parameter to tune the overall electrochemical properties of this material. Cations with high polarizing power and small ionic radius are interesting candidates to tune the interlayer: they can modify the water organization and promote proton diffusion through the hydrogen bond network via this Grotthus mechanism. To do so, we aim to tune the nanoconfined water located in the layer of the Lepidocrocite type TiO2 (further named: L-TiO2) to upsurge the proton mobility.

Zinc ions are small (0.74 Å) and highly polarizable (Z/r² = 3.65 Å -2 ), which makes them suitable for bonding with dipolar molecules, such as water in the interlayer and interacting with water molecules to influence their structure. A previous study by Kang et al. delved into the utilization of a precursor, ZnCl2, to achieve the chemical intercalation of zinc ions within the interlayer water of a material. 6 They have successfully chemically demonstrated that the lepidocrocite structure is conserved under zinc incorporation and that chloride remains close to the Zn 2+ . 6 From thermal analysis, and 1 H NMR they reached the following general formula: Within the scope of our research, we chose to follow a similar methodology but made a key adjustment in the precursor solution. Specifically, we opted for zinc acetate instead of the previously employed zinc chloride. This intentional choice was made to eliminate the presence of counter ions, thereby ensuring the exclusive presence of stabilized zinc cations. According to Density Functional Theory (DFT) calculations, zinc ions within the water interlayer were

ZnyCl2y Ti2-x□xO4-4x(OH)4x • nH2O
Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 76 the interlayer space and can coordinate with a greater number of water molecules (Figure

3.1.b).

Accordingly, we adopted this strategy.

Experimental: The synthesis procedure for both Lepidocrocite and Zn-Cl2 titanate materials has been previously reported. 6 To briefly outline the process for Zn- The obtained gel was washed with absolute ethanol recovered by centrifugation and dried for 4 hours at 80°C. Products were then ground manually using a pestle and pillar until a fine powder was obtained.

The synthesis used herein consists of a sol-gel process to obtain a lepidocrocite type TiO2. The structural nature of the lepidocrocite structure was characterized through high-energy X-ray total scattering and described as previously shown (space group: Immm). 5 Interestingly to prevent the transformation into anatase, specific parameters were set, including a hydrolysis ratio (h = nH2O/nTi) of 3.33 and a constant temperature of 90 °C. As it has been already discussed, the lepidocrocite structure consists of TiO6 edge-sharing octahedra, separated by a single layer of crystal water. The titanate layers remain electrically neutral, with no cations in the interlayer. The synthesis at low temperatures leads to the formation of hydroxo groups instead of oxo groups. 12 This in turn induces a positive charge due to the charge disparity between O 2-and OH -groups. S3.3). The existence of peaks corresponding to the elongated vibration mode of carbon-carbon bonds resembles the peaks observed in the Zn Acetate precursor. Nevertheless, the intensity of these peaks is notably diminished, consistent with the anticipated elimination of these groups during the synthesis process.

II.3. Bragg & total scattering diffraction

Having achieved the successful incorporation of zinc ions into L-TiO2 established, we now proceed to explore the consequences of this integration on the material's structure.

Examining the scanning electron microscopy (SEM-FEG) image provided in the Supplementary Information (Figure S3.4) for L-TiO2, the corresponding image for 20% Zn-Ace titanate displays a strikingly similar morphology. Specifically, this morphology comprises aggregated spherical nanoparticles, devoid of any distinct crystal-like arrangements. In Figure S3.5 is summarized the XRD patterns for L-TiO2 as well as samples synthesized with varying Zn/Ti ratios of 10%, 20%, and 30%. Notably, the introduction of zinc maintains a resemblance in the XRD pattern to that of L-TiO2, characterized by the absence of discernible Bragg peaks.

However, the structural changes confronted challenges associated with the intrinsically disordered nature of the crystallographic structure, the uncertain location and speciation of Zn 2+ ions within the structure, and the dynamics of interlayer interactions. To overcome these difficulties, we studied the structural characteristics using the X-ray pair distribution function x ~ 0.5 for reported Lepidocrocite type TiO2, 5 to x ~ 0.61 for 20% Zn-Ace titanate (for L-TiO2,

x ~ 0.5). Thus, the higher the Zn content, the greater the prevalence of vacancies.

Regarding the shift in the 3 rd peak, where Ti-Ti corner-sharing pairs are shortened from 3.66 to 3.52 Å, several explanations are plausible. This shift may result from the increasing concentration of Zn 2+ , which could potentially introduce clusters of zinc. 6 Another possibility is that cations induce a bending effect in the layer structure, as the positions of the 1 st and 2 nd peaks remain unchanged (Figure 3.3.b). Refinement of the structure using a real-space approach indicates that this peak is not fully accommodated by the lepidocrocite structure, suggesting that it is induced by the presence of Zn ions (Figure 3.3).

III. Transport properties

To get a better understanding of the impact of Zn ion on bulk ionic conductivity, we performed broadband dielectric spectroscopy (50-10 10 Hz) to measure the complex permittivity 3.1).

Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 83 this lower protonic conductivity couples to a low activation energy Ev might imply a shorter range but high mobility of proton hopping. Furthermore, the attempt frequency of water rotation presents an interesting insight into the behavior of nanoconfined water within the material. A slower relaxation time (7 × 10 11 Hz, Table 3.1) can be explained by either a higher ordering effect of the nanoconfined water or a longer distance between the hydrogen bonds. The proton remains in a specific location for a longer period before moving to another site, although the transition time between sites is notably shorter. Additionally, a significant distinction becomes apparent when compared to the previously reported Lepidocrocite-type TiO2 material.. 6 In the case of the Zn-Ace titanate, the movement of proton hopping and water molecules appear correlated, as evidenced by the similar attempt frequency values of 3 and 7 × 10 11 Hz, respectively. This correlation of oscillation between H2O and proton within a site is particularly significant concerning the Grotthus mechanism, where this movement can be likened to Newton's cradle's motion occurring for proton hopping. Such local correlated motion can contribute to a higher mobility of protons within the material, a crucial factor for facilitating efficient proton conduction. 7 Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 85 Hz with an activation energy of 0.14 eV. On the other hand, for ZnCl2, the attempt frequency is higher, standing at 5.4 × 10 9 Hz with an activation energy of 0.25 eV. 6 These differences in mobility indicate that there is less participation of hydrogen bonds within Zn-Ace titanate, likely due to the higher ordering effect present in the interlayer. This suggests that chloride ions hinder the zinc ions from forming higher coordination states with water, This excess mobility of protons within Zn-Ace titanate holds significant implications, especially when the active material is immersed in an aqueous electrolyte and subjected to an applied potential. This increased proton mobility facilitates faster proton transfer within the material, which can potentially confer advantages in improving the overall performance of the material when serving as an electrode in an electrochemical system. In situations where the material is exposed to an aqueous electrolyte and an applied potential, the enhanced proton mobility can lead to faster intercalation and de-intercalation of protons within the material's structure. Consequently, this can lead to enhances charge and discharge rates, as well as an overall improvement in electrochemical efficiency. The synergy between reduced hydrogen bond interactions and increased proton mobility creates a favorable environment for more efficient proton conduction and enhanced electrochemical performance of Zn-Ace titanate across a range of applications, particularly in the domain of aqueous batteries.

IV. Electrochemistry properties

IV.1. Electrode preparation and choice of electrolyte

In our study, we conducted electrochemical experiments on the layered titanate material using a slurry-based electrode configuration. The experimental procedure closely follows the methodology illustrated in Figure 3.6. Specifically, we employed a drop-casting approach to coat the layered titanate material, represented by doped or undoped L-TiO2 materials, onto a titanium (Ti) foil substrate. The slurry used for this purpose consisted of 17% weight carbon black mixed with 65% weight in tetrahydrofuran (THF) solvent. Additionally, 18% weight of Nafion 117 ® was incorporated into the slurry to effectively coat the aggregates and improve electrode performance. The resulting slurry was deposited onto a Ti foil, and the mass of the material ranged from 0.7 to 0.9 mg. To ensure a controlled experimental setup and isolate the exposed electrode area, a surface area of 0.36 cm² was defined using a layer of nail varnish.

This also served to cover the Ti foil that would come into contact with the electrolyte, maintaining accurate electrochemical measurements. Following the methodology proposed by Notably, the quasi-steady-state current (iC) reached at potentials around -0.6 V (i.e., E > ECB = conduction band potential) directly correlates with the electrode's double-layer capacitance. The property of TiO2, featuring a pH-dependent conduction band, leads to poor electronic conductivity until the potential surpasses the ECB. [16][17][18] Consequently, a transition from an insulator to a conductor occurs when the potential approximates ECB, indicating the gradual occupation of electronic states within the band gap until E < ECB. Before ECB is attained, the accumulation of negative charge within TiO2 is balanced by the electrosorption of cations from the electrolyte at the electrode surface. When E << ECB, the current density notably intensifies, accompanied by the emergence of a broad faradic wave spanning from -0.6 to -1.2 V. This pronounced increase in current density, compared to the response in the unbuffered electrolyte, signifies the onset of a faradaic intercalation process (in this case, discharge). This phenomenon can be attributed to the faradaic reduction of Ti IV ions concurrent with the intercalation of protons.

To ensure the selectivity of proton intercalation within the L-TiO2 material, we conducted a comprehensive study with varying pH levels. Considering that TiO2 exhibits a pHdependent potential conduction band, it should theoretically adhere to the Nernst equation if the proton intercalation process is intricately linked to electron transfer. The Nernst equation describes the equilibrium potential for an electrode reaction that involves an exchange of electrons and ions. In our case, as protons are involved in the intercalation process within L-TiO2, changes in the pH of the electrolyte can directly influence the potential at which the intercalation occurs. By systematically altering the pH of the electrolyte, we can assess how the proton intercalation potential varies with changes in pH. 4,14,16,17,19 E Ti 4+ /Ti 3+ = E Ti 4+ /Ti 3+ 0 -0.059pH (eq 3.4) signifies that the storage mechanism within the L-TiO2 material operates exclusively through a proton-coupled electron transfer mechanism, effectively excluding other ionic species present within the electrolyte solution. Furthermore, by analyzing the intercept of the linear regression, we determined the potential at which TiO2 insertion becomes fully conductive. This potential, measured to be around -0.52 V, concurs with our observations from cyclic voltammetry, where a significant exponential increase in current density was observed around approximately -0.6 V.

Chapter
To investigate the proton intercalation within Zn-Ace titanate, we performed cyclic voltammetry under the same conditions as those used for L-TiO2. The resulting cyclic voltammogram exhibited a distinct and pronounced contrast when compared to that of L-TiO2, as presented in Figure 3.7.c. The faradaic contribution associated with proton insertion in Zn-Ace titanate revealed well-defined broad waves, suggesting a higher degree of reversibility compared to L-TiO2. At specific potential ranges (0 to -1.2 V), the calculated capacity based on the oxidizing current was found to be 127 mAh.g -1 for L-TiO2 and 146 mAh.g -1 for Zn-Ace titanate, respectively. This higher capacity value for Zn-Ace titanate demonstrates its greater propensity for intercalating protons.

Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 91 the TiO2/electrolyte interface. To reinforce our observation of the reduced occurrence of the HER, we calculated the HER formation at the identified cut-off potential displaying the highest degree of reversibility. By isolating the irreversible HER contribution, we could ascertain its effect on the columbic efficiency. The HER capacity (QHER) was determined by subtracting the integrated anodic and cathodic current densities as outlined in equation 3.5:

Q HER = Q Discharge -Q Charge (eq. 3.5) 𝐶E = Q charge Q discharge * 100 (eq. 3.6)
where Qcharge corresponds to proton de-intercalation and Qdischarge corresponds to proton intercalation within the TiO2 electrode. The results of this analysis, as depicted in ) and Anatase (132 mAh.g -1 ). 4 The theoretical capacity of Zn-Ace titanate, calculated based on its chemical formula Ti1.39□0.61O1.54(OH)2.46 • 0.64H2O 0.27Zn(OH)2, and based on Ti 4+ /Ti 3+ redox couple is 217 mAh.g -1 . The reversible capacity, which corresponds to a stoichiometry of x ≈ 0.83, indicates a substantial reduction of almost all Ti redox sites, affirming the bulk insertion of a proton within this material. This value is consistent with x ≈ 0.77 for lithium insertion in layered lepidocrocite type TiO2. 5 This increase in the working potential window and capacity holds promise for enhancing energy density under battery conditions (density energy (Wh.g -1 ) = potential window (V) × capacity (Ah.g -1 )).

Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 94 temperature in K. C 0 Ti is the maximal concentration of the redox species in mol.cm -3 , For L-TiO2, the molar mass is 149.6 g.mol -1 (based on the chemical formula) and the lattice volume obtained through PDF refinement is 194.71 × 10 -24 cm 3 . This yields a maximal concentration of C 0 Ti = 0.008 mol.cm -3 . For Zn-Ace titanate, the molar mass is 171.3 g.mol -1 with a lattice volume of 191.56 × 10 -24 cm -3 , resulting in a similar concentration of C 0 Ti = 0.008 mol.cm -3 .

The estimated diffusion coefficients calculated from equation 3.9 are DH+, L-TiO2 = 1.7 × 10 -13 cm 2 .s -1 and DH+, Zn-Ace titanate = 1.02 × 10 -12 cm 2 .s -1 for L-TiO2 and Zn-Ace titanate, respectively. Interestingly, the calculated diffusion coefficient for Zn-Ace titanate is higher than that of L-TiO2, indicating a more efficient diffusion process in the presence of zinc. These values closely align with the diffusion coefficients found for lithium insertion in lepidocrocite type TiO2, as reported by Ma et al. 5 Notably, these diffusion coefficients are significantly higher than those found for H + diffusion in nanoparticles of amorphous and anatase TiO2 (10 -15 cm 2 .s -1 ). 4,13 Interestingly, these findings are consistent with the observations obtained through BDS analysis, confirming that the presence of zinc within the interlayer facilitates proton migration and diffusion.

IV.5. Galvanostatic experiments vs CV

To assess the cyclability of the titanate materials, Galvanostatic Charge/Discharge (GCD) cycling experiments were performed. Initially, these experiments were carried out under strong accumulation conditions, with the hydrogen evolution reaction (HER) plateau being the limiting factor. Upon the first discharge (Figure 3.10.a), an impressive deintercalated capacity of 212 mAh.g -1 was achieved, surpassing the value obtained from cyclic voltammetry (181 mAh.g -1 ) and aligning with the theoretical full capacity of Zn-Ace titanate (217 mAh.g -1 ).

Interestingly, the HER plateau was not observed during the intercalation process, but the total 230 mAh.g -1 capacity was not reached, suggesting that the HER occurred in another competitive way or some proton are retained within the compound. The maximum potential achieved during the initial discharge was -1.31 V. In contrast, L-TiO2 managed to deintercalate 165 mAh.g -1 ,

with the HER plateau being effectively observed at -1. To address the need for improved cyclability, a strategy inspired by Balland et al. was adopted, involving mild accumulation conditions. 4 This approach involved chronopotentiometry experiments in which proton intercalation was constrained by fixed discharge times (36 s) and a current density of 15.4 A.g -1 . This current density corresponded to an inserted capacity of 154 mAh.g -1 , translating to a stoichiometry value of x ≈ 0.71 for Zn-Ace titanate and x ≈ 0.58 for L-TiO2. In this protocol, the charging process was allowed to proceed without a fixed time constraint, reaching a potential limitation up to 0 V. Notably, this approach deliberately avoided the occurrence of the HER, as demonstrated in first 50 cycles and exhibited 29% greater capacity retention than L-TiO2 after 100 cycles. These performances are higher than the cyclability reported found in titanate layered with 33 mAh.g -1 Ti3O7(H3O)2 over 50 cycles or higher than ~70 mAh.g -1 found in amorphous TiO2 over 55 cycles. 4,18 This remarkable cyclability highlights the advantages of Zn-Ace titanate in electrochemical energy storage applications.

Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 97

V. Conclusion

The sol-gel synthesis method utilized in this study demonstrated the successful incorporation of zinc ions and the formation of the desired structures (without counter anions).

The structural analysis through the X-Ray pair distribution function highlighted the preservation of the lepidocrocite-type TiO2 structure upon zinc incorporation.

To gain a comprehensive understanding of the influence of Zn ions on bulk ionic conductivity, a detailed investigation was carried out using broadband dielectric spectroscopy.

Specifically, the bulk protonic conductivity of Zn-Ace titanate was approximately 10 times higher than that of the reported Lepidocrocite type TiO2 counterpart. 6 A comparison with ZnCl2 titanate synthesized by Kang et al. indicated that while Zn-Ace titanate's bulk protonic conductivity did not surpass that of ZnCl2 titanate, it displayed distinctive characteristics. Zn-Ace titanate exhibited a higher proton hopping frequency with lower activation energy in comparison to ZnCl2 titanate. This lower activation energy, coupled with a moderate protonic conductivity, suggested the potential for shorter-range yet highly mobile proton hopping behavior.

The electrochemical characterization demonstrated a significant improvement regarding the proton intercalation in Zn-Ace titanate compared to L-TiO2. The Zn-Ace titanate electrode exhibited favorable properties for proton intercalation, offering higher capacity, enhanced reversibility, and potential extension compared to L-TiO2. Scan rate studies unveiled that both L-TiO2 and Zn-Ace titanate followed an intercalation process governed by proton diffusion.

Diffusion coefficients estimated show a higher value in Zn-Ace titanate than in L-TiO2 showing the straight benefit of zinc ion. Under galvanostatic charge/discharge conditions, Zn-Ace titanate demonstrated improved cyclability compared to L-TiO2.

Hence, the influence of zinc ions on proton diffusion has led to enhanced cyclability and extended potential ranges. These findings contribute to the understanding of proton insertion mechanisms and offer valuable insights for the design of high-performance energy storage materials for proton ion aqueous batteries.

Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 98

In the future, further investigations into the chemical intercalation of other cations could be envisaged, with the aim of controlling the structure of the water in the crystals and obtaining specific properties in a finely tuned way. Accordingly, we investigate the use of a defective TiO2 material, that has been synthesized using a low-temperature sol-gel method (Vac-TiO2, see Chapter II. II. Synthesis and structure). We have demonstrated that this unique titanate compound, Vac-TiO2, exhibits an anatase framework with a high content of cationic vacancies that are stabilized by OH groups. We have studied the proton transport properties by dielectric spectroscopy and we have shown an interesting increase of the grain bulk conductivity of 1.34 × 10 -4 S.m -1 compared to stoichiometric anatase with 1.86 × 10 -5 S.m -1 . Up-to-date, the study of such titanate compounds regarding proton intercalation has never been done. In this chapter, we look at the electrochemical signature of Vac-TiO2, focusing on proton insertion behaviors.

Chapter IV. Defective TiO2: pushing the limit of HER 109 II.

Electrochemistry properties II.1. Intercalation process

First, we attempt to study the proton intercalation process in the Vac-TiO2 sample. To comprehensively characterize the electrochemical signature, we prepared electrodes of Vac-TiO2 using the same protocol as described previously. Briefly, 65% wt of Vac-TiO2 was mixed with 17% of Carbon Black in THF, and 18% wt of Nafion 117 ® was added to the mixture.

The obtained slurry is drop cast on top of a Ti foil and the surface area is controlled with nail varnish (0.36 cm² exposed). As mentioned in the previous chapter, the current density measured by Cyclic Voltammetry (CV) is the current normalized by the geometrical surface area. All potentials are expressed regarding Ag/AgCl, KCl Sat' (+0.197 V vs SHE).

To differentiate between the faradaic and capacitive contributions of proton intercalation within the Vac-TiO2 sample, we adopted the same procedure outlined in Chapter III. In brief, we employed two distinct electrolytes:

i/ A low proton-concentration electrolyte (1 M KCl, pH ≤ 7) to evaluate the capacitive contribution and, ii/ A 1 M acetate buffer electrolyte (CH3COOH/CH3COOK, pH = 5) to study the faradaic contribution, involving the (de)intercalation of ions within the electrode structure. 7,8 In Figure 4.2.a, the dashed curves present the Vac-TiO2 and pure Anatase response in the KCl electrolyte, revealing a nearly rectangular shape characteristic of a double-layer capacitance. The calculated capacity values were 6.5 mAh.g -1 (0.24 mF.g -1 ) and 18.9 mAh.g -1 (0.7 mF.g -1 ) for pure Anatase and Vac-TiO2, respectively. This minor contribution suggests the absence of insertion phenomena within this potential range when protons are not available. To identify the working potential window of Vac-TiO2, a cut-off analysis was performed using 1 M acetate buffer electrolyte. As depicted in 𝑇𝑖 𝐼𝑉 𝑂 2 (𝑠𝑢𝑟𝑓𝑎𝑐𝑒) + 𝑥𝐻 + + 𝑥𝑒 -↔ 𝑇𝑖 𝐼𝐼𝐼 𝑂𝑂𝐻 (𝑠𝑢𝑟𝑓𝑎𝑐𝑒) + 𝐻 2 𝑂 (eq. 4.1) 𝑇𝑖 𝐼𝐼𝐼 𝑂𝑂𝐻 (𝑠𝑢𝑟𝑓𝑎𝑐𝑒) + 𝑇𝑖 𝐼𝑉 𝑂 2 (𝑏𝑢𝑙𝑘) ↔ 𝑇𝑖 𝐼𝑉 𝑂 2 (𝑠𝑢𝑟𝑓𝑎𝑐𝑒) + 𝑇𝑖 𝐼𝐼𝐼 𝑂𝑂𝐻 (𝑏𝑢𝑙𝑘) (eq. 4.2)

Remarkably, the current response exhibits significant perturbations at potentials below V. Notably, a maximum of capacity of 259 mAh.g -1 for proton de-intercalation was achieved at -1.5 V. This remarkable value not only surpasses that observed for pure Anatase (187 mAh.g -1 ) but also those found by Makivic et al. 9 Based on the chemical formula obtained by TGA for Vac-TiO2 (Chapter II. II.3. Composition Analysis using Thermogravimetric Analysis (TGA) and PDF) Ti0.76□0.24O1.03(OH)0.97, this reversible capacity corresponds to a stoichiometry of x ≈ 0.88 which is almost the entire reduction of all Ti redox site, attesting the bulk insertion of a proton within this sample. Additionally, this value compares well with reported Li + insertion in nanosized TiO2 and also results achieved on L-TiO2. 10 In comparison, the 187 mAh.g -1 capacity observed for pure Anatase corresponds to x ≈ 0.58, which is close to the one encounter with Li-ion insertion inside bulk anatase with x ≈ 0.5 Li + . 11,12 The shapes and potential positions of the CV obtained with Vac-TiO2, manifest two broad, reversible faradaic peaks at a midpoint potential of -0.85 V and -1.21 V. Remarkably, as shown in Figure 4.3 the latter potential aligns with the reversible sharp peak observed in pure Chapter IV. Defective TiO2: pushing the limit of HER 111 Anatase voltammograms. As a result, we attribute this faradaic contribution to the insertion of protons within the anatase lattice: Ti3-O⸱⸱⸱H. Based on the findings from PDF, TGA, NMR spectroscopy, and XRD Rietveld refinement analyses which indicated the presence of vacancies (Chapter II), our initial hypothesis is that the peak centered at -0.85 V corresponds to the insertion of H + within the vacancy site. These sites are characterized by under-coordinated O 2- ions, typically involving Ti2□-O or Ti1□2-O. Importantly, when we compare these results with those from L-TiO2 results (Figure 4.3), where intercalation likely occurs within vacancies due to their abundance, we observe a distinct overlap between the broad faradaic waves encountered in L-TiO2 and those seen in the Vac-TiO2 system. This further supports the hypothesis of H + insertion within vacancies.

faradaic contribution at each cut-off, discharge (orange), charge (green), and HER (violet). mAh.g -1 of active titanate materials.

II.2. Proton couple electron Transfer

As in Chapter III, we performed a pH-dependent study to investigate the nature pf the charge carrier within the Vac-TiO2 material in both sites. Indeed, TiO2 has a pH-dependent potential conduction band and should follow the Nernst equation if the proton is coupled to the electron transfer (see eq. 4.4). 8,9,[13][14][15] E Ti 4+ /Ti 3+ = E Ti 4+ /Ti 3+ °-0.059pH (eq. 4.4)

From this equation 4.4, the slope per unit of pH corresponds to 0.059 when a Nernstian behavior is followed. In addition, the potential of the conduction band can be derived from the Chapter IV. Defective TiO2: pushing the limit of HER 114 point of intersection. We performed cyclic voltammetry at 20 mV.s -1 upon a wide range of pH, using proton donor electrolyte at 1 M, including HCl (pH 0), Chloroacetate buffer (pH 3.0), Acetate buffer (pH 5.0), HEPES buffer (pH 7.0), and Carbonate buffer (pH 9.3), (Figure S4.

2).

The linear regression from eq. 4 is made with the oxidizing interstitial peak in Vac-TiO2 Figure

4.4.

A slope of 0.061 V is obtained, consistent with the theoretical value predicted by the Nernstian pH-dependent equation in standard conditions. This observation implies that the storage mechanism involving lattice insertion within this titanate material operates through a proton-coupled electron transfer mechanism, therefore, excluding other ions present within the solution. In addition, the intercept linear regression gives the potential at which TiO2 lattice insertion becomes fully conductive and here is equivalent to -0.85 V. This result agrees with our observations from cyclic voltammetry on pure Anatase, further affirming that the intercalation process within Vac-TiO2, particularly regarding the second redox potential, follows the same underlying mechanism as observed in pure Anatase. However, it is noteworthy that the study of the pH-dependent on the vacancies' faradaic peak proved difficult to carry out due to the broad nature of the peaks. This complexity hindered the establishment of a specific potential for subsequent linear regression analysis. Nonetheless, in the case of L-TiO2 (Chapter

III. IV.2 Intercalation of H +

), a successful pH-dependent study was conducted, yielding a slope of 0.059 V. This result demonstrates that solely protons are intercalated within the cationic vacancies of L-TiO2. By assuming the structural proximity between the arrangement of vacancies in Vac-TiO2 and L-TiO2 samples, we hypothesize that the proton is the only ion intercalated in cationic vacancies in the Vac-TiO2 compound.

III. DFT calculations: probing proton intercalation mechanism

We turned to Density Functional Theory (DFT) calculations in collaboration with B. J.

Morgan at University of Bath (UK) to reinforce our findings concerning proton intercalation within vacancies. From a thermodynamic perspective, Regarding the open framework of pure Anatase, the proton finds its intercalation site nestled within an interstitial position, forming a bond with oxygen that is already 3-fold coordinated with Ti IV . It is similar to the process encounter of lithium intercalation into anatase. 3,6,12,16 Notably, six distinct possibilities for proton insertion within the anatase lattice exist. The calculated intercalation energy for this configuration is determined to be -1.38 eV.ion -1 , which is close to the intercalation energy for Li + in anatase (-1.23 eV.ion -1 ). 3 According to the relaxed structure for Vac-TiO2

IV. Kinetics investigation: insights into proton diffusion mechanisms

To understand the kinetics mechanism of both intercalation sites within Vac-TiO2, a scan rate study was performed. The potential cut-off was fixed at -1.4 V (to limit the HER) and the scan rate was screened from 100 mV.s -1 to a minimum just before the initiation of HER, here 20 mV.s -1 (Figure S4.3). The logarithm of the maximal cathodic current peak for both interstitial and vacancy sites was plotted in function of the logarithm of the scan rate following the power law: 𝑖 = 𝑎𝑣 𝑏 (eq. 4.5)

with i the current in mA.cm -2 , and v the scan rate in mv.s -1 . a and b are coefficient values.

By fitting the power law from eq. 4.5, we can have access to the b-value. Where a b-value of 0.5 corresponds to a rate controlled by diffusion, and a b-value of 1 is a surface-controlled rate.

The log-log plotting (Figure 4.6) reveals a small yet distinct difference between these two sites.

Linear regression fits for the two peaks at potentials of -0.83 V and -1.14 V, respectively, for vacancies and interstitial sites (Figure S4.3). Notably, the b-value obtained for the interstitial site, standing at 0.64, closely reflects an intercalation process governed by proton diffusion within the anatase lattice of Vac-TiO2. This finding finds congruence with observations of Li + insertion into Anatase. 17 Interestingly, for the vacancy site, the b-value increased to 0.72, suggesting a more surface-influenced current response. Despite this increase, we hypothesize that proton diffusion remains through the vacancies, as the range of b-values typically associated with capacitive mechanisms have a b-value between 0.8 and 1. 18 It is worth noting that the selected potential position can induce significant variation, especially regarding the vacancy site and its broad faradaic wave. Additionally, electrodes are far from model electrodes and therefore, some contributions such as ohmic drop can participate and may impact the results, especially at a high scan rate. 19 Where Ip is the peak current of the redox reaction (A), n is the number of electrons (here 1), F is Faraday's Constant (96485 C.mol -1 ), A is the surface area of the working electrode in (0.36 cm 2 ), v is the scan rate used to collect the voltammogram in V.s -1 , D is the diffusion coefficient of the redox species in cm 2 .s -1 , R is the Gas Constant 8.3144 J.mol -1 .K -1 , and T is the temperature in K. C is the concentration of the redox species in mol.cm -3 , in stoichiometric anatase TiO2 the maximum concentration of Ti is C = 0.048 mol.cm -3 . In defective Vac-TiO2, the active redox Ti is 0.76 (based on the chemical formula), thus yielding a maximal concentration of C = 0,03648 mol.cm -3 .

The estimated diffusion coefficients for Vac-TiO2 from equation 4.7 are reported in Table 4.1. We found DH+, interstitial = 2.02 × 10 -15 cm 2 .s -1 and DH+, vacancy = 1.71 × 10 -15 cm 2 .s -1 for proton insertion in interstitial and vacancy sites respectively. Interestingly, the calculated diffusion coefficient in the vacancy site is slightly smaller than the one found in the interstitial site, suggesting a slower diffusion process. These values are strongly different from the observed coefficient diffusion for L-TiO2 and Zn-Ace titanate with respectively 10 -13 and 10 -12 cm 2 .s -1 .

Nevertheless, these values closely align with those found for proton insertion in Anatase TiO2, as reported by Balland et al. 7,9 They are also similar to Li + coefficient diffusion found in anatase. 17,[20][21][22] Remarkably, in pure Anatase, we found a b-value of 0.53 (Figure S4.4) with a diffusion coefficient of DH+, pure Anatase = 1.5 × 10 -15 cm 2 .s -1 , which is very similar to the coefficient found in an interstitial site in Vac-TiO2. Interestingly, this suggests that the proton migration/diffusion occurs at the same rate regarding the interstitial site of both Vac-TiO2 and pure Anatase. 1.7×10 -13 1.02×10 -12 1.71×10 -15 2.02×10 -15 1.5×10 -15 10 -15

V. Mono or biphasic compound

Regarding the similar shape for proton intercalation between L-TiO2 and Vac-TiO2 To further explore this hypothesis, mixtures of pure Anatase (A) and L-TiO2 (L) were prepared in varying ratios (ranging from 90:10 to 10:90, A:L), followed by cut-off studies to mimic the electrochemical response of Vac-TiO2. Among these mixed ratios, the closest resemblance to the electrochemical signature of Vac-TiO2 was observed with the 7:3 (A:L) mixture. Chapter IV. Defective TiO2: pushing the limit of HER 122

VI. Cyclability analysis

The cyclability of the crystalline titanate material was further investigated by Galvanostatic Charge/Discharge (GCD). Initially, the study was conducted under conditions of strong accumulation limitation, where a fixed potential was the only imposed limiting factor.

In the initial discharge process, as the onset potential for the HER in galvanostatic conditions for Vac-TiO2 was not known, the intercalation process was allowed to proceed freely until the appearance of the HER plateau. Subsequently, a positive current was applied, establishing a potential limit for the subsequent GCD cycles, notably fixed at -1.4 V (Figure 4.8.a). Upon the first discharge, a deintercalated capacity of 275 mAh.g -1 was achieved, slightly exceeding the value obtained from cyclic voltammetry (259 mAh.g -1 ) and closely approaching the theoretical full capacity of Vac-TiO2 (294 mAh.g -1 ). Interestingly, the HER onset in galvanostatic conditions started at -1.4 V, unlike the shifted potential observed in cyclic voltammetry. This divergence can be attributed to the distinct operational modes: galvanostatic conditions maintain a fixed discharge current and varying potential, whereas cyclic voltammetry varies both potential and current. Therefore, the excess electron injected in Vac-TiO2 after full proton storage serves for the H2 formation. Thereby, the resulting Coulombic Efficiency (CE) of the first cycle is low with CE = 61%. Despite, a significant capacity fading at the 2 nd cycle, the potential limitation enables for a deintercalation capacity of 219 mAh.g -1 , with an improvement of the Coulombic efficiency to 91% observed. This trend of capacity fading was reproduced in subsequent cycles, reaching capacities of deintercalation of 195 and 106 mAh.g -1 for the 5 th and 20 th cycles, accompanied by CEs of 89% and 88%, respectively.

To address the need for improved cyclability, we revisited the strategy originally inspired by Balland et al., which involves mild accumulation conditions. 9 In line with this approach, we conducted chronopotentiometry experiments where proton intercalation was controlled through fixed discharge durations (36 s) and a current density of 15.4 A.g -1 , resulting in an inserted capacity of 154 mAh.g -1 (stoichiometry value of x = 0.52). In contrast, the charging process was unrestricted in terms of time and allowed to continue until maximum reversibility was achieved, with a potential limit set at 0 V. For an almost fully reversible process, the deintercalation is generally realized within 30 seconds. This protocol deliberately avoided the HER potential formation, as evidenced by Anatase. Additionally, the overall intercalation in Vac-TiO2 occurs at more cathodic potentials compared to pure Anatase, where this former competes with the HER region. As an illustration, for 60 mAh.g -1 of proton intercalated (discharge curve), the potential reaches -1.0 V for Vac-TiO2 and already reaches -1.2 V for pure anatase. The reversibility at the 2 nd cycle, here in such mild accumulation conditions, is greatly improved in both cases, with a CE = 96.8% and CE = 77.6% for Vac TiO2 and pure Anatase respectively. This demonstrates again the benefits of performing GCD under mild accumulation protocol.

The reversibility of the two samples was further explored through multiple cycles of GCDs. We initially expected better cyclability for Vac-TiO2, we attempt to explain this relatively poor performance regarding the nature of the counter electrode utilized during the experiment. The choice of the counter electrode in electrochemical experiments is crucial, as it can significantly influence the overall performance. In half-cell, a Pt wire was employed as the counter electrode and is well-known for its catalytic properties towards both the HER and the OER. Therefore, it should be noted that the Pt counter electrode may not be optimal for materials capable of attaining more negative potentials, for Vac-TiO2. During the cycling measurements, the current passing through the Pt wire, opposing the working electrode (TiO2), reaches a potential plateau indicative of H2 or O2 production. This behavior arises from the high adsorption affinity of Pt, which facilitates the evolution of gas species. Notably, according to the capability of Vac-TiO2 to intercalate as low as -1.42 V, and regarding the OER plateau stabilized at 1.89 V, the measured potential window reaches a value as high as 3.3 V at 154 mAh.g -1 for the 10 th cycle. At such potential, both the OER and HER processes occurring at the counter electrode exhibit enhanced kinetics (Figure 4.9).

In contrast, pure Anatase, achieves a potential of -1.29 V during the 10 th cycle for the same capacity, while the OER plateau on the Pt wire stabilizes at 1.5 V leading to a potential Chapter IV. Defective TiO2: pushing the limit of HER 124 window of 2.79 V at the end of discharge. Consequently, gas production within the pure Anatase in the half-cell is comparatively lower than that observed in the Vac-TiO2 half-cell. It is important to know that gas evolution leads to electrolyte consumption and degradation, thereby affecting the reversibility of the electrochemical system. 23 The ability for Vac-TiO2 to intercalate at a lower potential, leads to a higher degradation of the electrolyte and explains this gradual loss of capacity upon cycles. Therefore, with another counter electrode where a stable redox reaction occurs at a lower potential, we can expect that the cyclability will be greatly improved (Chapter V. IV.

Full cell cycling performances)

To evaluate the rate performance of Vac-TiO2 and Pure Anatase under mild accumulation conditions, we realized a series of GCDs across a range of current densities spanning from 1.54 to 46.1 A.g -1 . As illustrated in Figure 4.8.d, at lower current densities (i.e.

1.5 A.g -1 ), poor efficiency was observed, which is consistent with previous findings in literature for proton insertion at low rates. 9 This phenomenon can be attributed to the prevalence of the HER, which is known to be more favorable at slower discharge rates. The underlying mechanism of the reaction occurs primarily on the material's surface. At lower current densities, an extended duration is available for the adsorption and recombination of protons on the TiO2 surface, thereby facilitating the HER process. In contrast, an increase in efficiency was evident at higher current densities, reaching a peak of around 15.4 A.g -1 . This observation suggests a more effective insertion of protons at elevated rates compared to lower rates. Consequently, this kinetics competition at a higher rate gives less time for protons to undergo the recombination with other protons onto the TiO2 surface, hindering thus the HER reaction.

VII. Conclusion

In this investigation, we have extensively explored the electrochemical characteristics and mechanisms underlying proton intercalation within Vac-TiO2. This titanate material, synthesized using the low-temperature sol-gel approach, is promising not only for postponing the onset of HER but also for substantially increasing the proton capacity. The proton-coupled electron transfer mechanism, validated through pH-dependent investigations, DFT calculations, and kinetic analyses, establishes Vac-TiO2 as an exclusive host for proton intercalation within its structural framework. The increase in the H + diffusivity and rapid kinetics governing proton intercalation are closely linked to the small size of the proton ions.

Moreover, our research shows the remarkable cyclability of Vac-TiO2, outperforming pure Anatase over 50 cycles under mild accumulation conditions. This outcome underscores the material's potential to facilitate sustainable proton intercalation processes. Nonetheless, a persistent challenge remains in extending cycling durability, alongside the ongoing competition with the HER.

Looking ahead, our aim is to explore these titanate electrodes in full-cell configurations.

Figure S4.5: Raman signature of L-TiO2 Introduction

Recently, aqueous batteries have engendered increasing interest in grid-scale energy storage, primarily due to their distinctive advantages of safety and low cost. 1 They offer higher power capabilities but propose relatively low energy densities compared to conventional Li-ion batteries, which are known for their higher energy density, typically ranging from 100-300

Wh.kg -1 .

Among the aqueous-based batteries, Zn||MnO2 batteries have gained increasing attention this past few decades, regarding their high gravimetric capacity, reversibility, and lowcost design. 12,13 The MnO2 polymorph can show gravimetric capacity comprised between 300 and 600 mAh.g -1 when metallic Zn is used. Interestingly, Pan et al in 2016 reversibly performed 5000 cycles with Mn 2+ dissolved in a mild-based electrolyte. 12 Although this technology is already established and very promising, many improvements and questions remain. These include the growth of zinc dendrites, leading to short circuits, the capacity fading over time related to the discharge/charge mechanism of MnO2, and the self-discharge of the battery in certain conditions. [14][15][16] For the negative electrode, other chemistry has been tested including TiO2. The combination of TiO2 and MnO2 electrodes represents a promising option for energy storage applications, as these materials are abundant in the Earth's crust, cost-effective, and non-toxic metal oxides. 17,18 Compared to zinc, TiO2 goes with an insertion de-insertion mechanism during the charge/discharge process, avoiding the formation of dendrites and short circuits. Progress has recently been made in understanding the mechanisms of proton insertion into TiO2 electrodes. Evidence of rapid and massive proton insertion into TiO2 has been shown with either anatase TiO2 or amorphous TiO2. 9,[19][20][21] Up to date, a single report focused on the complete cell configuration using both anatase TiO2 for proton intercalation and MnO2 electrodeposition/dissolution on carbon cloth. 9 They succeeded in performing 10 cycles with energy density and power density of 143.2 Wh.kg -1 and 47.9 kW.kg -1 respectively. However, the poor cyclability encountered is linked to HER occurring at the negative electrode (TiO2), which Chapter V. Full cells 134 degrades not only the contact between the TiO2 nanoparticles but also between the films and the Ti current collector foil.

Importantly, the realization of the full potential of aqueous proton-ion cells is not without its challenges. One of the main obstacles is to achieve a high energy density as the electrochemical stability of water in water-based electrolytes leads to lower cell voltages. To overcome this limitation, this requires the exploration of new electrode materials and/or electrolyte formulations. 22 In this chapter, we have tested an aqueous MnO2||TiO2 battery with a weakly acidic buffered medium (1 M at pH = 5 of CH3COOH/CH3COOK), as electrolyte. Compared to other aqueous-based batteries, we have explored the design of novel TiO2 materials to postpone the HER and increase the energy density of the battery. The positive electrode is synthesized in situ using manganese salts (0.3 M (CH3COO)2Mn) in the electrolyte and carbon fiber paper as the current collector. The theoretical potential of this battery is 1.41 V since the standard potential of the developed titanate material is E° = -0.52 to -0.85 V and the standard potential of the conversion MnO2/Mn 2+ in mild acidic buffer electrolyte is E° = 0.56 V. 15 This should lead theoretically to an energy density of  472 Wh.g -1 (with TiO2 maximum capacity of 335 mAh.g -1 ) and a power of 47 kW.g -1 (for 36 seconds of charge/discharge).

II. Carbon nanofibers electrospun paper as current collector for MnO 2 plating/stripping II.1 Generalities of the electrospinning process

Electrospinning is a versatile and widely used processing technique for the fabrication of nanofibers web using electric force to draw charged threads of polymer solutions. This process allows the fabrication of continuous fibers with diameters in the order of a few nanometers to several micrometers. 26 The electrospinning setup typically consists of a high-voltage power supply, a syringe or spinneret containing the polymer solution, and a grounded collector (Figure 5.1.a). 27 Initially, a polymer solution or polymer melt is prepared by dissolving or melting a polymer in a suitable solvent or melt carrier, respectively. The polymer solution is then loaded into a syringe or placed in a reservoir connected to the spinneret.

process is typically carried out at room temperature unless a specific temperature is required to keep the polymer in a liquid state.

By manipulating these parameters, it is possible to achieve nanofibers with desired diameters, surface morphologies, and mechanical properties. Accordingly, electrospun nanofibers find ap²plications in various fields, including tissue engineering, filtration, energy storage, sensors, drug delivery, and protective clothing, owing to their high surface-to-volume ratio, porous structure, and tunable properties. [26][27][28][29][30][31]

II.2 Fabrication of Carbon Nano-Fibers

II.2.1 Generalities

Carbon nanofibers exhibit exceptional properties, including high mechanical strength, excellent thermal and electrical conductivity, and a large surface area. These characteristics make them highly desirable for a wide range of applications, such as energy storage, electronics, catalysis, and composite materials. 26,[31][32][33][34] A recent development in the literature involves the synthesis of conductive microfibers without the need for heat treatment. [35][36][37] This approach includes mixing a polymer solution with a conductive carbon material (e.g., carbon black) and subsequently electrospinning the composite. However, it is worth noting that the electronic conductivities of these microfibers are generally lower compared to those of pure carbon materials (10 1 to 10 3 S.cm -1 ) or electrospun carbon nanofibers followed by heat treatment (PAN-based carbon nanofiber 10 1 S.cm -1 ). 38,39 According to Jeremy-Luc Sanchez's studies, Polyacrylonitrile (PAN) was chosen as the polymer precursor ((C3H3N)x). 40 PAN is an acrylic resin exhibiting thermoplasticity properties with rigidity and toughness. It is resistant to most of the solvents and has slow combustion and low gas permeability. 41 Typically, the polymer is first dissolved in adequate solvents such as N,N-Dimethylformamide (DMF) and then spun into acrylic fibers. 32,42 These fibers are soft, flexible, and result in lightweight and lower cost. 41 Interestingly, these PAN fibers serve as precursor materials in the production of carbon and graphite fibers. 31,42 Chapter V. Full cells 137

In our electrospinning process, we have used the specific parameters that have been reported elsewhere. 15,43,44 The high voltage applied ranged between 17 to 18 kV, with a distance of 15 cm between the nozzle and the current collector. The current collector, in this case, was a static plate covered with aluminum foil. The flow rate was adjusted to 0.03 mL/min at room temperature, with a humidity level below 10%. The electrospun PAN fibers are illustrated in 32,45 During this stage, the polymer chains undergo chemical and structural rearrangements, cross-linking, and elimination of volatile components, resulting in the formation of a carbon-rich structure with improved thermal resistance and a characteristic brown color. 41,45 The final step is carbonization (Figure S5.1), where the stabilized fibers are subjected to high temperatures (typically above 800°C) in an inert or controlled atmosphere. 32,34,45 This process removes non-carbon elements, such as hydrogen, nitrogen, and oxygen, from the fibers, leading to the formation of carbon nanofibers with a graphitic structure. The conditions and duration of carbonization are crucial factors influencing the final properties of the carbon nanofibers, including their crystallinity and conductivity. For our electrospun carbon membrane, the carbonization step was carried out in a tubular furnace at 1200 °C for 1 hour under an argon flow, with a heating rate of 10 °C/min as optimized by previous work in our group (Figure 5.3.c). 40 Chapter V. Full cells 139 indicative of the presence of disordered structures associated with the sp3 hybridization of carbon atoms. In contrast, the second band, designated as the G band, demonstrates a peak at 1599 cm -1 , suggesting sp 2 hybridization of carbon atoms and thus the existence of ordered graphitic structures. Furthermore, the observed decrease in the ID/IG ratio during the heating process implies a concomitant reduction in structural defects. Those experiments confirm the presence of graphitic structure and structural defects in our CNFs.

These non-porous/microporous CNFs exhibit electrical conductivity in the range of 5-10 S.cm -1 , as reported previously. 42 To identify the electrochemical working potential window of the CNFs, we realized a cut-off study by Cyclic voltammetry (CV). realized in 1 M acetate buffer electrolyte with a scan rate of 10 mV.s -1 . In this set of experiments, it can be seen that the Oxygen Evolution Reaction (OER) appears at 1.27 V on the positive side and the Hydrogen Evolution Reaction (HER) at -0.83 V on the negative side.

To determine the Electrochemical active Surface Capacitance Area (ESCA), a scan rate study was conducted in a non-faradaic region, between 0.0 and 0.1 V (Figure 5.3.d). The CNFs exhibited a higher capacitive contribution as evidenced by the nearly rectangular shape obtained. By plotting the current (I) as a function of the scan rate, the slope of the curve allows us to estimate the ESCA. A value of 39.4 mF.cm -2 (equivalent to 13.7 F.g -1 ) is calculated. The specific surface area of CNFs was estimated to be 35 m².g -1 based on the measurement of fiber diameters (143 nm). 49,50 Areal capacitance values associated with such specific surface areas have been reported to be approximately 5 F.g -1 . 51 However, our CNFs exhibit higher areal capacitance values probably due to higher porosity contents.

Chapter V. Full cells 141 due to its ability to undergo reversible redox reactions in the presence of water-based electrolytes, as indicated by the Pourbaix diagram. 57 Interestingly, within a pH range of 3 to 5, a conversion mechanism is observed for MnO2, leading to the formation of Mn2O3 or MnOOH, or Mn 2+ ions. 57 In acidic conditions, the storage mechanism involved with MnO2 has been extensively studied and results in a conversion reaction based on electrodeposition-electrodissolution processes in a mild acidic electrolyte that can be described by a highly reversible protoncoupled reaction: 15,16,58 ❖ Charging process: where AH and A -are the weak Bronsted acid and base from the acetate buffer electrolyte.

To determine the optimal concentration of manganese salt in the electrolyte, CV was used to study the electrodeposition/electrodissolution of MnO2 using a 1 M acetate buffer electrolyte containing different concentrations of Mn(CH3COO)2 (hereafter referred to as MnAce) of 0.1, 0.3 or 0.5 M (Figure 5.4). The potential window selected changes from 0 to Chapter V. Full cells 142 0.8 V with a scan rate of 10 mV.s -1 . A geometric surface area of 0.785 cm² of CNFs in contact with the electrolyte has been used. First, the plating reaction corresponding to equation 2, is observed on the CV by a sharp peak in the oxidizing current while the dissolution of MnO2 into Mn 2+ follows equation 1. This is characterized by a broad faradaic wave in the reducing current.

According to the CV experiments, the optimal concentration of MnAce is found to be 0.3 M as at this concentration, the cyclic voltammogram exhibited the largest capacity, both in reduction (discharge) and oxidation (charge). Indeed, when the MnAce concentration is too low (e.g., 0.1 M), there is an insufficient amount of Mn 2+ available for electron transfer and it is then depleted. 58,59 Conversely, at high concentrations, the Mn 2+ migration rate is slower, and the electronic conductivity may be compromised due to a large ohmic drop (IR) coming from the semiconducting nature of the thick MnO2 layer deposited. 58,59 Additionally, we confirm that this process is reversible as indicated in Table 5.1. Furthermore, the electrodeposition/electrodissolution of MnO2 on CNFs electrode was studied by galvanostatic charge-discharge conditions. CNFs are placed inside the three electrodes cell containing 1 M acetate buffer with 0.3 M of MnAce. A current density of 10.0 A.g -1 was used and ~ 0.5 cm -2 of CNFs was dipped into the electrolyte (mCNFs = 0.108 mg). The electroplating process was carried out for 5 min, followed by 5 min of electrodissolution, corresponding to a total charge of 648 mC.cm -2 . As depicted in Figure 5.5.a., the potential immediately reached a maximal value of 0.69 V vs Ag/AgCl, which is attributed to the generation of MnO2 nuclei on the CNF. The potential is then stabilized to a remarkably stable value of 0.62 V vs. Ag/AgCl, characteristic of the growth mechanism. During the following discharge process, a single and stable plateau at a potential value of 0.55 V is observed and attributed to the electrodissolution of MnO2 to soluble Mn 2+ (see eq. 5.1), as previously reported in this electrolyte. 16,33 The highly stable potentials for both the deposition and dissolution confirm that the conductivity of CNFs is sufficient to maintain fast kinetics and the electrochemical reactions occur almost under thermodynamic control, without significant pH and Mn 2+ gradient at the electrode/electrolyte interface. 16,33,59 Finally, a Coulombic efficiency of 99.96%, corroborates the full reversibility of the process with the complete dissolution of MnO2 during the discharge step.

We then investigated the morphology and compositional characterization of the electrodeposited materials. The crystallinity of the MnO2 deposited was performed by X-Ray diffraction. The ex-situ XRD pattern of the electrodeposited MnO2 exhibits two broad peaks centered at 2θ = 36.29° and 2θ = 65.30°, suggesting that the MnO2 is poorly crystalline (Figure 5.5.b). 60 Then, the surface morphology was examined by SEM-FEG after electrodeposition/electrodissolution processes (see and have shown that it is proportional to the concentration of Mn 2+ in the electrolyte and results in a thicker MnO2 layer around the fibers. 33 Interestingly, it has been shown by XPS that the average oxidation state (AOS) of Mn is 3.7, suggesting the presence of 30% Mn III and 70%

Mn IV . This value compares well with the one reported by Mateo et al. on 2D ITO and 3D

GLAD-ITO electrodes. 16,59 Accordingly, we will refer as far as MnO2 is concerned to a molecular mass of 87 g/mol. According to the fact that 1.7 electrons per Mn center are involved Chapter V. Full cells 145 10 and 20 A.g -1 . Accordingly, we have investigated the optimal current density for the plating/stripping of MnO2 with current density ranging from 8 A.g -1 to 15 A.g -1 . Briefly, we investigated the cyclability of the various TiO2 electrodes at a current density of 15.4 A.g -1 with a fixed proton intercalation at 154 mAh.g -1 (mild accumulation) in a half-cell On first reading, this plateau could be attributed to H2 production. However, the charge curve shows almost complete proton de-insertion and thus diverges from HER production. We have therefore attributed this plateau to proton insertion into the diversity of insertion sites (c.f.

Chapter IV. II.1. Intercalation process).

A possible hypothesis could be formulated regarding an eventual phase transition that could lead to a structural transition from the tetragonal to the orthorhombic phase. This phase transition has been observed in anatase compounds upon lithium insertion and recently confirmed upon proton insertion. 61 However, further investigations are required to confirm this attribution. Interestingly, Vac-TiO2 exhibits the largest stable working potential window (1.4~1.9 V) with the lowest discharge hysteresis evolution (0.08 V), allowing an increase in the energy and power density (IV.4. Energy and power density).

Pure anatase exhibits the HER plateau within the range of -1.2 to -1.3 V (Figure 5.10.b), which is the range where the proton intercalation is observed for the other titanate compounds, confirming that the HER phenomenon is highly related to the material's structure. not be optimal for materials like Vac-TiO2 that can reach more negative potentials. The Pt counter electrode exhibits gas evolution due to its strong adsorption affinity, leading to potential plateaus indicative of H2 or O2 production. For Vac-TiO2, which can intercalate down to -1.42

V and has an OER plateau at 1.89 V, the potential window extends to 3.3 V during cycling, enhancing both OER and HER kinetics (Figure S5.3). In contrast, pure Anatase reaches a potential of -1.29 V during the 10th cycle, resulting in a narrower potential window of 2.79 V due to the OER plateau at 1.5 V. This difference in gas evolution affects electrolyte consumption and degradation, influencing system reversibility. This is why, in full cell conditions, the potential window for each titanate compound never exceeds 2 V. This is based on the stable plateau centered at a mid-potential of 0.58 V during the electrodeposition/dissolution of MnO2 (Figure 5.5.a). Thus, this may explain the improvement in cyclability when the full cell is performed.

IV.4. Energy and power density

The energy density of a battery refers to the amount of energy that can be stored per unit volume (Wh.L -1 ) or per unit mass (Wh.kg -1 ). It represents the energy storage capacity of the battery relative to its size or weight. Here, it will be expressed in Wh.kg -1 and it corresponds to the capacity multiplied by the working potential window.

E(Wh. kg -1 ) = potential difference (V) * capacity (Ah. kg -1 ) (eq. 5.6)

The power density of a battery refers to the rate at which it can deliver electrical power.

It represents the ability of a battery to provide high power output, which is crucial in applications that require rapid charging or discharging, or high-power bursts.

P (W. kg -1 ) = ∆E (Wh.kg -1 )
∆t (h) (eq. 5.7)

Power density is typically measured in watts per unit volume (W.L -1 ) or watts per unit mass (W.kg -1 ). Here the power will be expressed in W.kg -1 , and it corresponds to the energy density during a definite period of time, for example during discharge or charge.

The energy and power densities were estimated for our "TiO2" materials and the results are summarized in Consequently, we compare our work with the systems in the literature, using Ragone plot representation. In this representation, the energy density is plotted versus the power density, with the rate at which the battery delivers power. Each data on the Ragone plot represents a specific battery system, and the location of the point corresponds to the energy density and power density values of that particular system. In

, 7795-7866. , we have illustrated where our systems will be located compared to others, and especially compared to Zn-ion in aqueous configuration. As we can see, the TiO2||MnO2 system is located at the top of the right corner, showing the promise of this system for future application. Typically, according to the values attained for power and energy densities, our TiO2||MnO2 battery could be used for short-term energy applications or stationary batteries.

Materials

IV.3. Self-discharge

In order to investigate the irreversibility of proton (de)intercalation observed during cycling, a self-discharge study was conducted. This study aims to understand the competitive reactions that consume protons and hinder the complete reversibility of the proton intercalation process of the positive electrode during charge in the full cell configuration.

Previous research by Makivic et al. 61 has shown the occurrence of a self-discharge reaction in anatase and amorphous TiO2, represented by the following equation:

2 Ti III OOH k self → 2 Ti IV O 2 + H 2 (eq. 5.8)
This reaction shows that the mechanism is purely chemical.

To evaluate the self-discharge characteristics of the titanate electrodes, a dedicated procedure was implemented. 10 cycles of discharging and charging the electrode at a rate of 10 This result provides further confirmation of the advantageous role of vacancies in the TiO2 network. One hypothesis that can be put forward is that the presence of vacancies strengthens the adsorption of protons within the structure, thereby inhibiting or minimizing the recombination step of the HER. Furthermore, these findings underscore the importance of conducting cyclability experiments at high current densities. The HER that takes place during the discharge process, as well as the HER resulting from self-discharge during the charge step, are relatively sluggish reactions. However, when the experiments are performed at high current density, these slower reactions are bypassed or diminished. 20 The high C-rate enables higher gravimetric loading capacity and Coulombic efficiency to be achieved.

To solve the problem of self-discharge and minimize its faradaic contribution, additional measures should be considered. The application of surface coatings or the incorporation of additives into electrode materials could help to reduce self-discharge. In addition, these results reinforced the need to carry out the cyclability experiment with high current densities. Additionally, elements such as copper and phosphorus were detected, indicating the dissolution of the crocodile twizzle. These elements appeared to be randomly aggregated on the electrode surface rather than being intercalated.

On the positive side (Figure 5.14.e), the formation of MnO2 did not exhibit the morphology observed in half-cell experiments.

Indeed, previous studies in our group have shown that MnO2 deposition occurs uniformly along the CNFs. 33 However, here, MnO2 was observed to form dispersed spherical aggregates or "balls" with an average diameter of 30.5 µm. Within these MnO2 balls, longitudinal marks similar to former nanofibers were visible. The mechanism behind the formation of these MnO2 balls is still under investigation. One hypothesis is that the rapid kinetics imposed by the high current density upon such long cycling resulted in the quick growth of MnO2, leading to a thermodynamically stable shape with reduced surface area, similar to spherical aggregates. Another hypothesis is the presence of Nafion 117 ® dispersed in the electrolyte that can play the role of surfactant that can subsequently modify the morphology of MnO2.

V. Conclusion

We have successfully integrated TiO2 electrodes with different structures into MnO2||TiO2 batteries using a buffered, weakly acidic electrolyte. As a part of our investigation, we studied the MnO2 electrodissolution/electrodeposition mechanism at the positive electrode.

To achieve this, we employed the electrospinning technique of polymers followed by a subsequent high-temperature thermal treatment, resulting in the fabrication of carbon felts consisting of conductive carbon fibers with a diameter of 150 nm, thus forming a porous network. Subsequently, we conducted a rapid analysis using cyclic voltammetry and galvanostatic charge-discharge techniques. Our findings demonstrate that these carbon felts exhibit a thermodynamically favorable and remarkably reversible electrodissolution/electrodeposition process. Therefore, we studied the rate capabilities of this reaction to couple it with titanates electrodes operating, for better reversibility, at high current density. We found that even in such conditions, the electrodissolution/electrodeposition mechanism remains stable.

Then, we evaluated the electrochemical behavior of these oxides in half-cell configurations and observed good Coulombic efficiency for all three materials (L-TiO2, Zn-Ace titanate, and Vac-TiO2).

Finally, the electrodes integrated in full cell (MnO2||TiO2) exhibited several key advantages: excellent cycling stability enduring 1000 cycles, high power density, and reduced self-discharge rate in the case of Vac-TiO2, highlighting the advantage of the presence of vacancies in the TiO2 structure.

In perspective, we aim at further improving electrode engineering, particularly by optimizing the integration of Nafion into the electrode design, to enhance the battery's performance and overall efficiency.

General Conclusion and Outlooks

To overcome the inherent limitations stated by the constrained energy density of aqueous batteries, in this Ph.D. work, we have investigated the proton charge storage mechanism in nanostructured TiO2 compounds. We have investigated the complex interplay of material structures, and proton transport mechanisms in interaction with a mild acidic buffer electrolyte as a proton source, to propose more efficient, sustainable, and high-performance energy storage solutions.

The fundamental challenge intrinsic to aqueous batteries is the limited energy density.

The occurrence of parasitic gas production, arising from hydrogen and oxygen evolution reactions, diminished the overall battery performance. This is particularly true for dihydrogen formation, which generally occurs at the negative electrode. Among the short list of materials that can serve as negative electrodes under aqueous conditions, we have explored TiO2 as a promising candidate, exploiting its proton-coupled electron transfer mechanism for bulk proton insertion. Incorporating defects in 3D TiO2 structures revealed the possibility of achieving a high capacity of 259 mAh.g -1 , and tuning nanoconfined water in layered titanate materials improved proton diffusion kinetics to 1.02 × 10 -12 cm 2 .s -1 .

Initially, our research focused on the synthesis and morphological changes of TiO2 structures. We observed the transformation of amorphous layered lepidocrocite-type TiO2 into anatase through topotactic dehydration. By carefully adjusting the temperature, we were able to stabilize and investigate a metastable TiO2 intermediate phase characterized by cationic vacancies. This phase, referred to Vac-TiO2, was synthesized at a low temperature of 110 °C.

Our microstructural analysis confirmed the presence of 17% vacancies in the 3D open structure of Vac-TiO2. Moreover, broadband dielectric spectroscopy revealed enhanced proton ion conductivity in this sample (1.34×10 -4 S.m -1 ) compared with lamellar lepidocrocite-type TiO2 (8.1×10 -5 S.m -1 ) or stoichiometric TiO2 (1.86×10 -5 S.m -1 ).

Layered lepidocrocite-type TiO2 initially exhibited low ionic conductivity. To enhance its conductivity, we revisited the solvothermal synthesis method and introduced zinc ions into the interlayer spaces. These zinc ions can coordinate with up to three crystalline water molecules in the absence of counter-anions. This capability can significantly influence the water and hydrogen network, which in turn has an impact on proton behavior. While the layered structure remained largely unchanged, the incorporation of zinc ions, as demonstrated in Chapter 3, significantly increased overall protonic conductivity and improved proton hopping behavior. This increased proton mobility can be attributed to reduced energy barriers for proton jumps due to the presence of zinc ions with 0.49 and 0.18 eV for L-TiO2 and Zn-Ace titanate respectively. This translates directly into improved proton electrochemical properties of the protons, resulting in higher capacity to 212 mAh.g -1 , better reversibility, and a wider working potential window extending down to -1.4 V (vs Ag/AgCl, KCl sat').

The influence of vacancies in Vac-TiO2 on the electrochemical properties of the proton was investigated in Chapter 4. This material showed great promise in enhancing both proton insertion capacity, reaching up to 259 mAh.g -1 , and extending the potential window up to -1.6 V (vs Ag/AgCl, KCl sat'). The significant increase in capacity is directly linked to the presence of vacancies. However, further research is needed to investigate the mechanisms underlying the extension of the potential window. Additional studies should explore the optimal number of vacancies and their effects on the proton binding energy on the oxygen network.

Finally, our study led to the integration of TiO2 electrodes in MnO2||TiO2 batteries using a weakly acidic buffered electrolyte as a proton source. This positive electrode integration was achieved by fabricating carbon felts composed of conductive carbon fibers forming a porous network structure, employing the electrospinning technique of polymers and heat treatment in a controlled atmosphere at a very high temperature. We then conducted a comprehensive analysis using cyclic voltammetry and galvanostatic charge-discharge techniques. Our research revealed that these carbon felts facilitate a highly favorable and remarkably reversible electrodissolution/electrodeposition process for the Mn 2+ /MnO2 reaction. On this basis, we explored the rate capabilities of this electrochemical reaction, optimizing reversibility by coupling it with titanate electrodes operating at high current density. Remarkably, our research has shown that even under high-rate conditions, the electrodissolution/electrodeposition mechanism remains stable. We then integrated the titanate electrodes into full cells. These full cells offer several notable advantages, including remarkable stability over more than 1000 cycles and a noticeable energy density of around 99 and 105 Wh.kg -1 , with a power density of 9.9 and 10.5 kW.kg -1 for optimized materials, Zn-Ace titanate and Vac-TiO2. Notably, we also observed a faradaic contribution leading to a self-discharge reaction involving the chemical formation of H2, which had an impact on battery performance.

Collectively, this research highlights the critical role of material design, structural engineering, and electrochemical mechanisms in advanced energy storage devices. The study highlights the potential for improving proton conductivity by exploring interlayer water modulation in TiO2 structures. Additionally, the study of the chemical intercalation of other cations could lead to finely tuned materials with specific properties. The incorporation of cationic vacancies has proved to be an effective strategy for increasing proton capacity, and the optimum concentration needs to be investigated.

However, certain challenges remain. To minimize secondary reactions, the application of surface coatings such as Nafion or additives could help address cycle durability issues, including difficulties such as self-discharge, and improve overall performance. Additionally, the use of highly concentrated salt in the electrolyte is also a relevant strategy for minimizing the electrocatalysis of water and thus improving the efficiency of proton-ion batteries.

This work has provided valuable insights into the complexities of aqueous batteries, highlighting the challenges and opportunities for their development. These chapters have established the foundations for a deeper understanding of materials, structures, and mechanisms, paving the way for sustainable and high-performance proton ion batteries. The research results offer promising prospects for the future of energy storage systems that are both efficient and environmentally friendly.

General Conclusion and Outlooks

Schematic representation of MEMS chip used in TEM in-situ

The sample preparation involved grinding and dispersion in ethanol (EtOH), followed by drop-casting 1 drop (1.5 µl) of the solution onto the top heating chip. The closed cell was assembled using a second chip with a central SiNx membrane, forming a sealed environment with O-rings to maintain a high vacuum within the TEM column. A gas pressure of 1 bar was achieved using synthetic air under static conditions inside the cell. The heating process involved incremental steps from room temperature up to 500 °C, with a heating rate of 5 °C.min -1 . To minimize the influence of the electron beam, only one specific region was irradiated during the thermal ramps to monitor changes through video recording. Additionally, regions that were not observed during the in-situ experiment were characterized afterward to investigate possible changes induced by the electron beam.

A.5: Rietveld refinement

Rietveld refinement is a powerful technique used in the analysis of X-ray or neutron diffraction data to determine the crystal structure and refine the atomic positions in a crystalline material. This method combines the measured diffraction data with a structural model to obtain Where βh is the size contribution to the integral breadth of reflection h, are the real spherical harmonics with normalization as in [M. Jarvinen, J. Appl. Cryst. 26, 527 (1993)]. 3 The arguments are the polar angles of the vector h concerning the Cartesian crystallographic frame. After refinement of the coefficients, the program calculates the apparent size (in angstroms) along each reciprocal lattice vector if the IRF is provided in a separate file.

A.6: Scattering density: 4 The conventional equation used to compute the structure factor relies on a model that involves point scatters, with their positions in the x, y, and z dimensions influenced by thermal motion:

Fhkl = Σ n bn Nn exp (2πi (hxn + kyn + lzn)) exp (-Bn sin 2 θ / λ 2 )

Here, F represents a vector encompassing both the amplitude (F) and the phase (φ).

Theoretical foundations dictate that the structure factor constitutes the Fourier transform of the scattering density, denoted as ρ(x, y, z), which encompasses electron density for X-ray scattering and nuclear density for neutron scattering, across the entire unit cell. In other words, it is calculated by integrating: Conventional X-ray diffraction focuses on Bragg peaks in reciprocal space to glean structural details from materials. Yet, this approach demands structural periodicity and doesn't suit disordered systems or nanocrystals. 5 An alternative, the atomic pair distribution function (PDF), employs total scattering to analyze both Bragg and diffuse scattering data, offering insights into short-range order.

Fhkl = c ∫ 0 b ∫ 0 a ∫ 0 ρ(x,
The PDF necessitates a substantial momentum transfer Q for obtaining high-resolution data. Q is defined as a function of the wavelength λ and the angle θ, where 2θ represents the angle between the incident and diffracted X-rays: 𝑄 = 4𝜋𝑠𝑖𝑛𝜃 𝜆 Conventional X-ray diffractometers utilizing Cu Kα radiation (λ = 1.54 Å) can only achieve a Q value of approximately 8 Å -1 , which is inadequate for PDF analysis. 6 As a result, experiments necessitate the use of high-energy synchrotron X-rays. Typically, Q values in the range of 20 to 40 Å -1 yield satisfactory data. 6 For our study, measurements were conducted using high-energy X-rays (λ = 0.2128 Å) at the 11-IDB beamline located at the Advanced Photon Source of Argonne National Laboratory.

The experimental atomic PDF G(r) can be derived through the Fourier transform of the total scattering structure function S(Q), which is obtained after correcting the experimental diffraction intensity I(Q): The atomic PDF can also be described using the atomic pair density ρ0, the atomic number density ρ(r), and the radial distance between a pair of atoms r:

𝐺(𝑟) = 4𝜋𝑟[𝜌(𝑟) -𝜌 0 ]
G(r) is represented by peaks that indicate the number of atomic pairs at a specific distance r, within a spherical shell of unit thickness, formed by two atoms. 8

List of Figures

Figure 1 . 1 :

 11 Figure 1.1: ESW of pure water and compatible electrode material adapted from Kim, H. et al. 6

Figure 1 . 2 :

 12 Figure 1.2: (a) Elemental abundance in the Earth's crust. (b) The comparison of molar masses of different ionic charge carriers. (c) The radius comparison of various ions, where a naked proton is too small (0.89 fm) to be shown. (d) A schematic of the Grotthuss mechanism. (e) A timeline of representative proton battery research. Adapted from Xu, Y. 2021. 1

Figure 1 . 4 :

 14 Figure 1.4: Solvation sheath of salt in water-electrolyte (left) and water in salt electrolyte (right). Adapted from Suo 2015 97

  where a proton can hop from one molecule to another (Figure1.2.d).When a water molecule donates a proton to an adjacent water molecule, a hydronium ion is formed. The next step involves the hopping of the proton from the hydronium ion to a neighboring water molecule. This proton transfer is facilitated by the shared hydrogen bonds between adjacent water molecules. This way of conduction can have a significant role in proton insertion reversibility during electrochemistry, and such material has not yet been explored for proton ion insertion.

Figure 1 . 5 :

 15 Figure 1.5: (left) Cationic defective fluorinated TiO 2 anatase with calculated intercalation energy of Li + , Mg 2+ , and Al 3+ . Adapted from Koketsu 2020 106 ; (right) tuning of a layered lepidocrocite type TiO 2 with zinc ions. Adapted from Kang 2022[START_REF] Wang | Hydrogen Production from Water Electrolysis: Role of Catalysts[END_REF]

3. 33 (

 33 see synthesis protocol in Annexes A.1). The precipitation was achieved under solvothermal conditions at temperatures of 90, 110, and 150 °C for 12 hours. The samples obtained at these temperatures were investigated using powder X-Ray Diffraction (XRD) to analyze their structural variations. At 90 °C (Figure 2.1.a), the XRD pattern displayed diffuse scattered intensities without characteristic Bragg peaks, indicating the amorphous nature of the compounds. At 110 °C, the presence of Bragg peaks revealed crystallization with a tetragonal symmetry characteristic of TiO2 anatase, and the absence of diffuse scattered intensities indicated complete crystallization. As the temperature increased to 150 °C, the Bragg peaks became more intense and sharper, signifying the growth of anatase crystals. To confirm the structural variation with temperature, High-Resolution Transmission Electron Microscopy (HRTEM) was employed (Figure 2.1.a, Figure 2.1.b, and Figure 2.1.c).

Figure 2 . 2 .

 22 c), resembling the zigzag configuration observed in anatase. However, it contains two inequivalent O 2-sites of the same multiplicity, 2-fold and 4-fold coordinated to Ti 4+ cations. On the other hand, the anatase crystal structure is tetragonal (Figure 2.2.d), featuring a three-dimensional arrangement of TiO6 octahedra. These octahedra are interconnected through an edge-sharing configuration, wherein each O 2-ion is coordinated with three neighboring Ti 4+ ions. Notably, the first two peaks in the PDF curve, corresponding to Ti-O bonds within the TiO6 octahedra and Ti-Ti sharing the edges, are similar in all compounds, demonstrating the structural similarity between the titanate layer in Lepidocrocite type-TiO2 (hereafter named L-TiO2) and the three-dimensional framework of anatase.

  charge compensation. With the titanate layers being electrically neutral, each vacancy is stabilized by four OH groups. The chemical composition of L-TiO2 within the experimental margin error is Ti1.48□0.52O1.94(OH)2.06 • 0.69H2O, which is close compared to the reported layered Lepidocrocite type TiO2 with Ti1.5□0.5O2(OH)2 • 0.55H2O.14

Figure 2 . 3 :

 23 Figure 2.3: Thermogravimetric curve and calorimetric analysis of L-TiO 2 Interestingly, the PDF refinement of L-TiO2 confirms the presence of both oxide and hydroxide close to Ti 4+ , as observed in the first peak. The Gaussian fit of this peak (Figure S2.3) distinguishes two sets of distances, 1.881 and 2.038 Å corresponding to Ti-O and Ti-OH interatomic distances, respectively. The approximately equal peak areas suggest the equiproportion of the two bonds, which is consistent with the established chemical formula.

Figure 2 . 5

 25 Figure2.5 presents the thermal X-ray powder diffraction pattern of L-TiO2, displayed as a two-dimensional mapping color view (Annexes A.3). The diffraction pattern does not reveal any crystalline phases at temperatures up to approximately 300 °C. However, at 325 °C, the first Bragg peak corresponding to the anatase structure, specifically the (101) lattice peak, emerges. As the temperature rises further, the intensity of the anatase pattern is progressively strengthens. It is worth noting that up to 500 °C no transition to a thermodynamic Rutile or Brookite phase is notable, which is generally formed around 600 °C for Rutile with pure synthetic titanium oxide and between 400~800 °C for Brookite.33,34

Figure 2 . 5 :

 25 Figure 2.5: Thermal X-Ray Diffraction pattern of L-TiO 2

andFigure 2 . 7 :

 27 Figure 2.7: TEM images with (a) in situ gas phase of L-TiO 2 , and (b) post-mortem HRTEM of L-TiO 2.

intervals ranging from 1

 1 to 72 hours. The X-ray diffraction patterns of the obtained precipitates are shown in Figure 2.8.a. Interestingly, for up to 4 hours, the precipitate remains amorphous, and it takes 6 hours for the thermal activation to start the crystallization, resulting in the appearance of initial broad and weak Bragg peaks. Subsequently, as the reaction proceeds for Chapter II. Synthesis and characterization of titanate nanomaterials via soft chemistry method 51 III. Microstructure analysis of Vac-TiO 2

  Vac-TiO2 exhibited a slightly larger lattice volume (a*b*c) with 136.49 Å 3 compared to 136.13 Å 3 in Pure Anatase (Figure S2.6). This variation is mainly due to longer a and b lengths in Vac-TiO2 compared to Pure Anatase. On the other hand, Pure Anatase exhibited a longer -c-length and shorter a/b lengths.

14 )

 14 By plotting β* = βcos(θ)/λ as a function of d* = 1/d = (2sin(θ)/λ), we can construct the Average Size and Strain plot (ASS) or the Williamson-Hall (WH) plot. In Figure 2.10.a the WH plots of Vac-TiO2 displays a highly dispersed distribution of the breadth peak as a function of the two-theta angle in both materials, indicating the presence of size and strain factors that depend on the crystallographic planes. This -hkl-dependent breadthens allows us to identify two crystallographic directions, (h00) and (00l). We can observe a noticeable slope in both directions, indicating the participation of strain. Conversely, for Pure Anatase (Figure S2.7),

Figure 2 .

 2 13.a), the obtained plot shows the existence of a second relaxation P2 attributed to proton hopping (Figure 2.13.b). After subtracting contributions P1 and P2, the obtained plot highlights the Chapter II. Synthesis and characterization of titanate nanomaterials via soft chemistry method 60 relaxation P3 (Figure 2.13.c) which is thus attributed to water rotation. By fitting P2 and P3

  synthesized three starting titanate compounds via the sol-gel method. An amorphous layered lepidocrocite type TiO2 characterized by stacked layers of TiO6 octahedra separated by crystal water molecules. Through thermal treatment, a topotactic dehydration process occurred, leading to layer collapse and rearrangement, ultimately transforming L-TiO2 into anatase. The crystallization mechanism involves the precipitation of L-TiO2, followed by the departure of adsorbed and crystal water and the condensation of the layers to adopt the anatase structure. At an intermediate temperature (110 °C), the resulting Vac-TiO2 compound maintains a partial occupancy of vacancies, indicating the conservation of defects during the transformation from a 2D layered structure to a 3D open framework, and an almost defect-free pure Anatase synthesized at a higher temperature (150 °C).

Figure S2. 3 :Figure S2. 4 : 66 Figure S2. 5 :

 34665 Figure S2.3: Gaussian fits of the first peak of the PDF data on L-TiO2

Figure S2. 7 :

 7 Figure S2.7: WH plot and ASS plot of pure Anatase

  . The compositional analysis of the resulting ZnCl2 titanate compound yielded the formulation: Zn0.10Cl0.20Ti1.27□0.73O1.11(OH)2.89 • 0.16H2O. In addition, the Energy X-ray Adsorption Spectroscopy (EXAS) indicates that Zn 2+ ions are closely coordinated with both oxide and chloride ions (which suggests a coordination with one other crystal water possible). This unique coordination environment has yielded a remarkable 450fold enhancement in the bulk proton conductivity of the material. The pivotal finding of the study resides in the profound influence of zinc ions on the hydrogen bond network, thereby facilitating proton hopping mechanisms. One important thing to consider is that the study encountered difficulty in attributing all of the observed effects to the presence of Zn 2+ ions alone. This issue emerged due to the presence of chloride anions within the structure, which may have influenced the observed outcomes. The interplay between the introduced zinc ions and the associated chloride anions introduces complexity to the observed effects, necessitating further investigations to clarify the distinct and combined impacts of these constituents.

(

  PDF) technique. As presented in Figure3.3, the PDF data for both L-TiO2 and Zn-Ace titanate share similar peak positions, corroborating the structural affinity between these two compounds. Moreover, the coherence length-associated with crystallinity-does not exceed Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 79 10 Å for both L-TiO2 and Zn-Ace titanate compounds. This shortened coherence length is indicative of the concise order intrinsic to these materials. In consequence, we can confidently affirm that the lepidocrocite type TiO2 structure remains even upon the incorporation of zinc ions.

Figure 3 . 3 : 80 (

 3380 Figure 3.3: Pair distribution function analysis for L-TiO 2 (black line), and Zn-Ace titanate from 10 to 30% (dark red to light red line). The introduction of Zn 2+ ions into the material induces discernible although minor modifications. Focusing on 20% Zn-Ace titanate, specific changes in the 1 st and 3 rd peaks of the histogram are evident, corresponding to Ti-O(H) and Ti-Ti sharing corners within the TiO6 octahedra. Notably, the intensity of the 1 st peak (~ 1.98 Å) increases upon zinc incorporation. This peak can be further deconvoluted into two separate peaks via Gaussian fit, into Ti-O and Ti-OH at approximately 1.881 and 2.038 Å, respectively (Figure S3.6). In the L-TiO2 sample, the proportion of these two interatomic distances is nearly equivalent. However, for 20% Zn-Ace titanate, the fits reveal a proportion of Ti-OH and Zn-(OH)2. This observation aligns well with the chemical formula derived for Zn-Ace titanate, indicating a greater abundance of OH - ions in comparison to pure L-TiO2 or the Lepidocrocite type titanate reported by Ma et al. 5

ε 1 )

 1 (ω) and conductivity s(ω) linked by s(ω) = iωε0ε(ω) where ε0 is the vacuum permittivity (ε0  8.854×10-12 F.m -1 ) and ω the angular frequency (ω = 2πν with ν the frequency).Kang et al. have already demonstrated that the proton is the charge carrier on Lepidocrocite type TiO2 and ZnCl2 titanates and they are provided from the self-dissociation of H2O molecules following a Grotthus diffusion mechanism.6 Nyquist plots of complex permittivity ε" vs. ε' show that the decomposition of the dielectric spectra highlighted five relaxations (Figure3.4). These last ones are fitted with Cole-Cole functions represented by circular arcs centered below the ε'-axis in the complex plane (Figure3.4). The complex permittivity of Zn-Ace is thus the sum of Cole-Cole functions given by:ε (ω)=ε ∞ + ∑ [Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 81 with m = 1 to 5 numbering in order of the relaxations Pm from the lowest (i.e. P1) to the highest (i.e. P5) frequencies. ε is the lattice permittivity of the compound, Δεm the dielectric strength (i.e. width of the relaxation on the ε'-axis), τm the relaxation time, and αm a fitting parameter (0 ≤ αm ≤ 1). The relaxation time τm is a measure of the timescale of dipolar and charge carriers' motions. Consequently, the relaxation frequency νm = (2πτm) -1 is higher, the smaller the object concerned (i.e. from macroscopic to interatomic sizes) and the faster the motions. The exponent αm indicates the degree to which the system deviates from the Debye model and the existence of a distribution of relaxation times (frequencies) that is wider the further αm deviates from zero.As the sample is a compact set of grains, three sources of polarizations result from charge (protons) accumulations at different interfaces (grain boundaries, sample surface, sample/metal interface). In these three cases, their dynamics are linked to the mobility of the charges (protons). The low-frequency part of the Nyquist plot of permittivity (Figure3.4.a) is wellfitted by a giant relaxation P1 whose characteristic frequency ν1 is equal to 3 Hz at room temperature. This relaxation P1 is attributed to the existence of an ionic double-layer at the silver paint/sample interface. After subtracting the contribution P1 (Figure 3.4.b), the obtained plot shows the existence of a second relaxation P2 due to the sample polarization with a characteristic frequency ν2 equal to 1.7 × 10 4 Hz at RT. After subtracting contributions P1 and P2, the obtained plot highlights the relaxation P3 (Figure 3.4.c) which is thus attributed to grain (bulk) polarization with a characteristic frequency ν3 equal to 6 × 10 5 Hz at RT. The last two relaxations P4 and P5 (Figure 3.4.d & Figure 3.4.e) thus result from local motions within the grain (bulk) and due to proton hopping and water rotation, respectively (Table

  Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 86 which can affect the proton hopping in the interlayer. The lack of counter-ions and the full zinc coordination with crystal water in the interlayers is supported by the chemical composition studies.

  Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 88 KCl and acetate electrolytes are depicted in Figure3.7.a, with the dashed curves representing each case. In the KCl electrolyte, the CV shape demonstrates a nearly rectangular profile characteristic of a double-layer capacitance. The calculated capacity values was measured at 1.5 mAh.g -1 (equivalent to 0.56 µF.g -1 ). This minor contribution implies the absence of insertion phenomena within this potential range, considering the limited availability of protons. To identify the occurrence of proton intercalation in L-TiO2, CV experiments were carried out using the 1 M acetate buffer electrolyte. As illustrated in Figure3.7.b, the current response is significantly altered, demonstrating a distinct transition from an insulating to a conducting state.

89 As Equation 3 . 4

 8934 III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries suggests, a Nernstian behavior is expected when the slope per unit of pH corresponds to the theoretical value of 0.059, a fundamental constant related to Faraday's law of electrolysis. Additionally, the potential of the conduction band can be determined from the point of intersection. Our experimental approach consisted of performing cyclic voltammetry at a scan rate of 20 mV.s -1 over a wide range of pH levels. The electrolytes used spanned from HCl for pH 0 to Carbonate buffer at pH 9.3, with intermediary pH values covered by Chloroacetate buffer (pH 3.0), Acetate buffer (pH 5.0), and HEPES buffer (pH 7.0). All of these electrolytes act as proton donors, serving as a source of protons for the intercalation process. The linear regression of the data from Equation3.4 was performed using the oxidizing interstitial peak in L-TiO2, as depicted in Figure S3.7. Remarkably, the obtained slope was found to be 0.059 V.pH -1 unit, aligning precisely with the theoretical value predicted by the Nernstian pH-dependent equation under standard conditions Figure 3.7.b. This observation

Figure 3 .

 3 8.b, demonstrated that the Zn-Ace titanate electrode exhibited minimal HER contribution even at notably negative potentials such as -1.4 V. Notably, at a potential of -1.4 V, a maximum capacity of 181 mAh.g -1 was achieved for proton de-intercalation, accompanied by a columbic efficiency of approximately 93% Figure 3.8.d. This value is drastically higher than the one found in L-TiO2 (125 mAh.g -1 , CE = 90%, Figure 3.8.a & Figure 3.8.c) and is higher to that observed for amorphous TiO2 (144 mAh.g -1

4 V

 4 at the end of discharge. The ability of Zn-Ace titanate to facilitate proton intercalation led to a shift in the discharge curve towards more cathodic potential, thus moving away from the HER potential. Consequently, the Coulombic Efficiency (CE) during the first cycle reached 92% for Zn-Ace titanate whereas it Chapter III. Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 95 was 72% for L-TiO2. Nevertheless, at these elevated capacities, the electrode's cyclability was markedly compromised, with rapid and abrupt capacity degradation becoming evident in subsequent cycles. (Figure 3.10.c).

Figure 3 .

 3 10.b by the potential values reached at the end of chronoamperometry: -1.16 V for Zn-Ace titanate and -1.22 V for L-TiO2. The intercalation process in Zn-Ace titanate occurred slightly at more cathodic potentials compared to L-TiO2. The reversibility shown at the 2nd cycle, under such mild accumulation conditions, was greatly improved for both materials, with Coulombic efficiencies of 97.4% and 89.6% for Zn-Ace titanate and L-TiO2, respectively. This underscores the benefits of performing GCD under mild accumulation protocols. The reversibility of both materials was further explored through multiple cycles of GCDs. Figure 3.10.d illustrates the cyclability under mild accumulation conditions. While L-TiO2 exhibited moderate cyclability with rapid and consistent capacity fading, Zn-Ace titanate demonstrated exceptional cyclability. It maintained Coulombic efficiency above 90% for the

Figure S3. 3 :Figure S3. 7 :Figure S3. 8 :

 378 Figure S3.3: Infrared spectrums of L-TiO 2 (red curve), Zn-Ace titanate 20% (black curve), and Zn-Acetate precursor (light blue curve)

Figure 4 . 1 :

 41 Figure 4.1: Structural representation of intercalation site (left to right) defective free, single vacancy, and double vacancy anatase. (Bottom) DFT calculation of intercalation energy Li + , Mg 2+ , and Al 3+ . Adapted from: D.Dambournet Acc. Chem. Res. 2022, 55, 696-706.4

Figure 4 . 2 .

 42 b, the current response is strongly modified, manifesting the familiar transition from an insulating to a conducting state. As observed earlier, the quasisteady-state current (iC) attained at potentials exceeding -0.5 V (i.e., E > ECB = conduction band potential) is directly related to the electrode's double-layer capacitance. When E << ECB, the current density substantially increases, accompanied by the appearance of a broad faradaic wave ranging from -0.65 to -1.2 V. Additionally, a shorter faradaic wave emerges within the range of -1.2 to -1.6 V. This increase in current density compared to that obtained in the Chapter IV. Defective TiO2: pushing the limit of HER 110 unbuffered electrolyte primarily signifies the occurrence of a faradaic intercalation process (here = discharge). This phenomenon can be attributed to the faradaic reduction of Ti IV concomitant with the intercalation of protons, as described by equations 4.1 and 4.2.7,8

- 1 . 6 V 3 )

 163 , which is notably lower than what was observed for pure Anatase (-1.4 V, Figure S4.1).This shows the straight benefit of this Vac-TiO2 electrode toward the direct irreversible reduction of water to H2 at the TiO2/electrolyte interface. To reinforce this observation toward the carry-over of HER, we estimated the formation of HER formation at each cut-off. By considering that the reduction in capacity is exclusively attributed to the irreversible HER contribution, the HER capacity (QHER) was calculated through the subtraction of both integrated anodic and cathodic current densities: QHER = QDischarge -QCharge (eq. 4.Qcharge corresponds to proton de-intercalation and Qdischarge to proton intercalation within the TiO2 electrode. The results, summarized in Figure 4.2.c, highlight the minimal HER contribution within the Vac-TiO2 electrode, even at remarkably negative potentials such as -1.6

Figure 4 . 3 :

 43 Figure 4.3: CV performed at 20 mv.s -1 in 1 M acetate buffer electrolyte of L-TiO 2 (black curve), Vac-TiO 2 (red curve), and pure Anatase (blue curve).

Figure 4 . 4 :

 44 Figure 4.4: Quantification of the Nernstian relationship with pH, where the E interstital oxidizing peak of Vac-TiO 2 varies linearly with pH.

Figure 4 . 5 .

 45 a reveals the simulated relaxed structures of both Vac-TiO2 and pure Anatase upon intercalating a single proton.

Figure 4 . 5 .

 45 b, the proton demonstrates a clear affinity for oxygen located within a cationic vacancy. The associated intercalation energy is computed as -1.66 eV.ion -1 , which is more favorable than the one found in pure Anatase. This result is linked to the TiO2 architecture, characterized by TiO6 octahedra with a 3-fold coordinated O 2-. The introduction of monovalent hydroxyl groups, resulting from the

Figure 4 . 6 :

 46 Figure 4.6: Log-log plot of Vac-TiO 2 showing the linear regression to obtain the b-value of interstitials site (black square) and vacancies sites (red dot)Therefore, we attempt to determine the H + ion diffusion in both interstitial and vacancy sites within Vac-TiO2. To roughly estimate the apparent solid-state proton diffusion coefficient, we used the Randles-Sevcik equation assuming ideal conditions. The Randles-Sevcik equation describes the relationship between the current and the scan rate potential:

(Figure 4 . 2 .

 42 b), a notable discussion arises to investigate the potential presence of a biphasic composition in Vac-TiO2. To study the potential contribution of an amorphous phase to the proton intercalation mechanism, several analyses were conducted. Firstly, refinement of PDF analysis on Vac-TiO2 was performed to examine the presence of an amorphous signature, the results indicated that Vac-TiO2 could be described only with a single phase of anatase (ChapterII. II.2. Local structure analysis using Pair-Distribution Function (PDF)). Additionally, theChapter IV. Defective TiO2: pushing the limit of HER 120 absence of diffuse scattering in the XRD pattern corroborates the complete crystallization of the material.

Figure 4 . 7 .

 47 a effectively presents the electrochemical responses, revealing the simultaneous presence of interstitial and vacancy faradaic contributions in this particular mixture. It is noteworthy, however, that disparities emerged when comparing this mixture to Vac-TiO2, particularly in the context of the hydrogen evolution reaction (HER) contribution. In particular, HER was observed immediately after proton insertion in the pure Anatase phase, in parallel with the observations in the cut-off study of pure Anatase. Moreover, the interstitial contribution within the mixture exhibited a sharper profile compared to the broader potential range characterizing proton interstitial insertion in Vac-TiO2. These distinctions suggest that while the proton intercalation mechanism in Vac-TiO2 is closely linked to pure Anatase, discernible disparities highlight Vac-TiO2's distinctive nature. Raman spectroscopy was also employed to compare the structural characteristics of Vac-TiO2 and pure Anatase. By focusing on the first and high-sensitivity normalized Eg peak (143 cm -1 ), a striking contrast between Vac-TiO2 and the 7:3 (A:L) mixture in terms of intensity and FWHM values can be observed (Figure 4.7.b). The Raman signature of the mixture closely aligns with that of pure Anatase, whereas Vac-TiO2 exhibits distinct intensity and FWHM values. This observation harmonizes with the absence of Raman bands associated with Lepidocrocite, which are represented by broad features (Figure S4.5). This collective evidence further reinforces the conclusion that Vac-TiO2 embodies a distinctive material identity.

Figure 4 .

 4 8.b. The potential reached -1.25 V for Vac-TiO2 without significant HER plateau formation. Despite this protocol, in the Chapter IV. Defective TiO2: pushing the limit of HER 123 case of pure Anatase, the discharge curve reaches -1.3 V and coincides with HER. This distinction highlights the favorable proton intercalation in Vac-TiO2, in contrast to pure

Figure 4 . 8 .

 48 c illustrates the cyclability under mild accumulation conditions. Vac-TiO2 demonstrated exceptional cyclability, with Coulombic efficiency (CE) reaching 97.1% by the 6 th cycle and consistently staying above 90% up to the 48 th cycle before a gradual capacity fading. Interestingly, pure Anatase exhibited a moderate level of reversibility, maintaining a CE of 75% for over 50 cycles before displaying a gradual decline in capacity retention.

Figure 4 . 9 :

 49 Figure 4.9: HER and OER catalyst on Pt wire (red curve) in galvanostatic mild accumulation condition for Vac-TiO 2 (left) and pure Anatase (right)

Figure 5 .

 5 Figure 5.3.a.

Figure 5 . 3 .

 53 c shows the CV experiments of CNFs in half-cell (counter: Platinum (Pt) wire, ref: AgAgCl, KCl sat')

oo 1 )❖o

 1 Electrons were given by the generator to the positive electrode, here the CNFs, o Simultaneously, the accumulation of negative charge induces the attraction of protons to the electrode surface, The combination of protons and electrons leads to the reduction of MnO2 species present at the electrode, resulting in the electrodissolution of MnO2 to Mn 2+ : 16 MnO2 + 4 AH + 2 e -→ Mn 2+ + 4 A -+ 2 H2O (eq. 5.Discharging process: Electrons are provided by the positive electrode to the negative electrode o Simultaneously, protons are attracted to the negative electrode from the dissociation of water molecules o It results in the electrodeposition of Mn 2+ to MnO2: Mn 2+ + 4 A -+ 2 H2O → MnO2 + 4 AH + 2 e -(eq. 5.2)

Figure 5 . 4 :

 54 Figure 5.4: impact of MnAce concentration on electro-deposition/dissolution realized by cyclic voltammetry from 0 to 0.8 V at 10 mV.s -1

Figure 5 . 5 .

 55 c and Figure 5.5.d). As we can see, MnO2 with a sheet-like structure uniformly coats the CNFs. After complete discharge, we observe a complete dissolution of the MnO2 as the initial pristine morphology of the CNFs is achieved. Singh et al. have studied the electrodeposition/electrodissolution of MnO2 on CNF

Figure 5 . 6 illustrates

 56 the capacity of electrodissolution as a function of current density with a fixed electrodeposition of 100 mAh.g -1 . This figure shows that in the range of 8 to 15 A.g -1 , the electrodeposition/electrodissolution reaction is highly reversible with a CE close to 100%. This set of experiments indicates that this range of current densities is compatible with the high kinetics reaction of H + in TiO2. However, note that at current densities higher than 16 A.g -1 , the dissolution reaction efficiency decreases with a CE  90%; an accumulation of MnO2 at the CNFs is observed (FigureS5.2).33 III. Electrochemical properties of Ti-based compounds in half-cell configurationTo realize full-cell and half-cell, we follow the same process of electrode fabrication seen in the previous chapter (Chapter III. IV.1. Electrode preparation and choice of electrolyte).

Figure 5 . 6 : 1 Chapter V. Full cells 146 condition in 1 M 5 . 7 .

 561146157 Figure 5.6: Rate capability of MnO 2 plating/stripping from 8 to 15 A.g -1

Figure 5 . 7 :IV. 1 . 5 . 8 .

 57158 Figure 5.7: Cyclability performances of L-TiO 2 (black sphere), Zn-Ace titanate (green sphere), Pure Anatase (blue sphere), Vac-TiO 2 (red sphere) and with fixed intercalation of 154 mAh.g -1 (grey sphere).

Figure 5 . 9 :Figure 5 . 10 :

 59510 Figure 5.9: GCD cyclability of Full-cell with (a) L-TiO 2 , (b) Zn-Ace titanate. mAh.g -1 of active titanate materials

Figure 5 . 11 :

 511 Figure 5.11: Cyclability performances of L-TiO 2 (black sphere), Zn-Ace titanate (green sphere), Pure Anatase (blue sphere), Vac-TiO 2 (red sphere) and with fixed intercalation of 100 mAh.g -1 (grey sphere). mAh.g -1 of active titanate materials

Figure 5 . 12 :

 512 Ragone plot with Pure anatase (blue star), L-TiO 2 (black star), Zn-Ace titanate (green star) and Vac-TiO 2 (red star). Adapted from Jia, X.; Liu, C.; Neale, Z. G.; Yang, J.; Cao, G. Chem. Rev. 2020, 120

Figure 5 . 12 :

 512 Figure 5.12: Ragone plot with Pure anatase (blue star), L-TiO 2 (black star), Zn-Ace titanate (green star) and Vac-TiO 2 (red star). Adapted from Jia, X.; Liu, C.; Neale, Z. G.; Yang, J.; Cao, G. Chem. Rev. 2020, 120 (15), 7795-7866. Kg of active titanate materials

Figure 5 .

 5 Figure5.13 presents the charge capacity as a function of the resting time for the studied titanate materials. As we can observe, an increase in the resting period leads to a higher capacity loss for all samples. The rate of capacity loss is relatively similar among the different compounds, except for pure Anatase, which experiences the most rapid decrease and reaches a capacity of 0 mAh.g -1 after only 5 minutes of rest. In contrast, the other titanate compounds demonstrate better proton retention, with approximately half of the initial capacity being lost after 20 minutes. Notably, Vac-TiO2 exhibits superior performance, retaining 65% of the initial capacity after a 20 minute resting period at the open circuit voltage. These findings underscore the enhanced stability and proton retention capabilities of Vac-TiO2 compared to the other titanate compounds.

Figure 5 . 13 :

 513 Figure 5.13: Self-discharge study conduct on Pure Anatase (black square), L-TiO 2 (red dot), Zn-Ace titanate (green upper triangle), Vac-TiO 2 (blue lower triangle). mAh.g -1 of active titanate materials

Figure S5. 2 :Figure S5. 3 :

 23 Figure S5.2: High rate of MnO 2 plating/stripping

  accurate information about the atomic positions, thermal vibrations, and other parameters describing the crystal structure. The goal is to minimize the difference between the observed diffraction pattern and the calculated diffraction pattern based on the proposed structural model.Rietveld refinement considers various factors that affect the diffraction pattern, including instrumental effects, sample characteristics, and background contributions. It can handle complex crystal structures, including those with multiple phases, disorders, preferred orientations, and anisotropic growth The sample was well-grounded and deposited in significant quantity onto a substrate holder, and the XRD acquisition covered a wide range of angles (20-120 °2θ) over 12 hours.According to Rodríguez-Carvaj et al. (2002),2 concerning anisotropic size broadening it is possible to use a very general phenomenological model, using the Scherrer formula, that considers the size broadening can be written as a linear combination of spherical harmonics (SPH). The explicit formula for the SPH treatment of size broadening is the following:

A. 7 :

 7 y,z) exp (2πi (hx + ky + lz)) dx dy dz Crystallographers often employ atomic coordinates and thermal displacement parameters as a practical representation (referred to as a model) of the scattering density ρ(x, y, z). In a scenario where we possess F and seek to derive ρ(x, y, z), we can employ the inverse Fourier transform to deduce ρ. In this context, the integral can be substituted with a summation over all reflections, as illustrated below: ρ(x,y,z) = Σ hkl Fhkl exp (-2πi (hx + ky + lz)) / V Synchrotron X-ray pair distribution function (PDF)

7

 7

Figure 1 :

 1 Figure 1: (a) World total final consumption by source from 1971 to 2018. (b) World electricity

1 Figure 1 . 1 : 11 Figure 1 . 2 : 14 Figure 1 . 3 : 20 Figure 1 . 4 : 21 Figure 1 . 5 :

 1111112141320142115 Figure 1: (a) World total final consumption by source from 1971 to 2018. (b) World electricity generation by source from 1971 to 2018. Adapted from IEA (International Energy Agency) 20201 .......................................................................................................................................... 1 Figure 1.1: ESW of pure water and compatible electrode material adapted from Kim, H. et al. .................................................................................................................................................. 11 Figure 1.2: (a) Elemental abundance in the Earth's crust. (b) The comparison of molar masses of different ionic charge carriers. (c) The radius comparison of various ions, where a naked proton is too small (0.89 fm) to be shown. (d) A schematic of the Grotthuss mechanism. (e) A timeline of representative proton battery research.28 .............................................................. 14 Figure 1.3: (a) Theoretical volcano plot of favorable reaction. (b) Volcano plot of current density as a function of hydrogen biding Gibbs energy for some metal oxide. Adapted from Ganguly 201580; Laursen 2012,81 .......................................................................................... 20 Figure 1.4: Solvation sheath of salt in water-electrolyte (left) and water in salt electrolyte (right). Adapted from Suo 201597 ........................................................................................... 21 Figure 1.5: (left) Cationic defective fluorinated TiO2 anatase with calculated intercalation energy of Li + , Mg 2+ , and Al 3+ . Adapted from Koketsu 2020106; (right) tuning of a layered lepidocrocite type TiO2 with zinc ions. Adapted from Kang 2022 94 ..................................... 23

Figure 2 . 1 : 41 Figure 2 . 2 :

 214122 Figure 2.1: (a) XRD pattern of T = 90 °C in black curve, T = 110 °C in red curve and T = 150 °C in blue curve. TEM images of titanate at (b) T = 90 °C, (c) T = 110 °C, and (d) T = 150 °C. .................................................................................................................................................. 41 Figure 2.2: (a) PDFs of L-TiO2 in black curve, Vac-TiO2 in red curve, and pure Anatase in blue curve. (b) PDFs refinement of L-TiO2 (top), Vac-TiO2 (middle), and pure Anatase (bottom). Structural representation of (c) L-TiO2 and (d) Anatase .......................................................... 43

Figure 2 . 3 :

 23 Figure 2.3: Thermogravimetric curve and calorimetric analysis of L-TiO2 ............................ 44

Figure 2 . 4 :

 24 Figure 2.4: (a) Thermogravimetric analysis of L-TiO2 in black curve, Vac-TiO2 in red curve, and pure Anatase in blue curve. (b) Structural representation of Vac-TiO2 ............................ 45

Figure 2 . 5 :

 25 Figure 2.5: Thermal X-Ray Diffraction pattern of L-TiO2 ..................................................... 46

Figure 2 . 6 :

 26 Figure 2.6: SEM-FEG images of (a) L-TiO2, (b) Vac-TiO2, and (c) pure Anatase ................ 48

Figure 2 . 7 :

 27 Figure 2.7: TEM images with (a) in situ gas phase of L-TiO2, and (b) post-mortem HRTEM of L-TiO2. ..................................................................................................................................... 49

Figure 2 . 8 : 50 Figure 2 . 9 : 52 Figure 2 . 10 :

 28502952210 Figure 2.8: (a) Diffraction pattern of Vac-TiO2 after various synthesis times, (b) FWMH of Vac-TiO2 along (101), (004) and (200) crystallographic directions ........................................ 50 Figure 2.9: Refinement of Vac-TiO2 with (a) anisotropic orientation factors and (b) without anisotropic orientation factors with a focus on the Bragg peak (004). .................................... 52 Figure 2.10: (a) WH plot of Vac-TiO2, (b) ASS plot of Vac-TiO2 ......................................... 55

Figure 2 . 11 : 56 Figure 2 . 12 : 57 Figure 2 . 13 : 61 Figure 3 . 1 : 75 Figure 3 . 2 : 77 Figure 3 . 3 : 79 Figure 3 . 4 : 82 Figure 3 . 5 : 84 Figure 3 . 6 : 87 Figure 3 . 7 : 90 Figure 3 . 8 : 92 Figure 3 . 9 : 93 Figure 3 . 10 :

 211562125721361317532773379348235843687379038923993310 Figure 2.11: (a) TEM image of Vac-TiO2 (b) 3D and 2D crystallite shape visualization obtained by GFourier Suite software for Vac-TiO2 ................................................................................ 56 Figure 2.12: (a) Electronic map density of Vac-TiO2 (b) structural representation of Vac-TiO2 with the transversal plan location of the electronic map density. ............................................ 57 Figure 2.13: Nyquist plots of the complex permittivity ε" vs. ε' for Vac-TiO2 at room temperature: a) low-frequency part of the entire plot and evidence of the relaxation P1 (giant circular arc); b) evidence of the relaxation P2 (circular arc) from the plot obtained on subtracting the contribution P1; c) evidence of the relaxation P3 (circular arc) from the plot obtained on subtracting contributions P1 and P2. ........................................................................................ 61 Figure 3.1: Schematic representation of lepidocrocite type TiO2. a) L-TiO2 material. b) Lepidocrocite type TiO2 with Zn 2+ in the interlayer coordinated with crystal water ............... 75 Figure 3.2: (a) Quantification of zinc by WDXRF. (b) Thermogravimetric analysis of L-TiO2 (Black) and Zn-Ace titanate 10, 20, 30% (blue, red, and green respectively). The experiment was realized with HI-RESTM TGA from 20 to 650 °C under N2. .......................................... 77 Figure 3.3: Pair distribution function analysis for L-TiO2 (black line), and Zn-Ace titanate from 10 to 30% (dark red to light red line). ...................................................................................... 79 Figure 3.4: Nyquist plots of the complex permittivity ε" vs. ε' for Zn-Ace titanate at room temperature: a) low-frequency part of the entire plot and evidence of the relaxation P1 (giant circular arc); b) evidence of the relaxation P2 (circular arc) from the plot obtained on subtracting the contribution P1; c) evidence of the relaxation P3 (circular arc) from the plot obtained on subtracting contributions P1 and P2; d) evidence of the relaxation P4 (circular arc) from the plot obtained on subtracting contributions P1 to P3; e) evidence of relaxation P5 (circular arc) from the plot obtained on subtracting contributions P1 to P4. ................................................. 82 Figure 3.5: (a) Bulk (grain) conductivity times temperature σT vs. inverse temperature 10 3 /T for Zn-Ace titanate. (b) Dielectric relaxation frequencies for proton hopping (•) and water rotation (•) in Zn-Ace titanate. ................................................................................................. 84 Figure 3.6: Method of electrode fabrication ............................................................................ 87 Figure 3.7: (a) CV of L-TiO2 in 1 M KCl (dash curve) and 1 M acetate buffer electrolyte (black curve) at 10 mV.s -1 . (b) CV of L-TiO2 and Zn-Ace titanate in 1 M acetate buffer electrolyte.

Figure 4 . 1 :

 41 Figure 4.1: Structural representation of intercalation site (left to right) defective free, single vacancy, and double vacancy anatase. (Bottom) DFT calculation of intercalation energy Li + , Mg 2+ , and Al 3+ . Adapted from: D.Dambournet Acc. Chem. Res. 2022, 55, 696-706.4 ....... 108

Figure 4 . 2 :Figure 4 . 3 :Figure 4 . 4 :

 424344 Figure 4.2: (a) CV recorded at 20 mv.s -1 of pure Anatase (black dashed curve) and Vac-TiO2 (red dashed curve) in 1 M KCl pH < 7 electrolyte. (b) Cut-off study of Vac-TiO2 realized at 20 mv.s -1 in 1 M acetate buffer electrolyte at pH = 5.0. (c) Capacities of each faradaic contribution at each cut-off, discharge (orange), charge (green), and HER (violet). mAh.g -1 of active titanate materials. ................................................................................................................................ 112 Figure 4.3: CV performed at 20 mv.s -1 in 1 M acetate buffer electrolyte of L-TiO2 (black curve), Vac-TiO2 (red curve), and pure Anatase (blue curve). .......................................................... 113 Figure 4.4: Quantification of the Nernstian relationship with pH, where the Einterstital oxidizing peak of Vac-TiO2 varies linearly with pH. ............................................................................. 115

Figure 4 . 5 :

 45 Figure 4.5: Structural representation of DFT calculated energy for proton intercalation in (a) pure Anatase and (b) Vac-TiO2 .............................................................................................. 116

Figure 4 . 6 :

 46 Figure 4.6: Log-log plot of Vac-TiO2 showing the linear regression to obtain the b-value of interstitials site (black square) and vacancies sites (red dot) ................................................. 118

Figure 4 . 7 :

 47 Figure 4.7: (a) CV realized at 20 mv.s -1 of Vac-TiO2 (red curve) and a mixture of pure Anatase (A) and L-TiO2 (L) in a 7:3 A:L ratio, from 0 to -1.6 V. (b) Raman spectroscopy of pure Anatase (green curve), Vac-TiO2 (red curve) and a mixture of 7:3 A:L (blue curve) ......................... 121

Figure 4 . 8 :

 48 Figure 4.8: a) GCD of Vac-TiO2 under strong accumulation, no potential limitation at 1st cycle, then, a potential limitation of -1.4 V was applied for subsequent cycles. b) GCD 2nd cycle at

Figure 4 . 9 :Figure 5 . 1 :Figure 5 . 2 :Figure 5 . 3 :Figure 5 . 4 :Figure 5 . 5 :Figure 5 . 6 :Figure 5 . 7 :Figure 5 . 8 :Figure 5 . 9 :Figure 5 . 10 :Figure 5 . 11 :Figure 5 . 12 :

 49515253545556575859510511512 Figure 4.9: HER and OER catalyst on Pt wire (red curve) in galvanostatic mild accumulation condition for Vac-TiO2 (left) and pure Anatase (right) .......................................................... 126 Figure 5.1: (a) Schematic of the electrospinning process, (b) picture of the electrospinning machine .................................................................................................................................. 135 Figure 5.2: Pictures of carbon nanofibers stage elaboration with (a) after the electrospinning process, (b) after stabilization 280 °C 5h, and (c) after carbonization 1200 °C 1h................ 138 Figure 5.3: (a) XRD pattern of MnO2 deposition on CNFs. Cyclic-voltammetry of (b) CNFs cut off, and (c) scan rate study for the determination of EASA ............................................. 140 Figure 5.4: impact of MnAce concentration on electro-deposition/dissolution realized by cyclic voltammetry from 0 to 0.8 V at 10 mV.s -1 ............................................................................. 142 Figure 5.5: (a) Galvanostatic charge/discharge of MnO2 plating/stripping during 5 min for each step, (b) XRD pattern of MnO2 deposited after 5 min on CNFs, (c) SEM-FEG image of CNFs with 5 min of MnO2 deposited, (d) SEM-FEG image of CNFs after 5 min of MnO2 dissolution ................................................................................................................................................ 144 Figure 5.6: Rate capability of MnO2 plating/stripping from 8 to 15 A.g -1 ............................ 145 Figure 5.7: Cyclability performances of L-TiO2 (black sphere), Zn-Ace titanate (green sphere), Pure Anatase (blue sphere), Vac-TiO2 (red sphere) and with fixed intercalation of 154 mAh.g -1 (grey sphere). .......................................................................................................................... 146 Figure 5.8: Photography of the Full-cell set-up in classical three electrodes cell ................. 148 Figure 5.9: GCD cyclability of Full-cell with (a) L-TiO2, (b) Zn-Ace titanate. mAh.g -1 of active titanate materials .................................................................................................................... 150 Figure 5.10: GCD cyclability of Full-cell with (a) Vac-TiO2, and (b) Pure Anatase. mAh.g -1 of active titanate materials .......................................................................................................... 151 Figure 5.11: Cyclability performances of L-TiO2 (black sphere), Zn-Ace titanate (green sphere), Pure Anatase (blue sphere), Vac-TiO2 (red sphere) and with fixed intercalation of 100 mAh.g -1 (grey sphere). mAh.g -1 of active titanate materials .................................................. 152 Figure 5.12: Ragone plot with Pure anatase (blue star), L-TiO2 (black star), Zn-Ace titanate (green star) and Vac-TiO2 (red star). Adapted from Jia, X.; Liu, C.; Neale, Z. G.; Yang, J.; Cao, G. Chem. Rev. 2020, 120 (15), 7795-7866. Kg of active titanate materials ......................... 155

Figure 5 . 13 :Figure 5 . 14 :

 513514 Figure 5.13: Self-discharge study conduct on Pure Anatase (black square), L-TiO2 (red dot), Zn-Ace titanate (green upper triangle), Vac-TiO2 (blue lower triangle). mAh.g -1 of active titanate materials .................................................................................................................... 157 Figure 5.14:(a) SEM-FEG image of (a) clean Vac-TiO2 electrode, (b) Vac-TiO2 electrode after cycling. EDS-mapping of Vac-TiO2 (c) before GCD cycling experiment and (d) after GDC cycling. (e) SEM-FEG image of CNFs positive electrode after GCD cycling ...................... 159

Table 2 .2: Contributions, Relaxation Frequencies (at 293 K), Activation Energies, and Bulk Conductivities (at 293 K) of the L-TiO 2 , Vac-TiO 2 and pure Anatase. Relaxation frequency ν (Hz), activation energy E  (eV), attempt frequency ν 0 (Hz), dielectric strength 

 2

	Sample	Contribution	Proton hopping (P2)	Water rotation (P3)
		ν (Hz)	4.3×10 4	2.7×10 9
		E  (eV)	0.49	0.24
	L-TiO2 15	ν0 (Hz) 	9×10 12 23	4×10 13 3.5
		 (S.m -1 )		8.1×10 -5
		E  (eV)		0.51
		ν (Hz)	7.2×10 4	1.4×10 9
		E  (eV)	0.53	0.32
	Vac-TiO2	ν0 (Hz) 	8×10 13 22	4×10 14 3.3
		 (S.m -1 )		1.34×10 -4
		E  (eV)		0.52
		ν (Hz)	9.0×10 4	5.9×10 9
		E  (eV)	undefined	undefined
	pure Anatase	ν0 (Hz) 	undefined 20	undefined 3.2
		 (S.m -1 )		1.86×10 -5
		E  (eV)		0.56

, Bulk conductivity  (S.m -1 ), activation energy E  (eV). Undefined: couldn't be extracted from fit.

Table S2 .1: Structural parameters of XRD Rietveld refinement of pure Anatase

 S2

	Sample			pure Anatase
	Lattice parameters (Å)	a = b = 3.78551 (0.00002); c = 9.49941 (0.00005)
		Y00	Y20		Y40	Y44+
	Orientation factors	3.963	-2.663		-0.265	-0.527
		Y44-	Y60		Y64+	Y64-
		0.000	-0.796		0.000	0.000
	Refinements	Atom	X	Y	Z	Biso	Occ
	Model 1	Ti 4b	0.500 0.750 0.375 0.31999 0.125
		O 8e	0.500 0.750 0.167 0.34437	0.25
	(Without vacancies)		X² = 6.43, Bragg R-factor = 1.43%
	Model 2	Ti 4b	0.500 0.750 0.375 0.31999 0.12184
		O 8e	0.500 0.750 0.167 0.34437 0.250
	(With vacancies)

X² = 8.16, Bragg R-factor = 1.18% Vacancies 2.53%

  The two solutions were then mixed in a Teflon-lined container. After 10 min of stirring, the container was sealed in a stainless-steel autoclave and heated in an oven at 90 °C for 12 h.

	Ace titanate
	realized in collaboration with Ph.D. candidate M. Sorriaux at PHENIX (Paris, France): two
	solutions were prepared with (i) 4 mL of Ti(IV) isopropoxide (>97%, Sigma-Aldrich) dissolved
	in 10 mL of isopropanol (99.5% min, Alfa Aesar) and (ii) Zn-Acetate-dihydrate (Sigma-Aldrich
	>98%; Zn/Ti = 0.1, 0.2, and 0.3) dissolved in 15.19 mL of isopropanol and 0.81 mL of ultrapure
	water.

  To balance this charge, cationic vacancies are stabilized within Tuning the interlayer composition in hydrated lamellar TiO2 structure increases both capacity and potential window for aqueous proton batteries 78 primary distinction lying in the greater abundance of OH groups within the Zn-Ace titanate compound. Moreover, given that each zinc atom can potentially coordinate with up to three water molecules in the Zn-Ace titanate compound based on 0K DFT calculation,8 the count of free water molecules is drastically reduced. Assuming a 4-fold coordination at room temperature, Zn 2+ can coordinate with two crystal water molecules (in addition to (OH)2), which reduces the number of free water molecules to 0.1 H2O. It is well in agreement with the deduced chemical composition, where H2O/Zn = 2.3. Importantly, it is essential to highlight that although the majority of acetate groups were eliminated during the synthesis process, traces of residual acetate groups were still detected in the Zn-Ace titanate compound, as evidenced by infrared spectroscopy (Figure

	the	structure.	The	general	formula	of	lepidocrocite	type	TiO2	is
	Ti2-x□xO4-4x(OH)4x • nH2O (x  0.5 and n ∼ 0.55) with (□) as the titanium vacancies. This
	structural arrangement also generates undercoordinated oxygen sites, which can serve as host
	sites for proton intercalation.

II.2. Composition

Using Zn acetate dihydrate precursor, we prepared different materials with different molar ratios of Zn/Ti fixed at 0.1, 0.2, 0.3 during the synthesis. After the recovery of the powder Chapter III.

Table 3 .1: Contributions, Relaxation Frequencies (at 293 K), Activation Energies, and Bulk Conductivities (at 293 K) of the Lepidocrocite type TiO 2 , 6 and Zn-Ace titanate. Relaxation frequency ν (Hz), activation energy E  (eV), attempt frequency ν 0 (Hz), dielectric strength 

 3

	When comparing the crystal water mobility between Zn-Ace titanate and ZnCl2,
	intriguing patterns emerge. The attempt frequency for Zn-Ace titanate is notably lower,
	measuring at 2.8 × 10 9

, Bulk conductivity  (S.m -1 ), activation energy E  (eV).

Table 4 .1: Comparison of proton diffusion coefficient within titanate compounds. Samples L-TiO 2 Zn-Ace titanate Vac-TiO 2 pure Anatase Anatase 7,9 Vacancy Interstitial Diffusion coefficients (cm 2 .s -1 )

 4

Table 5 .1: Charge and discharge capacity in function of MnAce concentration Concentration Mn 2+ (M) Q discharge (mAh.cm -2 ) Q charge (mAh.cm -2 )

 5

	0.1	-0.0049	0.0058
	0.3	-0.0119	0.014
	0.5	-0.0088	0.011

Table 5 . 2 .

 52 For discussion, we have also reported for each titanate compound Chapter V. Full cells 154 their cyclability, the potential window, and their Coulombic Efficiency. As it can be seen, Vac-TiO2 exhibits the highest CE% (95.4%) and the highest power and energy density, 10.5 kW.kg-1 and 105 Wh.kg -1 respectively. These values are very promising for developing aqueous batteries with TiO2 as a negative electrode, especially regarding the high-power density. The other promising candidate is Zn-Ace titanate with a close value to Vac-TiO2 regarding energy, power, and Coulombic efficiency.

Table 5 .2: Cyclability, potential window, Coulombic efficiency, energy, and power density of titanate electrode in full cell conditions (Kg of active titanate materials).

 5

(112) Larcher, D.; Tarascon, J. M. Towards Greener and More Sustainable Batteries for Electrical Energy Storage. Nat. Chem. 2015, 7 (1), 19-29. https://doi.org/10.1038/nchem.2085.

Supporting information

Chapter IV. Defective TiO 2 : pushing the limit of HER

I. Introduction

Defect engineering has emerged as a fascinating strategy to improve the electrochemical properties of electrodes. 1,2 Such an approach has been demonstrated for one of these interesting materials, cationic deficient TiO2, which has been extensively studied. [3][4][5] More precisely, the presence of vacancies in fluorinated stabilized TiO2 has demonstrated a capacity increase vs.

Li + , Mg 2+ , and Al 3+ in non-aqueous electrolytes. The presence of cationic vacancies in anatase improves its ability to intercalate cations, leading to higher capacity and better diffusion and cycling stability. For example, Mg 2+ can reach a capacity of around 160 mAh.g -1 , Al 3+ can reach 90 mAh.g -1 (compared to ~20 mAh.g -1 and ~30 mAh.g -1 respectively for defect-free anatase), and Li + diffusion is improved to 10 -13 cm 2 .s -1 . 3,4 Additionally, DFT calculation has shown that vacancy favors the thermodynamic ion intercalation process as shown in Figure 4.1. According to Koudriachova et al., their first-principles simulations suggest that proton insertion doesn't occur at the ideal octahedral site but rather forms within a hydroxyl group. 6 Their study indicates that proton insertion in Anatase, beyond x = 0.0625, is unfavorable due to the formation of H2 caused by local distortions of the OH groups. Hence, the modification of the oxygen network plays a crucial role in promoting proton insertion into TiO2. Consequently, the formation of vacancies, such as Ti vacancies, allows for a much higher proton concentration while maintaining structural stability. Through defect engineering, our goal is to alter oxygens coordination via the introduction of Ti vacancies, thereby providing different adsorption sites for H + .

Supporting information

Chapter V. Full cells

Annexes

A.1: Synthesis protocol:

According to Ma et al., the synthesis used herein corresponds to the same synthesis protocol used for synthesizing a metastable phase of TiO2. A solution of 4 mL titanium (IV) isopropoxide (>97% Sigma-Aldrich) was mixed with 25.1 mL of isopropanol (99.5% min, Alfa Aesar) and 0.81 mL of ultrapure water at room temperature for 10 min in a Teflon container.

The hydrolysis ratio was fixed to h = 3.33. Mixtures were then set in an autoclave for 12h and the temperature was fixed at either, 90 °C, 110 °C and 150 °C. The obtained milky sol-gel was washed 3 times with absolute EtOH using centrifugation. The supernatant was removed and products were dried at 80 °C for 4h. Then, the dried gels were ground manually by using a pestle and mortar until a fine powder was obtained.

A.2: DSC and TGA:

The Differential Scanning Calorimetry analysis was carried out under an N2 atmosphere (10 mL/min), with a heating ramp of 10 °C/min from 25 °C to 600 °C. The TGA was performed with high-resolution HI-RESTM TGA and TA instruments between 25 °C and 600 °C under N2 (10 mL/min) with a variable heating rate to maximize the resolution. 1

A.3: Thermal X-ray powder diffraction:

The sample is in powder form and placed in a chamber with a specific atmosphere such as air, inert, or void as depicted. The basic principle of X-ray thermodiffraction involves exposing a crystalline sample to a source of X-rays while varying its temperature and measuring the intensity as a function of the angle of the X-ray beams that are diffracted by the crystal lattice. In our study, we conducted thermal X-ray diffraction experiments starting from 30 °C and gradually heating the sample to 500 °C. The heating rate employed was 5 °C/20 min, including 10 minutes of temperature stabilization and 10 minutes of data acquisition. These experiments were performed under air and atmospheric pressure conditions. Green batteries that are safe, made of readily available and ethically sourced resources, costeffective and easy to manufacture are needed. Lithium-ion technology, which uses Li-based materials and toxic flammable solvents, does not currently meet all of these criteria. With this in mind, we propose a complementary solution to Li-ion batteries in the form of an aqueous proton-ion battery. The primary challenges posed by aqueous batteries stem from their low energy density, which is the result of the restricted electrochemical stability window of water. This limitation narrows the field of candidates capable of intercalating protons without triggering parasitic reactions such as the hydrogen evolution reaction (HER). TiO 2 materials present an intriguing option for use as a negative electrode for proton insertion. The chemical structure synthesized in solvothermal condition has been modified to achieve three distinct compounds: a layered, amorphous lepidocrocite type titanate sample further optimized with zinc ions chemically incorporated into its interlayer, a cationic defective anatase, and a stoichiometric anatase. The electrochemical analysis of the zinc-optimized layered titanate and defective anatase samples, and their integration as negative electrodes in full-cell conditions, demonstrates promising characteristics with notably, an extended potential window and heightened reversible proton capacity. These results were correlated to the materials characterization and the effect on proton transport.

List of Tables

Keywords: Proton-ion, Aqueous battery, TiO 2 electrodes, Defects, Lamellar, HER

Développement d'une électrode négative à base de TiO2 par chimie douce pour batteries aqueuses protoniques

Résumé :

Des batteries vertes qui sont sûres, fabriquées à partir de ressources facilement disponibles et éthiquement sourcées, rentables et faciles à fabriquer sont nécessaires. La technologie lithium-ion, qui utilise des matériaux à base de Li et des solvants inflammables toxiques, ne répond pas actuellement à tous ces critères. Dans cette optique, nous proposons une solution complémentaire aux batteries Li-ion sous la forme d'une batterie proton-ion aqueuse. Les principaux défis posés par les batteries aqueuses proviennent de leur faible densité énergétique, dû à la fenêtre de stabilité électrochimique restreinte de l'eau. Cette limitation réduit le champ des candidats capables d'intercaler des protons sans déclencher des réactions parasites telles que la réaction d'évolution de l'hydrogène (HER). Les matériaux TiO 2 présentent une option intéressante pour être utilisés comme électrode négative pour l'insertion de protons. La structure chimique synthétisée en condition solvothermale a été modifiée pour obtenir trois composés distincts : un échantillon titanate de type lépidocrocite amorphe disposé en feuillet, qui a été optimisé par la suite avec des ions zinc chimiquement incorporé dans son interfeuillet, un anatase avec des défauts cationiques et un anatase stoechiométrique. L'analyse électrochimique des échantillons de titanate en feuillet optimisé au zinc et d'anatase défectueux, et leur intégration en tant qu'électrodes négatives dans des conditions de cellule complète, montre des caractéristiques prometteuses avec notamment, une fenêtre de potentiel étendue et une capacité réversible de protons accrue. Ces résultats ont été corrélés à la caractérisation des matériaux et à l'effet sur le transport des protons.

Mots-clés : Proton-ion, Batterie aqueuse, Électrodes TiO 2 , Défauts, Lamellaire, HER