1 Para mi familia.

To my family.

ACKNOWLEDGE MENTS

First, I would like to thank all the jury members for accepting to review my work and for the enriching discussion: Dr Raymundo-Piñero, Dr. Jansson, Dr Matei Ghimbeu, Dr Ronsse, Dr Vidal, Dr Eiss-Hortala, and Dr Moog.

I would like to thank my main supervisor, Dr Capucine Dupont, for all the support even before this thesis started. Thanks for introducing me to the biomass and the research world. I deeply appreciated all your support and advice.

I would like to thank my direct supervisor, Dr Loïc Simonin, for his constant guidance and for sharing with me all your knowledge of the Na-ion technology. I appreciated a lot working together during these three years.

Thank you, Capucine and Loïc for all the hard work at the final stage of this thesis and for helping me keep a positive mindset.

Moreover, I would like to thank all the amazing people I have the fortune to work with during these three years.

From the biomass team, I would like to acknowledge Snoop, Philippe and Sébastien for helping me build my biochar reactor. This research would not be possible without you. Thank you, Hélène, for all the analysis and your good energy. Muriel, Serge, Thiery, Maguelone, Guillaume and Françoise, and all the other team members, thank you for sharing your biomass knowledge with me and for forming me in the different techniques.

From the batteries team, I would like first to thank Virgine for being my very first mentor in the battery world. It was an honour to continue the research you started.

Thank you, Philippe, for teaching me so many things of the carbon science. Thank you, Jean-Fred, Jean-Francois, Sylvie, Djamel, Dan, Vasily, Thibaut, Didier, Sebastien Patoux, Léo, Gaëlle, and so many others I might forget, for all your help. Thank you, Sébastien Martinet, for teaching me the statistical analysis and all your work with the Minitab tool. Thank you, Sébastien Launois, for never losing hope in the new oven and for all your support! Thank you very much, Kim, for never giving up with my endless administrative problems and always being there for me! From other laboratories, I want to thank all the scientists who helped me characterise my materials. Thanks, Claude for forming me to the DRX and for all the interesting discussions while preparing the samples. Thank you, Denis, for teaching me all I know of Raman spectroscopy and for your help with the results' interpretations. Thank you, Arnaud, for letting me learn the SAXS basics and teaching me all the data treatment behind. Thank you, Michel Bardet, for the NMR analysis and explanations. Thank you, Gregory, for teaching me the SEM-EDX technique. Thank you, Olivier Sicardy, for teaching me TOPAS. And finally, thank you very much Anass, for forming me into the XPS science and for all your advice and guidance in the final stage of my research, I am glad I manage to learn so many things from you. Also, I want to acknowledge Maryline Roumanie, Michel Bailleux and Sebastien Vry from the DTNM laboratory for kindly letting me use their oven for synthesising my materials. I appreciate the support. I also would like to acknowledge Camelia, for being implied in my research since the early stages. Thank you for helping me improve constantly. I want to send special thanks to all the team in the IS2M institute that collaborated with us with the porosity measurements and the TEM observations. Thank you, Adrian, for all your help too. It was a pleasure to work with you these years.

Thanks to Stina, Eleonora and Eric for the mobility in Sweden, producing hydrochars. Thank you very much, Eleonora, for introducing me to the archery art.

And thanks Lionel for the covoitourage every Friday to practice it.

Thank you to the InnoEnergy KIC, especially to Elisabeth Schuster, for giving me the opportunity of being part of such an enriching program and allowed me to go to Delft in research mobility.

Thank you to all the people I met in Delft for all your help. Blanca, Arlex, David, and Berend. Thanks to all the IHE students, I had the pleasure to meet.

Finally, I need to thank all my PhD colleagues, that were struggling at the same time as me with our research. I am very happy to have shared many lab hours and conversations with you. Thanks, Benjamin for all your help with so many experiments and your welcome for the nice SEM images…Thank you Laurent, Xavier Fleury, Maxime, Thomas, Eddy, Pierre, Guigui, Claire, Dadou, Adrien, Alice, Xavier Judez, Arthur, Manon, Diana, Andrea, Mike, Morganne, Tilia, Florent, Maria, Clement and Oscar. And all the others that I might have forgotten.

I want to thank my friends that were a constant support for me during this time, Cata, Flo, Milan, Aleja, Anlly… Thank you, Themis, for all your patience.

Finally, I would like to thank my family for teaching me hard work and perseverance. Thanks for all your efforts during all these years. Gracias Orlando, Beatriz, Ana y David. Thanks to all my other family members that constantly supported my education, Maye, Cuyi, Sergio, Maria Teresa, Tio Wilson, Doris, Lili, Tio Hector, Turo, tia Mary, Tatavo and so many others.

To all of you, my endless thanks!

GENERAL INTRODUCTION

Our modern society enjoys today technological advances that could not even be dreamed a century ago. The introduction of Lithium-ion batteries (LIBs) in the late 90's, has had a major role in this technological revolution [START_REF] Deng | Li-ion batteries: basics, progress, and challenges[END_REF]. LIBs are widely used in portable electronics, such as mobile phones or laptops. More importantly, LIBs are considered as a key factor for the energy decarbonisation. As energy storage devices, they can be used as buffers for the intermittent renewable energy production. Additionally, LIBs are the most advanced technology used to power electrical vehicles. [START_REF] Deng | Li-ion batteries: basics, progress, and challenges[END_REF]- [START_REF] Etacheri | Challenges in the development of advanced Li-ion batteries: a review[END_REF].

Figure 1. LIB and SIB components and working principle [2], [5]. Author's illustration.

The basic LIB configuration consists of a positive electrode material, which provides the Li ions, and a negative electrode material, often graphite, which is the Li ions host. In between the electrodes, an electrolyte ensures the ions transfer, while a separator prevents the electrons flow between the electrodes. When an external current is applied, the battery undergoes charging, making the Li ions to travel from the positive to the negative electrode, through the electrolyte. When the electrical energy is required, the discharge of the battery implies the opposite ions flow, from the negative to the positive electrode, while the electrons flow through the external circuit. This mechanism is depicted in Figure 1.

Since their commercialisation in 1991, LIBs demand has not stopped growing. In 2000, Li-ion batteries provided 2 GWh to the energy storage market, while in 2018, they delivered around 160 GWh [START_REF] Pillot | The Rechargeable Battery Market and Main Trends 2018-2030 -AVICENNE ENERGY[END_REF]. LIBs demand and production are expected to keep rising, as electrical vehicles are thought to take the place of internal combustion engine vehicles, with manufacturers and governments encouraging this shift [START_REF] Pillot | The Rechargeable Battery Market and Main Trends 2018-2030 -AVICENNE ENERGY[END_REF].

Unfortunately, the massive production of LIBs relies on the consumption, and possible depletion, of non-renewable resources. In particular, Li, and Co in LIBs are considered as critical metals, given their technological relevance and possible threatened supply by natural, geopolitical or economic factors [7] [8]. Li and Co production for LIBs represents the mayor end use of these materials [7] [8].

Li is not considered as an abundant element, its exploitable reserves are estimated on around 16 million tons in 2017. Moreover, the known Li reserves are localized in South America, China and Australia. [START_REF] Ambrose | Understanding the future of lithium: Part 1, resource model[END_REF]. It has been estimated that the available Li resources might be able to supply an optimistic scenario of massive production of electrical vehicles until 2100 [START_REF] Ambrose | Understanding the future of lithium: Part 1, resource model[END_REF], [START_REF] Gruber | Global Lithium Availability[END_REF], [START_REF] Leader | The effect of critical material prices on the competitiveness of clean energy technologies[END_REF]. However, Li supply remains threaded by geopolitical and geographical factors, as its production is highly centralized. This could have a great influence on the versatility of its price in the future, which already has been observed to increase in the recent years [START_REF] Ambrose | Understanding the future of lithium: Part 1, resource model[END_REF].

Co, presents an even less optimistic scenario. Co is normally obtained as a coproduct of Cu, Ni or Fe mining, which increases its supply risk. Co production was recently estimated in 123-128 kilo tons, from which 69 kilo tons were mined in the Democratic Republic of the Congo. Co refining comprises 98 kilo tons, from which 45 kilo tons were processed in China. Moreover, the mining of Co in the Democratic Republic of the Congo has been associated with artisanal mining, human right violations and environmental negligence. The limited Co resources, together with its problematic geopolitical context, might represent high risk of an eventual supply disruption. Indeed, Co price has been observed to fluctuate several times in the last 40 year. [START_REF] Tkaczyk | Sustainability evaluation of essential critical raw materials: cobalt, niobium, tungsten and rare earth elements[END_REF], [START_REF] Leader | The effect of critical material prices on the competitiveness of clean energy technologies[END_REF], [START_REF] Van Den Brink | Identifying supply risks by mapping the cobalt supply chain[END_REF].

Besides Li and Co, LIBs often contain Ni, in the positive electrode and Cu in the negative current collector [START_REF] Durmus | Side by Side Battery Technologies with Lithium-Ion Based Batteries[END_REF], [START_REF] Leader | The effect of critical material prices on the competitiveness of clean energy technologies[END_REF]. The criticality of the latter elements is lower than for the former metals but their supply still needs to be assured. Additionally, Graphite negative electrode, is also mined or synthetically obtained from high temperature treatment of fossil-based precursors [START_REF] Pillot | The Rechargeable Battery Market and Main Trends 2018-2030 -AVICENNE ENERGY[END_REF].

It is clear that more attention should be given to the inventory and supply chain of raw materials for LIBs production. While the recycling process are still technologically immature, other options need to be implemented, to anticipate the demand of energy storage device production in the future.

In this scenario, Sodium-ion batteries (SIBs) can play a crucial role. SIBs work under the same principle as LIB, explained in Figure 1. SIBs and LIBs were actually studied in parallel back in 1980, but because of the higher performance of LIBs using graphite, research on SIB was left behind [START_REF] Hwang | Sodium-ion batteries: present and future[END_REF]. The Na element is larger and heavier than Li, therefore, delivering theoretical energy densities that are lower in SIBs than in LIBs.

In a practical scenario, what is important is the capacity delivered by the positive electrode material, and SIBs have the theoretical potential to approach 90% of the LIBs material capacities [START_REF] Eftekhari | Sodium-ion batteries: New opportunities beyond energy storage by lithium[END_REF]. Today, the sustainability concerns towards LIBs have brought back SIB's research. SIBs development could alleviate the demand of LIBs in targeted applications such as the stationary energy storage of renewable energy sources [START_REF] Durmus | Side by Side Battery Technologies with Lithium-Ion Based Batteries[END_REF], [START_REF] Pu | Recent Progress in Rechargeable Sodium-Ion Batteries: toward High-Power Applications[END_REF].

The main attractiveness of SIBs is related to the abundance, accessibility and low cost of their raw materials. Na is 1000 times more abundant than Li. Na reserves are estimated to 23 billion tons of soda ash, only in the USA [START_REF] Hwang | Sodium-ion batteries: present and future[END_REF]. Na is evenly distributed on the Earth crust and on the ocean, so its supply is not geographically centralized. Therefore, Na carbonate is less expensive than the Li carbonate (133 USD/ton vs. 12000 UDS/ton in 2017 [START_REF] Ambrose | Understanding the future of lithium: Part 1, resource model[END_REF], [START_REF]Soda Ash Statistics and Information[END_REF]).

Regarding the positive electrode materials in SIBs, Li, Co or Ni are not required, which reduces the critical metal content, and potentially, the cost of the positive material. Moreover, one of the biggest advantages of SIBs, is the possibility to replace the Cu negative current collector with Al. The replacement of Cu by Al could imply a cost reduction of 8% in the full battery, which is actually more important than the reduction cost given by the positive electrode material [START_REF] Eftekhari | Sodium-ion batteries: New opportunities beyond energy storage by lithium[END_REF]. As

Al is also lighter than Cu, the energy density of the SIBs improved by this replacement.

Furthermore, as the Na storage capacity of graphite is limited, other negative electrode needs to be implemented for SIBs [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF]. Hard carbon is the most promising option, which can be obtained from renewable organic sources such as biomass, in contrast to the fossil-derived graphite in LIBs. This fact not only influences the sustainability of SIBs, but also opens a completely new scenario for biomassderived materials.

Biomass materials derived from terrestrial vegetation are often referred to as lignocellulosic biomass. The annual production of lignocellulosic biomass is estimated in 181.5 billion tons, from which 4.6 billion tons are originated from agricultural residues [START_REF] Dahmen | Integrated lignocellulosic value chains in a growing bioeconomy: Status quo and perspectives[END_REF]. From these agricultural residues, only 25% have found some valorization pathway [START_REF] Dahmen | Integrated lignocellulosic value chains in a growing bioeconomy: Status quo and perspectives[END_REF]. This means that there are more than 1 billion tons of agricultural residues available to be converted into hard carbons for SIBs.

Lignocellulosic is highly versatile towards tailoring the desired materials properties. Its use as electrode material in energy storage devices has been gaining therefore increasing attention in recent years [START_REF] Gao | A cross-disciplinary overview of naturally derived materials for electrochemical energy storage[END_REF]- [START_REF] Feng | Upgrading agricultural biomass for sustainable energy storage: Bioprocessing, electrochemistry, mechanism[END_REF].

Despite its potential, the optimised use of lignocellulosic biomass in the hard carbon production suffers from a bottleneck: the lack of an interdisciplinary approach, going from the biomass field to the material or electrochemical fields.

Indeed, attention has been given up to now nearly only to the final material properties and electrochemical performances, with very little attention given to the biomass quality and its conversion process. This is why the present research work proposes to study in a systematic way the synthesis of hard carbon negative electrode materials from lignocellulosic biomass, as an attempt to better understand the links between biomass properties, thermal process conditions and hard carbon properties as well as electrochemical performances in SIBs.

For that purpose, the first chapter of this manuscript introduces the reader to the interdisciplinary context and concepts related. Chapter II gathers a detailed stateof-the-art review on biomass derived hard carbons for SIBs and identifies the relevant existing approaches as well as the research gaps. Chapter III presents the materials and methods used in this study.

As the main objective of this study is to identify the role of biomass composition on hard carbon properties and performance in Na-ion batteries, three experimental axes were followed.

First in chapter IV, the final pyrolysis temperature effect is analysed over the physicochemical and electrochemical properties of hard carbons produced from 4 reference lignocellulosic biomasses. From here, it was suspected that biomass composition, in particular its inorganic fraction, might had a more important role on the hard carbon formation, than the thermal treatment, at the range of temperatures used.

To confirm this hypothesis, in Chapter V a larger number of biomass precursors were used, presenting significantly different chemical compositions. Hard carbons were synthesised following the same procedure as before, but at the same final pyrolysis temperature of 1400°C. Correlations between the biomass composition and the final hard carbon properties and performances are discussed, confirming the critical role of the biomass inorganic content, which would seem to promote higher irreversible loss in the electrodes.

Finally, to deepen the understanding of this irreversible loss, the SEI formation and the trapping of Na ions in the hard carbon matrix is analysed in Chapter IV.

The conclusions and future studies arising from this research work are highlighted and presented at the end of this manuscript.

CHAPTE R I. CONTEXT AND OBJECTIVES DEFINITION 1. Introduction

The commercial development of Sodium-ion batteries (SIBs) requires overcoming several challenges to achieve competitive energy and power densities. Among several factors, a lot of research is focused on optimising the electrode materials.

While there is a wide variety of positive electrodes materials with acceptable performance, the negative electrode materials still need to be improved in terms of irreversible loss, specific capacity, capacity retention and rate capability [START_REF] Wang | Emerging non-lithium ion batteries[END_REF], [START_REF] Chen | Readiness Level of Sodium-Ion Battery Technology: A Materials Review[END_REF].

Most of the SIB's developments are based on the insights gained during the Lithium-ion battery (LIB) development. However, the most common negative electrode material used in LIBs, graphite, cannot be used in SIBs. Li ions can intercalate into a graphitic lattice forming LiC6 graphite intercalation compound, with a theoretical specific capacity of 372 mAh.g -1 [START_REF] Yabuuchi | Research Development on Sodium-Ion Batteries[END_REF]. However, Na ions do not interact in the same way with graphite, given a competition between the ionization of the metal atom and the ion-substrate coupling, which renders the Na ion intercalation energetically unfavourable [START_REF] Rosli | The Study of Band Structure of Graphite Intercalation Compound Containing Sodium Calculated Using Density Functional Theory[END_REF], [START_REF] Liu | Origin of low sodium capacity in graphite and generally weak substrate binding of Na and Mg among alkali and alkaline earth metals[END_REF]. Therefore, graphite-Na intercalation compounds deliver specific capacities as low as 35 mAh.g -1 because of the formation of NaC48 and NaC64 compounds [START_REF] Yabuuchi | Research Development on Sodium-Ion Batteries[END_REF], [START_REF] Liu | Origin of low sodium capacity in graphite and generally weak substrate binding of Na and Mg among alkali and alkaline earth metals[END_REF], [START_REF] Asher | Lamellar Compound of Sodium with Graphite[END_REF].

Therefore, other negative electrode materials need to be implemented for SIBs.

The most common negative electrode options can be classified as, insertion materials (carbonaceous or titanium-based), conversion materials (metal oxides or sulphides, transition metal oxides/sulphides), alloying materials ( such as Sn, Bi, Si, Ge, As, Sb and P) [START_REF] Hwang | Sodium-ion batteries: present and future[END_REF], [START_REF] Yabuuchi | Research Development on Sodium-Ion Batteries[END_REF], [START_REF] Wang | Advanced Nanostructured Anode Materials for Sodium-Ion Batteries[END_REF], and organic molecules based on carbonyl compounds [START_REF] Hwang | Sodium-ion batteries: present and future[END_REF]. Figure 2, presents the capacity delivered by the afford mention materials as a function of their working potential. It is worth mentioning that for negative electrode materials, the working potential is expected to be as low as possible, to increase the potential difference with the positive material and therefore, to increase the energy density of a full system.

Because of their low volumetric change, acceptable specific capacity, low working potential, and good cycling stability, the carbonaceous materials, especially hard carbon materials, are considered until now, as the most favourable option for the first generation of SIBs [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF].

As seen from the red square in Figure 2, hard carbons can deliver specific capacities around 300 mAh.g -1 , working at potentials under 0.3V. A SIB full cell using commercial hard carbon from Kureha (240 mAh.g -1 of specific capacity) and a positive electrode of NaNi0.5Mn0.5O2 (120 mAh.g -1 of specific capacity), can deliver an energy density of 224 Wh.Kg -1 (based on active material mass), which represents 60% of the energy density of a LIB containing LiFePO4 or LiCoO2 materials as positive electrodes, and graphite as negative electrode. [START_REF] Komaba | Electrochemical Na Insertion and Solid Electrolyte Interphase for Hard-Carbon Electrodes and Application to Na-Ion Batteries[END_REF], [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF]. This result is encouraging, as commercial hard carbon does not exhibit the best performance in SIBs [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF]. Increasing the hard carbon specific capacity is under constant research, as well as the positive electrode materials, electrolyte formulations, electrodes configurations, among others. Therefore, reaching closer energy densities to the one presented by LIBs is conceivable [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF].

What is a hard carbon material?

Hard carbon is the common denomination for non-graphitic and non-graphitisable carbonaceous materials obtained from reticulated organic precursors by pyrolysis at temperatures that are often higher than 1000°C. The structure of hard carbon is extremely complex, and no universally accepted model exist to describe it.

Overall, hard carbon structure is a complex composite matrix of locally organised domains and pores. The local organised domains are composed of few staked graphene layers of short dimensions, often called turbostratic domains. These domains are cross-linked forming nanovoids, often called ultramicropores (average pore size <1nm). [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF], [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF]. Hard carbon is mainly composed of sp 2 C, and to a lower extend, defects and heteroatoms such as O, H or N. The nature of the cross-linking species is not well defined yet but is thought to be related to O-C-O covalent bonds [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF]. Despite that hard carbon is obtained at high temperatures, its structure is still very different from that of graphitic carbons or natural graphite. In hard carbons, there will always be a lack of long-range order, which avoids the formation of a threedimensional graphitic lattice. The graphene layers in the turbostratic domains of the hard carbons are randomly stacked and are thought to be non-planar [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF].

The most common hard carbon precursors are [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF], [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF], [START_REF] Irisarri | Review-Hard Carbon Negative Electrode Materials for Sodium-Ion Batteries[END_REF]:

-Synthetic polymers: such as petroleum pitch, or phenolic resins.

-Sugars: such as glucose and sucrose.

-Bio-polymers: such as cellulose, or lignin.

Recently, biomass had also attracted attention as a hard carbon precursor for SIBs [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF], [START_REF] Górka | Recent Progress in Design of Biomass-Derived Hard Carbons for Sodium Ion Batteries[END_REF].

Negative electrode -Anode

Specific Capacity / mAh.g -1

Hard carbon final properties are strongly linked to the initial precursor properties, especially to their composition, which can also determine the ability of the precursor to graphitize or not [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF].

In general, the presence of heteroatoms, especially O and S seemed to prevent the long-range graphitization of organic precursors [START_REF] Bonnamy | Chapter 4 -Transmission Electron Microscopy[END_REF], [START_REF] Franklin | Crystallite growth in graphitizing and non-graphitizing carbons[END_REF]. It has been observed that avoiding a fluid phase during thermal treatment also results in non-graphitic carbon. The prevention of the fluid phase seemed to be directly related to higher O contents and lower H contents in the precursors [START_REF] Kipling | The pore structure and surface area of high-temperature polymer carbons[END_REF] [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF]. The formation of micro and ultramicropores at early stages of thermal treatment is also thought to prevent the graphene sheets alignment and further graphitization of the materials [START_REF] Franklin | Crystallite growth in graphitizing and non-graphitizing carbons[END_REF].

The presence of other heteroatoms such as Cl or F, seem to promote hard carbon formation too, as in the case of the resulting material from the pyrolysis of polyvinylidene chloride (PVDC) or polyvinylidene fluoride PVDF precursors [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF], [START_REF] Franklin | Crystallite growth in graphitizing and non-graphitizing carbons[END_REF].

It should be mentioned that the initial degree of cross-linking in the precursor could also determine its ability to graphitise, despite of its low heteroatoms content [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF]. Precursors such as phenolic resins or lignin, with low O/C ratios, present highly reticulate structures, which result in non-graphitisable carbons. This crosslinking can be also achieved by pre-oxidation stages, such as in the case of PAN derived hard carbons.

A schematic representation of the evolution of carbon structures with temperature, according to the O/C ratio of the precursors is given in Figure 3. Notice that if hard carbons are treated at 3000°C, a glassy carbon structure will be achieved. This structure is characterised by a high volume of wide closed micropores.

Furthermore, Table 1 gathers the electrochemical performance and some physicochemical properties of hard carbons derived from sugars, biopolymers, and synthetic polymers. Notice that reversible capacities as high as 345 mAg.g-1 have been achieved for hard carbons derived from phenolic resins, with an outstanding coulombic efficiency of 95%. This confirms that the improvement of hard carbon performance is possible. Despite the potential of hard carbons as negative electrodes for SIBs, they are still more expensive than the graphite used in LIBs. The hard carbon price is around 20 USD.Kg -1 , while graphite price (natural or synthetic) ranges between 6 and 10 USD.Kg -1 [START_REF] Pillot | The Rechargeable Battery Market and Main Trends 2018-2030 -AVICENNE ENERGY[END_REF]. The main reasons for the elevated cost of hard carbon are related to the synthesis conditions and the cost of the raw materials used to produce them.

Natural graphite does not require high thermal treatments, while hard carbon production requires the thermal treatment of the precursor at temperatures above 1000°C [START_REF] Liedel | Sustainable Battery Materials from Biomass[END_REF]. Additionally, if hard carbon is obtained from synthetic polymers, the synthesis of the precursor itself implies an additional cost.

Therefore, if the main attractiveness of SIBs is their lower cost concerning LIBs, cheaper hard carbon precursors need to be used to lower the final hard carbon price. This is the main reason why biomass feedstock has gained attention as a precursor for hard carbon. Biomass as hard carbon precursor is also more advantageous than pure biopolymers or sugars, as these substances also imply previous preparation.

Biomass and biomass wastes are potentially less expensive than the other precursors cited before. Moreover, they are widely accessible, non-toxic and renewable. All these features improve the sustainability of the hard carbon process and might reduce the environmental impact of SIBs [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF]. g. NR: non-reported

Table 1. Electrochemical performances and properties of hard carbons derived from sugars and fossil precursors.

What is Biomass?

In the absolute, biomass can refer to any organic material derived from vegetation (land-base or aquatic) or animals, which are alive or were alive in the recent past (non-fossil) [START_REF] Basu | Biomass Gasification, Pyrolysis and Torrefaction: Practical Design and Theory[END_REF]. Therefore, biomass includes all the types of plants, animals, and even microorganisms, and the biodegradable by-products that they can generate.

Special attention is given to the lignocellulosic biomass, which is mainly composed of cellulose, hemicelluloses and lignin macromolecules. Because of its low nutritional value for humans, lignocellulosic biomass is widely used in energy and chemical production [START_REF] Dahmen | Integrated lignocellulosic value chains in a growing bioeconomy: Status quo and perspectives[END_REF].

Lignocellulosic biomass includes all the land-based vegetation, such as wood and herbaceous species. Wood comprises both softwoods and hardwoods species. The by-products of wood (forestry) and herbaceous (mostly agricultural) crops are also included, which are obtained after harvesting or processing the feedstock. Most of the agricultural and forestry residues include [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF]:

-Forestry: barks, branches, leaves, bushes, chips, pellets, sawdust, sawmill.

-Agricultural: straws, stalks, fibres, pits, grains, seeds, cobs, bagasse, pulps, cakes, residual food, among others.

From here, the focus is given to the lignocellulosic biomass and the term biomass will be used to refer to it.

Biomass composition

Lignocellulosic biomass is the result of a photosynthetic process, in which a plant absorbs sun radiation, water minerals and CO2 to synthesised macromolecules, made of basic units of sugars and aromatic compounds [START_REF] Lago | The Role of Bioenergy in the Emerging Bioeconomy: Resources, Technologies, Sustainability and Policy[END_REF]. The three main macromolecules, cellulose, hemicellulose and lignin, arrange to form the secondary cell wall of plants. Cellulose organises to form microfibrils and these fibrils regroup to form macrofibrils. Hemicelluloses keep the cellulose fibres together and lignin provides strength and rigidity to this matrix. [START_REF] Rubin | Genomics of cellulosic biofuels[END_REF]. Several interactions are known to exist between the macromolecules, to preserve the architecture of the cellular arrangement, such as covalent cross-linking or non-covalent forces. The proportion of each macromolecule will vary in between the different plant cells and between other plant species [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF]. A schematic representation of the biomass structural arrangement is presented in Figure 4.

Lignocellulosic biomass also contains minor non-structural compounds such as extractives or fluid phases. Moreover, lignocellulosic biomass contains an inorganic fraction, that can be originated from natural processes, such as mineral absorption, or by technogenic procceses, such as contamination during harvesting or transportation stages. [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF].

Consequently, lignocellulosic biomass is mainly composed of carbon (C, 40 -60%), oxygen (O, 30 -50%), and hydrogen (H, 3 -10%) [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF], which constitute the macromolecules. Other elements that include the necessary macronutrients for the plant growth are present in minor concentration, and are N, P, K, Ca, S, and Mg.

Biomass could also present other inorganic species such as Si, Na, Cl, Mn, Fe, Zn, Cu, Mo, Ni, and Se. [START_REF] Lago | The Role of Bioenergy in the Emerging Bioeconomy: Resources, Technologies, Sustainability and Policy[END_REF]. Other inorganic elements can be also present as traces.

Cellulose

Cellulose is a linear polymer of β-D glucose units linked by the 1,4-glycosidic bond. [START_REF] Meents | The cell biology of secondary cell wall biosynthesis[END_REF]. Cellulose degree of polymerization goes up to 10000 and the linear polymeric chains can aggregate into crystalline microfibrils of nanometric diameter by means of intra and intermolecular H-bonds and Van der Waals forces. Amorphous regions on cellulose microfibrils are due to discontinuity of the linear arrangement [START_REF] Gomez | Sustainable liquid biofuels from biomass: the writing's on the walls[END_REF].

Interestingly, cellulose is the most abundant biopolymer on earth, and high molecular weight cellulose can only be obtained from vegetal biomass [START_REF] Kim | Characterization of crystalline cellulose in biomass: Basic principles, applications, and limitations of XRD, NMR, IR, Raman, and SFG[END_REF].

Hemicelluloses

Hemicelluloses refer to the groups of heterogeneous and amorphous polymers composed of different sugars. Hemicelluloses have a lesser degree of polymerization than cellulose. Sugars in hemicelluloses polymers are mainly composed of linear backbones of pentoses such as xylose or arabinose, and hexoses such as glucose, mannose, or galactose. However, the final polymeric arrangement is not linear, presenting continuous side chains of different molecules such as glucuronic acid, acetate and methyl groups, and even arabinose or galactose ramified molecules [START_REF] Meents | The cell biology of secondary cell wall biosynthesis[END_REF]. Composition of hemicellulose varies strongly from one specie to another and between cell types. Hardwood hemicelluloses for example, are mainly composed by xylan backbones presenting glucoronic acid ramifications, while softwoods hemicelluloses contain mainly mannan backbones. Grasses are mainly composed of xylan with ramified substitutions of arabinofuranose [START_REF] Meents | The cell biology of secondary cell wall biosynthesis[END_REF] [START_REF] Hatfield | Grass Cell Walls: A Story of Cross-Linking[END_REF].

Hemicellulose arrangements interact closely with cellulose microfibrils by H-bonds to form a cohesive network and to prevent microfibrils to be in direct contact, ensuring a plasticizing effect of the cell wall [START_REF] Gomez | Sustainable liquid biofuels from biomass: the writing's on the walls[END_REF].

Lignin

Lignin is a complex racemic heteropolymer of 4-hydroxyphenylpropanoids units derived mainly from three monolignols: p-coumaryl alcohol, conipheryl alcohol and sinapyl alcohol. These monomers polymerize by free-radical coupling mechanisms forming the building blocks of lignin: p-hydroxyphenyl (H-), guaiacyl (G-) and syringyl (S-), respectively [START_REF] Meents | The cell biology of secondary cell wall biosynthesis[END_REF] [58], as can be seen in Figure 4. These blocks differ in their methoxyl, oxygen or carbon substitutions in the o-position regarding the phenyl oxygen and are linked to each other by several ether and C-C linkages such as β-O-4, β-5, or β-β, among others [START_REF] Ralph | Lignins: Natural polymers from oxidative coupling of 4hydroxyphenyl-propanoids[END_REF]. The final lignin composition and structure varies from one specie to another. It will depend on the proportion of each building block, the presence of other compounds in the polymerization process (such as tricin, hydroxylstilbenes or other non-phenolic compounds) [START_REF] Morreel | Profiling of Oligolignols Reveals Monolignol Coupling Conditions in Lignifying Poplar Xylem[END_REF] [60], and the chemical bonding of other oxygenated groups that esterified within the final lignin polymer (such as alcohols, carbonyl and carboxylic acid functions) [START_REF] Ralph | Lignins: Natural polymers from oxidative coupling of 4hydroxyphenyl-propanoids[END_REF], [START_REF] Collard | A review on pyrolysis of biomass constituents: Mechanisms and composition of the products obtained from the conversion of cellulose, hemicelluloses and lignin[END_REF], [START_REF] Terrett | Covalent interactions between lignin and hemicelluloses in plant secondary cell walls[END_REF].

Lignin and hemicelluloses polysaccharides are cross-linked by covalent bonds [START_REF] Terrett | Covalent interactions between lignin and hemicelluloses in plant secondary cell walls[END_REF],

sealing the polysaccharides in the cell wall and providing waterproof, mechanical reinforcement and a strong barrier to microbial digestion [START_REF] Gomez | Sustainable liquid biofuels from biomass: the writing's on the walls[END_REF].

Extractives

Extractives are compounds that are not part of the structure of the cellular wall and that can be extracted from the lignocellulosic biomass by polar or non-polar solvents. Vassilev et al. [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF] enumerated the main extractive compounds as organic species such as saccharides or carbohydrates, proteins, hydrocarbons, oils, aromatics, lipids, fats, starches, phenols, and waxes. Other substances such as chlorophyll, resins, terpenes, terpenoids, acetyls, uronic acids, sterols, organic acids, glucosides, alkaloids, gums, tannins, flavonoids, dyes, mucilage and saponins, could also be found as biomass extractives. Inorganic species can also be present as extractives.

Inorganic matter

Inorganic elements in biomass could be any element other than C, O, H, N or S.

The inorganic species absorbed by the plant can be fixed in different forms [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF], [START_REF] Vassilev | Ash contents and ash-forming elements of biomass and their significance for solid biofuel combustion[END_REF]:

-Bound to the organic matter (covalent and ionic bonds).

-Ion-exchangeable forms such as organic salts or carboxylic acids.

-Discrete organic materials, such as metal agglomerations.

-Organic or inorganic mineral matter.

The inorganic species in biomass are of particular interest in thermochemical conversion processes as they can form ash, which might generate operational problems. Herein, ash is defined as the non-combustible inorganic residue that remains after complete or incomplete combustion of biomass.

Most of the inorganic ash-forming species in biomass are present in the form of crystalline minerals, such as silicates, sulphates, sulphites, sulphides, phosphates, oxyhydroxydes, carbonates, nitrates, chlorides, etc. Semi-crystalline compounds or mineraloids can be also present, such as silicates, phosphates, hydroxides, chlorides, etc. Additionally, some amorphous compounds such as silicates can be present too. [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF].

Inorganic matter in biomass varies strongly in-between the different biomass species. Ash content in lignocellulosic biomass was observed to range between 0.1 to 42% on dry basis [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF].

4. Biomass pyrolysis as the main synthetic route to produce carbon materials.

Overview

Pyrolysis is the degradation of organic material by the action of heat under an inert atmosphere. Pyrolysis of biomass produces three main products, a solid phase called biochar, a liquid phase and a gas phase. The yield and composition of each product will directly depend on the operational conditions and the biomass feedstock composition. If biochar is produced at temperatures over 1000°C, it can be considered as a hard carbon [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF].

Biochar definition is often ambiguous in literature and is regularly related to agricultural applications. Herein, it is defined as the solid residue of pyrolysis, regardless of its application, and as such, the biochar term will be used all along in this manuscript.

To favour the biochar yield, generally slow (or conventional) pyrolysis is used. Slow heating rates (<60°C/min), long residence times (> 5min) and temperatures between 400 and 500°C, have been observed to increase the biochar yield [START_REF] Tripathi | Effect of process parameters on production of biochar from biomass waste through pyrolysis: A review[END_REF]. Slow heating rates allow a controlled diffusion of the generated volatile matter through the forming biochar; therefore, the macrostructure of the original biomass is preserved after the thermal treatment [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF]. Moreover, slow heating rates and long residence time of the volatiles give the system enough time for biochar-formation reactions to proceed [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF]. The presence of alkali metals, such as Ca, K and Na, can also catalyse the biochar formation [START_REF] Collard | A review on pyrolysis of biomass constituents: Mechanisms and composition of the products obtained from the conversion of cellulose, hemicelluloses and lignin[END_REF], [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF].

Pyrolysis mechanisms

Pyrolysis mechanisms are very complex, as many reactions take place simultaneously. The mass loss during slow pyrolysis is commonly followed by thermogravimetric analysis (TGA), from which the rate of degradation can be evaluated. The range of temperature in which a given biomass loss its highest mass, is highly variable. At the same pyrolysis conditions, this temperature strongly depends on the macromolecular and inorganic composition of the feedstock, as well as, on the existent interaction between its different components.

Overall, slow pyrolysis proceeds by primary and secondary reactions, as follows [START_REF] Collard | A review on pyrolysis of biomass constituents: Mechanisms and composition of the products obtained from the conversion of cellulose, hemicelluloses and lignin[END_REF], [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF]:

-With initial heating, dehydration of biomass will take place, liberating water vapour and light extractive species. properties that the primary char [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF].

Biochar to hard carbon formation

The biochar formation could be considered as a progressive evolution from unaltered plant material, towards aromatic clusters that finally form turbostratic crystals at temperatures around 700°C [START_REF] Keiluweit | Dynamic Molecular Structure of Plant Biomass-Derived Black Carbon (Biochar)[END_REF]- [START_REF] Mafu | Chemical and structural characterization of char development during lignocellulosic biomass pyrolysis[END_REF], as can be seen from Figure 5a and b. The resulting material can be seen as an early stage hard carbon. Biochar elemental composition gets richer in C and losses O, N, S and H as the pyrolysis temperature rises. At the same time, the ash are mainly non-volatile and concentrate on the solid product.

Biochars obtained between 700 and 1050°C (approximately) are microporous with average pore sizes of 3 to 10 Å. These pores can form voids in the aromatic planes, producing defective sites, such as free radicals, or dangling bonds [START_REF] Bourke | Do All Carbonized Charcoals Have the Same Chemical Structure? 2. A Model of the Chemical Structure of Carbonized Charcoal[END_REF].

It is known that the specific surface area of the biochar tends to increase up to a critical temperature (between 600 and 1200°C according to the biomass precursor), while there is a decrease of the solid density. This behaviour is related to the porosity formed by the escape of the species during volatilization. The generated pores can exhibit different sizes, from mesopores to ultra-micropores [START_REF] Weber | Properties of biochar[END_REF], often quantified by adsorption tests. After the critical temperature, the biochar specific surface area starts decreasing, while its density remains almost constant. This behaviour is related to an internal coalescence of the pore walls with temperature.

[29], [START_REF] Zhang | Investigating the structure of non-graphitising carbons using electron energy loss spectroscopy in the transmission electron microscope[END_REF], [START_REF] Matei Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF], [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF], [START_REF] Buiel | Model of micropore closure in hard carbon prepared from sucrose[END_REF].

At temperatures over 1000°C, the turbostratic domains account for 70% of the biochar structure. With further thermal treatment, these domains keep growing while they become less defective. The graphene layers stack and get closer to each other, increasing the average size of the closed pores. At this point, the hard carbon structure has been achieved. [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF], [START_REF] Matei Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF], [START_REF] Weber | Properties of biochar[END_REF]. This final stage is schematised in Figure 5c.

The hard carbon composition, structure, defects, porosity (both accessible and closed), and surface chemistry, are strongly related to the final pyrolysis temperature, as presented in Table 2. A detail discussion of the hard carbon properties variation with temperature is presented in Chapter II. He measurements [START_REF] Kipling | The pore structure and surface area of high-temperature polymer carbons[END_REF] The highest possible Highly dependent on the temperature, precursor and gas of measurement.

Morphology

For N2 : > 300 m 2 .g -1 ( for T<1000°C) until limit of detection (T>1400°C)

For CO2: > 400 m 2 .g -1 ( for T<1200°C) until limit of detection (T>1400°C) [START_REF] Zhang | Investigating the structure of non-graphitising carbons using electron energy loss spectroscopy in the transmission electron microscope[END_REF], [START_REF] Matei Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF], [START_REF] Zhang | Correlation Between Microstructure and Na Storage Behavior in Hard Carbon[END_REF], [START_REF] Beda | Hard carbons derived from green phenolic resins for Na-ion batteries[END_REF], [START_REF] Izanzar | Hard carbons issued from date palm as efficient anode materials for sodium-ion batteries[END_REF] For N2 SSA > 40 m 2 .g -1

No specification can be proposed for CO2 SSA

Closed porosity

From SAXS observations and density measurements.

Average pore size increase with temperature but is always < 2nm [START_REF] Kipling | The pore structure and surface area of high-temperature polymer carbons[END_REF], [START_REF] Simone | Hard carbon derived from cellulose as anode for sodium ion batteries: Dependence of electrochemical properties on structure[END_REF],

[77], [START_REF] Dou | Research Update: Hard carbon with closed pores from pectin-free apple pomace waste for Naion batteries[END_REF] As high as possible

Surface composition

Depending on the precursor and temperature. sp 2 C [START_REF] Matei Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF], [START_REF] Matei Ghimbeu | Valorizing low cost and renewable lignin as hard carbon for Naion batteries: Impact of lignin grade[END_REF],

[79]- [START_REF] Conder | Chitin and Chitosan-Structurally Related Precursors of Dissimilar Hard Carbons for Na-Ion Battery[END_REF] Depending on the precursor.

O functional groups and inorganic impurities.

O > 10 at% at the extreme surface.

Objectives of the thesis

The synthesis of biomass-derived hard carbon is at the interface of two distinctive scientific fields, which implies different languages and approaches. On the one hand, there is the understanding of the biomass feedstock properties and the optimisation of the processing techniques to obtain the desired products. On the other hand, there are the electrochemical requirements that the produced materials should meet to be able to host the highest amount of Na ions in a SIB's electrode during a given lifetime. Therefore, the optimal biomass-derived hard carbon electrode cannot be conceived without an interdisciplinary approach. This is why this research work aims to provide an interdisciplinary study of biomass-derived hard carbons for SIBs, where equal weight is given to the biomass and the electrochemistry fields.

It was presented that biomass could be used as precursor for hard carbons, as it has a high O content, an important-crosslinking degree and a macroporous hierarchical macrostructure. After high-temperature slow pyrolysis, these precursors will produce non-graphitic materials with preserved morphologies, able to store Na ions. However, it was also presented that biomass can comprise a wide range of different species, that will differ in chemical composition and macrostructures. These differences will certainly lead to diverse hard carbon characteristics, which are going to affect the Na ions storage on the material.

Therefore, the main objective of this research work is to study the impact of the biomass precursor properties on the hard carbon final properties and to identify the critical factors that will affect the hard carbon formation from a biomass precursor.

To address this objective, an initial literature review on hard carbons derived from biomass was done, to identify the existing scientific gaps. This review is presented in the upcoming chapter. From the conclusions of this review, an experimental approach is proposed to achieve the previous objective.

40

CHAPTE R II. STATE -OF-THE -A RT OF BIOMASS DERIVED HARD CARBONS

Herein, the available knowledge regarding the hard carbon production from biomass feedstock is explored. In particular, the feedstock and the synthetic routes used by the different authors are presented and discussed in this chapter.

An extensive literature review of biomass-derived hard carbons for SIBs was performed, leading to the selection of 62 publications, for discussing their approaches and results. To the best of our knowledge, the selected publications provide an exhaustive overview of the biomass precursors and hard carbons synthesis conditions used for SIBs application.

The reviewed data is presented in Table 3 at the end of the chapter. From now on, the term feedstock is replaced by the term precursor. The term feedstock is usually used in the thermal process field to designate biomass, while the term precursor is used in the material science field. Here, for the sake of consistency, only the term precursor will be used.

From the reviewed data, we aim to determine if the role of the biomass composition on the hard carbon properties has been already identified. Moreover, the biomassderived hard carbon synthesis procedure is evaluated to determine which synthesis conditions influence the most the hard carbon properties and performance, and whether or not the synthesis procedure is optimised. Finally, the performance of the biomass-derived hard carbon is evaluated, to determine if it achieves the requirements for SIBs application.

General trends

All the studies reported in Table 3 were performed at a laboratory scale, in where few grams or even milligrams of the final hard carbons were obtained.

Apart from the large diversity of biomass precursors reported in CarboxyMethyl Cellulose (CMC), Na alginate, sodium polyacrilate (PANa), and polyacrylic acid. Not all the studies used a carbon additive on the composite electrodes. Different current densities were used to cycle the systems, ranging from 10 to 500 mA.g -1 . All these set-ups are reported in Table 3.

It is also possible that studies sometimes used different models or fitting techniques to calculate structural, porosity or chemical parameters.

This wide set of variables makes the results very difficult to compare. To perform a fair comparison, in which the role of the biomass precursor is fully evidenced, it is clear that we need a systematic study in which the synthesis routes and the analysis of the results are carried out in the exact same conditions. However, general trends were identified from the literature data, and are discussed below.

Biomass precursors diversity

Selection criteria

The most common hard carbon biomass precursors are agricultural non-valorised residues. These precursors were chosen because of their abundance, low cost, renewable nature, and to avoid their disposal that causes environmental pollution.

In addition to these criteria, some precursors were selected given their theoretical content of cellulose, hemicelluloses and lignin [START_REF] Dou | Research Update: Hard carbon with closed pores from pectin-free apple pomace waste for Naion batteries[END_REF], [START_REF] Ding | Carbon Nanosheet Frameworks Derived from Peat Moss as High Performance Sodium Ion Battery Anodes[END_REF]- [START_REF] Romero-Cano | Surface functionalization to abate the irreversible capacity of hard carbons derived from grapefruit peels for sodium-ion batteries[END_REF]. It seems that authors assumed that the cross-linking of these macromolecules in the biomass precursor would lead to the hard carbon structure. However, as will be discussed in the next section, none of these studies seemed to quantify the composition of the biomass precursor.

Similarly, other precursors were selected because of their N [START_REF] Selvamani | Garlic peel derived high capacity hierarchical N-doped porous carbon anode for sodium/lithium ion cell[END_REF], [START_REF] Li | Heteroatom-doped hierarchically porous carbons derived from cucumber stem as highperformance anodes for sodium-ion batteries[END_REF], or F [START_REF] Wang | Fluorine-Doped Carbon Particles Derived from Lotus Petioles as High-Performance Anode Materials for Sodium-Ion Batteries[END_REF],

content, arguing that naturally doped hard carbons can be obtained. Still no quantification of these elements in the biomass precursor was performed.

It is worth noticing that most of the reviewed publications gave a limited description of the biomass origin. Only a few studies reported where the biomass precursors were obtained and their harvesting conditions. The initial physical conditions of the precursors, such as shape, particle size, or moisture, were also barely reported. The latter information is useful to decide the most adapted pretreatments and synthesis methods to process the biomass. Therefore, the performed synthesis conditions might not be optimized for the biomasses initial properties.

Biomass precursor composition

As seen in chapter I, the hard carbon properties are strongly dependent on the precursor properties, especially on their composition. Parameters such as the C to O ratio, the C to H ratio, the proportion of the different macromolecules, the level of cross-linking, the amount and composition of the extractives, and the concentration and composition of the inorganic fraction, could lead to very different hard carbon characteristics. Variations can be seen in terms of structure, surface chemistry, porosity development, and C purity of the resulting hard carbons, which are thought to result in different electrochemical performances.

It is therefore surprising that none of the 62 analysed precursors of Table 3 were characterised in terms of macromolecular and or inorganic composition. Only four studies quantified the C, H, N and S elemental composition of the precursors [START_REF] Ding | Carbon Nanosheet Frameworks Derived from Peat Moss as High Performance Sodium Ion Battery Anodes[END_REF][85][102] [START_REF] Romero-Cano | Surface functionalization to abate the irreversible capacity of hard carbons derived from grapefruit peels for sodium-ion batteries[END_REF]. Some studies reported the macromolecular composition of the precursor that was found in the literature. However, variability is known to exist in between precursors of the same type [START_REF] González Martínez | Assessing the impact of woody and agricultural biomass variability on its behaviour in torrefaction through Principal Component Analysis[END_REF], and therefore, to report a literature's composition for another precursor could provide a mistaken composition of the analysed sample.

Macromolecular composition

It is generally assumed in hard carbon's literature that cellulose, hemicellulose and lignin are the principal components of biomass, together with some minor inorganic elements, lipids and proteins. It is also widely assumed that the structural organization of the final hard carbon is dependent on the proportion between cellulose, hemicelluloses and lignin in the precursor.

The ratio of cellulose, hemicelluloses and lignin varies significantly from one biomass precursor to another and even for the same biomass if a different part of the plant is analysed. Moreover, the basic sugar or phenolic units in the macromolecules could present very different chemical composition, according to the biomass type. This variety translates into a distribution of chemical bonds that might reinforce or weaken the crosslinking of the precursors among the thermal treatment. Therefore, the macromolecular composition could indeed affect the graphitization degree of hard carbon.

However, there is no information in the literature regarding a relationship between the macromolecular composition and the hard carbon's final structure.

The few existing studies regarding this subject were held under conditions that might not represent the interaction between the macromolecules on given biomass.

On the one hand, there is the enzymatic fractionation of a series of biomass precursors into different extracted samples, containing combination or single macromolecules [START_REF] Feng | Chemical-enzymatic fractionation to unlock the potential of biomass-derived carbon materials for sodium ion batteries[END_REF]. While on the other hand, extended H3PO4 activations were done before the hard carbon synthesis [START_REF] Dou | Pectin, Hemicellulose, or Lignin? Impact of the Biowaste Source on the Performance of Hard Carbons for Sodium-Ion Batteries[END_REF]. In both cases, biomass precursors were are agricultural residues such as brewer spend grain, grape pomace, walnut shell, apple waste, corn cob and peanut shells. Therefore, the behaviour of other herbaceous and woody precursors is not considered, and the results might not represent all the biomass types, as the macromolecular composition is known to vary strongly between biomass types (see Chapter I).

It would seem, however, that for the mentioned agricultural residues, the hemicelluloses content and the lignin to cellulose ratio play a major role on the graphitization degree and structural defects of the resulting materials. The porosity and the surface chemistry of the resulted hard carbons also seemed to be impacted by the different macromolecular compositions.

Hemicelluloses seemed to decrease the d002 interlayer space of the hard carbon.

Feng and co-workers [START_REF] Feng | Chemical-enzymatic fractionation to unlock the potential of biomass-derived carbon materials for sodium ion batteries[END_REF] suggested that as hemicelluloses decompose at lower temperatures than cellulose and lignin during pyrolysis, an earlier re-organization of the carbon sheets might take place, leading to lower d002 values. Additionally, hemicelluloses are heterogeneous polymers (xylose + glucomanose + arabinose, among others monomers), while cellulose is a homogeneous polymer only composed of glucose monomers. This could imply that the re-arrangement towards sp 2 aromatic clusters results in lesser defective sites for a non-hemicellulose-derived hard carbon, as was confirmed by their lower ID/IG ratio and their highest sp 2 -C concentration.

It was also observed that lignin to cellulose ratios around 2:1 seemed to produce highly defective carbons with acceptable performance. In contrast, low lignin to cellulose ratios or not lignin at all, led to hard carbons with closer graphene layers, high porosity and an important quantity of sp 3 C concentration, of less competitive electrochemical performances. [START_REF] Feng | Chemical-enzymatic fractionation to unlock the potential of biomass-derived carbon materials for sodium ion batteries[END_REF] Regarding the porosity and surface area, the results suggested that high lignin values seemed to produce hard carbons with a larger mesoporous volume. For lower temperature carbonaceous materials, such as biochars, it is well known that the surface area and the porosity development can be influenced by the biomass macromolecular composition [START_REF] Tripathi | Effect of process parameters on production of biochar from biomass waste through pyrolysis: A review[END_REF], [START_REF] Weber | Properties of biochar[END_REF], [START_REF] Tomczyk | Biochar physicochemical properties: pyrolysis temperature and feedstock kind effects[END_REF]. However, studies regarding this aspect for hard carbons produced at higher temperatures are not available.

Considering that above 1000°C, the pores present in the carbonaceous solid start to collapse or merge, it is possible that the effect of the macromolecular composition in the surface area and porosity formation is not significant under these hightemperature conditions.

It should be remembered that enzymatic fractionation and H3PO4 treatments of the precursor might destroy the interaction between the macromolecules in the biomass cellular arrangement that are known to exist and to affect the thermal degradation and char properties [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF], [START_REF] Collard | A review on pyrolysis of biomass constituents: Mechanisms and composition of the products obtained from the conversion of cellulose, hemicelluloses and lignin[END_REF], [START_REF] Jönsson | Pretreatment of lignocellulose: Formation of inhibitory by-products and strategies for minimizing their effects[END_REF]. Therefore, more representative studies are needed to provide conclusive insights into the role of the macromolecular composition and hard carbon performances.

Inorganic composition

As for the macromolecular characterization, inorganic species in biomass seemed to be rarely quantified. Inorganic matter in the biomass precursor is considered an impurity and is normally removed by washing the precursor with different solutions, as can be seen from Table 3. Inorganic matter can also be removed from intermediate biochars or the final hard carbons.

Because inorganic matter is either unidentified or directly removed, there are no available insights on its effect on the hard carbon development or properties. It is only assumed that it has a detrimental effect on the final hard carbon performances.

It is known that a biomass precursor can possess up to 40% of its dried weigh on inorganic matter (see Chapter I). It is also known that the chemical composition of this inorganic fraction will affect the pyrolysis reactions and therefore, the hard carbon formation. Alkali metals such as Ca or K are known to catalyse biochar formation reactions [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF] and therefore, they could increase the hard carbon yield.

Species such as K, Zn, Mg, P or S compounds have been observed to catalyse porosity formation reactions, [START_REF] Feng | Chemical-enzymatic fractionation to unlock the potential of biomass-derived carbon materials for sodium ion batteries[END_REF], [START_REF] Stratford | Intrinsic activation: The relationship between biomass inorganic content and porosity formation during pyrolysis[END_REF]. No studies were found where these parameters were studied for biomass-derived hard carbons formation.

Some other species such as Ca compounds have been observed to promote catalytic graphitization on precursors that normally lead to non-graphitic carbons [START_REF] Ōya | Phenomena of catalytic graphitization[END_REF]. As before, no systematic studies have been carried out on the catalytic graphitization of biomass precursors by the effects of their inorganic compounds.

The available studies are mainly focused on the effect of the acid wash on the electrochemical performances. Some authors have tried to investigate the effect of the inorganic matter removal from raw biomass [START_REF] Dahbi | Synthesis of hard carbon from argan shells for Na-ion batteries[END_REF] and an intermediated biochar [START_REF] Zhang | Pinecone biomass-derived hard carbon anodes for highperformance sodium-ion batteries[END_REF], in the resulting structure, porosity and electrochemical performance of the hard carbon. In both cases, acid treatments with HCl solutions were performed. K was found to be the main inorganic specie on the samples, whose concentration was observed to decrease significantly (almost completely) after the washing of the samples. The electrochemical performance of the washed samples was observed to be better than the one of the raw samples, in terms of reversible capacity, capacity retention and rate capability, as can be seen from Table 3, samples [START_REF] Zhang | Pinecone biomass-derived hard carbon anodes for highperformance sodium-ion batteries[END_REF], [START_REF] Dahbi | Synthesis of hard carbon from argan shells for Na-ion batteries[END_REF]. Moreover, it was observed that the washed samples presented better performances at lower pyrolysis temperatures that the raw samples. Both studies attributed this improvement to the cleaning of active sites on the washed samples, which were occupied by inorganic matter in the raw samples.

However, as for the enzymatic fractionation and H3PO4 treatments discussed before, the HCl washing could also modify the carbonaceous matrix and therefore, the better electrochemical performance cannot be fully attributed to the removal of the inorganic matter. The thermal degradation kinetics were observed to varied for argan shell samples that were HCl-washed, presenting a higher thermal resistance [START_REF] Dahbi | Synthesis of hard carbon from argan shells for Na-ion batteries[END_REF]. The structural and porosity properties of the washed and unwashed derived hard carbons were also observed to vary. Washed samples seemed to present lower d002 distances and ID/IG ratios, while a significant increase of their N2 SSA was seen. A comparison between the afford mention properties, as well as the reversible capacity and ICE, for the washed and raw argan shell derived materials can be seen in Figure 6.

Therefore, the acid wash approach is not able to determine the role of the inorganic matter on the hard carbon properties, and more studies, in which the carbonaceous matrix is not modified, are needed to fill in this gap. It is worth mentioning that the role of the inorganic fraction in biomass does not only affect the hard carbon formation and final structure, but it could also have a role in the electrochemical phenomena during cycling. It is possible to think that if inorganic elements remain as chemically bonded compounds on the hard carbon surface after pyrolysis, they could modify the surface properties, by either creating isolating regions [START_REF] Conder | Chitin and Chitosan-Structurally Related Precursors of Dissimilar Hard Carbons for Na-Ion Battery[END_REF], or interacting within the electrolyte species [START_REF] Matei Ghimbeu | Valorizing low cost and renewable lignin as hard carbon for Naion batteries: Impact of lignin grade[END_REF], depending on the chemical nature of the remaining compounds. If inorganic species remain in the hard carbon bulk, they could not only decrease the active sites for Na ions storage, but they could also decrease the electronic and ionic conductivity of the carbon substrate [START_REF] Zhang | Pinecone biomass-derived hard carbon anodes for highperformance sodium-ion batteries[END_REF], [START_REF] Dahbi | Synthesis of hard carbon from argan shells for Na-ion batteries[END_REF]. If the remaining inorganic species is electrochemically inactive, it will contribute to a dead weigh in the electrode composition, therefore, decreasing the energy density of the electrode [START_REF] Adamson | Peat-derived hard carbon electrodes with superior capacity for sodium-ion batteries[END_REF]. For all these reasons, the contribution of the specific inorganic species from biomass, on each of the previously describe phenomena, needs to be determined.

Hard carbon synthesis conditions

A whole set of experimental conditions can be tuned to produce hard carbon, that could produce different hard carbon properties, and therefore, electrochemical performances. The final temperature of pyrolysis, the pyrolysis atmosphere, the pyrolysis residence time, the heating rate, and the gas flow rate are known to affect the pyrolysis mechanisms [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF] and therefore, they can modify the physicochemical properties of the resulting hard carbons. Additionally, the biomass precursor can be subjected to pre-treatments before to the pyrolysis process, such as the already mention acid wash, a pre-pyrolysis at lower temperatures, a stabilization under air at lower temperatures, an activation, and hydrothermal carbonization. In some cases, the resulting hard carbon is also subjected to post-treatment, such as heteroatoms doping, acid/basic wash, and or activation.

All kinds of combinations of synthesis conditions have been tested in literature as can be seen from Table 3. Consequently, it was observed that from the same precursor, different physicochemical and electrochemical properties are reported.

Notice for example, the case of the different hard carbons derived from pomelo peels when treated at different conditions. The reversible capacity of the pomelo peel derived hard carbons varied from 170 to 430 mAh.g -1 . The ICE of the same materials varied from 27 to 69 %, while their rate capability at 1 A.g -1 , varied from 50 to 172 mAh.g -1 . These results confirmed the critical role of the synthesis conditions on the hard carbon properties and performance.

In the upcoming sections, the main pyrolysis conditions and pre-treatments are discussed, regarding their impact on the final hard carbon performances.

Pyrolysis temperature

To vary the pyrolysis temperature during the hard carbon synthesis yields in very different hard carbon properties in terms of structure, surface area, surface chemistry, porosity and final C purity, from the very same precursor, delivering different electrochemical performances. This is why almost all the focus on literature is based on the analysis of the hard carbon properties among thermal treatment, as can be seen from Table 3. This section presents a detailed analysis of hard carbons produced by direct or two-step pyrolysis, without any pre-or posttreatment, in order to be as comparable as possible between the different reported results. This might also allow identifying if different precursors behave differently under similar pyrolysis conditions.

Interestingly, there seem to be no criteria for the final pyrolysis temperature choice, and temperatures as low as 600°C and up to 1600°C have been used. Only few approaches in literature have used thermogravimetric analysis (TGA) to choose the range of pyrolysis temperatures to produce hard carbon [START_REF] Izanzar | Hard carbons issued from date palm as efficient anode materials for sodium-ion batteries[END_REF], [START_REF] Wang | Kelp-derived hard carbons as advanced anode materials for sodium-ion batteries[END_REF], [START_REF] Jiang | Biomass carbon micro/nano-structures derived from ramie fibers and corncobs as anode materials for lithium-ion and sodium-ion batteries[END_REF]. A TGA analysis allows determining the temperature at which the biomass precursor reaches a constant mass loss. From this point, mainly, a re-organization of the structure is taking place. Therefore, it is interesting to synthesise hard carbons above this temperature. This point is observed to be dependent on the precursor type, which of course relates to the precursor composition.

Hard carbon composition

The final pyrolysis temperature was observed to determine the final hard carbon composition. This way, a material synthesised at 700°C is more likely to contain higher amounts of heteroatoms, than a hard carbon synthesised at 1600°C. This can be seen from Figure 7, with the increase of the C content and the decrease of the O content, for apricot shells, pectin-free apple waste, and corn straw piths derived hard carbon, as well as the treated banana peels and peat moss derived hard carbons (washed with KOH and HCl). Elemental composition of directly pyrolysed hard carbons is also rarely reported.

Hard carbon yield

The hard carbon yield is also strongly dependent on the final pyrolysis temperature. It was observed that the yield is influenced by the early volatilization of the precursors and the final volatilization related to the structural organization.

Biochars derived from different woods and herbaceous precursors exhibited their major mass loss at temperatures between 200 and 800°C [START_REF] Weber | Properties of biochar[END_REF]. Therefore, the final hard carbon yield mainly depends on this first step. Then, the yield can keep decreasing through minor volatilization. Once again, it should be mentioned that the main volatilization temperature strongly depends on the biomass composition [START_REF] Weber | Properties of biochar[END_REF].

Small variations of the yield are observed after a given temperature. Notice for example, that for the directly pyrolysed mangosteen shell precursor [START_REF] Wang | Low-Cost and High-Performance Hard Carbon Anode Materials for Sodium-Ion Batteries[END_REF], the hard carbon yield varied from 55% at 800°C to 35% at 1300°C. Then, only 5% of the yield is lost between 1300°C and 1600°C.

Despite the industrial relevance of the hard carbon yield, it was observed that it is not often reported in the literature. Moreover, when reported, the basis of the calculation (moisture-free or moisture-ash-free) is not always specified. 52

Hard carbon structure

The hard carbon structure is quite complex and still, not completely understood.

Often, hard carbon structure is studied using X-ray Diffraction (XRD). Even though this technique is generally used for characterising well-crystallised compounds, it provides many insights into the hard carbon structure. The distorted turbostratic domains in hard carbon produce asymmetrical and very broad diffraction peaks of 00l and hk(0) planes, as seen from Figure 8a. No hkl planes exist on the hard carbon XRD pattern, as these planes are characteristic of a homogenous long-range stacking of the graphene layers that do not exist in hard carbon.

With increasing the final pyrolysis temperature, the hard carbon peaks in the XRD pattern become narrower, and the 002 peak shifts towards higher angles indicating a smaller interlayer space between the graphene layers (d002 parameter) of the turbostratic domains. The diminution of the d002 interlayer space has been observed for several directly pyrolysed hard carbons, as reported in Figure 9a.

This behaviour is also an indication of lateral growth of the turbostratic domains.

The d002 interlayer distance of hard carbon is normally higher than 0.344 nm, which is still significantly higher than the d002 distance for graphite (0.335 nm) [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF], [START_REF] Inagaki | Materials Science and Engineering of Carbon: Characterization[END_REF]. Therefore, the d002 parameter is often used as a measure of the degree of organization reached in the structure.

XRD observations are often complemented with Transmission Electron Microscopy (TEM) and Raman spectroscopy observations.

High-resolution TEM images give qualitative information about the local ordering of the hard carbon structure. The average interlayer distance of the graphene layers can also be calculated by using TEM observations [START_REF] Bommier | Internal structure -Na storage mechanisms -Electrochemical performance relations in carbons[END_REF]. Figure 8c andd present the TEM images for mangosteen shell derived hard carbons obtained after two hours pyrolysis at 800 and 1600°C. While for the 800°C sample, short turbostratic domains can be seen, for the 1600°C sample, an increase of the number of stacked graphene layers is seen, as well as larger domains. The Raman spectrum of hard carbon is characterized by two peaks, the so-called, D and G bands, as can be seen from Figure 8b. The D band is related to the presence of defects in the graphene layers, while the G band is assigned to the bond streaking of any pair of sp 2 atoms (both rings and chains) [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF], [START_REF] Bommier | Internal structure -Na storage mechanisms -Electrochemical performance relations in carbons[END_REF]. To quantify the level of structural order, normally a ratio between the D and the G bands is calculated, the ID/IG ratio (or also IG/ID ratio). This ratio can be calculated from the absolute peaks heights or the integrated area of the peaks.

The ID/IG ratio is therefore, an indication of the defectiveness of hard carbon material. In most of the hard carbon literature, the ID/IG is found to decrease with temperature, as reported in Figure 9b. The latter indicating that lesser defects are present at higher temperatures, and that the hard carbon structure is more organized.

The lower d002 spaces and ID/IG ratios are indications of a higher structural organization, achieved at higher temperatures for all the hard carbons, despite their origin. However, the absolute values varied in between different precursors.

These variations could be attributed to the different precursor's composition and physical properties. However, as mention at the beginning of the section, the different models and fitting techniques to calculate the structural parameters makes difficult to compare the results between each other, and therefore to conclude, regarding different precursors.

Hard carbon porosity and surface area

The porosity and specific surface area (SSA) in hard carbons could strongly affect the electrochemical performance of the electrode material [START_REF] Irisarri | Review-Hard Carbon Negative Electrode Materials for Sodium-Ion Batteries[END_REF]. For SIB's application, a hard carbon should exhibit a low SSA to limit the irreversible losses, such as the solid electrolyte interface formation, which is thought to occur all over the hard carbon surface [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF]. Moreover, hard carbons might exhibit pores of different sizes, ranging from mesopores to ultramicropores, which can affect differently the Na ions storage [START_REF] Matei Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF].

The accessible porosity volume is normally quantified by gas adsorption/desorption isotherms, as it is typical in porous carbon materials [START_REF] Correa | Biobased Functional Carbon Materials: Production, Characterization, and Applications-A Review[END_REF]. It should be mentioned that there is no accepted model of porosity on hard carbons, and thus, the calculations derived from the isotherm data should be treated with precaution [START_REF] Correa | Biobased Functional Carbon Materials: Production, Characterization, and Applications-A Review[END_REF].

Interestingly, not all the gases used in adsorption isotherms can quantify the same type of pores. Therefore, the final pore volumes and SSAs might vary for the same material, depending on the adsorption gas used. Ghimbeu et al. [START_REF] Matei Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF] have demonstrated that some pores can be closed to the N2 probe molecule but open or accessible to the CO2 molecule. This behaviour originates from the smaller size of the CO2 molecule and the faster kinetics of the test, that allows reaching narrower pores, such as ultramicropores (pore size > 0.7 nm) [START_REF] Matei Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF]. Hard carbons would seem to present higher CO2 SSA than N2 SSA at a given temperature, evidencing a significant presence of ultramicropores in the material.

Moreover, because of the structural reorganisation and pore coalescence with high thermal treatment, hard carbons might also exhibit internal or segregated micropores that are not reachable by the CO2 molecule [START_REF] Buiel | Model of micropore closure in hard carbon prepared from sucrose[END_REF]. These types of pores are often referred as to closed pores.

From this point of view, to observe the totality of the material's pores, regardless of the probe molecule accessibility, scattering techniques such as small-angle Xray diffraction (SAXS) are used [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF]. SAXS technique uses the difference between the electronic density of the carbonaceous phase and a void phase, to quantify the porosity volume and the size of the pores. Naturally, this technique is sensitive to both accessible and closed porosity. A complete description of the physical fundamentals of SAXS technique is given in Chapter III, section 4.4.

It should be mentioned that the porosity in a hard carbon material can comprise up to 23% of its volume, even if its surface area is as low as 1 m 2 .g -1 [START_REF] Dou | Research Update: Hard carbon with closed pores from pectin-free apple pomace waste for Naion batteries[END_REF]. Therefore, the characterisation of the non-accessible pores should not be neglected.

The accessible specific surface area quantified by N2 isotherms (N2 SSA) of the series of biomass-derived hard carbons, was observed to significantly decrease for all the materials with increasing the pyrolysis temperature, as evidenced in Figure 9c (notice the logarithmic y-axis). It was also seen that the N2 SSA varies strongly from one precursor to another at low temperatures. While tamarin seeds derived hard carbon only exhibited 50 m 2 .g -1 of N2 SSA at 800°C, mangosteen shell derived hard carbon exhibited 540 m 2 .g -1 at the same temperature 1 . While sugar cane bagasse hard carbons exhibited N2 SSA under 10 m 2 .g -1 at all the synthesis temperatures. These behaviours evidence again, the impact of the biomass composition on the hard carbon properties, in particular over the developed N2 SSA. However, as their macromolecular and inorganic compositions are not reported, correlations cannot be proposed. 1 Both values were calculated using the BET model. The average pore size of the reported hard carbons seemed to vary with temperature as well. A closer look into the pore size distribution evolution with temperature for argan shell derived hard carbons showed that there seems to be an increase of the average pore size with increasing the temperature. Dahbi et al. [START_REF] Dahbi | Synthesis of hard carbon from argan shells for Na-ion batteries[END_REF] found that the average pore size calculated by BET measurements for the argan shell derived hard carbons, increased from 2.2 to 7.6 nm from 800 to 1300°C, respectively. Moreover, they also found that the ultramicropores present in the hard carbon structure, increase their average radius with increasing the pyrolysis temperature. The SAXS measurements for the same samples are shown in Figure 9d. It can be seen that the intensity of the curves in the microporous region (Q values between 0.1 and 0.5) increase with temperature, evidencing an increase of the contrast between the carbon and the void phase in the materials, in this case, the ultramicropores. The SAXS-calculated average pore radius increased from 0.7 nm at 800 °C to 0.9 nm at 1300°C.

The presence of ultramicropores, quantified by CO2 isotherms, in biomass derived hard carbons is often ignored, as evidenced by the lack of this type of measurements from the reported data in Table 3. The CO2 measurements over hard carbons derived from cellulose [START_REF] Matei Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF], or phenolic resins [START_REF] Beda | Hard carbons derived from green phenolic resins for Na-ion batteries[END_REF], would seem to indicate that the CO2 SSA also decreases with the increase of the final pyrolysis temperature of the hard carbon, from 1000 to 1700°C. Therefore, a similar behaviour could be expected for biomass derived hard carbons.

Hard carbon surface composition

The surface chemistry of hard carbon is of crucial importance as it is in direct contact with the external media, in particular, with the Na ions and the electrolyte species. Therefore, different surface chemistries can affect the electronic and wettability properties of the materials [START_REF] Zhang | Correlation Between Microstructure and Na Storage Behavior in Hard Carbon[END_REF].

The surface chemistry of hard carbons is less studied than its structure and its specific surface area. Techniques like X-ray photoelectron spectroscopy (XPS) [START_REF] Dou | Research Update: Hard carbon with closed pores from pectin-free apple pomace waste for Naion batteries[END_REF],

[80], [START_REF] Lotfabad | High-Density Sodium and Lithium Ion Battery Anodes from Banana Peels[END_REF], and temperature-programmed desorption coupled with mass spectrometry (TPD-MS) [START_REF] Matei Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF], [START_REF] Matei Ghimbeu | Valorizing low cost and renewable lignin as hard carbon for Naion batteries: Impact of lignin grade[END_REF], [START_REF] Zhang | Correlation Between Microstructure and Na Storage Behavior in Hard Carbon[END_REF], are the most useful techniques to analyse the hard carbon surface composition. Energy Dispersive X-ray spectroscopy (EDX)

have been used as well in some reports [START_REF] Li | Controlled synthesis of macroscopic three-dimensional hollow reticulate hard carbon as long-life anode materials for Na-ion batteries[END_REF], [START_REF] Zhu | A green route to synthesize low-cost and high-performance hard carbon as promising sodium-ion battery anodes from sorghum stalk waste[END_REF].

The surface chemistry of carbon materials is dominated by oxygenated functional groups [START_REF] Inagaki | Materials Science and Engineering of Carbon: Characterization[END_REF], such as carbonyls (C=O), anhydrides (O=C-O) and ethers (C-OR). Ncontaining species, such as pyridine, pyrolic and quaternary N have been also

identified [START_REF] Zhang | Correlation Between Microstructure and Na Storage Behavior in Hard Carbon[END_REF]. These observations were confirmed by TPD-MS experiments. It is commonly observed that the content of the surface functional groups seemed to decrease for samples obtained at higher pyrolysis temperatures, while the content of sp 2 C increased [START_REF] Chen | Nonignorable Influence of Oxygen in Hard Carbon for Sodium Ion Storage[END_REF], [START_REF] Wang | Kelp-derived hard carbons as advanced anode materials for sodium-ion batteries[END_REF].

Electrochemical performance

Na ion storage mechanisms

When a hard carbon electrode is tested, it is normally cycled against a metallic Na counter electrode, in a half-cell configuration. The cycling of a half-cell at a constant current density generates the profile shown in Figure 10. First, a reduction occurs at the hard carbon electrode (Na uptake or sodiation), by applying a negative current, that makes the potential of the system decrease, until a cut-off potential. The sodiation achieved in this step is referred to as the reduction capacity. When the cut-off potential is reached, the hard carbon electrode undergoes oxidation and de-sodiation. The Na-ions leave the hard carbon matrix by means of inverting the current sense, which makes the potential of the system to increase again. In this stage, the oxidation capacity is quantified and accounts for the number of Na ions that can be recovered. Therefore, it is also called reversible capacity. Between reduction and oxidation, some electric charges are consumed in irreversible phenomena. This irreversibility is quantified by the Coulombic Efficiency, CE (or Initial Columbic Efficiency, ICE for the first cycle), which is the ratio between the reversible and the reduction capacities.

Notice from Figure 10, that the evolution of the capacity with the potential seems to have two main contributions. First, at potentials over 0.15 V, there is a constant variation of the capacity with the potential. This portion of the capacity is referred to as the sloping capacity. Under 0.15 V, the capacity variation with the potential follows a flatter behaviour, which is commonly referred to as the plateau capacity in literature. This marked differentiation in the potential profile evidenced that different mechanisms are responsible for the Na storage in the hard carbon material.

However, the Na storage mechanisms involved in the slope and plateau regions are a subject of dynamic discussion, and, so far, there is no common agreement on this subject.

The complexity of the hard carbon structure is the main reason for its unacknowledged electrochemical behaviour. Additionally, the different local or bulk techniques used to characterize this structure and the electrochemical behaviour can lead to ambiguous results. Additionally, it is not if each of the possible mechanisms occurred sequentially or if some of them occur simultaneously, adding to the complexity of the interpretation. Initially, Stevens and Dahn [START_REF] Stevens | High Capacity Anode Materials for Rechargeable Sodium-Ion Batteries[END_REF], [START_REF] Stevens | The Mechanisms of Lithium and Sodium Insertion in Carbon Materials[END_REF], proposed that the sloping region of the curve relates to the insertion of the Na ions into the graphene layers while the plateau region is related to the filling of the closed ultramicropores. This model is often called insertion-adsorption or insertion-filling and is schematized in Figure 12a. The authors based their proposition on in-situ XRD and SAXS experiments over hard carbons derived from glucose at 1000°C. It was observed a decrease of the 002 peak intensity during the decrease of potential in the slope region. This decrease of intensity was attributed to the presence of inserted Na ions in the graphene layers. In addition, a decrease of the scattered intensity in SAXS curves was observed in the plateau region, suggesting that the voids are being filled.

Since then, several studies have confirmed this model using similar techniques [START_REF] Irisarri | Review-Hard Carbon Negative Electrode Materials for Sodium-Ion Batteries[END_REF], [START_REF] Simone | Hard carbon derived from cellulose as anode for sodium ion batteries: Dependence of electrochemical properties on structure[END_REF], [START_REF] Gotoh | NMR study for electrochemically inserted Na in hard carbon electrode of sodium ion battery[END_REF]. However, several studies have demonstrated that there might be an influence of the heteroatoms, functional groups and defects concentration of hard carbons on the intercalation of the Na ions into the graphitic domains.

Therefore, these factors could contribute to extending the sloping region. [START_REF] Bommier | New Mechanistic Insights on Na-Ion Storage in Nongraphitizable Carbon[END_REF], [START_REF] Zhang | Correlation Between Microstructure and Na Storage Behavior in Hard Carbon[END_REF],

[70], [START_REF] Morikawa | Mechanism of Sodium Storage in Hard Carbon: An X-Ray Scattering Analysis[END_REF], [START_REF] Cao | Sodium Ion Insertion in Hollow Carbon Nanowires for Battery Applications[END_REF]. Additionally, other studies have proposed that the slope capacity is actually completely related to capacitive or pseudo-capacitive phenomena. While the insertion of the Na ions into the turbostratic domains generates the plateau region. [START_REF] Bommier | New Mechanistic Insights on Na-Ion Storage in Nongraphitizable Carbon[END_REF], [START_REF] Ding | Carbon Nanosheet Frameworks Derived from Peat Moss as High Performance Sodium Ion Battery Anodes[END_REF], [START_REF] Wang | Controlling pseudographtic domain dimension of dandelion derived biomass carbon for excellent sodium-ion storage[END_REF], [START_REF] Lotfabad | High-Density Sodium and Lithium Ion Battery Anodes from Banana Peels[END_REF]. Cao et al. [START_REF] Cao | Sodium Ion Insertion in Hollow Carbon Nanowires for Battery Applications[END_REF] first suggested that the sloping region corresponds to the charge transfer of Na ions at the surface of the graphene layers, while the plateau region was related to the sodium ions insertion/de-insertion into and from the stacked graphene layers. They used cyclic voltamperometry (CV) curves and theoretical energy binding simulations to support their hypothesis. This model is then called adsorption -insertion. Also, Bommier et al. [START_REF] Bommier | New Mechanistic Insights on Na-Ion Storage in Nongraphitizable Carbon[END_REF] then suggested that in the sloping voltage zone, the Na ions are stored by the defective carbon sites, given that the sloping region capacity directly decreases with the decreasing of defects concentration in carbon samples. Furthermore, they were able to observe a reversible expansion and contraction of the 002 peak by ex-situ XRD. This indicated that in the plateau region, the Na ions might intercalate into graphene layers. Bommier and co-workers also performed kinetic studies through galvanostatic intermittent titration (GITT) in which the diffusion of Na ions could be analyzed as a function of the potential. The results confirmed that at high potentials (over 0.2 V), diffusion of Na ions is considerably faster than in the plateau zone, so the ions are stored at easily accessible sites, present in the surface of the turbostratic domains. When reaching the plateau zone, Na ions must overcome the repulsive charge of the ions already present in the surface and enter in-between the graphene layers, so the diffusion decreases, as can be seen from Figure 11. Notice, however, that at 0.05V the diffusivity starts increasing again, which suggest that faster phenomena than diffusion is occurring. Authors suggested that this behaviour is most probably related to the filling of pores by Na clusters, as proposed by Stevens and Dahn. While others denied this mechanism by in-situ XRD, ex-situ 23 Na nuclear magnetic resonance (NMR), electron paramagnetic resonance (EPR) observations and simulations of the Na-ion binding energy [START_REF] Qiu | Manipulating Adsorption-Insertion Mechanisms in Nanostructured Carbon Materials for High-Efficiency Sodium Ion Storage[END_REF].

Morikawa et al. [START_REF] Morikawa | Mechanism of Sodium Storage in Hard Carbon: An X-Ray Scattering Analysis[END_REF] also agreed with the adsorption-insertion model. They highlighted that the predominant contribution of each storage mechanism is dependent on the final pyrolysis temperature, as the hard carbon structure evolves with thermal treatment. Their results are summarized in Figure 12b.

It is evident that many discrepancies exist regarding the Na-ion storage mechanisms, and that different techniques lead to different proposed models.

Despite this fact, and regardless of the potential at which they occur, some insights regarding the Na insertion mechanisms can be highlighted:

-There is capacitive storage, which is related to the interaction of the Na ions with the surface of the turbostratic domains, where the electrolyte can reach (accessible pores, defects, heteroatoms). This contribution is more important for hard carbon synthesized at lower temperatures as these hard carbons present higher concentration of capacitive sites.

62

-There is also adsorption of the Na ions in the internal active sites, such as defects, heteroatoms or functional groups. The internal sites are not in contact with the electrolyte; therefore, diffusion of the Na ions is required to achieve this mechanism.

-The intercalation of the Na ions into the graphene layers of the turbostratic domains seems to be affected by the d002 interlayer space, and the presence of defects and heteroatoms. -The role of the ultramicropores in the hard carbon structure is not well understood yet, and therefore its contribution to Na clustering is not confirmed or rejected. If this phenomenon is favourable, it would be also governed by solid diffusion trough the pore walls.

-The plateau capacity increases with increasing the final pyrolysis temperature for a series of hard carbons obtained under the same conditions. Which could suggest that it is enhanced for more highly organized structures.

Electrochemical performance of the biomass-based hard carbon

The electrochemical performance was observed to vary as a function of temperature, as a result of the structural and surface parameters variation with temperature, that were describe before. For some directly pyrolysed samples, the reported reduction and reversible capacities, the ICE and the rate capability, are gathered and presented in Figure 13.

Notice from Figure 13a, that there is no clear tendency regarding the effect of temperature over the initial reduction capacity. For the different precursor-derived hard carbons, a combined effect of the different degree of organization together with the different surface areas, and surface compositions, are responsible for the heterogeneous behaviour of their reduction capacity.

However, when the reversible capacity is analysed, there is a clear tendency of a higher reversible capacity value with temperature. Interestingly, notice from Figure 13b, that some samples exhibited an inflexion point of the reversible capacity, which starts decreasing between 1200 and 1400°C. Therefore, it seems to be an optimal pyrolysis temperature, which provides the higher amount of active sites for the reversible storage of the Na ions in the hard carbon material. As a matter of fact, controlling the size of the turbostratic domain has been reported to be a critical factor for obtaining high capacity hard carbon anodes for SIBs application [START_REF] Wang | Controlling pseudographtic domain dimension of dandelion derived biomass carbon for excellent sodium-ion storage[END_REF], [START_REF] Zhu | High-yield humic acid-based hard carbons as promising anode materials for sodium-ion batteries[END_REF], [START_REF] Zhu | Expanded biomass-derived hard carbon with ultra-stable performance in sodium-ion batteries[END_REF], [START_REF] Alvin | Extended flat voltage profile of hard carbon synthesized using a two-step carbonization approach as an anode in sodium ion batteries[END_REF]. It can be thought that if the graphene layers are too close to each other at temperatures over 1200°C, the Na insertion into the graphitic domains will be limited. Moreover, if the defects concentration decreases with temperature, this will also reduce active sites for Na storage. On the contrary, if the temperature is low and the resulted hard carbon has not well developed turbostratic domains, there will not be a significant contribution to the capacity. Lower temperatures also imply hard carbons with higher non-C residual elements, such as O, N, S, or inorganic element from the precursor ash that might affect the performance.

To better understand this compromise, it is useful to analyse the plateau and the slope regions of the reversible capacity. The plateau and slope capacity contributions for the different hard carbons obtained from direct pyrolysis are plotted in Figure 14. The increase of the plateau region capacity with increasing temperature seemed to be correlated to the increase of the structural organization of the samples (decrease of the d002 and ID/IG ratios), which also exhibited an inflexion point between 1200 and 1400°C. Therefore, this seems to agree with a possible Na insertion into the turbostratic domains of the hard carbons. Concerning the slope capacity, its variation with temperature seemed less strong than for the plateau capacity, for certain derived hard carbons. However, there seemed to be a negative tendency with temperature. As the surface contributions decrease with increasing temperature, it would seem that the slope region is indeed affected by these contributions. In addition, defects concentration decrease with temperature, probably contributing to the decrease of the sloping region.

Notice that for most of the hard carbons, the slope region contribution is more important to the reversible capacity than the plateau region, for the hard carbons obtained between 800 and 1000°C. This is in agreement with the lower degree of organization and higher N2 SSA and impurities, producing higher capacitive contributions to the reversible capacity in the final hard carbons.

The previous observations would also align with an adsorption-insertion mechanism model.

It has been observed that two-step pyrolysis could increase the plateau capacity of the derived hard carbons, as well as their rate capability [START_REF] Adamson | Peat-derived hard carbon electrodes with superior capacity for sodium-ion batteries[END_REF][123], as can be evidence in Figure 15. It would seem that the increase of the plateau reversible capacity is related to a better-developed turbostratic domain, evidenced by the larger size of the turbostratic domain and to the lower d002 interlayer space, as well as a lower quantity of sp 3 -C in the two-step pyrolysis sample. However, the surface area and the microporosity volume did not seem to be affected by the pre-pyrolysis treatment. It was also observed that the improvement of the plateau capacity is strongly dependent on the first stage pyrolysis temperature. For a series of hard carbons derived from peat moss for example [START_REF] Adamson | Peat-derived hard carbon electrodes with superior capacity for sodium-ion batteries[END_REF], the optimal pre-pyrolysis temperature was found to be 450°C. Higher pre-pyrolysis temperature, of over 600°C, induced higher graphitization degree on the resulting hard carbons, therefore, limiting the Na ions uptake.

Pyrolysis heating rate

Slow heating rates favour the hard carbon yield, which is in agreement with the slow heating rates used to pyrolyse all the studied samples, between 1 and 13 °C.min -1 . Given this small variation, it is hard to assess its role on the hard carbon properties. However, it is worth mentioning that slow heating rates also ensured the preservation of the macrostructure of the biomass precursor after the thermal treatment.

Several authors confirmed this phenomena [START_REF] Li | Controlled synthesis of macroscopic three-dimensional hollow reticulate hard carbon as long-life anode materials for Na-ion batteries[END_REF], [START_REF] Ding | Carbon Nanosheet Frameworks Derived from Peat Moss as High Performance Sodium Ion Battery Anodes[END_REF], [START_REF] Jiang | Biomass carbon micro/nano-structures derived from ramie fibers and corncobs as anode materials for lithium-ion and sodium-ion batteries[END_REF], [START_REF] Wang | Controlling pseudographtic domain dimension of dandelion derived biomass carbon for excellent sodium-ion storage[END_REF], [START_REF] Zhu | Expanded biomass-derived hard carbon with ultra-stable performance in sodium-ion batteries[END_REF], [START_REF] Zhu | Hard carbon derived from corn straw piths as anode materials for sodium ion batteries[END_REF], [START_REF] Adamson | Peat-derived hard carbon electrodes with superior capacity for sodium-ion batteries[END_REF],

[124]- [START_REF] Zhang | High capacity hard carbon derived from lotus stem as anode for sodium ion batteries[END_REF], and took advantage of this fact to produced hierarchical hard carbons.

The hierarchical channels created by the natural macro porosity of the precursor can enhance the electrolyte percolation of the electrode, reduce the diffusion channels of the Na ions and the electrons, and hence improve the ionic and electronic conductivity of the systems. These parameters might be translated into a better reversible capacity and rate capability of the electrodes.

The final morphology and microstructure of some selected hard carbons are presented in Figure 16.

Pyrolysis atmosphere and residence time

Pyrolysis atmosphere

Ar was identified as the most used pyrolysis atmosphere; however, several studies also used N2, as can be seen from Table 3. Few authors had explored the use of reductive atmospheres containing H2 to reduce the surface functional groups [START_REF] Zhu | A porous biomass-derived anode for high-performance sodium-ion batteries[END_REF].

While H2S diluted atmosphere has been used to produced S doped hard carbons [START_REF] Wang | Fabrication of Microporous Sulfur-Doped Carbon Microtubes for High-Performance Sodium-Ion Batteries[END_REF].

As for the other synthesis conditions, the effect of the pyrolysis atmosphere cannot be evaluated if the other parameters are variated. A single study was found where three different atmospheres were used and the final hard carbon characteristics were compared. Marino et al. [START_REF] Marino | Biowaste Lignin-Based Carbonaceous Materials as Anodes for Na-Ion Batteries[END_REF] used Ar, N2 and Ar with 5% of H2 atmospheres to produce hard carbons from almond shells.

Marino and co-workers found that the resulting hard carbons pyrolysed in Ar and in N2 atmospheres exhibited very similar structures and porosity characteristics, as well as electrochemical performances, as could be seen from Table 3. While the hard carbon pyrolysed under the reductive atmosphere (Ar with 5% of H2) presented less content of defects, a higher accessible surface area and closed ultramicropes, as revealed by Raman, N2 isotherms and SAXS observations, respectively. The electrochemical performance of the latter hard carbons was the lowest one of the series of hard carbons, with 128 mAh.g -1 of reversible capacity (at 30 mA.g -1 ) and 47% of ICE. This sample also exhibited larger solid electrolyte interface (SEI) formation and lower plateau capacity, than the N2 and Ar pyrolysed samples, which most probably evidenced and interaction between the reduced surface functional groups, and the higher surface area with the Na ions.

Interestingly, it was observed that when the reduction atmosphere was used as a post-treatment to already synthesised hard carbons, it could actually increase the performance of the hard carbon electrode. Hard carbons derived from apricot shells [START_REF] Zhu | A porous biomass-derived anode for high-performance sodium-ion batteries[END_REF] were reduced under an Ar/H2 atmosphere of 93:7 v/v composition, during 2 hours at 800°C. The resulting samples exhibited also extended accessible surface areas, however, the reversible capacity was observed to be around 400 mAh.g -1 (at 25 mA.g -1 ) with an ICE of 79%.

Pyrolysis residence time

As mentioned before, the pyrolysis residence time was varied between 30 minutes to 6 hours. Most of the studies kept this parameter constant while varying the final pyrolysis temperature, the pre or post-treatments or the pyrolysis atmosphere, making its analysis quite difficult. Only one study explored the effect of the pyrolysis residence time on the hard carbon properties [START_REF] Wang | Low-Cost and High-Performance Hard Carbon Anode Materials for Sodium-Ion Batteries[END_REF]. It was found that with increasing the residence time, the organisation degree of the hard carbon increases, leading to less defective structures. However, the electrochemical performances indicated that there is an optimal time of pyrolysis that led to the best reversible capacity and retention over cycling. For mangosteen shells synthesised at 1500°C, this value was of 2 hours. This single study is not enough to conclude on the residence time effect. However, it is logical to think that longer residence times at a given temperature allow the solid matrix to re-accommodate towards a more organised configuration, than if lower residence times are used.

Pyrolysis pre-treatments

Activation

According to Table 3, activation seemed to be a rather common treatment, before pyrolysis, to increase the accessible specific surface area (N2 SSA) of the hard carbons. Different authors [START_REF] Wu | Apple-Biowaste-Derived Hard Carbon as a Powerful Anode Material for Na-Ion Batteries[END_REF], [START_REF] Xiang | Activated hard carbon from orange peel for lithium/sodium ion battery anode with long cycle life[END_REF]- [START_REF] Hasa | Nanostructured Na-ion and Li-ion anodes for battery application: A comparative overview[END_REF] suggested that the increase of the N2 SSA could lead to an improvement of the electrolyte percolation and the Na ions diffusion into the hard carbon bulk. Higher N2 SSA could create more active sites for Na ions interactions, thus improving the Na storage capacity and the rate capability of the electrodes.

Both, physical and chemical activations have been reported as biomass pretreatments to produce hard carbons. Physical activation have been made using air as an activation agent at 300°C [START_REF] Ding | Carbon Nanosheet Frameworks Derived from Peat Moss as High Performance Sodium Ion Battery Anodes[END_REF], [START_REF] Lotfabad | High-Density Sodium and Lithium Ion Battery Anodes from Banana Peels[END_REF]. Chemical activation have been made using several activation agents, such as KOH [START_REF] Wang | Biomass derived hierarchical porous carbons as high-performance anodes for sodium-ion batteries[END_REF], [START_REF] Xiang | Activated hard carbon from orange peel for lithium/sodium ion battery anode with long cycle life[END_REF], [START_REF] Lv | Peanut shell derived hard carbon as ultralong cycling anodes for lithium and sodium batteries[END_REF], NaOH [START_REF] Jiang | Biomass carbon micro/nano-structures derived from ramie fibers and corncobs as anode materials for lithium-ion and sodium-ion batteries[END_REF] ,

H3PO4 [START_REF] Dou | Pectin, Hemicellulose, or Lignin? Impact of the Biowaste Source on the Performance of Hard Carbons for Sodium-Ion Batteries[END_REF], [START_REF] Wu | Apple-Biowaste-Derived Hard Carbon as a Powerful Anode Material for Na-Ion Batteries[END_REF], [START_REF] Hong | Biomass derived hard carbon used as a high performance anode material for sodium ion batteries[END_REF], [START_REF] Dou | Impact of the Acid Treatment on Lignocellulosic Biomass Hard Carbon for Sodium-Ion Battery Anodes[END_REF], or ZnCl2 [START_REF] Yu | Old-loofah-derived hard carbon for long cyclicity anode in sodium ion battery[END_REF]. Hydrothermal carbonization under different activations agents have also been done to activate and or purified the resulting hard carbon and will be discussed in the upcoming section.

It was observed that the N2 SSA of the activated hard carbons was indeed, significantly higher than the N2 SSA of the same non-activated derived hard carbons, at temperatures below 1000°C. With further thermal treatment, the SSA of the activated and non-activated samples were very close, because of the pore coalescence at higher temperatures. It would also seem that the activation process produced a higher fraction of microporous volume than for the non-activated hard carbons [130] [133].

Besides, of the higher N2 SSA, the structure of the activated hard carbon was also affected by the activation process, as observed from the different d002 distances, ID/IG ratios and SAXS curves, between the activated and non-activated samples [START_REF] Adamson | Peat-derived hard carbon electrodes with superior capacity for sodium-ion batteries[END_REF], [START_REF] Lotfabad | High-Density Sodium and Lithium Ion Battery Anodes from Banana Peels[END_REF], [START_REF] Xiang | Activated hard carbon from orange peel for lithium/sodium ion battery anode with long cycle life[END_REF], [START_REF] Hong | Biomass derived hard carbon used as a high performance anode material for sodium ion batteries[END_REF].

Overall, the effect of the activation treatment on the hard carbon properties seems to be more important at pyrolysis temperatures lower than 1000°C. Dou et al. [START_REF] Dou | Impact of the Acid Treatment on Lignocellulosic Biomass Hard Carbon for Sodium-Ion Battery Anodes[END_REF] suggested that activation treatment could be an effective process to lower the pyrolysis temperature and obtain good hard carbon properties in terms of structural parameters. Up to some extent, the structural degradation of the macromolecules in biomass by the activation agent could compensate for a lower temperature of pyrolysis, to obtain similar structural and surface properties of the final hard carbons.

Regarding the electrochemical performances, it was seen that the first reduction capacity of all the activated samples is higher than for the non-activated samples.

Reduction capacity ranged between 200 and 650 mAh.g -1 for different activated hard carbons reported in Table 3, [START_REF] Wang | Fluorine-Doped Carbon Particles Derived from Lotus Petioles as High-Performance Anode Materials for Sodium-Ion Batteries[END_REF], [START_REF] Adamson | Peat-derived hard carbon electrodes with superior capacity for sodium-ion batteries[END_REF], [START_REF] Lotfabad | High-Density Sodium and Lithium Ion Battery Anodes from Banana Peels[END_REF], [START_REF] Xiang | Activated hard carbon from orange peel for lithium/sodium ion battery anode with long cycle life[END_REF], [START_REF] Hong | Biomass derived hard carbon used as a high performance anode material for sodium ion batteries[END_REF]. The higher reduction capacity is explained by the increase in the SSA of the activated hard carbons. However, the obtained reversible capacities varied from 100 to 300 mAh.g - Moreover, it would seem that chemical activation strongly influences the electrochemical profile of the activated hard carbons, reducing the plateau capacity of the samples [START_REF] Xiang | Activated hard carbon from orange peel for lithium/sodium ion battery anode with long cycle life[END_REF]. While physical activation on air did not seem to affect the plateau capacity of the samples [START_REF] Ding | Carbon Nanosheet Frameworks Derived from Peat Moss as High Performance Sodium Ion Battery Anodes[END_REF], [START_REF] Lotfabad | High-Density Sodium and Lithium Ion Battery Anodes from Banana Peels[END_REF].

Finally, it was observed that the reversible capacity at high current density (1A.g - 1 ) was better for all the activated hard carbons, especially for the hard carbons obtained at temperatures lower than 1000°C. Therefore, the activation treatments did seem to enhance the rate capability thanks to the faster Na ion diffusion into the hard carbon bulk, generated but the increased surface area. However, it was observed that the plateau capacity disappear when higher current densities are applied. There was mainly a contribution of the slope capacity to the recovered reversible capacity at these conditions.

Despite the discussed results, it is thought that more extensive studies are needed to determined, if the higher reversible capacities and rate capabilities of the activated hard carbons, compensate their rather low ICE.

Hydrothermal carbonization

As seen in Table 3, several studies used hydrothermal carbonization (HTC) treatments to produce an intermediated hydrochar that was further pyrolysed. As biomass undergoes different degradation pathways during HTC treatment that during pyrolysis, it is expected that a pre-treated HTC hard carbon exhibited different properties than a directly pyrolysed hard carbon.

In an HTC process, the biomass precursor is subjected to degradation in a closed pressurised system in the presence of a liquid reaction media. For the studied hard carbons, the HTC treatment temperatures ranged from 180 to 250 °C, while the residence time varied from some minutes to more than 24 hours [START_REF] Wang | A review of the hydrothermal carbonization of biomass waste for hydrochar formation: Process conditions, fundamentals, and physicochemical properties[END_REF]. The maximum pressure achieved by the system was not reported for any of the studies.

HTC was performed using different reaction media such as water [START_REF] Cao | Rape seed shuck derived-lamellar hard carbon as anodes for sodium-ion batteries[END_REF], [START_REF] Wang | Facile hydrothermal treatment route of reed straw-derived hard carbon for high performance sodium ion battery[END_REF], [START_REF] Cong | Hydrothermally assisted transformation of corn stalk wastes into high-performance hard carbon anode for sodium-ion batteries[END_REF],

H2SO4 [START_REF] Li | Controlled synthesis of macroscopic three-dimensional hollow reticulate hard carbon as long-life anode materials for Na-ion batteries[END_REF], [START_REF] Li | Controlling the Thickness of Disordered Turbostratic Nanodomains in Hard Carbon with Enhanced Sodium Storage Performance[END_REF], NaOH [START_REF] Xu | Curly hard carbon derived from pistachio shells as highperformance anode materials for sodium-ion batteries[END_REF]- [START_REF] Nie | Hierarchical Porous Carbon Anode Materials Derived from Rice Husks with High Capacity and Long Cycling Stability for Sodium-Ion Batteries[END_REF], and KOH [START_REF] Li | Heteroatom-doped hierarchically porous carbons derived from cucumber stem as highperformance anodes for sodium-ion batteries[END_REF]. Technically speaking, HTC terminology is reserved for treatments in water as reaction media. However, in here, HTC terminology will also be applied for the other reaction media.

It was observed that the main effects of the HTC treatment on the hard carbon properties, are related to the development of a specific morphology, an activation of the carbonaceous subtract and a removal of the inorganic matter from the precursor.

HTC treatment has been widely used to tailor the morphology of the resulting hard carbon particles into reticulate hollow particles [START_REF] Li | Controlled synthesis of macroscopic three-dimensional hollow reticulate hard carbon as long-life anode materials for Na-ion batteries[END_REF], microspheres [START_REF] Li | Controlled synthesis of macroscopic three-dimensional hollow reticulate hard carbon as long-life anode materials for Na-ion batteries[END_REF], [START_REF] Li | Controlling the Thickness of Disordered Turbostratic Nanodomains in Hard Carbon with Enhanced Sodium Storage Performance[END_REF], curly particles [START_REF] Xu | Curly hard carbon derived from pistachio shells as highperformance anode materials for sodium-ion batteries[END_REF], lath shape particles [START_REF] Li | Controlled synthesis of macroscopic three-dimensional hollow reticulate hard carbon as long-life anode materials for Na-ion batteries[END_REF] or lamellar particles [START_REF] Cao | Rape seed shuck derived-lamellar hard carbon as anodes for sodium-ion batteries[END_REF]. Hollow particles and microspheres were obtained after 24 hours of HTC treatment, while the other morphologies were obtained between 2 to 6 hours of HTC treatment. Some of the afford mention samples are presented in the SEM micrographs of Figure 17.

These morphologies were observed to enhance the electrolyte percolation, allowing better Na ions diffusion, improving the reversible capacity, rate capability and capacity retention over cycling, of the studied samples. The hydrochars obtained from pollen grains [START_REF] Li | Controlled synthesis of macroscopic three-dimensional hollow reticulate hard carbon as long-life anode materials for Na-ion batteries[END_REF], pomelo peels [START_REF] Li | Controlling the Thickness of Disordered Turbostratic Nanodomains in Hard Carbon with Enhanced Sodium Storage Performance[END_REF], and pistachio shells [START_REF] Xu | Curly hard carbon derived from pistachio shells as highperformance anode materials for sodium-ion batteries[END_REF] that were further pyrolysed at temperatures between 600 and 1000°C, exhibited the highest reversible capacities, rate capabilities and capacity retention, when compared to the hydrochars pyrolysed at higher temperatures, as can be seen from Table 3.

However, the best ICE was of 62%, which is not outstanding.

The HTC activation was observed to be done using KOH as reaction media. The hard carbons obtained after an HTC treatment with KOH [START_REF] Li | Heteroatom-doped hierarchically porous carbons derived from cucumber stem as highperformance anodes for sodium-ion batteries[END_REF], [START_REF] Wang | Biomass derived hierarchical porous carbons as high-performance anodes for sodium-ion batteries[END_REF], exhibited N2 SSA higher than 1000 m 2 .g -1 . While the other HTC studied samples exhibited surface areas ranging from 10 to 420 m 2 .g -1 .

As for the activation pre-treatments discussed in the previous section, the reversible capacity of the HTC-activated samples seemed to be higher than the non-treated samples. However, the ICE of the samples was observed to be as low as 30%. The reversible capacity of the resulting HTC treated hard carbons is in the range of the activated or directly pyrolysed hard carbons. It is however possible that the same reversible capacities could be achieved at lower pyrolysis temperatures for the HTC treated hard carbons. More efforts are still needed to improve the ICE of the resulting HTC treated samples. A significant improvement on the rate capability of the HTC samples is also seen in all reaction media, and can be associated to the better ionic diffusion and electrolyte accessibility, given the enhanced morphological futures resulting from HTC process.

Finally, some other authors used HTC in water media as a process to reduce the inorganic "impurities" of the biomass precursor [START_REF] Cong | Hydrothermally assisted transformation of corn stalk wastes into high-performance hard carbon anode for sodium-ion batteries[END_REF]. The results showed that the hard carbons without the HTC treatment had a residual ash content of 8 %, while the HTC treated hard carbons presented residual contents as low as 1.5%.

Therefore, HTC could be used, indeed, as a purification process to reduce the inorganic content in the final hard carbons. Moreover, the authors found that the resulting hard carbons presented higher thermal stability than the non-treated sample, under air combustion. These HTC-treated samples exhibited higher reversible, ICE and rate capabilities than the non-treated samples, as can be seen from Table 3, see sample [START_REF] Cong | Hydrothermally assisted transformation of corn stalk wastes into high-performance hard carbon anode for sodium-ion batteries[END_REF].

As water HTC treated samples seem to exhibit better ICE than the other HTC treated samples, it could be thought that the reaction media has an important role on the irreversible loss of the resulting hard carbons.

Conclusions and approach

It was seen that the biomass-derived hard carbon studies are focus on investigating the final properties of the obtained material, while little attention was given to the properties of the precursor. The precursor choice was observed to be related to its "waste" value or availability, regardless of its quality. The studied data evidence a sort of "trial and error" approach in which an available organic waste was taken to produce a hard carbon, and test its feasibility as an electrode for SIBs. Therefore, there are no concrete insights regarding the role of the lignocellulosic biomass composition on the hard carbon properties.

This context is evidenced by the fact that the chemical composition of the biomass precursors used was never quantified. Additionally, a huge amount of synthesis conditions were tested. The combined effect of varying the un-characterised precursors and the synthesis conditions prevents identifying the effect of each parameter on the final hard carbon properties.

Some general trends can however be derived from the existing studies. The pyrolysis temperature seemed to be a key parameter to control the final hard carbon properties. Higher pyrolysis temperatures seemed to develop hard carbons with higher C content and organised structure. Activation pre-pyrolysis treatment increased the accessible surface area of the precursors consistently, while hydrothermal treatments were useful to remove inorganic matter and tailored the desired morphology. Both pre-treatments seemed to improve the rate capability of the hard carbons. However, more efforts are needed to improve the ICE efficiency of these resulting hard carbons. The optimal combination of synthesis conditions has not been explored either.

Based on these observations, it was decided that to elucidate the role of the biomass composition in the hard carbon formation, only the precursors would be variated, while the other parameters were kept constant. It was clear that a systematic study in where only the composition of the precursor is varied, while all the other conditions are kept constant is needed, to identify its role on the hard carbon properties.

Only a preliminary study on the effect of the pyrolysis temperature was done to confirm the tendencies of literature over different biomasses, tested at the same conditions. No acid wash or activation was performed on the precursors, as they were observed to modify the carbonaceous matrix too.

Consequently, the research approach was divided on exploring three main axes:

1. Confirm the effect of the pyrolysis temperature over the hard carbon properties derived from different precursors. A first study was held in where four biomass precursors were selected and pyrolysed at three different temperatures of 1000, 1200 and 1400°C.

Identify the role of the biomass precursor composition on the hard carbon

properties. A series of 25 precursors representative of different biomass types, with variated compositions, were selected. All the precursors were treated at the exact same conditions, under the same characterisation techniques and interpretation tools. They were electrochemically analysed under the same set-up.

3. Understand the origin of the irreversible loss on hard carbons derived from biomass precursors, by ex-situ analysis of the surface composition.

CHAPTE R I I I. MATERIALS AND EXPERIMENTAL

METHODS

This chapter presents:

-The description of the biomass precursors used along with the research and the techniques used to quantify their composition.

-The synthesis procedures for obtaining hard carbon by pyrolysis, as well as the characterisation techniques used to analyse the final physicochemical and electrochemical properties of the hard carbons.

Description of the biomass precursors

Twenty-five different biomass precursors were used to synthesise hard carbons by pyrolysis. The precursors were chosen taking into account their availability and diversity in terms of organic and inorganic composition. Moreover, some of the selected precursors were non-valorised wastes that represent a potential precursor for biochar or carbonaceous materials production. The selected precursors comprise several woody and grasses species, and agricultural residues and byproducts such as cereal straws, husks, cobs, shells and seeds. All the biomasses were selected from CEA's internal databases in which each biomass precursor composition was previously reported. After collecting the physical samples, the composition of each sample was verified with updated measurements available in the laboratory or by performing new composition measurements, as described in section 2.

Sampling

All the biomass precursors were prepared following the NF EN ISO 14780 standard. This standard protocol is as follows: a fraction of the biomass sample is divided into four quarters of a circle. The opposite quarters are selected, while the other two opposite are removed. The selected quarters are mixed and divided again into four quarters, until the targeted amount of sample is reach, herein of 100 g.

Drying

Drying was done to prevent deterioration during storage. All the woody precursors were dried using heated air at 60°C for 24 hours, given their high moisture content in the as-received samples, of over 50%. No drying was necessary for the other samples given their low moisture content of under 15%.

Storage

After sampling and drying of the woody biomass precursors, the biomass samples were stored in sealed plastic bags or plastic containers for further manipulation.

All the samples were store in an indoor storage facility at ambient temperature with open ventilation. The storage conditions and the drying of the high moisture samples have been observed to assure conservation of the biomass samples properties in time [START_REF] Darr | Biomass storage: an update on industrial solutions for baled biomass feedstocks[END_REF].

Grinding

Before pyrolysis for biochar and hard carbon production, all the biomass precursors were ground to a particle size between 300-800 μm, using a variable speed rotor mill Pulverisette 14 of Fritsch, at 700 rpm. For the walnut shell precursor, the sample was first pre-ground in a Universal cutting mill Retsch, equipped with a 1.5 mm grid.

The bagasse, corn husk, coconut husk, tamarin seed and papaya seed were received as ground particles with sizes between 500 and 800 μm, from the IHE2 Institute.

Biomass composition characterization

Biomass precursors underwent the compositional analysis gathered in Table 5. *weight moisture-free basis (wmf%)

Table 5. Compositional analysis for biomass precursors.

Inorganic elemental composition

The content of calcium (Ca), iron (Fe), magnesium (Mg), phosphorus (P), potassium, (K), sodium (Na) and silicon (Si) were quantified following the European standard for solid biofuels -NF EN ISO 16967 (2015). The ground sample was first mineralized in strong acid solutions. Inductively Coupled Plasma Atomic Emission Spectroscopy (ICP-AES) was used to quantify the inorganic elements. Silicon was first molten with lithium tetraborate and then mineralized.

The obtained values were expressed mg of inorganic element per kg of dried biomass (mg.kg -1 ). These analyses were carried out by the laboratory of environmental analysis, SOCOR.

Macromolecular and sugar monomers composition

The quantification of extractives, cellulose, hemicelluloses, and lignin was performed in the French Institute of Technology for Forest-based and Furniture Sectors (FCBA). International standards and internal methods of the Institute were followed. For the extractives content, samples were placed in an accelerated solvent extractor, during two cycles, each at 1500 psi. For the first cycle, water was used at 110°C, followed by acetone at 95°C for the second cycle. From the obtained extractions, the solvents were evaporated; the weighted residue was reported as the extractives content of each sample in wmf%. Lignin content was measured following the Klason method in the standard TAPPI T222 om-83-1988 (from the Technical Association of the Pulp and Paper Industry).

A fraction of the obtained samples from the previous extraction process (extractives free sample) was dissolved in sulphuric acid, and the insoluble precipitate was then recovered and weighted as the lignin content of each biomass. This value was then corrected to subtract the ash amount (which is also insoluble). Final lignin content is express in wmf%. For the cellulose and hemicelluloses content, an internal method was developed from the TAPPI standard T249 cm-85.

Another fraction of the residual extracted sample (extractives free sample) is completely hydrolysed using sulphuric acid to obtain the monomer sugars: glucose, xylose, mannose, galactose, and arabinose, and the acetyl groups present in hemicelluloses backbone or ramifications. The concentration of each sugar was quantified by ionic chromatography equipped with an amperometric detection and the acetic acid liberated from the hydrolysis of acetyl groups by High-Performance Liquid Chromatography (HPLC). It is important to notice that glucose is present in both cellulose and hemicelluloses macromolecules. Therefore, literature ratios of glucose content in different types of biomass are used to estimate the glucose amount in each macromolecule, following the methodology of [START_REF] Da | Characterisation of the Most Representative Agricultural and Forestry Biomasses in France for Gasification[END_REF]. The other sugar monomers and acetyl groups are only present in hemicelluloses. To present the results, the sugar monomers were normalised to the total monomer concentration. The final cellulose and hemicelluloses content is expressed in a weight moisture-free basis (wmf%).

Hard carbon synthesis by pyrolysis

A two-stage pyrolysis set-up was used to synthesise the hard carbons of this study.

The first stage until 450°C was done to remove the major fraction of volatile species of the biomass precursors, obtaining intermediate biochars. The biochars underwent a second thermal treatment until temperatures over 1000°C to achieve the hard carbon structure.

First, approximately 30g of each ground biomass sample was placed in a stainless steel mesh and introduced in a home-made pyrolysis reactor. The reactor consisted of a stainless steel container of 7.5 cm diameter and a 5 cm height, sealed with a graphite gasket. Figure 18a schematises the reactor configuration. The bottom of the reactor supplied an N2 flow while the outlet of the gases was placed in the cover. Heating coils were placed before the N2 inlet and after the gas outlet pipes to ensure a homogeneous temperature of the inlet and outlet gases, respectively.

Five thermocouples were installed in the reactor, one in the nitrogen inlet, three in the bottom of the gasket, and one in the gas outlet, to monitor the pyrolysis temperature, as can be seen from Figure 18a. The whole reactor configuration was placed inside an oven (from Thermolyne with a 33-F30430CM regulator). This configuration ensures that no tar condensation will take place in the gas outlet, and therefore, that there is any risk of blockage. N2 flow was supplied at a rate of 130 L.h -1 at ambient conditions. The thermal treatment is described in Figure 18c and consisted of fast heating until 200°C at 10°C.min -1 , followed by a plateau of 115 min to eliminate moisture and reach a homogeneous heating transfer towards the reactor. Then, the samples were heated at 1°C.min -1 until a final temperature of 450°C, which was held for 20 min. After the thermal treatment, the samples were let to cool down overnight under nitrogen flow.

The obtained biochar was weighted and a fraction of approximately 2 grams was used for hard carbon preparation at high temperatures. The biochar was placed in an alumina crucible of 1 cm wide, 1 cm height, and 15 cm long. The following equations were used to calculate the biochar and hard carbon yields: Where all the masses are in a moisture-free basis (wmf%). A yield in a weight moisture-ash-free basis (wmaf%) was calculated for the biochars, using the following equation:

The ash content in biochar was calculated as described in equation ( 5), in the next section.

The hard carbon yield was calculated as the overall recovered mass after the two steps of pyrolysis. First, a hard carbon yield from the biochar was calculated similarly to equation [START_REF] Deng | Li-ion batteries: basics, progress, and challenges[END_REF]. In this way:

It was no possible to quantify a yield in a wmaf% basis for the hard carbons, as the ash content in the hard carbon was unknown and it was not possible to measure it.

The uncertainty related to the yield calculations arose from the uncertainty of the weight measurements and the possible mass loss in the walls of the reactor and the tubular furnace. The reactor and the crucible where filled always in the same way, using the same mass of sample, therefore, assuming that the mass loss in the equipment was similar for all the samples, a relative uncertainty of 5% was given for each of the pyrolysis steps, following the procedure used in [START_REF] Dahou | Contribution to the understanding of the role of inorganic elements in biomass steam gasification[END_REF]. The final relative uncertainty for the overall hard carbon yield results in 7%.

Five biomass precursors, pine, beechwood, ashwood, miscanthus and wheat straw-1, were used to study the reproducibility of the low-temperature pyrolysis treatment. Each precursor underwent pyrolysis up to 450°C three times, and the yield values were compared, resulting in differences of less than 7% in between the three measurements.

Finally, the obtained samples were identified as follows:

Sample Name

Biochar of an X biomass precursor BC_X-450

Hard carbon at a temperature T in °C, for an X biomass precursor HC_X_T

The two-stage pyrolysis configuration was also chosen to avoid tar condensation over the biochar particles and the blocking of the tubular furnace outlet. In the first place, biomass precursors have a high amount of volatile species that can condensate in the cold zones of the tubular furnace, and therefore, the risk of blocking the gas outlet is imminent. On the other hand, if the gas outlet is blocked, there is no Ar flow. Blocking the outlet of the tubular furnace will activate the relief valve of the furnace. Therefore, pyrolysis would take place in an atmosphere saturated of volatiles coming from biomass thermal degradation, which would enhance secondary reactions such as tar condensation in the biochar particles and affect the yield, surface area and functionalities of the obtained biochars. The use of the homemade reactor with a high N2 flow ensures that minimal tar condensation is happening. Furthermore, pyrolysis until 450°C ensures the complete degradation of cellulose and hemicelluloses and partial degradation of lignin; therefore, the biochars volatile content is lower than for the raw biomass, and there is a low risk of blocking the tubular furnace in the second pyrolysis stage.

Physicochemical characterisation techniques for hard carbon analysis

Given the complexity of hard carbon materials properties, several characterisation techniques are needed to create a general view of the hard carbon specifications regarding structural, porosity, and surface properties. In this way, the techniques chosen for this work are thought to provide the necessary information for further correlation with each hard carbon precursor and electrochemical behaviour. The techniques are summarised in Table 6.

Table 6. Description and justification of the physicochemical characterisations preformed in this study to the obtained hard carbons.

Property Characterization

Technique Objective

Chemical composition

Elemental analysis* To quantify and follow the CHNS content of the HCs.

Particle morphology and size

Scanning electron microscopy

To identify morphological changes and verify the particle size.

Raman spectroscopy

To identify and quantify defects in the turbostratic structure.

Porosity and surface area

Gas isotherms. N2 and CO2.

To quantify the surface area and volumes, and the pore size distribution of the HCs.

CO2 is a complement to N2 results. *techniques performed on the biochar samples too.

Inaccessible porosity

Elemental composition of biochar and hard carbon

The elemental composition in C, H, N, and S content was measured following the same standard and equipment as for the raw biomass quantification. See Table 5.

For the biochars, it is possible to calculate their ash fraction by a mass balance, as ash is assumed to be preserved and concentrated in the solid product until 450°C.

In this way, the ash content of the biochars was calculated by the equation:

For hard carbons, the O content cannot be calculated by difference because ash content is unknown, as some inorganic species are thought to volatilise through the high thermal treatment [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF]. In contrast, some inorganic elements concentrate on the carbon matrix. Ash content for hard carbons was not possible to quantify by the same procedure as the raw biomass because it needs significant mass quantities of the sample that were not available. For the hard carbons produce for chapter III, the oxygen content was measured directly in the elemental analyzer Vario El cube from ELEMENTAR®.

Hard carbon morphology and surface composition

Scanning Electron Microscopy (SEM) was used to observe the morphology and size The presented conditions of EDX analysis assure a depth from one to several microns.

Hard carbon structure

In this work, three techniques were used to characterise the hard carbon structure, X-ray diffraction (XRD), Transmission electron microscopy (TEM) and Raman spectroscopy.

XRD measurements

XRD uses a monochromatic X-ray beam to irradiate a sample. The radiation interacts with the electrons in the atomic arrangements, producing coherent and incoherent scattered radiation. The reflected coherent radiation carries information about the periodic structural organization of the sample. The scattered angle θ is related to long-range ordering dimensions, while the intensity at this angle is related to the location of the electrons in the given order [START_REF] Kvick | X-Ray Diffraction, Materials Science Applications[END_REF]. The Bragg's law describes this phenomenon:

Where:

-λ is the wavelength of the incident radiation, -d is the space between atomic layer arrangement, -hkl are the miller indices for the atomic planes in the sample, they are specific to the material and indicate the plane that is being diffracted;

-and θ is the angle at which coherent scatter occurs.

The obtained XRD pattern is plotted as a scattered intensity as a function of the scattered angle in 2θ. XRD is normally used for crystalline materials, such as graphite, where the atomic layers are well defined and stack in parallel long-range arrangements. However, given that a hard carbon material exhibits turbostratic arrangements of graphene layers, it is possible to use XRD to study this arrangement. Diffraction peaks for hard carbon are therefore broadened features of some of the 2H-graphite diffractions peaks. The first peak around 2θ = 25°, is related to the diffraction of the graphene basal planes (002). This peak overlays with a background signal at low angles, given by the particular microstructure of hard carbon [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF]. A second asymmetric peak near 2θ = 43°, denotes the (100) or [START_REF] Leader | The effect of critical material prices on the competitiveness of clean energy technologies[END_REF] diffraction plane perpendicular to the graphene layers. Iwashita [START_REF] Iwashita | Chapter 2 -X-ray Powder Diffraction[END_REF] consider that this plane should be denoted as [START_REF] Leader | The effect of critical material prices on the competitiveness of clean energy technologies[END_REF] rather than (100), given that no threedimensional order is present in non-graphitic carbon. Moreover, no (101) lines are observed for hard carbon as they are also characteristic of materials with regular stacking order [START_REF] Iwashita | Chapter 2 -X-ray Powder Diffraction[END_REF].

The average interlayer distance between graphene layers (d002) in the turbostratic domains of hard carbon can be calculated using Equation 6. The size of the turbostratic crystal in terms of depth (Lc) and length (La), can be deduced from the full-width at half-maximum parameter (B) of the (002) and ( 10) peaks, respectively, by using the Scherrer equation:

Where:

-k takes values of 0.9 for Lc and 1.84 for La.

XRD measurements for the hard carbons were performed using a D8 Advance Bruker AXS diffractometer with a copper cathode (λ Kα=1,5418 Å). The 2θ angle was varied from 5° to 80° using a step size of 0,05° and a counting time of 1s. When present, crystalline inorganic phases were identified using the PDF4+ database, supplied by the ICDD 3 . Phase identification and background subtraction of the XRD patterns were made using EVA software. TOPAS software was used for the fitting of the hard carbon patterns to determine the average d002 and full-width at half-maximum parameters. The patterns were first corrected for background scattering from the raw data. Each peak was fitted separately using a baseline of polynomial order, n=3. Peaks were modelled using a pseudo-Voight function.

Before XRD measurements, hard carbon samples were ground manually using a mortar to reduce the particle size and assure a homogeneous distribution of the powder in the sample holder.

TEM observations

TEM is used as a complement of XRD observations, to verify the presence of localized graphitic domains and the ultramicropores formation in the arrangement of hard carbons. In contrast to SEM technique, TEM microscopes make use of a high-energy electrons beam (from 60 kV to 3 MeV). This high energetic beam, together with a very thin sample (from 5 nm to 0,5 μm for 100 kV beam), allows the electrons to pass through the target sample an obtain the atomic scale information. The resolution of TEM observations can be as high as 0,2 -0,5 nm. It will depend on the energy of the beam, the sample thickness, and the atomic mass of the sample elemental composition [START_REF] Reimer | Transmission Electron Microscopy: Physics of Image Formation and Microanalysis[END_REF]. For this work, TEM observations, coupled with EDX analysis, were performed in the laboratories of the Institute of Material Science of Mulhouse (IS2M), using a JEOL ARM-200F TEM microscope working at 200kV.

Raman measurements

Finally, Raman spectroscopy is a powerful tool to complete the structural characterisation of hard carbon, as it gives qualitative information about the degree of order and defects of the hard carbon sample. This technique uses a laser beam to irradiate the sample. The sample can scatter the light, in the same frequency as the laser beam or, shifted by the vibrational frequency of the sample molecules. The shifted light is the Raman scattering, and thus, the final Raman spectrum is a signature of the measured vibrational energy of the molecule as a function of the shifted light from the incident beam [START_REF] Ferraro | Introductory Raman Spectroscopy[END_REF]. When applied to carbon materials, the applied laser has the frequency of visible light, which resonates with the π electrons of the carbon arrangements. Therefore, only the vibrations from the sp 2 -hybridised carbons can be detected. However, indirect information can be obtained for the sp 3 -hybridised carbons, which are commonly associated with the defects or low organisation degree in carbon allotropes [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF].

A typical Raman spectrum for hard carbon materials shows two well-defined bands. The G band, which originates from the in-plane starching of the sp 2 C-C bonds, therefore, its apparition does not necessarily imply a six-fold ring configuration of carbon atoms. This band can appear between 1500 and 1630 cm -1 .

The other band is the so-called D band, which is found around 1350 cm -1 . This band corresponds to the breathing modes of graphitic arrangements, and only appears in disordered carbons; in this case, the intensity of the D band is strongly related to the presence of six-fold aromatic rings. [START_REF] Ferreri | Interpretation of Raman spectra of disordered and amorphous carbon[END_REF] The Raman measurements were done in a Renishaw Raman spectrometer, using a laser of 532 nm, at a power of 1mW, approximately, to avoid sample deterioration. The hard carbon sample was deposited on a glass holder and the Raman configuration was set by using an optical microscope with a magnification of x50. Different points of the hard carbon sample were tested, and each scan was held for 2 minutes. The obtained Raman spectra were fitted using WiRE 4.0 software. Herein, a five-band deconvolution model was used, as previously suggested in the doctoral work of Guizani [START_REF] Guizani | Effects of CO2 on the biomass pyro-gasification in High Heating Rate and Low Heating Rate conditions[END_REF] and recently applied to hard carbon analysis by Marino et al. [START_REF] Marino | Biowaste Lignin-Based Carbonaceous Materials as Anodes for Na-Ion Batteries[END_REF]. A detail description of the model is done in Chapter III. The Raman observations were done at least at 3 different locations of the hard carbon surface, to assure the reproducibility of the results.

Hard carbon surface area and porosity

Gas adsorption is the most used technique to quantify the porosity and surface area of porous materials. The accessible surface area of a porous material comprises the internal surface of the pore walls and the external surfaces. In this technique, a gas molecule is adsorbed in the surface area of the material using the physisorption phenomenon. The quantity of adsorbed gas molecules is quantified as a function of the relative pressure P/P0 of the probe gas, where P0 is the vapour pressure of the gas. This process is carried out at a constant temperature, and the final curves are often called adsorption-desorption isotherms (when the probe gas is desorbed from the solid surface). The shape of the isotherms is characteristic of different pore distributions, given that the interactions of the gas molecules is different according to the pore size. The IUPAC 4 classifies the pores into different categories, according to the pore width as follows [START_REF] Lowell | Characterization of Porous Solids and Powders: Surface Area, Pore Size and Density[END_REF]:

-Macropores > 50 nm, -Mesopores: between 2 and 50 nm, -Micropores < 2 nm. Ultrmicropres are micropores with a size < 7 Å [START_REF] Correa | Biobased Functional Carbon Materials: Production, Characterization, and Applications-A Review[END_REF].

The different isotherms are also classified by the IUPAC [START_REF] Lowell | Characterization of Porous Solids and Powders: Surface Area, Pore Size and Density[END_REF] and showed in Figure 19. Type I isotherm is characteristic of highly-microporous samples. Type II, is characteristic of non-porous or microporous samples. Type III is the typical shape for the inexistent, or very weak, interaction between the probe gas and the sample, or a powerful interaction between the probe gas molecules. Type IV is related to mesoporous distribution with a hysteresis caused by the condensation of the gas on the surface. Type V is also related to pore condensation and weak interaction between the probe gas and the sample. Finally, Type VI is a particular case in which the surface of the sample is not porous, but there is multilevel adsorption on the surface.

In order to determine the specific surface area of the material, a model must be used to determine the number of adsorbed layers of the gas. Different models exist that take into account thermodynamic and kinetic mechanisms, such as Langmuir, Brunauer-Emmett-Teller (BET), t-plot, Dubinin and Radushkevich (DR), and Dubinin and Astakhov (DA). On the other hand, molecular simulations approaches are also used, such as density functional theory (DFT) [START_REF] Correa | Biobased Functional Carbon Materials: Production, Characterization, and Applications-A Review[END_REF]. 4 International Union of Pure and Applied Chemistry.

109 The enhanced diffusion kinetics of the CO2 molecule allows it to reach smaller pores. When accessible ultramicropores are present in the hard carbon structure, surface area and volumes obtained by CO2 isotherms can be considerably higher

when measured with CO2 isotherms, than with N2 isotherms.

However, the coalescence of the open pores with thermal treatment gives place to very small pores, for which the probe molecules of the adsorption measurements are inaccessible. These pores are quantified by techniques such as X-ray or neutron scattering. In this work, Small Angle X-ray scattering (SAXS) is used. In this technique, an incident X-ray beam is diffused through the sample and the signal is then received in a detector. After signal treatment, the intensity is plotted as a function of the scattering vector Q:

Where:

-λ is the wavelength of the X-ray source, and -2θ is the angle between the incident beam and the diffused signal, in radians.

The intensity of the received signal is going to be proportional to the electronic density difference between the carbon matrix and the empty voids created by the If the scattering continues towards higher Q values, it will reach the domain of the X-ray diffraction measurement. Therefore, the characteristic peaks of graphene layers diffraction can be seen.

In this work, the SAXS measurements were conducted until Q values of 1, then corresponding to the beginning of zone 3. 112 vector, Silver Behenate (C22H43AgO2) was used. After the measurements, a radial intensity distribution is obtained. This information is normalized for the acquisition time and the transmission intensity. Furthermore, the radial distribution was integrated using the MPA WINNT software. One-dimensional data consists now in intensity as a function of pixels size, which is then converted to the Q vector using the Silver Behenate parameters [START_REF] Simone | Hard carbon derived from cellulose as anode for sodium ion batteries: Dependence of electrochemical properties on structure[END_REF]. Finally, the intensity is normalized to the sample mass.

To analyse the final data, the obtained SAXS pattern was plotted in a logarithmic scale to balance the difference in intensity magnitude. In this work, the pattern was fitted using the non-linear curve fit functionality of OriginPro software with the following model: Where:

-A is the factor describing the weight of the contribution of the particle, -y can vary from 3 to 4 depending on the sample nature, -Q is the scattering vector, -B is the factor of the micropores contribution, -Rg is the radius of gyration, which is a physical parameter related to the average distance from the centre of mass of an object and its rotation axe. It is considered as an average dimensional parameter of the micropores, and -C is the factor of weight for the structural contribution.

The notion of Rg is often used to analyse the micropore size variations as a reference value.

Hard carbon surface composition and functionalities

In addition to EDX, the surface of selected hard carbons was studied by X-ray photoelectron spectroscopy (XPS).

XPS observations were carried out on two selected hard carbons, on both pristine and cycled materials, in order to study the Solid Electrolyte Interface (SEI) formation and the irreversible losses associated with surface phenomena. XPS is based on the photoelectric effect, where a core electron of an atom or ion is ejected by the effect of a photon, in a single step process [START_REF] Van Der Heide | X-ray Photoelectron Spectroscopy: An introduction to Principles and Practices[END_REF]. In XPS observations, highenergy X-ray photons are directed to a sample, and the energy of the photon is entirely transferred to the core electrons, which are emitted at a certain kinetic energy. This kinetic energy is measure by an analyser and is a function of the binding energy of the electron, which in parallel is dependent on the chemical element and its environment. The binding energy is calculated from the following expression [START_REF] Ugarte | In operando study of lithium batteries by coupling X-ray photoelectron spectroscopy and impedance spectroscopy[END_REF]:

Where:

-BE is the binding energy, -hv is the energy of the incident photons, -KE is the kinetic energy of the emitted electron, and -Φ is the work function of the equipment. XPS can analyse the extreme surface of a material, up to 10 nm depth. This is particularly important, as it is on the extreme surface where the interaction of the material with the medium is taking place. Characterization techniques with deeper ranges, such as EDX, do not perceive the precise nature of the extreme surface that might differ considerably from that of the bulk of the material.

Herein, XPS was used to analyse hard carbons in their pristine form, as composite electrodes, and cycled at different states of discharge. The XPS observations were made in a Versaprobe II ULVAC-PHI spectrometer, equipped with a monochromatic beam with an X-ray source of Al, (Al-Kα 1486.6 eV). The beam was configured to focus on the samples at 100 μm of diameter and 25 W of power. Highresolution spectral analyses were performed using a pass energy of 23.9 eV, which corresponds to a resolution of 0.6 eV. All XPS measurements were carried out in an ultra-high vacuum chamber (7.10 -8 mbar). The binding energy calibration was performed using the photoemission lines of gold Au 4f7/2 (83.94 eV), silver Ag 3d5/2 (368.2 eV), and cupper Cu 2p3/2 (932.6 eV). The core-level binding energies were recorded within an error of ±0.1 eV. All the resulting spectra were corrected using the C 1s peak for sp 2 /sp 3 carbon found in the pristine electrodes. The spectra treatment was made in Multipack Software, using a Shirley background and a combination of 80% Gaussian and 20% Lorentzian distributions for the bands.

The pristine hard carbon samples, as well as the uncycled electrodes, were directly placed into a sample holder by a double face cupper tape.

For the cycled electrodes, half-cells were previously opened in a glove box using a coin-cell disassembling tool from HOHSEN. The recovered electrodes were washed in 1 mL of Dimethyl carbonate (DMC) solvent for 30 seconds. This was repeated three times for each electrode. After washing, electrodes were dried in the glovebox's SAS under vacuum. A section of the dried electrode was cut and scotched to the double face cupper tape. The tape support was introduced into a closed transfer vessel inside the glove box, and further transport to be inserted in the XPS equipment.

XPS measurements and data treatment validation were carried out in the DTNM department of the CEA-Grenoble.

Electrodes preparation and electrochemical characterization

Composite electrodes must be first prepared to evaluate the electrochemical behaviour of the hard carbons. A composite electrode commonly has three components: the active material, which is responsible for the Na storage, in this case, the hard carbon. Then, a conductive additive is added to the active material to ensure the correct electronic conductivity of the electrode. Finally, the active material and the conductive additive are held together by a binder, which also allows the electrode mix to adhere to a current collector properly. Herein, the electrode mix is composed of 80% of hard carbon, 10% of conductive additive carbon black super C65 from IMERYS, and 10% of polyvinilydene fluoride binder (PVDF 5130 from Solvay), initially dissolved in a N-methylpyrrolidone solvent (NMP from Sigma Aldrich). Hard carbon samples were first ground manually using a mortar and sieved through a mesh size of 100 μm. The ground hard carbon and the carbon black were then mixed using Cyclohexane. Once the solvent was evaporated, the PVDF solution was added. Furthermore, pure NMP was added until the achieving of a consistent and homogeneous slurry. The slurry was coated onto the current collector, which is an Al foil (25 μm thick) by using a doctor blade of 100 μm of aperture. The aluminium foil was held straight by using a vacuum table. The coated slurry was then placed into an oven (equipped with an extraction circuit) overnight at 55°C to evaporate the NMP solvent. Disks of 14 mm diameter were cut from the resulting electrode tape, using a punching machine. The electrodes disks were further pressed to reduce the thickness of the electrodes assuring higher volumetric energy, proper electronic percolation, and adhesion of the electrode to the current collector. The pressure applied for all the samples was of 5 tones.cm -2 . The electrodes were then weighted, and their final thickness was measured using a precision micrometre, Digital Micrometre of Testing Machines, Inc (TMI)®. Finally, the electrodes underwent drying under dynamic vacuum at 80°C and 48 hours, into a Büchi® oven, to eliminate the remaining moisture. The electrodes were lastly transferred to a glove box for the half-cell assembly.

Half-cells assembly was achieved inside a glove box filled with Ar, with O and H2O contents below 5 ppm. A two-electrode configuration was used for the half-cells test. In this set-up, the composite hard carbon electrode is the working electrode, which is tested against a counter electrode, which plays the role of a reference electrode at the same time. Metallic Na is used as a counter and reference electrode in this work. Until now, metallic Na is the preferred option to be used as a counter electrode and a reference electrode in this configuration. In theory, metallic Na has a small overpotential and exhibits a stable voltage during cycling [START_REF] Talaie | Methods and Protocols for Electrochemical Energy Storage Materials Research[END_REF]. However, there are no commercially available metallic Na tapes of a homogeneous thickness, as for the Li counterpart. Therefore, before the half-cell assembly, the metallic Na electrode needs to be prepared. Special care is needed at this step, as metallic Na is extremely reactive (interacts faster with the glove box atmosphere than metallic Li). Moreover, the thickness of the Na electrode and its final surface will strongly affect the polarization of the cell and might influence the measurements at low voltage in high current rates, as it is demonstrated by Conder et at. [START_REF] Conder | How reliable is the Na metal as a counter electrode in Na-ion half cells?[END_REF]. The following procedure was established to prepare the Na electrode: Na was obtained from a Na cube of 1 cm 3 (Sigma Aldrich 99.9% of purity). In order to avoid complete oxidation, the Na cubes were stored in oil. This oil was carefully washed with Pentane. An external layer of the cube surface was removed to expose the shiny metallic Na with no oxidation on it. The second layer of approximately 2 mm was cut from the pristine metallic Na and placed over a stainless-steel support of 16 mm diameter. Na layers were cut and manipulated using a scalpel and plastic pincer, previously cleaned with DMC. The Na layer was then rolled using an also cleaned plastic roller with no texture. The Na was rolled until achieving a homogeneous, flat surface of approximately half a millimetre. The excess of Na was constantly removed using the scalpel. No traces of Na should rest in the backside of the support or in the lateral space of it, as it can produce circuit cuts. Finally, the surface of the Na electrode was swept carefully with a brush to remove the white-rose oxides formed during its preparation. Na electrodes were stored in a glass vessel while the half-cell is assembled to avoid further surface oxidation. The stability and repeatability of the half-cells rely almost entirely on this step.

Once the Na electrode was ready, the half-cells are assembled in a coin cell setup (CR2032 : 20 mm of diameter, 3,2 mm height), following the organization presented in Figure 21: The can, the cap and the spring are made from stainless steel. A plastic gasket was used to ensure the airtightness of the coin cell. The assembly started from the can, placing the plastic gasket, and the hard carbon electrode, the separators the electrolyte, the metallic Na electrode, the spring, and finally the cap. Viledon separator is a fibrous material used to enhance the electrolyte wettability in the hard carbon electrode. While Celgard 2400 is a polypropylene monolayer membrane used to allow the preferential transport of ionic species and avoid circuit cuts. The electrolyte used in all the coin cells is a preparation consisting of 1 M Na Hexafluorophosphate (NaPF6) solution in 1:1 volume of Ethylene Carbonate (EC) with Dimethyl Carbonate (DMC) and an addition of 1.5 wt% of Fluoroethylene Carbonate (FEC). The electrolyte volume added was 150 μL. The coin cells were finally closed and sealed using a crimping tool from HOHSEN and clean with a wet tissue with isopropanol to remove the excess of electrolyte. After taking the cells out of the glove box, a sorely wash with ethanol was made to remove the electrolyte traces completely.

Galvanostatic cycling test was the technique used to characterize all the coin cells in this work. This technique consists of applying a constant current to charge and discharge the half-cell. The complete process of charge and discharge is called a cycle; thus charge and discharge sparely constitute half-cycles. The current density used to cycle the half-cell is referred to as the C-rate, which is defined as the expected capacity of the active material (Q) divided by the number of hours in which the material is going to be fully charge or discharge. Q can be calculated using the Faraday law, presented in equation 11:

Where:

-Q is the theoretical capacity of the active material, -n is the number of electrons exchange during (dis)charge.

-F is the Faraday constant F= 96500 C.mol -1 , and -M is the molar mass of the active material in g.mol -1 Therefore, the C-rate is expressed as: Herein a C-rate of C/10 was used, with C = 372 mAh.g -1 corresponding to Na insertion into carbon with the formula NaC6, by analogy to the lithium-graphite system. This arbitrary value is used because of the lack of information to properly calculate the theoretical capacity of hard carbon (number of ions exchanged and exact molecular mass). A C/10 rate is commonly used in literature, as it is slow enough to allow a correct quantification of the material's capacity and avoids parallel-undesired phenomena. In practical cases, as capacities are mostly below this value, real C-rates are a little bit above C/10.

The test started with an initial rest of 30 min. Then a negative C-rate was applied, and the half-cells started their discharge (voltage diminution). In this step, the hard carbon material was being reduced while hosting the Na ions (Na + ) coming from the Na electrode. A cut-off bottom voltage of 10 mV was set, in which the halfcell was set to a floating step at a constant voltage. This step was held until the Crate reaches C/100 or until 15 hours. This floating step ensures a complete Na uptake from the hard carbon electrode. After floating, the current direction was inverted. The C-rate is positive and charging started, increasing the voltage of the half-cell, while the hard carbon was being oxidized, releasing the Na + back to the Na electrode. A top voltage limit is also set, here of 3V for Chapter IV and V and 1.5 V for Chapter VI. Notice here that by using a Na counter electrode, Na + supply is not a limiting parameter, and the hard carbon capacity can be quantified with acceptable accuracy.

The capacity obtained from the first reduction step 5 is quantified in mAh.g -1 (of active material). It is the total amount of Na + stored in the hard carbon matrix/defects/surface. The capacity obtained from the first oxidation step will be referred to as the reversible capacity, as it comprises the amount of Na + that can leave back the hard carbon electrode. The difference between the first reduction capacity and the reversible capacity is the irreversibility of the hard carbon material expressed in mAh.g -1 .The ratio of the reversible capacity over the first reduction capacity is called Initial Columbic Efficiency (ICE). At each cycle, the ratio of the oxidation capacity over the reduction capacity is called the CE.

Galvanostatic cycling tests were performed using an Arbin Instrument facility operated by MITS PRO software. At least two coin-cells were tested for each hard carbon sample to analyse the reproducibility of the measurements. For the particular cycling of the samples analysed in Chapter VI, a VMP potentiostat from Biologic, commanded by EC-LAB software, was used.

121

CHAPTE R IV. STUDY OF THE HA RD CARBON FORMATION DURING THERMAL TREATMENT FOR FOUR REFE RENCE BIOMAS S PRE CURSORS

This chapter aims to elucidate the role of the thermal treatment operating conditions and the biomass chemical composition, in the hard carbon formation process. The attention was focus on the pyrolysis temperature while keeping a constant heating rate and pyrolysis atmosphere. The variation of the pyrolysis temperature under a slow heating rate leads to different hard carbon structuration and defines the porosity and the surface of the hard carbons. In this way, very different hard carbon properties can be obtained from the same precursor if it is treated at different pyrolysis temperatures. The heating rate was kept constant at low values of 1°C.min -1 until 450° and 3.33°C.min -1 until the final temperature, to assure a high solid yield of hard carbon and preserve the particle morphology.

Several studies have been published where different pyrolysis temperatures were used to synthesised hard carbons from biomass, as discussed in Chapter II.

Additionally, there are also a wide number of published studies dealing with hard carbon from numerous biomass precursors under fixed synthesis conditions. However, to the best of our knowledge, a transversal study taking into account the influence of both factors on the hard carbon properties is still unavailable.

Herein, four reference biomass precursors were selected to produce hard carbons at different pyrolysis temperatures of 1000, 1200, and 1400°C. The biomass precursors were pine, beechwood, miscanthus and wheat straw (1), described in Chapter III. They were characterised in terms of their elemental and macromolecular composition. Biochar and further hard carbons were synthesised following the procedure described in Chapter III. The obtained hard carbons were analysed to characterise their structure and surface properties.

Finally, the Na storage capacity was evaluated for the series of hard carbons to quantify their electrochemical performance, at a C/10 rate. The correlations between the biomass precursor, the pyrolysis temperature and the electrochemical performance, were analysed. By means of the presented results, the main factors influencing the hard carbon physicochemical properties were identified.

The obtained biochars and hard carbons were denoted as presented in Table 7: Table 7. Biochar and hard carbon abbreviations.

Sample Name

Biochar of an X biomass precursor BC_X-450

Hard carbon at a temperature T in °C, for an X biomass precursor HC_X_T

Chemical composition of the biomass precursors 1.1 C, H, N, O elemental composition

Table 8 presents the chemical composition of the four biomass precursors. As expected [START_REF] Da | Characterisation of the Most Representative Agricultural and Forestry Biomasses in France for Gasification[END_REF], the elemental analysis indicated that the four biomasses were mainly composed of C containing between 45 and 52 wmf%, and of O with values between 39 to 44 wmf%. H was the third most abundant element (around 6 wmf%), followed by N, which was present in a lower magnitude than the latter elements (from 0.1 to 0.8 wmf%).

The O/C ratio for all the biomasses was over 0.8, while their H/C ratio was under 0.1. Franklin [START_REF] Franklin | Crystallite growth in graphitizing and non-graphitizing carbons[END_REF] and further Saurel et at. [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF] have shown that high O/C ratios and low H/C ratios normally suggest that the precursors are suitable for developing the cross-linked turbostratic structure of hard carbon. Therefore, with the high thermal treatment, the four selected precursors should lead to a disordered carbon structure, containing turbostratic domains, as described in Chapter I, Table 2.

Moreover, a high amount of C in the biomass precursor may ensure a richer carbon content in the final hard carbon, thus enhancing the capacity and conductivity of the active material. However, pine possessed a higher amount of extractives than the other wood, beechwood. As extractives are released during early pyrolysis, they do not contribute to the final carbon yield. Notice that indeed, woods presented lower extractives than miscanthus and wheat straw. The higher extractives 123 content for the latter precursors might contribute to a decrease in their final carbon yield.

Macromolecular and sugars composition

The four-biomass precursors presented mainly glucose sugars, followed by xylose and mannose monomers, which indicate that cellulose and hemicelluloses are the main components of these biomasses. Indeed, the macromolecular composition for all the biomass precursors was in agreement with the reported literature [START_REF] Da | Characterisation of the Most Representative Agricultural and Forestry Biomasses in France for Gasification[END_REF].

Cellulose content ranged between 33 and 45 wmf%, while hemicelluloses and lignin contents were both between 20 and 28 wmf%, for all the biomasses. The woody biomasses, namely pine and beechwood, showed a higher amount of hemicelluloses and lignin in the structure than miscanthus and wheat straw.

Normally, rigidity in woods implies a higher lignin content than less rigid plants, such as grasses like miscanthus and wheat straw [START_REF] Mckendry | Energy production from biomass (part 1): overview of biomass[END_REF]. Biomasses containing high lignin contents are known to exhibit higher char yields during slow pyrolysis [START_REF] Hashimoto | Correlations of kinetic parameters in biomass pyrolysis with solid residue yield and lignin content[END_REF].

Therefore, because of the lower extractives content and the higher lignin content, higher biochar yields are expected from the woody precursors than from miscanthus and wheat straw precursors.

Beechwood exhibited higher holocelluloses (cellulose + hemicelluloses) and lower lignin content than pine. This observation is in agreement with the reported literature where normally, hardwoods present higher holocelluloses content and lower lignin contents than softwoods [START_REF] Asmadi | Characteristics of softwood and hardwood pyrolysis in an ampoule reactor[END_REF]. Moreover, it is important to remember that softwoods and hardwoods vary mainly on their cellular arrangements. While a softwood possesses very basic configurations, hardwood has a more complex arrangement also presenting transportation vessels, or pores [START_REF] Ross | Wood handbook : wood as an engineering material[END_REF].

Besides the cellular arrangement, there is also an important molecular distinction between softwoods and hardwoods. Softwood lignin is mainly constituted of guaiacyl-units, while hardwood lignin has both guaiacyl and syringyl units [START_REF] Huang | Relationships between Hemicellulose Composition and Lignin Structure in Woods[END_REF].

Moreover, the hemicelluloses in softwoods such as pine are known to possess a higher fraction of mannose, which is a six-carbon unit sugar. While for hardwood such as beechwood, xylose, a five-carbon unit sugar, is the predominant sugar of hemicelluloses [START_REF] Huang | Relationships between Hemicellulose Composition and Lignin Structure in Woods[END_REF]. This was evidenced by the content of mannose and xylose of pine and beechwood precursors, reported in Table 8.

These structural and molecular differences between pine and beechwood might be translated into different thermal degradation behaviours. They could also lead to different composition and structure of the derived hard carbons.

Miscanthus and wheat straw biomasses presented similar macromolecular distribution. As seen from Chapter III, they both belong to the taxonomic grass family. The cell wall configuration of grasses is known to be a complex arrangement that differs from the cell wall configuration of trees [START_REF] Hatfield | Grass Cell Walls: A Story of Cross-Linking[END_REF]. This difference might not be evident in the macromolecules content, as seen from their similar composition in Table 8, but in the constitutive units of those macromolecules. In particular, the xylose substitutions and interactions with lignin building blocks are the distinctive characteristics of the grass cell walls.

Xylose in grass is substituted with arabinofuranose (higher arabinose content than wood in Table 8), ester-like hydroxycinnamates and glucoronic acid. Moreover, there is evidence of a synergy between the substituted xylose in hemicelluloses and the lignin in grass, in where the addition of ferulic acid and p-coumaric acid in xylose, forms a cross-linked arrangement with the lignin units [START_REF] Hatfield | Grass Cell Walls: A Story of Cross-Linking[END_REF]. The extended degree of cross-linking in the cellular walls of these precursors might lead to a more cross-linked structure of the biochars and the resulting hard carbons [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF].

Ash and inorganic composition

Table 8 evidenced that miscanthus and wheat straw precursors presented the highest content of ash, which was mainly composed of Si, as observed in the literature [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF]. Si in grass come mostly from the silicic acid absorbed from the soil, which further transforms into amorphous or poorly crystalline hydrated silica, denominated opal-A (SiO2.nH2O) [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF] [START_REF] Strömberg | Functions of phytoliths in vascular plants: an evolutionary perspective[END_REF]. The presence of silica is associated with the formation of biological bodies called phytoliths, which are important for the growing, mechanical strength, rigidity, protection and cooling of the plants [START_REF] Yapuchura | Observation of the transformation of silica phytoliths into SiC and SiO2 particles in biomass-derived carbons by using SEM/EDS, Raman spectroscopy, and XRD[END_REF]. Nevertheless, silica could also come from external contributions such as dust or transportation pollution [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF]. Given the close interaction between Si and C in the phytoliths, the formation of a SiC phase is expected at high-temperature thermal treatment (<1400°C) in a reductive atmosphere [START_REF] Qadri | Nanotubes, nanobelts, nanowires, and nanorods of silicon carbide from the wheat husks[END_REF]. The morphology, surface area, ionic and electronic conductivity of this phase will most probably affect the Na uptake capacity of the hard carbon electrode.

Additionally, miscanthus and wheat straw also presented important contents of K (3364 and 8823 mg.Kg -1 , respectively) and Ca (4082 and 2539 mg.Kg -1 , respectively). These values were higher than the ones for pine (590 mg.Kg -1 of K and 2128 mg.Kg -1 of Ca) and beechwood ( 1123 of K and 1496 mg.Kg -1 of Ca).

Wheat straw, in particular, presented a higher concentration of other inorganic elements such as S, P, Mg, and Na than the other biomasses. The concentration of these elements added up to 3709 mg.Kg -1 in wheat straw, while the addition of the same elements for the other biomasses was less than 1954 mg.Kg -1 .

Regarding pine and beechwood ash, they were mainly composed of Ca, together with K and Si, in the case of pine. These observations were also in agreement with the reported literature for wood ash composition [START_REF] Dupont | Kinetic modelling of steam gasification of various woody biomass chars: Influence of inorganic elements[END_REF]. In woods, Si can also form phytoliths and is found mainly in the bark as a protection element [START_REF] Piperno | Phytoliths: A Comprehensive Guide for Archaeologists and Paleoecologists[END_REF]. Naturally, Si can also be provided to the sample by external factors, such as dust or transportation pollution [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF]. Therefore, as for miscanthus and wheat straw, it could be possible to observe the formation of a SiC phase after thermal treatment for this sample.

Notice that beechwood has a higher content of K (1123 mg.Kg -1 ) and a lower content of Ca (1496 mg.Kg -1 ) than pine (590 mg.Kg -1 of K and 2128 mg.Kg -1 ).

Beechwood also presented the lowest content of Si for all the biomasses. Inorganic compounds present in biomass are known to affect the pyrolysis processes, and the final properties of the biochars [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF], [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF]. Consequently, they can also influence the resulting hard carbon properties. In particular, Na, K and Ca can participate as catalytic agents enhancing char formation and solid yield [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF]. Moreover, Ca, Mg, or Fe can enhance the graphitization of non-graphitic carbons during the thermal treatment [START_REF] Van Den Brink | Identifying supply risks by mapping the cobalt supply chain[END_REF] [START_REF] Eftekhari | Sodium-ion batteries: New opportunities beyond energy storage by lithium[END_REF], therefore producing hard carbons with a lower Na uptake capacity. K and Na could also have a negative effect, because Na and K containing compounds can act as activation agents for the carbonaceous structure, producing higher surface areas, thus promoting a higher Solid Electrolyte Interface (SEI) formation and irreversible loss during the first cycles [START_REF] Matei Ghimbeu | Valorizing low cost and renewable lignin as hard carbon for Naion batteries: Impact of lignin grade[END_REF]. The role of Si in biomass pyrolysis, char formation or hard carbon structure, is not well-identified in the literature and will be discussed later in this chapter. It is expected, however, that Si compounds in biomass produce impurities in the hard carbon structure, given its low volatility at the chosen pyrolysis 127 temperatures. Previous studies have shown that phytoliths in biomass can lead to the formation of SiC compounds at temperatures of 1400°C [START_REF] Chiew | A review on the synthesis of SiC from plant-based biomasses[END_REF].

Notice that the identification of the effects of the different composition parameters on the biochar and hard carbon formation could not be as straight forward as one could desire. Several parameters are expected to affect the thermal degradation process simultaneously. While a higher lignin content is expected to enhance the char yield, the effect of inorganic catalytic elements could also enhance the char yield for low lignin-containing biomasses. Even lignin specific composition can have an stronger influence in the solid yield that its absolute quantity, if variations are not significant. Also, while a highly cross-linked hard carbon structure is expected from grass biomasses, their high Ca content can lead to graphitic domains, which have the opposite effect towards Na storage.

Biochar and hard carbon composition evolution with temperature

Biochar composition and yield

The C, H, N and O contents of all the biochars are reported in Table 9. Until 450°C, all the biomass ash is thought to be preserved and concentrated in the biochar. Therefore, a mass balance was done to calculate the ash concentration of the biochar (Eq. 5 of Chapter III). This allowed expressing the biochar yield on a weight moisture-ash free basis (wmaf%), as presented in Eq. 2 of Chapter III. Biochar yields in wmaf% are also reported in Table 9.

The O content in the biochar was directly measured using the elemental analyser.

However, O content of the biochar was also calculated by difference from the mass balance of the biochar composition, for comparison.

The O results showed for the biochar that the measured values were higher than the values obtained by the mass balance of the components. This relates to the fact that the measured O also could quantify the O linked to the ash components (inorganic O), as they can be in the form of oxides, or carbonates [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF]. While the O content obtained from the mass balance relates only to the O linked to the organic matrix (Organic O). Biochar yields in wmaf% were very close for all the biochars, thus slightly higher for the woody ones. As expected, higher moisture-ash-free yields were obtained for the biochars derived from biomasses with higher lignin contents and lower extractives contents [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF].

Figure 22 shows the evolution of C and O contents from the biomass precursors until the final hard carbons at each temperature, passing by the 450°C biochar.

The O results were reported as the calculated values from the mass balance for the biomass and the biochars, and as the measured values for the hard carbons.

The evolution of the O volatilization among thermal treatment should be analysed with precaution, given that the values reported are obtained from different approaches. Even though these values are not comparable, they are used here as reference values to quantify the total O amount in the hard carbons, because over 1000°C, it was no longer possible to calculate or measured the ash fraction. The charring reaction is enhanced by the slow heating rate of the thermal treatment. Therefore, C concentration increases while heteroatoms (O, H, N and S) are released in the gas and the tar products. While the extractive compounds are also volatilized through primary pyrolysis, the inorganic fraction is fully preserved at 450°C [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF], then explaining the lower C content of miscanthus and wheat straw biochars.

Hard carbons composition and yield

Overall, with the increase of the pyrolysis temperature from 1000 to 1400°C, an increase of the C content and a decrease of the O and H content was seen for all the resulting hard carbons, as evidenced in Figure 22.

Hard carbons obtained from pine, beechwood and miscanthus at 1000°C gained 20%, 18% and 16% of C content, respectively, in their structures, compared to the biochar after the first thermal treatment at 450 °C. Hard carbon from wheat straw at 1000°C only gained 12% of C content. From 1000°C to 1400°C, pine and beechwood derived hard carbons only gained 5% of additional C, presenting a final carbon content of over 98 wmf%. While miscanthus and wheat straw derived hard carbons gained 7% and 18% of C, with 93 wmf% and 83 wmf% of final C content, respectively, from 1000 to 1400°C. Miscanthus and wheat straw derived hard carbons concentrated more C at higher temperatures than pine and beechwood derived hard carbons. However, their final C content was lower.

Moreover, wheat straw derived materials, presented a more important C concentration from 1000 to 1400°C ( of 17%) than from 450 to 1000°C ( of 12%), in contrast to the behaviour of all the other samples. This could be explained by the late volatilization of inorganic compounds trapped in the carbonaceous matrix in these samples, at higher temperatures.

The lower C content of miscanthus and wheat straw derived hard carbons could also be related to the catalytic effect of the inorganic compounds in the precursors.

Considering the large concentration of K in miscanthus, and more markedly in wheat straw biomass precursors, one can suggest that an intrinsic activation of the carbonaceous matrix could proceed in the biochars from these two biomasses, and that this could generate porosity and therefore, consume C, lost in the form of CO or CO2 gases [START_REF] Shen | Catalytic pyrolysis of biomass with potassium compounds for Co-production of high-quality biofuels and porous carbons[END_REF].

The total O content (organic O + inorganic O) seemed to decrease for all biomass types between the 450°C-biochars and the hard carbons obtained at 1000 °C, as can be seen from Figure 22. However, once again, the decrease in the total O content was over 62 % for woody derived hard carbons, while for the miscanthus and the wheat straw derived hard carbons it was of only 38 and 20%, respectively.

At 1400°C, the total O content for all the hard carbons decreased again of about 60%, when compared to the O content at 1000°C. As for C evolution, miscanthus and wheat straw hard carbons presented a higher O content reduction between 1000 and 1400°C. Finally, the total O content for all the hard carbons obtained at 1400°C was very similar and range between 2 and 3.5 wmf%.

Overall, biochars from miscanthus and wheat straw, presented relative smaller variations of their C and O contents than the woody derived samples, between 450°C and 1000°C. This could lead to two hypotheses:

-Because miscanthus and wheat straw had a higher ash fraction than woody biomasses, the concentration of this ash fraction during thermal treatment superpose to the concentration of C and the decrease of O. Only when the concentrated ash fraction started to volatilise after 1000°C, the C and O variations are more important.

-It is also possible that the cross-linking difference between the cellular walls of grasses (miscanthus and wheat straw) and woody biomasses, required higher amounts of energy to be fully decomposed, and therefore, the degradation of the structural components is not completed at 1000°C. As discussed before, wheat straw and miscanthus cell walls present a more complex and cross-linked arrangement, within more chemical functionalities than the one of pine and or beechwood, therefore, requiring higher amounts of energy to decompose.

Beyond 1000°C, the interaction of the inorganic compounds and the carbonaceous substrate may be stronger, because of the volatilization of the concentrated inorganic fraction. This can promote side reactions such as intrinsic activation by the action of K compounds or the formation of more thermally stable inorganic phases such as carbides, which might contribute to significant fluctuations in the hard carbon composition. If C is consumed in the formation of another thermally stable phase such as SiC in wheat straw biochar, this will contribute to the low pure C content of the HC_wheat-straw-1400.

The overall yields for the hard carbons obtained at 1400°C are reported in Table 9.

These yields can only be expressed in a moisture-free basis, as it was not possible to quantify the ash content precisely. Notice that the yields for all the hard carbons were very close, between 18.7 and 22.5 wmf%. These values confirmed that the major mass loss occurred in the first pyrolysis step and that nearly 80% of the mass was preserved at the high-temperature pyrolysis step.

Finally, notice from Figure 23 that the sharp decrease of the H/C ratio for the hard carbons obtained at 1400°C, when compared to the 450°C-biochars, confirmed the increase of the aromatization degree in the materials.

Figure 23. H/C vs. O/C molar ratios chart for the biomass precursors, the biochars and the hard carbons obtained at 1400°C (HC 1400). This diagram is a

useful tool to analyse the degree of aromatization of carbon samples.

Structure analysis of the hard carbons

XRD results

The structure of all the hard carbons was first analysed using the X-ray diffraction (XRD) technique. Figure 24 and Figure 25 present the diffraction patterns of the woody and the grass materials, respectively. The biomass and biochar samples were also analysed by XRD, to follow the structural evolution of the samples. All the biomass precursors presented a typical XRD pattern of cellulose as shown in Figure 24 and Figure 25, with two broad diffraction peak lines of Iβ-cellulose, (110) around 2θ-15°, and (200) around 2θ-22° [START_REF] Kim | Characterization of crystalline cellulose in biomass: Basic principles, applications, and limitations of XRD, NMR, IR, Raman, and SFG[END_REF]. Hemicelluloses and lignin do not present crystalline phases and only contributed to the diffuse scattering in the patterns [START_REF] Poulain | Wide-angle X-ray scattering combined with pair distribution function analysis of pyrolyzed wood[END_REF].

According to the database match, crystalline SiO2 (PDF 00-046-1045) peaks of weak intensity were detected for miscanthus and wheat straw biomasses, in agreement to their high Si content measured by ICP. Interestingly, SiO2 phase was also detected for beechwood, but no for pine, which presented almost two times more Si content than beechwood. The latter could indicate that SiO2 in pine is rather amorphous, or that SiO2 in beechwood, miscanthus and wheat straw is present in both amorphous and crystalline phases. As discussed in section 1, when there is Si in biomass because of metabolic activities, it is normally present as opal, a hydrated silica with no long-range order of crystals. Opal can arrange into opal-A and opal-CT in biomass [START_REF] Smith | Opal, cristobalite, and tridymite: Noncrystallinity versus crystallinity, nomenclature of the silica minerals and bibliography[END_REF]. Their XRD signature is shown in Figure 1 of Annex I. Opal-A exhibits a XRD signature with a diffuse scattering in the 2θ angle region 20°-30°, which could be the case for the pine SiO2 phase. While opal-CT can present a broad peak near 2θ -22° with a shoulder at lower angles and could correspond to the SiO2 phase identified for beechwood, miscanthus and wheat straw.

In addition to Si, all the biomasses also presented high quantities of K (590-8823 mg.Kg -1 ) and Ca (1496-4082 mg.Kg -1 ), according to the ICP measurements.

Crystalline phases of compounds containing these elements were not detected in the biomasses XRD patterns, most probably indicating that they have an amorphous/disordered configuration.

When the biomass precursors were converted to biochars at 450°C, the cellulose signature was lost as the result of its degradation, which was almost completed at this temperature [START_REF] Poulain | Wide-angle X-ray scattering combined with pair distribution function analysis of pyrolyzed wood[END_REF], [START_REF] Chen | The structure evolution of biochar from biomass pyrolysis and its correlation with gas pollutant adsorption performance[END_REF]. A very broad feature stands over the background signal, around 2θ -22°, which stands for the scattering of the poorly organized aromatic layers in the biochars. This structural evolution was in agreement to previously reported data for biochars obtained from different biomass precursors reported by Keiluweit et. al [START_REF] Keiluweit | Dynamic Molecular Structure of Plant Biomass-Derived Black Carbon (Biochar)[END_REF], Chen et al. and Yang et al. [START_REF] Yang | Biomass-Based Pyrolytic Polygeneration System for Bamboo Industry Waste: Evolution of the Char Structure and the Pyrolysis Mechanism[END_REF].

The SiO2 phase was now identified in all the biochars, which can relate to the original SiO2 present in the biomass precursor or to a phase transformation of a fraction of the disordered opal-A to crystalline SiO2 with the thermal treatment.

Calcium carbonate CaCO3 (PDF 04-012-0489) was also detected in miscanthus biochar but not in the wheat straw, despite the higher Ca concentration of the latter biomass, suggesting that Ca species would not be present in the same form in these biomasses.

From 1000°C, the materials began to show the hard carbon structure, which was evidenced by the apparition of two broad peaks in the XRD patterns, of all the hard carbons. These peaks around 2θ-23° and 41°, correspond to the (002) and the [START_REF] Li | Heteroatom-doped hierarchically porous carbons derived from cucumber stem as highperformance anodes for sodium-ion batteries[END_REF] diffraction planes of graphite, respectively, as shown in Figure 24a and Figure 25a.

The apparition of these characteristic hard carbons peaks has been reported for biomass-derived hard carbons at 1000°C, as seen from Chapter II.

With further thermal treatment, hard carbons are expected to stack the initially formed graphene layers at 1000°C, creating ultramicropores and all types of structural defects. The local alignment of the graphene layers should produce a shift of the 002 peak towards higher angles. However, it can be seen from Figure 24 and Figure 25, that the evolution of the XRD patterns with temperature is not the same for all the samples. Moreover, mineral phases are detected all along with the thermal treatment for the miscanthus and wheat straw derived hard carbons.

Pine and beechwood derived hard carbons exhibited the typical hard carbon XRD pattern, at 1000, 1200 and 1400°C, as shown in Figure 24a andb. The carbonaceous structure of the woody hard carbons seemed to be quite stable upon the chosen temperatures of synthesis. Pine hard carbons, in particular, did not exhibit a significant shift of the (002) peak, as observed from Figure 24a. The latter indicated that there was no a significant decrease of the d002 interlayer space of the pine hard carbons with increasing the pyrolysis temperature, as it normally commonly see for biomass derived hard carbons (see Chapter II). For the beechwood derived hard carbons, the 002 peak slightly shifted towards higher angles, therefore indicating a diminution of the d002 interlayer space.

The d002 values for the obtained had carbons were calculated and plotted in Figure 26. It should be remembered that these values represent a mean of the graphene interlayer distance in the turbostratic domains, given the disordered structure of the hard carbons. Moreover, because of the inexistence of the three-dimensional graphite peak in the XRD pattern, there is no reference to confirm that the (002

)
peak shift is given by structural difference between the samples and not by a difference in the sample height during the experiment, despite of the extreme care when performing the XRD experiences. Herein the values were compared relatively between the different samples.

From Figure 26 it is seen that pine hard carbon obtained at 1000°C exhibited the lowest interlayer space of the series of pine hard carbons. The difference between the wider and the narrower calculated d002 was only 9 pm, which could be probably affected by the experimental error. For beechwood hard carbons, the d002 interlayer space was observed to decrease with the thermal treatment. This diminution with temperature is smooth but constant and accounts for 24 pm of difference between the 1000°C the 1400°C beechwood hard carbons.

These results suggested that the pine carbonaceous structure is rather stable already at 1000°C and did not undergo further structural organization of its structure with higher temperatures, as when compared to beechwood-derived structures. This behaviour is in agreement with the rather stable chemical composition of pine hard carbons from 1000°C. Moreover, the less cross-linked macromolecular arrangement in softwoods like pine when compared to hardwoods like beechwood, could also be implied in this structural and chemical stability.

For woody hard carbons, only SiO2 peaks were still detected in the 1000 and 1200°C hard carbons. SiO2 was no longer detected at 1400°C, probably indicating reactivity towards another phase.

Figure 26. d002 interlayer space for all the hard carbons.

The XRD patterns for miscanthus and wheat straw derived hard carbons differed from those of pine and beechwood derived hard carbons, especially at 1400°C.

From 1000 to 1200°C, the typical hard carbon patterns were observed. The (002) peak did not shift considerably for miscanthus hard carbons while it moved towards higher angles for wheat straw hard carbons, as seen from Figure 25. The (100) peak evolved similarly for both samples. At the same temperatures of 1000 to 1200°C, a higher contribution of mineral mater was seen for the miscanthus and the wheat straw derived hard carbons, than for the woody derived hard carbons.

The 1000 and 1200°C hard carbons from miscanthus and wheat straw exhibited SiO2 and hydrated potassium carbonate, K2CO3 *1,5H2O 7 (PDF 00-011-0655)

phases. The K2CO3 found in the hard carbons most probably originates from the thermal decomposition of the K containing compounds in the biomass precursors, such as KOH. The formation of K2CO3 from KOH during thermal treatment, as well as the decomposition of the K2CO3 phase at temperatures over 700°C, can lead to pore formation, by the mechanisms described in equations 1-5. 7 Notice that the identified K2CO3 is hydrated; however, this water is most probably not a result of the thermal treatment, but a consequence of the adsorption of ambient moisture by the sample. At 1400°C, the XRD patterns for miscanthus and wheat straw hard carbons presented very different signatures than for the woody derived hard carbons. First, the (002) peak was asymmetric and presented an embedded peak near 2θ -26°.

The presence of an embedded peak in the (002) broad feature suggested that there was a significant contribution of long-order graphitic-like domains within the turbostratic domains [START_REF] Iwashita | Chapter 2 -X-ray Powder Diffraction[END_REF]. This localized graphitization could be a result of a catalytic graphitization by the action of the inorganic species present in the carbonaceous matrix. Previous studies have found that some Ca compounds could enhance significantly, the graphitization of a carbonaceous material [START_REF] Ōya | Phenomena of catalytic graphitization[END_REF], [START_REF] Hirano | Cooperative accelerating effect of calcium carbonate and gaseous nitrogen on graphitization of carbon[END_REF].

Given the high concentration of Ca in miscanthus and even higher Ca concentration in wheat straw biomasses, catalytic graphitization may happen in these materials at 1400°C. However, additional observations, such as TEM, are needed to confirm the presence of graphitic-like domains, as discussed later in section 3.2.

Additionally, the 1400°C-XRD patterns of miscanthus and wheat straw hard carbons exhibited a major inorganic phase consisting of silicon carbide (SiC) compounds, as expected from the high Si content of the biomass precursors. These peaks were more intense for wheat straw hard carbon than for miscanthus hard carbon, most probably given the higher Si content in wheat straw precursor.

According to the diffraction database, the detected peaks could correspond to Si3CaC4 (PDF 01-82-3652), SiC (Moissanite, PDF 00-029-1129), or SiCN (PDF 01-

074-2308).

SiC is the product of a carbothermal reduction reaction between a source of C and SiO2 at high temperatures as described from equations 6 to 11 [START_REF] Chiew | A review on the synthesis of SiC from plant-based biomasses[END_REF].

Normally, industrial SiC is obtained from synthetic sources of Si and SiO2 at temperatures around 2000°C [START_REF] Chiew | A review on the synthesis of SiC from plant-based biomasses[END_REF]. However, in biomass precursors, the Si and the C sources are strongly linked together in the Opal bodies (phytoliths), which allows the formation of SiC at lower temperatures, from 1200 to 1800°C, by solid state reaction mechanisms (equations 6-7) [START_REF] Chiew | A review on the synthesis of SiC from plant-based biomasses[END_REF]. Therefore, it is possible to conclude that SiC was formed in our miscanthus and wheat straw hard carbons obtained at 1400°C.

TEM results

To confirm the presence of long-range order zones in the hard carbons derived from miscanthus and wheats straw precursors, we performed TEM observations on all the hard carbons obtained at 1400°C, which are reported in Figure 27.

For pine and beechwood hard carbons, the graphene layers were mainly organised in a turbostratic fashion, with no long-range order. It was possible to observe the ultramicropores formed within the nano-voids between the turbostratic domains.

Nevertheless, wider TEM observations revealed a more organized zone for pine hard carbon, as shown in Figure 28a and b, with zones presenting several graphene layers stacked regularly. Given that XRD observations did not evidence the presence of this phase in the pine pattern, these zones were probably localised over certain particles and were not develop homogeneously for the entire pine sample.

Pine derived hard carbon also seemed to present spherical nanoparticles within the carbon matrix, most probably coming from SiO2 or SiC compounds. As for the long-order phase, these particles seemed to be heterogeneously localised within the hard carbon structure and are not representative of the entire hard carbon sample, as they are not identified in the XRD pattern.

Figure 27. TEM observations for a) pine b) beechwood c) miscanthus and d) wheat straw derived hard carbons at 1400°C.

For miscanthus and wheat straw derived hard carbons, Figure 27c andd, respectively, confirmed that there is a localized zone of long-order graphitic-like domains, together with the turbostratic arrangements and the ultramicropores.

The graphitic domains seemed to be composed of 10 stacked graphene layers of more than 30 nm long. These observations support the presence of the (002) sharp embedded peak, as a graphitic-like domain in these hard carbons. Moreover, a wider observation of these samples also confirmed the presence of a noncarbonaceous phase, presenting morphologies of wires and spherical nanoparticles, as shown in Figure 28c andd. A TEM-EDX analysis over these particles evidenced that they have a high concentration of Si, therefore, confirming the formation of SiC compounds as whiskers and spherical particles from the thermal treatment of the raw miscanthus and wheat straw biomasses.

b) TEM image of a particle of pine derived hard carbon presenting spherical nanoparticles. c) TEM image over the SiC whiskers and spherical SiC nanoparticles for miscanthus and d) wheat straw derived hard carbons. The table under the TEM images reports the atomic elemental composition of the analysed

zones for c and d.

Raman results

Finally, Raman spectroscopy observations complemented the structural analysis of the obtained hard carbons, the Raman spectra of all the hard carbons are grouped in Figure 29.

All the spectra exhibited a D band around 1350 cm -1 and a G band around 1590 cm -1 . Qualitatively, the intensity of the D band was seen to be as strong as the G band intensity, confirming the disordered configuration of the graphene layers and the presence of a considerable amount of defects [START_REF] Sheng | Char structure characterised by Raman spectroscopy and its correlations with combustion reactivity[END_REF]. The main type of defects in hard carbons can be attributed to in-plane imperfections of the lattice, such as mono-or di-vacancies or stone-wales defects [START_REF] Tsai | Ab initio study of sodium intercalation into disordered carbon[END_REF]. In addition, the edges of the graphene sheets and the heteroatoms impurities could also contribute to the defects concentration of hard carbons [START_REF] Sadezky | Raman microspectroscopy of soot and related carbonaceous materials: Spectral analysis and structural information[END_REF].

To quantify the structural evolution of the series of hard carbons, the spectra were deconvoluted using five bands as suggested by Sadezky et al. [START_REF] Sadezky | Raman microspectroscopy of soot and related carbonaceous materials: Spectral analysis and structural information[END_REF], for fitting Raman spectra of disordered turbostratic carbons.

In addition to the D and the G band, three more bands were needed to improve the fit and are thought to be related with the contribution of disordered features in the structure, namely, D2, D3 and D4 bands. The D band near 1350 cm -1 corresponded herein to the D1 band. The D4 band was introduced around 1200 cm -1 and is seen in very disordered carbons, probably evidencing sp 2 -sp 3 sites or polyene-like structures [START_REF] Sheng | Char structure characterised by Raman spectroscopy and its correlations with combustion reactivity[END_REF]. The D3 band was located around 1500 cm -1 , it is a broad band that occurs in poorly organised structures, and it is commonly related to amorphous bonded carbon [START_REF] Sheng | Char structure characterised by Raman spectroscopy and its correlations with combustion reactivity[END_REF]. Finally, the D2 band was located near 1620 cm -1 and is always present if the D1 band is present for a given material. Its origin is in the same vibrations as for the G band, but accounts for graphene layers that are not directly stacked between other graphene layers [START_REF] Sheng | Char structure characterised by Raman spectroscopy and its correlations with combustion reactivity[END_REF].

The best fit was obtained with a combination of Gaussian and Lorentzian bands.

Once the bands were deconvoluted, the area of each band was used to analyse the ratios between bands in the spectra. Figure 30 shows the area evolution of all the bands with respect to the total area of the bands (AT) and to the area of the G band (AG). There was no significant evolution of the AG/AT ratio with the thermal treatment for any of the hard carbons, as can be seen from Figure 30e. This could indicate that there is no significant increase of the turbostratic domains formed from 1000°C to 1400°C in the materials.

The AD1/AT ratio slightly evolved with temperature. It seemed to decrease with the thermal treatment (Figure 30a), except for the pine-derived hard carbons, for which it was rather stable. The decrease of the AD1/AT ratio with temperature confirmed that fewer defects were present at higher temperatures. This defects diminution can be attributed to less lattice defects, less defective graphene edges, and lower impurities in the sample. The latter was confirmed by the increasing content in C with temperature, observed by elemental analysis for all the hard carbons discussed before.

Figure 9f shows how the AD1/AG ratio significantly decreased with temperature for beechwood and miscanthus hard carbons, while it was rather stable for pine and wheat straw hard carbons, the latter presenting the lowest ratios. Beechwood and miscanthus derived hard carbons followed the expected behaviour for the AD1/AG ratio of a hard carbon with thermal treatment (see Chapter II). The latter indicates that the structure increased its organisation. Given that the G band contribution did not vary with temperature, while the D1 band contribution decreased with temperature for these hard carbons, the higher structural organization is gained by remediation of the defective sites rather than by the growth of the turbostratic domains.

Figure 30. Area ratios for a) AD1/AT, b) AD2/AT, c) AD3/AT, d) AD4/AT, e) AG/AT, f)AD1/AG, g) AD2/AG, h) AD3/AG and i) AD4/AG for all the obtained hard carbon samples.

The AD2/AT and AD2/AG ratios did not evolve with temperature, confirming that the turbostratic domains are rather stable for all the hard carbons. The AD3/AT and AD3/AG ratios only varied for the wheat straw derived hard carbons, both increasing from 1000 to 1200 °C thermal treatment. This observation would suggest that the contribution of amorphous C species in the structure of wheat straw hard carbon is strong until 1200°C. Amorphous C contribution can be related to organic molecules or fragments of functional groups [START_REF] Sheng | Char structure characterised by Raman spectroscopy and its correlations with combustion reactivity[END_REF], which would be consisted to the thermal stability of the wheat straw samples discussed before.

HC_pine

HC_beechwood HC_miscanthus HC_wheat-straw

a) b) c) d) e) f) g) h) i)
Finally, AD4/AT and AD4/AG only seemed to vary for the pine hard carbons; however, the incertitude for both ratios at 1000°C was also significant and it would be inaccurate to conclude on this observation.

To summarise, the lack of variation of the D and G bands for the pine derived hard carbons confirmed the XRD observations, where it was suspected that these materials presented a rather stable structure that did not evolve significantly from 1000 to 1400°C. Moreover, the constant decrease with the temperature of the AD1/AG ratio for beechwood and miscanthus hard carbons, together with their constant value of the AG/AT ratio, indicated that the increase in their organization degree with temperature is related to a decrease of the imperfections, rather than to an expansion of the turbostratic domains. Finally, the lowest AD1/AG ration of HC_wheat-straw_1400 is an indication of its higher organisation degree when compared with the other obtained hard carbons.

Morphology and surface analysis of the hard carbons

Morphology

The morphology evolution of the hard carbon particles at high temperature was studied by SEM. Figure 31, Figure 32, Figure 33, and Figure 34 showed the SEM images for the hard carbons obtained from pine, beechwood, miscanthus and wheat straw, respectively, at all the temperatures.

It was observed for the pine and beechwood derived hard carbons that the morphology of the particles seemed to be preserved at all the three temperatures of pyrolysis. The slow heating rate of the thermal treatment is known to enhance the morphology preservation of the particles, given the that the slow heating rates and the natural vessels in biomass, allow the volatiles to transport through the carbon matrix without major morphological alterations [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF]. This behaviour has been observed for biomass-derived hard carbons obtained at 1300°C [128] [113]. It was not possible to performed SEM images on the biomass precursors or on the biochar samples. However, the particle size, the fibrous nature of the particles and the presence of macropores (most probably coming from the wood vessels), seemed to indicate the conservation of these futures from the biomass precursors. The surface of the pine and beechwood hard carbons seemed rugose, with smaller particles immersed over it. The EDX analysis reported in Table 10 showed that mainly C and O were present in the surface of these samples. Very small peaks of K and Ca were detected on the surface, which were consistent with the K and Ca content of the parent biomass. The small-immersed particles might be related to this minor inorganic presence; however, it is not possible to confirm this hypothesis with the performed observations. For the miscanthus and wheat straw derived hard carbons it was observed that the particles seemed larger, thinner and more curved than the woody derived hard carbon particles, at all the temperatures of pyrolysis, as can be seen from Figure 33 and Figure 34.

Unlike the woody hard carbons, the miscanthus and wheat straw derived hard carbons surface and morphology was observed to evolve with thermal treatment. This evolution was related to the SiC formation observed in the XRD patterns. At 1000°C, there was no particular evidence of an inorganic phase formation over the carbonaceous surface for miscanthus hard carbon, see Figure 33a. While wheat straw hard carbon obtained at 1000°C presented localised zones in the surface with spherical agglomerates, that seemed to originate from a liquid intermediated phase of SiO2, as has been described by Dahou et al. [START_REF] Dahou | Contribution to the understanding of the role of inorganic elements in biomass steam gasification[END_REF]. At 1200°C these agglomerates seemed to formed all over the surface of miscanthus hard carbons and were related to the intermediate state of SiO2, prior to the formation of SiC [START_REF] Yapuchura | Observation of the transformation of silica phytoliths into SiC and SiO2 particles in biomass-derived carbons by using SEM/EDS, Raman spectroscopy, and XRD[END_REF], as it is clearly seen from Figure 33b (higher magnification zoom).

Between 1200°C and 1400°C the SiO2 agglomerates reacted with the hard carbon surface forming SiC compounds in the form of whiskers over the particles surface, following the equations 6-11, presented in section 3.1. The whiskers formation was identified at the surface of the hard carbon derived from miscanthus at 1400°C.

For wheat straw hard carbons, the whiskers were identified from 1200°C. The whiskers formation seemed to be more significant for wheat straw hard carbon obtained at 1400°C, which is in agreement with its higher Si concentration and higher intensity of the SiC peaks in the XRD patterns.

The composition of the whiskers was confirmed by an EDX cartography over the surface of HC_wheat-straw_1400, shown in Figure 35, where Si was detected in high amounts on the formed SiC whiskers locations. Ca also was extended over the carbonaceous surface, while over the surface without SiC whiskers, C was the main element present.

Figure 35. EDX cartography for HC_wheat-straw_1400 and Chemical composition by EDX analysis of HC_wheat-straw_1400 on two different zones of the surface.

To sum up, the SEM-EDX observations fully support the XRD measurements in which the SiC phases are identified and allow determining their morphology of whiskers over the carbonaceous particles.

Specific surface area

The specific surface area (SSA) and pore volume of all the hard carbons were quantified by N2 and CO2 adsorption isotherms, presented in Figure 36 and Figure 37, respectively. The N2 isotherms for pine, beechwood and miscanthus hard carbons presented the typical shape of microporous materials, which was clearer for the woody hard carbons obtained at 1000°C. On the contrary, wheat straw hard carbons at 1400°C exhibited the typical isotherm of mesoporous materials, presenting some hysteresis related to the condensation of the adsorbate on the surface. The CO2 isotherms for all the samples are characteristic of microporous distributions.

N2 SSA

The N2 SSA of the 450°C-biochars, was not significant, as exhibited by the singlepoint measurements of the N2 SSA of these materials. The obtained values were of 3.3, 0.7, 1.5 and 2.5 m 2 .g -1 for the biochars derived from pine, beechwood, miscanthus and wheat straw, respectively. These low N2 SSA values have been observed for biochars from different precursors and are related to the incomplete removal of volatile matter in the biochar [START_REF] Chen | The structure evolution of biochar from biomass pyrolysis and its correlation with gas pollutant adsorption performance[END_REF]. The evolution of the N2 SSA with the final pyrolysis temperature is reported in Figure 38, for all the hard carbons. Except for pine derived hard carbon at 1000°C, all the other materials presented N2 SSA lower than 20 m 2 .g -1 at 1000°C. With temperature, the N2 SSA of pine, beechwood and miscanthus hard carbons seemed to decrease, while for wheat straw hard carbons, it increased with the thermal treatment.

In carbon materials derived from biomass, micro and mesoporosity created upon pyrolysis, reach a maximum commonly between 600 and 800°C, for woody biomass [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF], [START_REF] Weber | Properties of biochar[END_REF] and around 900°C for other biomasses [START_REF] Chen | The structure evolution of biochar from biomass pyrolysis and its correlation with gas pollutant adsorption performance[END_REF], which is exhibited in large N2 derived SSA (over 100 m 2 .g -1 ). With further thermal treatment, the surface area decreases because of the tar re-condensation on the pore walls or by the coalescence of the pores given by the structural reorganisation [START_REF] Rodríguez-Reinoso | Activated carbons from lignocellulosic materials by chemical and/or physical activation: an overview[END_REF] obtained hard carbons at 1000°C from beechwood, miscanthus and wheat straw hard carbons. Given that at temperatures over 1000°C, almost all the volatiles have been removed from the solid structure, the coalescence of the pores is thought to come from a microstructural reorganization, rather than from a tar recondensation on the pore walls. Naturally, there is not a punctual temperature distinction between these two phenomena of porosity diminution among thermal treatment, and they can undergo simultaneously at some point during the hightemperature treatment.

HC_pine_1000 exhibited a significantly higher N2 SSA of 141 m 2 .g -1 ; when compared to the N2 SSA of other hard carbons at the same temperature, including the other woody hard carbon, HC_beechwood_1000 (16 m 2 .g -1 ). This could be explained by their different macromolecular arrangement in the plant cell wall, inducing different types of macro, meso, and microporosity in the biomass precursor [START_REF] Darmawan | Characterization of Lignocellulosic Biomass as Raw Material for the Production of Porous Carbon-based Materials[END_REF], which are further enhanced by thermal treatment and might remain visible even at 1000°C. Hence, Darmawan et al. [START_REF] Darmawan | Characterization of Lignocellulosic Biomass as Raw Material for the Production of Porous Carbon-based Materials[END_REF] suggested in their study that softwoods such as pine may have a higher potential to produce porous carbons, given the particular arrangement of cellulose microfibrils in the structure that degrades easily compared to those of hardwoods. Nevertheless, N2 SSA for pine hard carbons dropped significantly at 1200°C, presenting the same trends of N2 SSA values as beechwood.

In contrast, the N2 SSA of the wheat straw derived hard carbons increased with the temperature. This increase of the N2 SSA can be explained by the contribution of the SiC whiskers formation and the intrinsic activation by the K containing compounds. The SiC whiskers growth over the carbonaceous surface might contribute to an increase of the porous volume. As seen by XRD and SEM, these whiskers were mainly present at 1400°C, explaining why their effect over the measured N2 SSA is more important at 1400°C. Additionally, the presence of K2CO3 revealed by XRD indicated that an intrinsic activation of the carbonaceous structure has occurred in this material, by the mechanisms described before (equations 1-5). As the overall K content is almost 3 times higher for wheat straw than for miscanthus, the K concentration in miscanthus may not be sufficient to produce as much porosity as in the case of wheat straw samples.

Figure 38. Evolution of the a) N2 SSA, b) CO2 SSA, c) N2 volume and d) CO2

microporous volume with temperature for the series of hard carbon.

CO2 SSA

The CO2 SSA and volumes values for all the hard carbons are reported in Figure 38, as a function of the final pyrolysis temperature. At 1000°C, the CO2 SSA values of all the materials were significantly higher than the N2 values. This was in agreement with the better accessibility of the CO2 molecule to smaller pores (<1 nm). The smaller molecular size, together with the higher temperature of the measurement when compared to the N2 experiments, enhanced the adsorption kinetics of the CO2 molecule, therefore allowing it to reach smaller pores.

As can be seen from Figure 38, the CO2 SSA and volume decreased with temperature for all the hard carbons, even for the wheat straw derived hard carbons, for which the N2 SSA and volume increased with temperature. Therefore, for pine, beechwood and miscanthus derived hard carbons, the CO2 and the N2 SSAs and volumes followed the same trends with temperature. The decrease in temperature treatment for both N2 and CO2 isotherms. The values obtained for the N2 PSD showed two peaks for miscanthus and wheat straw derived hard carbons.

A first peak for ultramicropores was observed around 0.6-0.9 nm and a second peak of mesopores was observed between 2.3 and 3.3 nm for miscanthus and 23 nm for wheat straw hard carbons. For woody hard carbons, the N2 PSD was centred around 1 and 3.4 nm. For the CO2 PSD, there seemed to be a single mean value centred on 0.5 nm for all the hard carbons.

SAXS observations

To complement the porosity analysis of the hard carbons, the internal and external porosity were observed by SAXS measurements, and the patterns are shown in Figure 39. The patterns obtained from the 450°C-biochars indicated that at this temperature, there is no significant porosity development in the samples, confirming the observation obtained by single-point measurements of N2 SSA.

Interestingly, two different behaviours were obtained from the samples at higher temperatures, as for the XRD observations. For the woody derived hard carbons, the SAXS patterns presented the typical shape for microporous carbons, with the characteristic Guinier's domain near 0,1 Å -1 . Miscanthus hard carbons exhibited very short microporous domains, while they were almost inexistent for the wheat straw hard carbons. The atypical result for miscanthus and wheat straw materials can be directly attributed to their higher inorganic content, especially in Si. Indeed, Si electronic contrast is higher than C electronic contrast, therefore masking the diffusion of the C signal. These results might indicate that the Si compounds produced at high temperatures (crystalline SiO2 or SiC compounds) are also present in the bulk of the materials and not only over their surface, confirming the TEM observations for local zones. Additionally, the intensity at higher Q values ( defined in Chapter III) is higher for wheat straw derived hard carbons than for the other derived hard carbons, which indicated that there was a higher surface area contribution for the former hard carbons, in agreement with the N2 SSA obtained values.

Figure 39. SAXS patterns for a) pine, b) beechwood, c) miscanthus, and d) wheat straw. e) Radius of gyration calculated for pine and beechwood hard carbons.

However, all the patterns exhibited an increase in the scattered intensity with temperature, which came from the higher electronic density difference between the continuous phase of C (or the Si compounds) and the ultramicroporous phases, therefore indicating that ultramicropores would be larger or more abundant at higher temperatures. As this trend is opposite to the CO2 SSA and pore volumes, this suggests that the ultramicroporosity increase would come from inaccessible pores and not from the accessible microporous volume. This result would be consistent with the internal structural reorganization of the graphene layers, in 160 which the evolving spaces are of the magnitude of some angstroms, therefore not accessible to probing molecules, but detected by X-ray scattering.

SAXS fitting was done over the patterns of the woody derived hard carbons and the resulting Radius of gyration (Rg defined in Chapter III) were reported in Figure 39e. The obtained Rg values did not vary significantly in this range of temperatures. They remained around 0.5 and 0.7 nm, which was close to the PSD from the CO2 pore size distributions. This behaviour was consistent with the observations previously reported over hard carbons derived from cellulose, were Rg was rather constant between 1000 and 1400°C [START_REF] Simone | Hard carbon derived from cellulose as anode for sodium ion batteries: Dependence of electrochemical properties on structure[END_REF].

All the fitted patterns exhibited a Q behaviour between Q -3 and Q -3,5 , indicating that the interfaces of the particles were not well-defined. This is in agreement with the heterogeneous morphology of the obtained hard carbon particles at all the temperatures, as seen from the SEM images in section 4.1.

Analysis of the hard carbons electrochemical performance

First cycle analysis

The galvanostatic curves obtained for the series of hard carbons are presented in Figure 40, while the quantified electrochemical parameters are presented in Figure 41. The electrochemical parameters reported are the mean of the values for the two coin-cells that were cycled for each material. All the hard carbons at all the temperatures exhibited the characteristic reduction-oxidation profile of a hard carbon material, as was presented in Chapter II. This profile consisted of a constant sloping variation of the potential between 1,5 V and 0,15 V. Under 0,15V, a pseudo-plateau behaviour was responsible for most of the gained capacity. The capacity lost between the reduction and the oxidation steps, accounted for the irreversibility of the system.

Notice from Figure 40, that during the reduction step, a shoulder on the reduction profile was developed at potentials around 0.7 to 0.3 V. It was more evident for the HC_wheat-straw_1400. This shoulder evidenced the formation of the SEI layer.

The SEI originates from the degradation of the electrolyte during reduction at potentials near 0,7 to 0,3 V. At this potential window, the organic and inorganic components of the electrolyte are not stable and undergo degradation reactions to form a solid passivation layer, the SEI, over the hard carbon surface that also trap Na ions.

Figure 40. First cycle reduction-oxidation profiles for a) pine, b) beechwood, c) miscanthus and d) wheat straw derived hard carbons.

However, the formation of the SEI during the first cycle prevents further degradation of the electrolyte in the upcoming cycles, despite of the initial Na ions lost in its formation 8 . The SEI layer normally forms all over the carbonaceous surface, therefore a higher accessible SSA (normally N2 SSA) will induce a higher SEI formation, yielding to higher reduction capacities and irreversible loss [START_REF] Bommier | Predicting capacity of hard carbon anodes in sodium-ion batteries using porosity measurements[END_REF]. 8 A deeper analysis of the SEI formation mechanisms, properties and importance, as well as its development on the different hard carbons derived from biomass, is discussed in Chapter VI. From the profiles on Figure 40, it was possible to see that the temperature-induced a decrease of the irreversibility for the hard carbons obtained from pine, beechwood and miscanthus, while it seemed to increase the irreversibility of the wheat straw hard carbons. In the same way, the average oxidation potential seemed to decrease with temperature for the former samples, as can be evidence by the flattening of the red curves for the hard carbons obtained at 1400°.

In the above section of this chapter, we saw that the obtained hard carbons presented a variety of properties, which evolved according to their biomass precursor and with the thermal treatment. The chemical composition, the morphology, the surface area, the porosity volume, and the structure are the main properties that differentiate the hard carbon samples, and are assumed to be reflected in the delivered electrochemical performances of the different electrodes. As expected, the different hard carbons did not show the same trends versus temperature/precursor regarding their final reduction and oxidation capacity, average oxidation potential, and Initial Coulombic Efficiency (ICE), as it is clearly seen from Figure 41. An individual analysis of each of the former parameters is presented in the following subsections.

Reduction capacity

The reduction capacity is related to the total Na uptake capacity of the hard carbon electrode in all its active sites, i.e. turbostratic domains, ultramicropores, micropores, mesopores, defects, active surface area, etc. As mention before, the formation of the SEI layer over the hard carbon surface during the reduction of the electrode, also accounts for the quantified reduction capacity, as Na ions are consumed in its formation. Therefore, the higher the surface area, the more extended the SEI formation, and consequently, the reduction capacity is expected to be higher too.

The reduction capacity of pine, beechwood and miscanthus derived hard carbons presented small variations with temperature. While the reduction capacity for wheat straw hard carbons increased with the thermal treatment. The influence of temperature treatment was not clear in these results. Indeed, it is hard to assign a single physicochemical parameter from the measured properties to these reduction capacity variations, as the derived hard carbons present all very different behaviours.

Only the increase of the N2 SSA of wheat straw hard carbons could be related to the increase of the reduction capacity with temperature for these samples. As stated at the beginning of this section, higher SSA will produce more SEI and therefore account for more reduction capacity. In this same logic, the rather constant behaviour of the N2 SSA for the miscanthus derived hard carbons could explain their also constant reduction capacity values with temperature. Notice also, that the CO2 SSA did not seem to influence the delivered reduction capacity.

This could be related to the fact that CO2 SSA measures smaller, less accessible pores, in which the SEI probably did not form.

164

Regarding the influence of the different biomass precursors, different reduction capacity values were observed for the different hard carbons at the same temperatures, as can be seen from Figure 41. Overall, at each temperature, the values of the reduction capacity for the derived hard carbons were observed to decrease as follows: pine > beechwood > miscanthus > wheat straw. For example, at 1400°C, HC_pine_1400 presented 376 mAh.g -1 of reduction capacity.

HC_beechwood_1400 showed 353 mAh.g -1 . The reduction capacity of HC_miscanthus_1400 was around 340 mAh.g -1 . While HC_wheat straw_1400 delivered a reduction capacity of 334 mAh.g -1 .

It was, however, unexpected the rather low reduction capacity of HC_pine_1000, given its significant N2 SSA of 141 m 2 .g -1 , when compared to the SSA of the other hard carbons. This observation might be explained by the low value of d002 interlayer space for HC_pine_1000, which was the lowest of all samples at all the temperatures, accounting for 0.38 nm. Therefore, it was possible that the high SSA of pine hard carbon was counteracted by its low d002 interlayer space, therefore delivering a reduction capacity close to that of the beechwood sample.

The lower reduction capacity for miscanthus hard carbons could be attributed to its lower C purity. Finally, wheat straw hard carbons exhibited the lowest reduction capacity of all the series of hard carbons, especially at 1000 and 1200°C.

As for miscanthus hard carbons, the lower reduction capacity was most probably related to less active sites towards Na uptake because of the of its high impurities content, especially SiC content. Its increase in the reduction capacity can be related to the increase of its N2 SSA, as it is the only sample to have a considerable higher N2-measured porosity volume at higher temperatures, as explained before. The role of the SiC compounds in the Na ions uptake is not known; it is, however, possible that the SEI was also formed over the SiC whiskers, which would also contribute to an increase of the reduction capacity at 1400°C. However, this hypothesis needs to be further confirmed. Indeed, SiC compounds are insulator materials that might be inert towards the electrochemical mechanism of Na uptake under the conditions of the study.

Moreover, graphitic domains are known to have a very limited capacity towards Na ion intercalation [START_REF] Matei Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF]. It is worth to note that at higher temperatures, SiC and graphitic phases were more important for miscanthus and wheat straw hard carbons, so one would expect the reduction capacity to decrease, but instead, it increases for the wheat straw hard carbons. This would suggest that the effect of the N2 SSA has a stronger impact on the reduction capacity than the other detrimental characteristics, notably the graphitic domains, for wheat straw materials.

Oxidation or reversible capacity

The reversible capacity quantifies the amount of Na ions that are able to leave the hard carbon structure during the oxidation of the electrode after being stored in it.

The quantified capacity that is lost between the reduction and oxidation steps accounts for the irreversibility of the material, which originates mainly from the Na ions lost in the SEI formation or trapped in the structure.

The reversible capacity was then lower than the reduction capacity for all the hard carbons and followed the same trends for the pine, beechwood and miscanthus hard carbons. Woody hard carbons presented the highest values (around 290 -315 mAh.g -1 ), which slightly fluctuated with temperature. Miscanthus hard carbons showed the lowest values near 250 mAh.g -1 , which did not vary significantly with temperature either. Nevertheless, wheat straw hard carbons also showed a constant reversible capacity with temperature, in contrast to its augmentation of the reduction capacity at 1400°C, indicating that more irreversible capacity is generated in this material at 1400°C. Na ions may be trapped irreversibly inside the hard carbon structure (turbostratic or graphitic-like domains), or by defects or active sites, as mentioned in Chapter II. Additionally, the higher N2 SSA area of the wheat straw hard carbon at 1400°C could explain its lower reversible capacity because of a higher SEI formation over its surface.

However, pine hard carbon obtained at 1000°C showed N2 SSA more than three times bigger than the one of wheat straw hard carbon at 1400°C (141 vs 42 m 2 .g - wheat straw hard carbon it is again the combination of its graphitic like-domains, its lower C content, and the presence of the SiC phase, that contributed to their higher irreversibility at 1400°C.

Even if graphitic domains do not insert/intercalate significant amounts of Na ions, the stored Na ions in these domains might irreversibly be trapped into them. In the same way, the interaction between Na ions and the SiC whiskers is thought to be irreversible. Therefore, the latter phenomena might contribute to reducing the reversible capacity of the electrode and explaining why the reversible capacity did not evolve with temperature for wheat straw hard carbons.

The discussion above was confirmed by the values obtained for the ICE of the hard carbons, where wheat straw hard carbons are the only samples that presented a diminution of the ICE with temperature, as seen from Figure 41.

Slope and plateau capacity

Figure 42. Slope and plateau Capacities of the oxidation of the first cycle. The white capacity accounts for the irreversible capacity lost on the first cycle. Error bars correspond to the variance of the mean values for the two coin-cells, for each material.

As mention at the beginning of this section, a hard carbon galvanostatic curve is characterised by two distinctive zones. From 1,5 to 0,15V (approximately), the curve exhibited a linear decrease or increase of the capacity as a function of the potential. The capacity given by this zone is called, slope capacity. At voltages below 0,15V, the capacity fluctuation is limited, presenting a pseudo-plateau behaviour. The capacity given by this zone is often called plateau capacity. These two distinctive behaviours indicate that different mechanisms are involved in the Na ions uptake in the hard carbon electrode. However, as discussed in Chapter II, there is no general agreement in literature about the correct assignment of electrochemical or physical phenomena to these zones during cycling.

Herein, the slope and plateau capacity of the oxidation step, are the only electrochemical parameters that varied significantly with temperature, for almost all the obtained hard carbons. The slope capacity decreased with temperature and consequently, the plateau capacity increased, as reported in Figure 42. This was reflected in the diminution of the average oxidation potential for all the hard carbons except for wheat straw derived ones (see Figure 41d), in which the slope and plateau capacity variation is not as significant as for the other hard carbons from 1200 to 1400°C. The decrease in the average oxidation potential caused the flattening of the galvanostatic profiles with temperature for all the hard carbons, except for wheat straw hard carbon at 1400°C (see Figure 40).

If the plateau capacity is related to the Na storage within the turbostratic domains, the bigger these domains, the higher this capacity. Herein, the diminution of the CO2 SSA and volume with temperature suggests a densification of the turbostratic domains. As the graphene layers arranged towards bigger turbostratic "crystals", less accessible porosity is left in the hard carbon structure. This leads to potentially wider inaccessible pores, as the graphene layers come closer for some hard carbons under this range of temperature. This latter fact is consistent with the increase of the scattered intensity in the SAXS patterns seen for all the hard carbons with temperature. Even if miscanthus and wheat straw hard carbons presented actual graphitic zones, i.e. long-range stacked graphene layers, they also exhibit turbostratic domains as demonstrated by the TEM observations (see section 3.2), and therefore, the Na uptake from this disordered domains seemed to follow the same trends as for woody hard carbons.

The CO2 SSA, CO2 volume, and the defects concentration follow the same trend as the slope capacity for all the samples: they decreased with temperature. The diminution of the slope capacity could be then related to the decrease of the defects concentration in all the hard carbons, as shown from the Raman results and previously observed by Bommier et al. [START_REF] Bommier | New Mechanistic Insights on Na-Ion Storage in Nongraphitizable Carbon[END_REF]. Bommier and co-workers [START_REF] Bommier | New Mechanistic Insights on Na-Ion Storage in Nongraphitizable Carbon[END_REF] correlated the diminution of the slope capacity with temperature to the diminution of microstructural defects with temperature, observed by the diminution of the AD1/AG ratio. They attributed the slope capacity to the interaction of Na ions with a wide morphological variety of defects in the hard carbon structure, such as dangling and sp 3 bonded C at the edges of the pseudo-graphitic domains, mono and di-vacancies, stone wales defects and extreme graphene curvature. Herein, the results clearly followed the same trend for the hard carbons derived from woody biomass precursors, which presented the purest carbon structures. The miscanthus and wheat straw hard carbons also exhibited a diminution of the slope capacity with temperature, however, it is difficult to attribute a single structural parameter to the mechanisms of Na ions storage in these hard carbons, because of the parallel contribution that might have their graphitic domains, their porosity and their SiC whiskers.

Cycling Stability

Finally, all the samples exhibited a stable behaviour up to 100 cycles at C/10 rate, as can be seen from Figure 43.

Among all the materials, the hard carbons obtained at 1400°C presented the best retention capacity over cycling. Interestingly, miscanthus and wheat straw derived hard carbons exhibited very stable cycling in spite of their low capacity of Na storage. The coulombic efficiencies for all the samples also stabilised after the initial cycles, to reach values up to 100% during cycling. Notice from Figure 43, that the cycling stability curves for HC_pine_1000 and HC_beechwood_1000 exhibited some atypical reversible capacity values during cycling. For HC_pine_1000, the unusual points were observed at the first cycles, and could be related to a limited electrolyte wettability at the beginning of cycling that improves with time. In the case of HC_beechwood_1000, the aberrant points were observed after the first 20 cycles, they could be more likely associated with other phenomenon, such as preferential dendrite formation at the edges of the Na electrode, the presence of gas bubbles producing sudden disruptions during cycling, or the misalignment of the half-cells constituents. However, more insights regarding these aspects are needed. Notice also that for the HC_beechwood_1000 and HC_beechwood_1400, the CE exhibited values over 100% after the first cycle.

This behaviour was clearly not related to physical phenomena, and was associated with an inaccurate measurement from the Arbin hardware.

Conclusions

In this chapter, we demonstrated that under the selected range of temperatures, the chemical composition of biomass precursors played a major role in the final electrochemical performance of the obtained hard carbons. Most of the final properties of the hard carbons were more affected by the phenomena induced by their composition, especially their inorganic fraction, than by the thermal treatment between 1000 and 1400°C.

The evolution of the hard carbon properties with the thermal treatment for each hard carbon is synthesised in Table 11.

CHAPTE R V. ELUCIDATION OF THE

Classification of the biomass precursors

In this work, 25 biomass precursors were used, all presented in section 1 of Chapter III. These precursors were chosen according to their availability for future industrial implementation and diversity in terms of chemical composition.

To simplify the data interpretation, herein we classified the different precursors into three different categories, according to their botanical origin, organic and inorganic composition and if they are a residue or a dedicated crop. The final groups are described in Table 12. A code has been assigned to each biomass precursor in order to simplify notation.

Elemental composition of the biomass precursors and ash content 2.1 C, H, N, O composition

The C, H, N, and O content of all the precursors, is presented in Table 13. Most of the precursors presented similar C, H, N and O contents. C content for all the precursors was between 42 to 55 wmf%. The H content ranged between 5 and 7 wmf% and the N content between 0.1 and 5 wmf%, for all the precursors. The calculated O content ranged between 37 and 45 wmf% for all the biomass precursors.

Papaya seed presented the highest C, H and N content of all the studied samples with 55, 7, and 5 wmf%, respectively. In contrast, it showed the lowest O content of 26 wmf%. No literature was found on the elemental analysis of papaya seeds for comparison. The properties of the hard carbon derived from the papaya seed are expected to be different from the other precursors, especially regarding its structure, which might be less cross-linked given the lower O content and higher H content.

The N content is also high and around 3 wmf% for GIII-19: grape seed cake and GIII-25: tamarind seed. Literature reports N contents for grape seeds mixtures of around 2 wmf% for different samples harvested in Spain [START_REF] Ischia | Realization of a solar hydrothermal carbonization reactor: A zero-energy technology for waste biomass valorization[END_REF] and Italy [START_REF] Diaz | Low-Cost Activated Grape Seed-Derived Hydrochar through Hydrothermal Carbonization and Chemical Activation for Sulfamethoxazole Adsorption[END_REF], thus confirming the high N value of this biomass. There are no available elemental composition reports for tamarind seeds. abundant element in GII was K, while Ca content was lower than for GI, varying only from 801 to 4081 mg.Kg -1 . Some species such as grasses and cereals, that are common herbaceous or agricultural residues used in biomass valorisation, are known to be Si accumulators [START_REF] Neumann | Silicon in Plants[END_REF]. These plants can absorb silicic acid from the soil and to transform it into amorphous hydrated opal (SiO2.xH2O) by polymerisation. These silica bodies are referred as phytoliths, and are known to accomplish structural functions on the cellular arrangements [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF], [START_REF] Strömberg | Functions of phytoliths in vascular plants: an evolutionary perspective[END_REF].

For the GIII group of biomass precursors, K was the most abundant element, found in concentrations from 2695 to 13999 mg.Kg -1 , except for the case of GIII-19: grape seed cake, which presented a higher Ca content (6740 mg.Kg -1 of Ca vs. 4140 mg.Kg -1 of K). No general trend was found for the second most abundant element in GIII. Elements such as P, Mg, S and Na were found in concentrations higher than 2000 mg.Kg -1 for some precursors. The heterogeneity of the biomass precursors in GIII group is responsible for this wide inorganic composition distribution. It is worth noticing that GIII-23: papaya seed, presented the largest values of P (4877 mg.Kg -1 ), Mg (4255 mg.Kg -1 ) and S (3650 mg.Kg -1 ) of all the precursors, which could be related to the high protein and nutritional value of papaya fruit. Its high mineral content, together with its high C, H, and N content, explain their low organic O content. Whereas GIII-24: coconut husk presented the highest value of Na (5910 mg.Kg -1 ) among all the precursors.

Ca-rich precursors of GI group might develop more organized structures than the lower Ca content precursors, as seen in the previous chapter for miscanthus and wheat straw Ca-rich precursors when compared to pine and beechwood. GII Si-rich precursors are expected to form SiC phases at 1400 °C and probably higher N2 surface areas than GI precursors, because of the contribution to the SSA given by the SiC whiskers. The simultaneous effect of an intrinsic activation in GII derived hard carbons could also take place given their K concentration. GIII K-rich precursors might also exhibit high N2 and CO2 surface areas and a different structure than the other precursors, as K, besides promoting intrinsic activation, could intercalate between the graphene layers and affect their final spacing [START_REF] Shen | Catalytic pyrolysis of biomass with potassium compounds for Co-production of high-quality biofuels and porous carbons[END_REF].

All these hypotheses will be disputed with the further observations. which is also characteristic of some grasses and cereal species [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF], [START_REF] Da | Characterisation of the Most Representative Agricultural and Forestry Biomasses in France for Gasification[END_REF].

The cellulose content in the precursors was highly variable and presented a general variation from 8 to 47 wmf% between all the precursors, in agreement with the glucose variations discussed before. GI-3: SRC-willow and GI-4: SRF-poplar showed the highest values of cellulose in forestry biomasses with 44 and 47 wmf%, respectively. While GII-10: miscanthus and GII-16: bagasse showed values of 45 and 44 wmf% of cellulose, respectively. The lowest cellulose content was found for GIII-19: grape seed cake, containing only 8 wmf%. These findings are in agreement with the sugar composition discussed in the previous section, and support the particular composition of some biomass subgroups also reported by Vassilev et al. [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF].

It should be remembered that the hemicelluloses content calculation is based on empirical ratios, therefore, it would be more accurate to discuss the sugar composition as presented before, than the final hemicelluloses content. It is, however, possible to say that the GII-16: bagasse and GII-17: corn husk derived hard carbons, had a total holocelluloses (cellulose + hemicelluloses) content of 80 and 77 wmf%, respectively, which was the highest of all the biomasses precursors.

This high amount of polysaccharides may affect the biochar and hard carbon yield, as a higher yield is expected from higher lignin contents than holocellulose contents. Moreover, the hard carbon structure and porosity derived from the latter precursors could also be different from the one of the precursors containing high lignin content, such as GI-9 and GIII-19.

Extractives content for all the biomasses fluctuated as well. In forestry biomasses, extractives compounds are related to a high content of bark and foliage [START_REF] Da | Characterisation of the Most Representative Agricultural and Forestry Biomasses in France for Gasification[END_REF],

which explained the higher extractives content for GI-1: orchard apple, GI- In GIII group, GIII-24: coconut husk presented the highest extractives content of all the samples, of 28 wmf%. As for cereals, this high amount of extractives is typical of agricultural residues and might contain important quantities of nonstructural sugars or starch [START_REF] Da | Characterisation of the Most Representative Agricultural and Forestry Biomasses in France for Gasification[END_REF]. As highlighted in Chapter III, a higher extractives content might result in a lower mass yield, as most of the extractives components are volatile at low temperatures [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF].

For GIII-23: papaya seed and GIII-25 tamarind seed, the quantification of macromolecules in these samples was not successful. The morphology of the raw biomass, which was hard grounded powder of around 500 μm size, prevented the complete solubilisation of the particles for the hydrolysis treatment, as well as a proper interaction with the extraction solvents. Therefore, the obtained results

were not reliable and it was decided not to take them into account.

Some data regarding the macromolecular composition of the latter biomass precursors was found on literature. For papaya seed, data was available on their nutritional components, and reported by Santos et al. [START_REF] Santos | Chemical characterization of the flour of peel and seed from two papaya cultivars[END_REF]. Holocelluloses content can be estimated from these values and ranged between 20.7 and 22.9 wmf%.

Therefore, cellulose and hemicelluloses content of papaya seed seemed to be low when compared to the other precursors. Fibre content can be related to the lignin content [START_REF] Tang | Chapter Eighteen -Biological mechanisms of glycan-and glycosaminoglycan-based nutraceuticals[END_REF], and ranged from 7.8 to 8.8 wmf%. This data suggests that lignin content is very low on papaya seed, in contrast to the high lignin content of the grape seed cake, reported before. Papaya seeds are mainly composed of lipids (29.7

-28 wmf%) and proteins (28.6 -28.1 wmf%), which explain their low holocellulose and lignin contents. The high proteins content might explain the high N concentration of this sample.

According to the available literature, hemicelluloses in tamarind seed are mainly composed of galactoxyloglucan, a polymer of glucose, xylose and galactose units, which is often referred as tamarind polysaccharides, and can reach concentrations up to 65% [START_REF] Chawananorasest | Extraction and Characterization of Tamarind (Tamarind indica L.) Seed Polysaccharides (TSP) from Three Difference Sources[END_REF]. Galactoxyloglucan in tamarind seeds differs from other hemicelluloses components given that it does not play and structural role, but is formed and stored in the seed [START_REF] Zavyalov | Plant Polysaccharide Xyloglucan and Enzymes That Hydrolyze It (Review)[END_REF]. Lignin content was low as deduced from the total fibre value of 8.8 % reported in [START_REF] Bhattacharya | Some physical and engineering properties of tamarind (Tamarindus indica) seed[END_REF].

Physicochemical properties of the obtained hard carbons

Biochars yield and composition

The C, H, N, S, O, and ash content together with the mass yields of the 450°Cbiochars are reported in Table 15. Ash quantity for the 450°C-biochars was calculated as in chapter IV, from a mass balance of the ash after pyrolysis. O content in the 450°C-biochar was calculated by difference.

As shown in Table 15, the yields for the 450°C-biochars resulted in being between 18 and 42 wmaf% for all the precursors. Values over 35 wmaf% were observed for GI-9: scots pine bark, GIII-19: grape seed cake and GIII-24: coconut husk. The high yield for the GI-9 and GIII-19 can related to their high lignin content of 40 and 63 wmf%, respectively. The high yield for GIII-24 biochar could also be observed in other studies [START_REF] Suman | Pyrolysis of coconut husk biomass: Analysis of its biochar properties[END_REF], and it does not seem to be only related to its lignin content as it was of 26 wmf%. Na is known to enhance charring reactions promoting biochar yield [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF]. As the Na content in this precursor is the highest of all the precursors (5910 mg.Kg -1 ) this could explain its higher biochar yield, despite its moderate lignin content. N in biochar could be retained because of the N retention from the biomass precursor, N adsorption from the pyrolysis atmosphere, or N uptake from the generated unstable gaseous products of the thermal decomposition of the biomass, such as cyanide (HCN), ammonia (NH3) and isocyanuric acid (HNCO). [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF], [START_REF] Lang | Retention of Organic Elements during Solid Fuel Pyrolysis with Emphasis on the Peculiar Behavior of Nitrogen[END_REF].

The H content for all the precursors varied between 3 and 5 wmf%. A volatilisation of at least 50% of the initial H content was observed for all the biochars. This volatilization strongly indicates that the macromolecules have been degraded towards a matrix mainly composed of aromatic compounds.

Notice from Table 15 that the quantities of O and ash are not reported for the GIII-23: papaya seed 450-biochar. The calculated O quantity for this sample was of (-)9 wmf%, indicating that the mass balance is over 100%. The overestimated mass balance might suggest that the ash fraction is not fully preserved for the papaya seeds biochar at 450°C. This can be explained by the high S content in papaya seeds. S in biomass is thought to be chemically bonded to the organic fraction of the biomass in the proteins and to be present as an inorganic mineral. It has been seen that during thermal treatment, S is released in two stages, where organic S is release at a temperature under 500°C and inorganic S is mainly released over 900°C, as reported by Saleh et al. [START_REF] Saleh | Release of Chlorine and Sulfur during Biomass Torrefaction and Pyrolysis[END_REF]. The total S reported in Table 13 for the papaya seed precursor corresponds to the organic and inorganic fractions and therefore, is also measured by the ash fraction. Therefore, considering the high amounts of proteins in papaya seed, it is possible to think that the ash fraction is reduced because of the high organic S volatilization.

Hard carbons yield and composition

The C, H, N, S, O content together with the mass yields of the hard carbons are reported in Table 15.

Given that the ash fraction is no more preserved at 1400°C, and that direct O measurements were not possible, a grouped residual quantity notated as O+ash, was calculated for the hard carbons by the equation ( 1) and reported in Table 15.

The total hard carbon yield ranged from 14 to 28 wmf% among the different precursors. The highest yields of over 25 wmf% were obtained again for GII-15:

rice husk, GIII-19: grape seed cake and GIII-24: coconut husk, and also for GIII-22: walnut shell. The high mass yield of rice husk is related to the retention of its elevated ash fraction. However, ash fraction for the BC_ricehusk_450 decreased from 34 wmf% to a total content of O+ash of 25 wmf% in the HC-rice-husk_1400.

The previous observation indicated the volatilization of some of the inorganics present in the ash, most probably the volatilisation of O-containing species [START_REF] Anca-Couce | Reaction mechanisms and multi-scale modelling of lignocellulosic biomass pyrolysis[END_REF].

The high yield of GIII-22: walnut shell could be related to its high lignin content of 44 wmf%. This contribution did not seem significant for the 450°C-biochar yield, probably because of its low inorganic content that did not catalyse charring reactions at low temperatures. It is also possible that the lignin composition of walnut shell precursor was different from that of grape seed cake or coconut husk, and thus led to the formation of a biochar being more thermally stable.

As for the 450°C biochar, HC_bagasse_1400 presented the lowest total yield of all the precursors, of 14 wmf%. Bagasse precursor exhibited the highest holocellulose content (80 wmf%), together with low lignin (20 wmf%) and ash (1 wmf%) values, which could be correlated to its low yield.

GI and GIII derived hard carbons presented C contents between 89 and 98 wmf%, while GII derived hard carbons presented lower C contents between 74 and 95 wmf%. The main non-C related species of the hard carbons were the ones grouped in the O+Ash fraction, except for HC_SRF-poplar_1400, where the N content was around 3 wmf%. N content was also around 3 wmf% for HC_switchgrass-1400.

As described before, some biomass species might concentrate N in the carbonaceous structure during low-temperature treatment such as 450°C, however, at this elevated temperature of 1400°C, most of the N is expected to volatilize in the gas products [START_REF] Liu | Conversion mechanism of fuel-N during pyrolysis of biomass wastes[END_REF]. Therefore, these atypical high values of N content should be treated with precaution. It has been observed that N from the air can be adsorbed in C-rich materials, like coal, phenolic resins or activated carbon fibres [START_REF] Hansson | Formation of HNCO, HCN, and NH3 from the pyrolysis of bark and nitrogen-containing model compounds[END_REF], and its contribution can increase the N concentration measured by the analyser. As the hard carbon samples are stored in an air atmosphere, it is possible that N could be adsorbed in the carbon matrix.

Nevertheless, this was only observed for two samples out of twenty-five, which probably suggests that another source of N is contributing in the measured values for these samples. More insights are needed to fully explain this behaviour, as there is no available reported data on N content of poplar and switchgrass hard carbons at 1400°C.

The final hard carbon N content was low (from 0.04 to 0.2 wmf%) and even close to the detection limit, for the samples that exhibited the highest N content in the precursor and in the 450°C-biochar: GIII-19, GIII-23 and GIII-25. As the atmosphere for the second pyrolysis treatment between 450 and 1400 °C was Ar, no additional N could interact with the hard carbon samples, and there must therefore be volatilisation of the N species at high temperatures for these samples.

It has been reported that doping the hard carbon material with heteroatoms such as N, P or B can enhance the electrochemical performance of the electrode, by increasing the electronic conductivity of the electrode and provides more active sites for Na storage [START_REF] Li | Preparation of nitrogen-and phosphorous co-doped carbon microspheres and their superior performance as anode in sodium-ion batteries[END_REF]. Therefore, the final N concentration of HC_SRF-poplar_1400 and HC_switchgrass-1400 could enhance their reduction and reversible capacity, and charging potential.

The content of H decreased strongly for all the hard carbons when compared to the precursors (5 to 7 wmf% ) and 450°C-biochars (3 to 5 wmf%), ranging from 0.2 to 2 wmf%. The significant H volatilisation for most of the hard carbon samples is an indication of a higher aromatization degree, as seen in the previous chapter, and therefore, it is related to a richer turbostratic structure composed of aromatic carbon layers and less aliphatic ramifications. GIII-19 and GIII-22 showed the highest H content with 2wmf%. Their H/C ratio was of 0.03, and was the highest of all the hard carbons. Based on its particular composition, the HC_grape-seed-cake_1400 sample might exhibit a particular electrochemical performance. Given the low C content on this hard carbon, together with its high H and S contents and its lower aromatization, its Na uptake is expected to be reduced when compared to higher C content hard carbons.

Structure of the obtained hard carbons

XRD observations

The XRD patterns for GI, GII and GIII derived hard carbons are shown in Figure 44, Figure 46, and Figure 47, respectively. 191

The d002 interlayer space of all the hard carbons was calculated as described in chapter III, by fitting the XRD pattern in TOPAS software. The results are reported in Figure 45.

Figure 45. a) d002 interlayer distance for all the obtained hard carbons. b) d002' interlayer distance of the graphitic-like domains present in some of the obtained hard carbons.

As discussed in chapter IV, the d002 values should be used with precaution, as they represent a mean value for the graphene interlayer distance in the turbostratic domains. Moreover, because of the inexistence of the three-dimensional graphite peak in the XRD pattern, there is no reference to confirm that the (002) peak shift is given by a structural difference between the samples and not by a difference in the sample height during the experiment. Even if all the XRD measurements were performed with extreme care and at the same conditions every time, without a reference peak it is not accurate to use the resulted d002 interlayer space as absolute values.

GI-derived hard carbons

All the hard carbons derived from GI precursors exhibited the (002) and ( 100) broad peaks near 2θ 23 and 41°, respectively, as seen from Figure 44. The XRD patterns obtained from this group of hard carbons confirmed that the woody precursors are most likely to develop the characteristic turbostratic structure of hard carbons, which is also seen in almost pure precursors such as phenolic resins [START_REF] Beda | Hard carbons derived from green phenolic resins for Na-ion batteries[END_REF], cellulose [START_REF] Simone | Hard carbon derived from cellulose as anode for sodium ion batteries: Dependence of electrochemical properties on structure[END_REF], or lignin [START_REF] Matei Ghimbeu | Valorizing low cost and renewable lignin as hard carbon for Naion batteries: Impact of lignin grade[END_REF]. Moreover, inorganic phases were identified in the hard carbon patterns, which are of low intensity but seemed to be present homogeneously in the irradiated sample.

For GI-2 to GI-5 and GI-9 derived hard carbons, a peak near 31-2θ° was present but any suitable compound was found to correspond for this diffraction. GI-2 and GI-9 derived hard carbons presented low-intensity SiC phases of cubic structure.

The presence of these phases can be explained by the relatively high content of Si in GI-2: pine-forest-residues and GI-9: scots-pine-bark precursors of 3936 and 1823 mg.Kg -1 , respectively. However, GI-6: pine precursor showed a Si content of 2074 mg.Kg -1 and it did not show SiC related compound peaks in the XRD pattern, but only in the local TEM observations (see Chapter IV). It is possible that part of the Si was present as contamination and therefore, there was no close interaction between SiO2 and C during the thermal treatment. On the contrary, all the SiC in GI-9 hard carbon could be related to the presence of chemically bonded Si to the carbon matrix of the scots pine bark precursor.

The GI-1: orchard apple and GI-4: SRF-poplar precursors showed the highest concentrations of Ca of all the studied precursors (together with GIII-23). As seen in the previous chapter (see section 3.1) the high Ca concentration seemed to promote the development of a graphitic-like phase in the derived hard carbons, evidenced by the embedded peak near 2θ-26°, denoted as 002' in the top left zoom of Figure 44. GI-2 and GI-9 derived hard carbons also developed a broader embedded peak near 2θ-26°. Interestingly, GI-5: vineyard and GI-8: ashwood hard carbons did not present graphitic-like phases, despite the fact that they had higher Ca content of these precursors when compared to GI-2 and GI-9 precursors.

Ōya et al. [START_REF] Ōya | Phenomena of catalytic graphitization[END_REF] presented that Ca vapour enhances the catalytic graphitization of non-graphitizable precursors at low temperatures (under 2000°C), producing long graphitic-like domains of Lc > 90 nm. Ca vapour is thought to break the crosslinking of the carbon matrix that prevents the graphitization. However, when Ca compounds, such as CaO, CaCO3, or Ca(OH)2 are used as catalytic agents, the catalytic activity can be different from the activity of the constituent metal. The Moreover, Ōya and co-workers also observed that some Ca compounds such as CaF2, do not exhibit any catalytic activity. Ca in biomass can be present in different chemically bonded compounds, which could be liquid or crystalline phases. Ca can also originate from external sources such as contamination, and therefore, not be chemically bonded to the C atoms in the precursor structure. This chemical and physical differences of the Ca compounds could explain why GI-5: vineyard and GI-8: ashwood hard carbons did not show any graphitic embedded component in the XRD pattern, while less Ca containing precursors seemed to present one (GI-2 and GI-9 derived hard carbons). Further fundamental research is needed to understand the formation of graphitic phases by the Ca catalytic activity from biomass precursors.

For GI derived hard carbons, d002 ranged between 0,364 to 0,392 nm, see Figure 45a. The softwoods pine and Scots pine bark derived hard carbons seemed to present the highest d002 distances, of 0.383 and 0.392 nm, respectively.

The d002' values of the graphitic-like phases in GI-1, GI-2, GI-4, and GI-9 derived hard carbons, are also reported in Figure 45b. Notice that GI-1 and GI-4 derived hard carbons presented the shortest d002' spaces of 0,339 and 0,342 nm, respectively. For the GI-2 and GI-9 derived hard carbons, the d002' values were of 0.346 and 0.347 nm, respectively. If the embedded peaks for GI-2 and GI-9 are related to a graphitic-like phase in these hard carbons, the larger interlayer space of these samples could evidence different catalytic graphitization mechanisms, which, at the same time, might be influenced by their softwood chemical composition (rich in 6-C sugars).

A wider interlayer space in the hard carbon turbostratic structure could enhance the Na uptake of the electrode during reduction. However, for GI-2 and GI-9 hard carbons, this effect could be masked by the presence of the graphitic-like domains, with shorter interlayer spaces, in which Na uptake is not favourable.

GII-derived hard carbons

The XRD patterns for GII derived hard carbons are shown in Figure 46. They exhibited the shape of the miscanthus and wheat straw derived hard carbons analysed in Chapter IV section 3.1. They all presented an asymmetric broad (002)

peak with an embedded peak near 2θ -26°, related to a graphitic-like phase in these materials.

In parallel, all the hard carbons derived from GII precursors developed SiC phases at 1400°C. According to the database, these phases matched with the diffraction of compounds with 3C structures: Si3CaC4 (PDF 01-82-3652), SiC (Moissanite, PDF 00-029-1129), and SiCN (PDF 01-074-2308). The intensity of the SiC related phases is stronger for GII-11: wheat-straw-1 to GII-15: rice-husk derived hard carbons, especially for the latter. All these precursors presented elevated Si contents (ranging from 1 to 3 wmf%), with rice husk presenting over 6 wmf% of Si content in the raw biomass. The mechanism behind the SiC formation in these biomasses was explained in Chapter IV, section 3.1. The SiC formation seemed to follow the whiskers morphology observed for miscanthus and wheat straw derived hard carbons in the previous chapter. A SEM-EDX analysis over the surface of the hard carbon derived from rice husk (see Figure 1 of Annex II), confirmed the whiskers formation and the higher Si concentration over them.

Notice from Figure 46, that the diffraction of the ( 200) and ( 222) peaks of the cubic SiC structure, presents a low, or almost inexistent intensity, suggesting a preferential growing of the SiC crystals. To confirm this, a slower XRD scan with higher resolution was performed for the GII-15: rice husk hard carbon, which is reported in Figure 2 of Annex II. The XRD pattern revealed the existence of the ( 200) and ( 222) peaks, however, their intensity was still much lower than the [START_REF] Adamson | Peat-derived hard carbon electrodes with superior capacity for sodium-ion batteries[END_REF] family peaks intensity. These results cannot confirm a preferential growing of the SiC crystals, but could suggest that the (111) family of peaks are dominant in the SiC structure developed by the rice husk precursor.

As mentioned before, all the hard carbons derived from the GII precursors exhibited a graphitic-like phase, evidenced by the embedded peak near 2θ 26°.

Herein, the 2θ -26° peak showed higher intensity than for the embedded peak present in the GI derived hard carbons. Previously, we discussed that graphiticlike domains in hard carbons are most likely related to the Ca content, as some Ca compounds (CaO, CaCO3, Ca(OH)2) are known to be graphitization catalysts [START_REF] Ōya | Phenomena of catalytic graphitization[END_REF].

However, the precursors in GII had, overall, lower Ca content than those of GI, and still the formation of a graphitic-like phase was seen for all of them. These results suggest that not only Ca is responsible for an enhanced graphitization of non-graphitizable precursors. Si is also known to catalyse graphitization [START_REF] Ōya | Phenomena of catalytic graphitization[END_REF]. Si in the GII precursors is part of the phytoliths plant constituents, which are cellular bodies composed of amorphous hydrated silica (SiO2.nH2O), called opal-A.

Few insights exist in the literature regarding the effect of SiO2 towards catalytic graphitization of non-graphitizable precursors. SiO2 seems to be used to promote the graphene formation trough an intermediate phase of SiC. According to Gubernat et al. [START_REF] Gubernat | Catalytic graphene formation in coal tar pitch-derived carbon structure in the presence of SiO2 nanoparticles[END_REF], stacked graphene layers can grow over a substrate of SiC nanoparticles at temperatures over 2000°C. They used a coal tar-pitch precursor as the carbon source, mixed with silica nanoparticles. SiC formation from biomass precursor has been observed to proceed at considerable lower temperatures than the needed temperature when coal and silica are mixed to form SiC, 1400°C vs.

2000°C, as discussed in Chapter IV. It is, therefore, possible that the graphitic-like domains formed by the interaction with SiC particles present in the hard carbon could start at lower temperatures too. Overall, the results seemed to indicate that the graphitic-like phase formation in the GII-derived hard carbons is more related to the Si content than to the Ca content. However, it is also possible that there is a synergy between the Si and Ca compounds in these biomasses towards a catalytic graphitization of a carbonaceous substrate. More evidence would be needed to sustain these hypotheses. The d002 interlayer space of the GII derived hard carbons varied from 0,360 to 0,386 nm (see Figure 45a). The d002' interlayer distance for the graphitic-like domains fluctuated around 0,345 nm, and was close to the one of the graphitic-like component found in GI-2 and GI-9 (see Figure 45b). This observation could suggest a similar graphitization mechanism for these precursors, as GI-2 and GI-9 hard carbons also contained SiC phases.

GII derived hard carbons are expected to show poor electrochemical performance, as seen for miscanthus and wheat straw derived hard carbons in the previous chapter. The graphitic-like domains in these materials might limit the Na uptake and increase the irreversibility, as some Na ions can be trapped in the structure during the oxidation of the electrode. The presence of SiC phases, especially for GII-15, are expected to increase the N2 SSA of the materials considerably, and therefore to contribute to the irreversibility of the electrode, as seen for the wheat straw derived hard carbon in chapter IV. The electronic and ionic conductivities of these materials are also expected to deteriorate. Higher oxidation potentials from the GII derived hard carbons are also expected.

GIII-derived hard carbons

The XRD patterns for the GIII hard carbons are shown in Figure 47. As for the GIderived materials, the typical hard carbon structure can be seen, where the (002) and ( 10) broad peaks were observed near 2θ-23° and 41°, respectively.

All the hard carbons in this group, except for the one derived from GIII-22: walnutshell, presented peaks for inorganic phases containing, Ca, S, Mg and N. Only GIII-24: coconut-husk hard carbon presented minor intensity peaks related to SiC compounds. Coconut husk biomass had similar Si concentration than pine species, thus promoting the same SiC formation in this material.

CaO (PDF 01-070-5490) was found in GIII-19: grape-seed-cake hard carbon, whereas CaS (PDF 04-001-3758) was found for GIII-20: sunflower-seed-shell hard carbon. GIII-23: papaya-seed hard carbon presented intense peaks of NaS (PDF 01-080-4869) as well as CaO. Finally, GIII-25: tamarind-seed hard carbon presented MgO (PDF 00-045-0946) peaks in the structure.

Notice that despite the high K content of the GIII precursors, no thermally stable K phase was detected in the XRD patterns for these hard carbons. This tends to confirm the volatilisation of K compounds before 1400°C.

The diversity of other inorganic elements than Si in the GIII precursors was evidenced with the presence of the inorganic phases described before, in the hard carbon structure. As these inorganic phases might be potentially embedded in the carbon matrix, they might deteriorate the Na uptake of the material, as they could reduce the active sites for Na storage. They could as well affect the conductivity of the electrode, and if present in the hard carbon surface, they could affect the composition and potentially the stability of the solid electrolyte interface formation.

Figure 47.

XRD patterns for the hard carbons derived from GIII. CaS PDF: 04-001-3758, CaO PDF: 01-070-5490, MgO PDF: 00-045-0946, and NaS PDF: 01-080-4869.

A higher graphitization degree was expected for the GIII-23: papaya seed derived hard carbon, because of the limited O content of the precursor (around 26 wmf%), and its high Ca content (9557mg.Kg -1 ). However, this hard carbon did not show any high graphitic-like embedded peak as for GII derived hard carbons, but a lowintensity peak similar to the one of GI-1 derived hard carbon. This latter fact supports the hypothesis of different pathways for the catalytic graphitization in non-graphitizable precursors, when Si and Ca compounds are present. The d002 distance for the derived GIII hard carbons seemed to be close, between 0,366 and 0,379 nm (see Figure 45a). No distinct structural property could be observed by XRD for the GIII-25: tamarind seed derived hard carbons, despite the particular hemicelluloses composition of the precursor, when compared to the other precursor in GIII group. However, it seemed to present the shorter d002 interlayer distance of the GIII-derived hard carbons.

Correlations

The d002 values of the obtained hard carbons were not found to be directly correlated to any of the compositional parameters of the biomass precursors, as seen from Figure 3 and Figure 4 in Annex II. It would only seem that the H content might affect the decrease of the interlayer space, as it is supposed to promote a soft carbon structure, rather than a hard carbon structure. However, as mentioned before, these results should be analysed with precaution, given the high uncertainty of the value.

Summary of the XRD observations

It was observed that the hard carbons derived from precursors with low ash content of up to 6.7 wmf%, developed very similar turbostratic structures at 1400°C, in terms of size and interlayer space. This was seen by the rather close d002 for the GI and GIII derived hard carbons. There was however, a significant difference regarding the mineral impurities found for these hard carbons. While GI derived hard carbons presented very few impurities, most of the GIII derived hard carbon presented mineral phases such as CaO, CaS, MgO or NaS. These inorganic compounds might decrease the available active sites for Na uptake or affect the surface chemistry of the electrode, therefore modifying the delivered performance of the materials. The GII derived hard carbons come from precursors with higher ash contents of up to 15 wmf%, which are mainly composed of Si. This high Si content triggered the formation of a graphitic-like phases and cubic SiC phases, therefore decreasing the number of turbostratic domains in these materials, which is expected to decrease the Na uptake of the final electrodes. 200

Raman results

The Raman spectra of some selected hard carbons, together with their AD1/AG ratio, are shown in Figure 48. The samples were selected according to the remaining quantity of material. The studied hard carbons presented AD1/AG ratios between 1,6 and 2,4. This range is slightly wider than the range of the same ratio for pine, beechwood, miscanthus and wheat straw derived hard carbons at 1400°C, presented in the previous chapter. These observations suggest that the hard carbons observed in Raman developed similar defective concentrations. Despite the slight variation of the AD1/AG ratio, it was found to be related to the ash and the K content in the biomass precursor, for all the studied samples, as can be seen from Figure 49, while less strong relations were seen for the C, O and Si contents in the precursors. The decrease of the AD1/AG ratio with the ash, and in Notice also from Figure 49, that the Ca content in GI and GIII precursors would seem to decrease the ratio of the defects of the derived hard carbon, which would be consistent with the graphitization induced in these materials by the Ca compounds, discussed before. Figure 5 of Annex II reports the AD1/AG ratio as a function of the other biomass composition parameters, and exhibits the lack of a simple relationship between them and the AD1/AG ratio.

Figure 49. AD1/AG ratio vs a) Ash, b) C, c) Si, d) Ca and e) K contents in biomass

precursors.

Surface area and porosity of the obtained hard carbons

General observations

The N2 adsorption-desorption isotherms for the obtained hard carbons are shown in Figure 50, while the CO2 adsorption isotherms are shown in Figure 51.

A D1 /A G Si (mg.Kg -1 * 10 4 ) 0,0

A D1 /A G K (mg.Kg -1 * 0,2 0,4 0,6 0,8 1,0 1,5 1,8 2,1 2,4
GI GII GIII 0,0

Pore size distributions of N2 and CO2 isotherms

The pore size distribution plots derived from the N2 and CO2 isotherms can be found in Figure 6 and Figure 7 of Annex II. For N2 isotherms, most of the GIderived hard carbons presented pore sizes centred in two zones, from 0.5 to 4nm and from 22 to 24 nm. The latter indicates the presence of mesoporous in the 204 samples, in agreement with the observations from the adsorption-desorption isotherms, discussed before. The distribution of pores centred in 24 nm was more important for the GI-5: vineyard hard carbon, confirming its mesoporous nature.

For GII-derived hard carbons, only GII-14: triticale hard carbon presented pore sizes centred in the previous two zones. Interestingly, the other samples did not show important differential volumes at sizes over 20 nm. GII-15: rice husk hard carbon exhibited an important contribution of mesoporous centred on 5 nm size.

These results might suggest that the GII derived hard carbons present smaller average mesoporous size and lower mesoporous fractions than the GI hard carbons.

For the GIII-derived hard carbons, the pore sizes were also centred in two zones, from 0.5 to 3nm and from 23 to 24 nm, for GIII-20, GIII-23 and GIII-24 derived hard carbons, as expected from their rather mesoporous shape isotherms. GIII-24 derived hard carbon also showed an important distribution around 4 nm.

The presence of different pore sizes in the samples could affect the electrochemical performance of the electrodes, and therefore, the distinctive distributions presented before are expected to contribute to the final electrochemical performance of the materials. In particular, the presence of mesopores has been related to an improvement of the electrolyte's percolation and therefore, to the Na ions diffusion [START_REF] Xiang | Activated hard carbon from orange peel for lithium/sodium ion battery anode with long cycle life[END_REF]. It could be therefore possible, that GI and GII derived hard carbons, that exhibited important mesoporous fractions, present enhanced electrochemical performances.

The CO2 isotherms results indicated that all the hard carbons had a pore size centred between 0.5 and 1 nm, in agreement with the results presented in the previous chapter.

N2 SSA

The N2 SSA and volumes of all the hard carbons are reported in Table 16. GI and GIII derived hard carbons presented low N2 SSA up to 14 m 2 .g -1 , close to the detection limit of the analyser (10 m 2 .g -1 ). The N2 SSA for the GII derived hard carbons was higher and between 12 and 97 m 2 .g -1 . As expected, higher N2 SSA and volumes were obtained for the biomass precursors with higher ash contents, in particular Si-rich ash, as seen from Figure 52. It is difficult to assign an absolute contribution to the high N2 SSA in these hard carbons. However, it is possible to think that the SiC whiskers could also contribute to the micro or mesoporous volume of these hard carbons, as they also exhibit their own porosity. Reported N2 SSA of SiC whiskers found in literature ranged from 45 to 100 m 2 .g -1 [START_REF] Hao | SiC nanomaterials with different morphologies for photocatalytic hydrogen production under visible light irradiation[END_REF], [START_REF] Yang | High Surface Area Silicon Carbide Whiskers and Nanotubes Nanocast Using Mesoporous Silica[END_REF]. In this way, the higher the whiskers concentration on the hard carbon samples, the higher the N2 SSA and volume.

Herein indeed, the highest N2 SSA and volume were obtained for the rice husk hard carbon, which presented SiC whiskers all around the hard carbon particles, as seen in the SEM images in Figure 2 of Annex II. Interestingly, when the Si is removed from the biomass precursor, the final N2 SSA of the hard carbon materials is low and close to 10 m2.g -1 , as reported by Wang et al., for hard carbons derived from washed rice husk precursors [START_REF] Wang | Rice husk-derived hard carbons as high-performance anode materials for sodium-ion batteries[END_REF].

The previous observations might suggest that inorganic elements in biomass precursors play a critical role when high surface areas are observed in materials obtained at 1400°C, by the effect of the SiC whiskers porosity.

The C content seemed to be correlated with the N2 SSA, as can be seen from Figure 52a and c. This effect seems to be stronger for the C content in the hard carbon, than for the C content in the biomass precursor. This observation is consistent with the fact that a higher C content in hard carbon implies that fewer C atoms were consumed to produce pores.

The GII derived hard carbons presenting N2 SSA over 20 m 2 .g -1 are expected to lose more capacity and to show a lower columbic efficiency during the first oxidationreduction cycle than the other hard carbons, as seen from the previous chapter for wheat straw derived hard carbons.

The higher amount of K in the GIII-23, GIII-24 and GIII-25 biomass precursors did not seem to be correlated with higher N2 SSA in the derived hard carbons as seen in Figure 8 of Annex II. However, GIII-23, GIII-24 hard carbons were the only samples that exhibited a rather mesoporous isotherm and higher total N2 volumes.

The N2 SSA of the GIII derived hard carbons did not seem to be affected either by the high content of Na, S, P or Mg in the biomass, as seen from Figure 53.

Some of the GII biomass precursors also presented the highest K content of the series of precursors. As seen in the previous chapter, K species in biomass can induce an intrinsic activation of the carbonaceous matrix. The effect of this activation was seen in the materials synthesised at lower temperatures of 1000 and 1200°C. Herein, Figure 52g showed that there seemed to be no correlation between the K content and the final N2 SSA. Therefore, even if an intrinsic activation has occurred at lower temperatures than 1400°C, its effect is not preserved until this final temperature, as the pores could have been subjected coalescence.

208

The N2 SSA was not found to be correlated to the macromolecular composition, or the H, N, and Ca contents in the biomass precursors, as shown in Figure 8 of Annex II.

CO2 SSA

The CO2 SSA and volumes are reported in Table 16. For most of the obtained hard carbons, the CO2 SSA was close to the N2 SSA, and variated from <1 to 85 m 2 .g -1 ,

as already observed from the samples obtained at 1400°C in the previous chapter.

Clear exceptions were the hard carbons obtained from GII-14: triticale and GII-16:

bagasse, with CO2 SSA values of 199 and 338, m 2 .g -1 , respectively.

Higher CO2 SSA and volumes than for N2 measurements in these materials could indicate that several micropores were not reached by the N2 probe molecule, or that they presented an inkbottle entrance, blocking the N2 molecule for quantification.

The particular high CO2 SSA values for these two materials might be related to the composition particularities of these biomass precursors. GII-14 triticale, presented one of the highest content of extractives of all the precursors (26 wmf%).

As explained before, extractives in agricultural biomass such as triticale might contain significant quantities of non-soluble sugars and starch, which are volatile during thermal treatment, enhancing the development of porosity when leaving the carbonaceous matrix [START_REF] Tripathi | Effect of process parameters on production of biochar from biomass waste through pyrolysis: A review[END_REF]. This would also explain the higher CO2 SSA and volume for GIII-24: coconut husk hard carbon, as this precursor exhibited a high extractives concentration of 28 wmf%. Moreover, Na content in coconut husk might also contribute to porosity development by the same mechanisms explained in chapter III for K compounds. Intrinsic activation with Na during hard carbon synthesis has been observed when preparing hard carbons from commercial lignin alkali at 1200°C, presenting CO2 SSA of 378 m 2 .g -1 [START_REF] Matei Ghimbeu | Valorizing low cost and renewable lignin as hard carbon for Naion batteries: Impact of lignin grade[END_REF]. However, no general tendency was seen between the CO2 SSA and the holocelluloses or extractives contents of the other biomass precursors, as seen from Figure 9 of Annex II. The other composition parameters of the biomass precursors were not found to be correlated to the CO2 SSA, as it is also shown in Figure 9 of Annex II.

GII-16: bagasse precursor had the highest holocellulose content of 80 wmf%.

Holocellulose is responsible for most of the volatile species formed during pyrolysis and therefore, they are associated with the development of microporosity in biochars [START_REF] Tripathi | Effect of process parameters on production of biochar from biomass waste through pyrolysis: A review[END_REF]. However, when the degradation of the holocelluloses is complete, and the temperature rises to 1400°C, the microporosity that was formed is expected to decrease, because of the coalescence of the pores. The high CO2 SSA of the bagasse hard carbon suggests that it underwent less pore coalescence than the other materials, which could be the result of its probable very high microporosity at lower temperatures. This hypothesis needs to be confirmed by analysing the CO2 SSA of bagasse biochars and hard carbons synthesised at different temperatures, under the same conditions of synthesis used in here.

The effect of high CO2 SSA over the electrochemical performance of the hard carbons is not well established. However, as shown in the previous chapter, it could indirectly suggest the development of a highly turbostratic matrix in the material, able to host more Na ions reversibly. For the coconut husk and the triticale derived hard carbons, this effect could be masked by the performance deterioration given by the presence of SiC phases.

Electrochemical properties of the obtained hard carbons

To analyse the electrochemical performance of the hard carbons, at least two halfcells were assembled per hard carbon in order to analyse the repeatability of the results. For the first five cycles, the half-cells underwent cycling at a C/10 rate (with C = 372 mA.g -1 ). In order to study the behaviour of the electrodes at higher rates, the half-cells were cycled at a C rate for the next ninety cycles. Hard carbons derived from GI-6: pine, GI-7: beechwood, GI-8: ashwood, GII-10: miscanthus and GII-11: wheat-straw-1, were not cycled at C regime, therefore, the values for the C parameters are not reported.

The reduction-oxidation profiles of one half-cell for each of the hard carbons derived from GI, GII and GIII groups are reported in Figure 54, Figure 55, and heteroatoms or defects, or clustered onto the ultramicropores [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF]. Therefore, it was expected to see higher reduction capacity from hard carbons presenting higher C or N contents and from higher SSAs and defects. However, no trend could be seen for the reduction capacity among the different biomass-derived hard carbons.

As seen from Figure 57, neither the C content in the hard carbon nor the Si content were correlated to the reduction capacity. The reduction capacity was also plotted as a function of the other composition parameters of the biomass precursor, with no simple relationship found, as seen from Figure 10 and Figure 11 of Annex II. Furthermore, SEI formation occurred during the first reduction of the hard carbons, contributing to the measured reduction capacity. The SEI layer is the result of the degradation of the electrolyte salt and its organic solvents at potentials between 1 and 0.5 V vs. Na + /Na. The degradation products create new species that form a solid layer over the carbonaceous surface, trapping Na ions in it. Therefore, the higher the N2 SSA, the more marked the SEI formation and the higher Na capacity used in the SEI formation. However, no relationship was found

GI GII GIII c)

between the obtained N2 SSA or CO2 SSA of the samples and their final reduction capacity, as can be seen from Figure 57d ande.

The low N2 SSA of the GI and the GIII derived hard carbons was not expected to contribute to the reduction capacity, and therefore, it is consistent that the result did not show any specific trend. However, for GII derived hard carbons, it was expected to see higher reduction capacities than for the other two hard carbon groups, which is not the case.

Regarding the CO2 SSA, it did not seem to affect the reduction capacity of the GII derived hard carbons. The CO2 SSA of hard carbons has been rarely reported before, and its relation to the reduction capacity is not well-established. It is not clear if the measured CO2 SSA is available for the electrolyte percolation.

Therefore, if there are pores in which the electrolyte cannot reach, this will not enhance the reduction or even the oxidation capacity of the electrodes.

It is worth noting that the GII-derived hard carbons presented significantly lower C contents when compared to the GI and GII derived hard carbons, as can be seen from Table 15, ranging from 74 to 95 wmf%. It is therefore quite inaccurate to express their specific capacity in terms of the active mass of the electrode, when this "active mass" actually presented up to 26 % of inactive components, such as SiC. The first cycle reduction-oxidation profiles for the GII-derived hard carbons were then corrected for the C content in the active mass, and the results are shown in Figure 58.

As expected from the correction, there was an increase of the reduction capacity of the GII derived hard carbons, which ranged between 369 and 485 mAh.g -1 . The reduction capacity was not significantly affected for the GI and GIII derived hard carbons as they had C contents of over 89%. The relationship between the first reduction capacity and the C content in hard carbon, the N2 SSA, the CO2 SSA and the Si content in the biomass precursor were re-evaluated and are shown in Figure 59. There was an obvious increase of the reduction capacity of the GII derived hard carbons as the N2 SSA increased. In parallel, the increase of the corrected reduction capacity seemed to be correlated to the higher Si content in the GII precursors, while a lower content of C in the GII hard carbons seemed to promote higher reduction capacities.

These results might suggest that there is a balancing effect between the lower C content in the GII hard carbons and their higher N2 SSA, that final produces higher reduction capacity. Moreover, as seen from section 3.4.3, the SiC whiskers might significantly contribute to the measured N2 SSA of the GII hard carbons. As there is no evidence that Na ions can interact with the SiC whiskers, either by adsorption or by insertion, it is most likely that the contribution of the SiC whiskers to the higher reduction capacities is related to the SEI formation over their surface. The SiC whiskers are thought to be at the same potential of the electrode during cycling, as they are in close contact with the active material in the composite electrode. Therefore, the electrolyte wetting the whiskers may also be decomposed to form a passivation layer over them, during the reduction.

Oxidation or reversible capacity

From here on, the analysed capacities are presented based on the total mass of active material and not for the specific mass of C in the hard carbon (as done for the correction for the reduction capacity for GII derived hard carbons). It is considered that this specific capacity is more accurate in terms of electrode engineering, as this dead mass affects the energy density of a full cell.

The oxidation or reversible capacity obtained for the different hard carbons varied from 225.9 to 314.7 mAh.g -1 for the GI and the GIII derived hard carbons, representing a difference of 28%. The lowest reversible capacities were delivered by the GII hard carbons, which ranged between 170.3 to 283.5 mAh.g -1 as reported in Table 1 in Annex II.

The reversible capacity seemed to be correlated to ash, C, H and Si content in the biomass precursor, as well as C content and N2 SSA in the hard carbons, as can be The ash and C contents in the biomass precursor are the only parameters that seemed to affect the reversible capacity for all the hard carbons. Additionally, the Si and H content in the biomass precursor, and N2 SSA of the hard carbons seemed to affect the reversible capacity of the hard carbons derived from GII group.

We have seen that a lower ash content in the biomass precursor could imply less side catalytic effects to take place during pyrolysis. The major inorganic elements in each group of precursors were analysed as a function of the reversible capacity delivered from the hard carbons (see Figure 60). It would seem that Si, Ca and K contents in the biomass precursors affect the reversible capacity. Hence, the ash content in biomass seemed to be the best indicator of the overall influence of the inorganic matter over the reversible capacity.

It is clear that there is a strong influence over the recovered reversible capacity given by the particular characteristics of the GII derived hard carbons. As expected, the effect seemed to be more direct for the reversible capacity than for the reduction capacity, as the reversible capacity quantifies the Na ions that are able to leave the hard carbon structure during oxidation, and not the total amount of Na ions used in reduction, that can be stored by several mechanisms.

As discussed in the previous section, a higher Si content in the GII precursors seemed to produce higher N2 SSA in the hard carbon, which in parallel, seems related to a lower reversible capacity. The lower reversible capacity from these materials evidenced the Na ions lost in the SEI formation over the extensive surface. Additionally, as previously explained, lower C content in the hard carbon is an indication of less turbostratic domains, when compared to purer hard carbons, which means that the amount of Na ions recovered from insertion in these domains will be lower. Additionally, GII hard carbons all presented graphitic domains, which are supposed to exhibit low Na uptake, therefore, decreasing even more the recovered Na capacity of these materials.

As for the reduction capacity, the effect of the CO2 SSA did not show any particular tendency for any of the hard carbon groups. More insights are needed to understand its contribution to the hard carbons synthesised in here. As discussed for the reduction capacity, it is possible that the CO2 SSA does not quantify all the surface that the electrolyte can reach, and therefore, there is no straightforward relationship with the capacity of the hard carbon.

Slope and plateau capacity of the first cycle

The first cycle reversible capacity was separated into the slope and plateau capacity contributions, and the results are shown in Figure 61.

All the hard carbons, except for the GII-15: rice husk derived, showed a higher plateau than slope capacity. As discussed in the previous chapter, our results seemed to indicate that the plateau capacity is related to the insertion of the Na ions into the turbostratic domains, while the slope capacity is related to the interaction with the surface, the defects and the porosity. It was seen in the previous chapter, that the plateau capacity increased with the diminution of the N2 and the CO2 SSA, while the slope capacity was higher for higher SSA values.

However, when considering the 25 hard carbon samples, the relationship is not as straightforward as presented before.

Plateau Capacity

The plateau capacity seemed to be only correlated to the ash and Si content of the GII precursors and the N2 SSA of the hard carbons derived from this group, as can be seen from Figure 62. For the other hard carbons, no correlations could be observed between these parameters and the plateau capacity, probably because of their narrower variation. The CO2 SSA did not seem to affect the plateau capacity of any group of hard carbons.

The reduced plateau capacity of the GII derived hard carbons is in agreement with the reduce turbostratic domains in these samples, given their low C content and their graphitic components.

Interestingly, the plateau capacity of the GII derived hard carbons would seem to increase with the increase of the AD1/AG ratio, as seen from Figure 63. For the GI and GIII derived hard carbons, the plateau capacity does not seem to vary with this defects ratio. The atypical observation for the GII derived hard carbons would suggest that the defects present in the GII derived hard carbons are involved in the Na storage at low potentials, while the defects concentration does not affect the Na storage for the GI and the GII derived hard carbon at low potentials.

Therefore, the plateau capacity of the GII derived hard carbons could not only originate from the Na storage in the turbostratic domains, but also from the Na storage in the defective sites.

Slope Capacity

For the slope capacity, no correlations were found neither with the SSAs, nor the Si content of the biomass precursors or hard carbons of any group, as seen from Figure 62. Not even the extremely elevated CO2 SSA of the GII-16: bagasse derived hard carbon seemed to affect the slope capacity. This is worth mentioning because bagasse hard carbon presented high C content and no SiC whiskers were seen on its surface with SEM observations (see Figure 13 in Annex II) that could decrease the Na ions uptake. This observation might indicate that, the CO2 SSA does not affect the slope capacity or, that most of the measured CO2 SSA, at least in this sample, is not accessible by the electrolyte, and therefore, no interaction is produced with the Na ions that stores capacity.

Notice from Figure 62, that the ash content in biomass seemed to decrease the slope capacity of the GI and the GIII derived hard carbons. However, the Si content on these biomass precursors did not seem to affect the slope capacity of the derived hard carbon materials. This suggests that other inorganic compounds than Si in the precursors, could affect the Na storage at high potentials. Notice indeed, from Figure 63, that Ca and K concentration seemed to decrease the slope capacity, especially for GI and GIII derived hard carbons. K and Ca compounds were related before to a lower defects concentration in the precursors. Here, notice from Figure 63, that the increase of the AD1/AG ratio, also seemed to improve the slope capacity of the GI and GIII derived hard carbons. These observations strongly support the Na ions storage in defective sites at high potentials, for the GI and GIII derived hard carbons.

For the GII derived hard carbons, the slope capacity was observed to remain constant with the variation of the AD1/AG ratio. This behaviour might suggest that the defects concentration did not seem to affect the slope capacity. As the plateau capacity was observed to increase with the increase of the AD1/AG ratio for these precursors, there would seem that for the GII derived hard carbons, the role of the defects Na uptake is not favoured at the same potentials, as for the GI and GIII derived hard carbons.

Average oxidation potential

The average oxidation potential was analysed from the 5 th cycle, as it is considered stable and representative of the C/10 cycling rate. The obtained values ranged between 0,20 and 0,38 V for all the hard carbons. The highest values of average oxidation potentials, between 0,27 and 0,38 V, were obtained from the hard carbons derived from GII. As for the other electrochemical parameters, the average oxidation potential seemed to be correlated to the ash and Si content in biomass and C content in hard carbon for the GII derived hard carbons, but no clear tendency was seen for the other groups (see Figure 64). This behaviour could be explained by the lower C content of the GII precursors in comparison to the other hard carbons. As mentioned before, the higher the ash content, the lower the resulting C content in the hard carbons, as some elements in the ash, in particular Si, are non-volatile. This is also why there is a correlation between the higher Si content and the higher average oxidation potential.

The operational oxidation potential of a negative electrode is very important as it will determine the final energy density of a full cell [START_REF] Eftekhari | Low voltage anode materials for lithium-ion batteries[END_REF]. This parameter tends to be shadowed by the oxidation capacity of the hard carbons, and it is rarely reported. Table 17 regroups some average oxidation values for hard carbons derived from low ash content precursors at 1400°C, and from a commercial hard carbon of Kureha Corporation. These reference hard carbons present values between 0,28 and 0,31 V, which are even higher than most of the GI and GIII derived hard carbons (0,20 to 0,27 V).

GI GII GIII

half-cell could also account to lower capacities. Therefore, the results obtained here at a C rate should be analysed carefully, keeping this consideration in mind.

As for the C/10 rate cycling, the CE stabilised after some cycles reaching values of over 99% at the tenth cycle for all the hard carbons (see Table 2 of Annex II).

As a result of the higher polarisation during C rate cycling, the average oxidation potential also increased for all hard carbons, taking values from 0.26 to 0.51 V, as it is reported in Table 2 of Annex II. The highest average oxidation potentials were again observed for the GII derived hard carbons, which can be attributed to the high fraction of non-C components in theses electrodes, as discussed before. The mechanisms involved in the Na ions storage at a higher rate are supposed to be the same as at lower C rate. However, these mechanisms are limited by several phenomena related to the physicochemical properties of an electrode active material, in this case, the hard carbon, and the electrolyte, such as i) the electronic transport in the hard carbon as well as ii) the ionic transport in the electrolyte, through the hard carbon pores embedded with electrolyte and trough interfaces such as the SEI, and iii) the solid-state diffusion of the ions through the hard carbon and the electrochemical reactions at the hard carbon/electrolyte interface. [START_REF] Tian | Quantifying the factors limiting rate performance in battery electrodes[END_REF]. The separators in the half-cell also play a role in the rate performance of the electrode [START_REF] Tian | Quantifying the factors limiting rate performance in battery electrodes[END_REF].

Therefore, despite the lower reversible capacity at C cycling, the sixth oxidation capacity followed the same trends as for the first oxidation capacity. It was found to be correlated to the ash and the Si content in biomass precursor and the C content in hard carbon, as seen in Figure 65. However, the fraction of reversible capacity that each hard carbon is able to retain when cycling at higher rates is another parameter which relies on the physicochemical properties of each hard carbon, the composite electrode properties and the nature of the metallic Na electrode. This parameter is reported in Table 3 of Annex II, as 10 th Ox. Cap / 5 th Ox. Cap. The 10 th Ox. Cap / 5 th Ox. Cap parameter was plotted versus the ash and de Si content in biomass, the C content in the hard

GI GII GIII

carbon, the N2 SSA, the CO2 SSA and the d002 interlayer space of the hard carbons and shown in Figure 66, but no simple correlations were found. The macromolecular and complete elemental composition of the biomass precursors was also analysed against the 10 th Ox. Cap / 5 th Ox. Cap factor (see Figure 14 of Anne II), but any parameter seemed to correlate with it. The cycling stability of the series of hard carbons was evaluated during 90 cycles for the C cycling, as reported in Figure 67. Notice that some samples in each group seemed to present a sudden and sharp decrease of the reversible capacity.

Moreover, notice that despite the lower reversible capacity exhibited by GII-15: rice husk derived hard carbon, it was very stable upon cycling at a C rate. In fact, this hard carbon presented a high retention capacity after 90 cycles at C, of 98 %, which represented a capacity fading of only 0.04 mAh.g -1 .cycle -1 . GII-12 and GII-13 derived hard carbon also presented low capacity fade of 0.04 and 0.08 mAh.g - 1 .cycle -1 , respectively. However, notice from Figure 68, that no correlation could be seen between the capacity fading and the composition of the biomass precursors or hard carbons and the surface areas (additional composition parameters are presented in Figure 15 of Annex II, where no clear correlations were found). The Capacity retention between the 10 th and the 95 th cycle was analysed as a function of the same biomass and hard carbon properties, in Figure 16 of Annex II, and as for the capacity fading, no simple correlations were observed.

The performance at high rates are linked with several phenomena, and it is rather complex to correlate it to one single hard carbon property [START_REF] Tian | Quantifying the factors limiting rate performance in battery electrodes[END_REF]. Herein, we are only considering the hard carbon properties contribution, without considering some macroscopic properties of the electrode, such as its wettability towards the electrolyte or its thickness, among several others. Moreover, the metallic Na electrode also affects the performance of the half-cell cyclability at higher rates. Not only because of its higher overpotential, but also, because of its ageing during cycling. All these factors might explain the lack of relationship between the biomass and hard carbon properties and the C cycling performance. Regarding the ageing of the metallic Na electrode, it is known that it can also form an SEI during cycling. This SEI formation over the metallic Na electrode could promote two phenomena. First, it can create a "dead Na" interface that consumes the electrolyte. With cycling, this generates the electrolyte drying leading the halfcell to die, in analogy to the phenomena observed in Li-metal half-cells [START_REF] Zheng | A high-performance sodiumion battery enhanced by macadamia shell derived hard carbon anode[END_REF].

Moreover, if the surface of the Na electrode is not very homogeneous, the formed SEI will also be unevenly distributed over the surface. This will generate a

GI GII GIII

preferential Na deposition during the platting of the Na electrode. There will be a preferential ion flux over the surface protuberances, which will enhance the growth of dendrites. If a dendrite grows enough to penetrate the separator and reach the hard carbon electrode, a short circuit will occur and the cell will undergo an exothermic failure. Dendrites formation are commonly associated with high current rates, however, they have also been identified at low to moderate current rates [START_REF] Lee | Sodium Metal Anodes: Emerging Solutions to Dendrite Growth[END_REF]. Figure 69a, presents a schematic explanation of Na dendrites growth and dead N formation during cycling. The Na induced half-cell dead could explain the abrupt reversible capacity drop observed during cycling for certain hard carbons such as GI-2, GI-9, GII-14, GIII-20, GIII-22 and GIII-25, presented in Figure 67.

To studied if a possible correlation between dendrite formation and half cell dead existed, the last cycle of all the assembled half cells was plotted as a function of the current density in terms of mA.cm -2 . Higher surface current density should increase the probability of dendrites formation or dead Na formation, and therefore induced a premature cell failure. Notice from Figure 69b that there is no correlation between the surface current density and the last cycle of the half-cells.

Therefore, higher surface current densities did not seem to affect the cells failure.

Other factors might be related to the dendrite formation, most probaly coming from the surface nature of the matellic Na counter electrode.

Conclusions

With an extended number of 25 hard carbons derived from different biomasses, it was possible to confirm that ash content in biomass, in particular ash rich in Si, Ca, and K compounds, play a critical role in the development of the hard carbon properties at 1400°C, and their performance at a C/10 cycling. Most of the correlations were only seen for the hard carbons derived from GII precursors, given their high inorganic content. Moreover, no simple relation was seen between the macromolecular composition and the final electrochemical parameters. It is possible that the high-temperature treatment masks the contribution of the macromolecules on the hard carbon formation.

The identified correlations between the selected hard carbon electrochemical parameters and the biomass/hard carbon compositional properties are summarised in Table 18. were held before start cycling. This time was determined by following the impedance increase as a function of time, for HC_walnut-shell_1400, as shown in Figure 2 of Annex III.

Eight representative samples from the woody and grass precursor groups presented in Chapter V were selected for the iterative cycling, as follows:

-Woody derived hard carbons (GI): HC_beechwood_1400, HC_ Orchard apple_1400, HC_pine, and HC_pine-forest-residues_1400.

-Grass derived hard carbons (GII): HC_rice husk_1400, HC_ wheat straw -1_1400, HC_miscanthus_1400, and HC_ bagasse_1400.

Only HC_walnut-shell_1400 sample was selected from the residues group (GIII),

given its good performance and the rather homogeneous structure. Other samples from this group were not tested given the high heterogeneity in the precursors, and the different mineral composition present on the hard carbons.

Figure 70 presents the full voltage profile of hard carbon after each reductionoxidation step. For comparison, a continuous galvanostatic cycle was reported (red curves in all the profiles). Notice that all the continuous cycling profiles followed the iterative step-by-step discharge closely, indicating that the stepwise discharge is representative of the phenomena that occur during the first galvanostatic cycle at the electrodes. Only the HC_beechwood_1400 hard carbon seemed to present higher irreversibility when cycled iteratively, most probably indicating that the two electrodes are not homogeneous.

The capacity lost in each step was quantified as described in Equation 1, and plotted as a function of the Depth of Discharge (DoD) as described in Equation 2,

at each cut-off potential, and is presented in Figure 71. In general, all the electrodes presented an irreversible loss after the 0.01V cut-off voltage, indicating that during the floating stage, there is also irreversible phenomena that consume Na ions. In order to correlate the different irreversible losses to the physicochemical characteristics of the studied electrodes, a summary of the main parameters is given in Table 19, together with the irreversible loss, presented as the percentage of the total discharge capacity at each cut-off potential ( Equation 3).

Notice that at 0.7 V all the electrodes presented 100% of irreversible capacity. This indicates that no Na ions are recovered at this point and probably the sodiation mechanism has not started yet. The generated irreversible capacity is most probably related to the Na ions that are being consumed to form the SEI layer over the hard carbon surface. At a discharge cut-off voltage of 0.3V, the electrodes started gaining reversible capacity (see Table 19), yet, the irreversible capacity is still higher than the reversible capacity. Notice, however, that some electrodes exhibited less irreversibility, of under 70%, while others presented very high irreversibility, up to 93% at this point. This behaviour would seem to be related to the N2 SSA of the electrodes, for the HC carbons derived from miscanthus, wheat-straw-1, rice husk and bagasse precursors. HC_orchard-apple_1400 electrode also exhibited high irreversibility at 0.3V of cut-off potential, despite of its low N2 SSA. It would seem that for this sample, the higher irreversibility could be related to the narrower d002 and d002' interlayer spaces, as reported in Table 19.

At further discharge up to 0.01 V, the irreversible percentage of the discharge capacity decreases for all the electrodes, and the reversible capacity becomes more significant for the hard carbons derived from pine, walnut shell, pine forest residues, and bagasse, while for the others, the irreversible loss is still more important that the generated reversible capacity.

After the floating step, most of the electrodes only exhibited 20-30% of irreversible lost. Clear exceptions were the HC_wheat-straw-1_1400, HC_orchard-apple_1400, and HC_rice husk_1400 with 41.6, 35.9, and 59.2 % of irreversible loss, respectively.

Hence, some hypothesis can be driven:

-The SEI formation seems to be mainly completed at 0.3 V for samples such as woods and walnut shell derived hard carbons, presenting high C contents and low N2 SSAs. While it seemed that, the SEI is still forming between 0.3 V and 0.01V for the HC_wheat-straw-1_1400, and the HC_rice-husk_1400 samples. The latter electrodes have low C content, high content of SiC and an elevated accessible surface area (42 -97 m 2 .g -1 ). These factors might be responsible for the prolonged SEI formation on these materials.

-However, it is also possible that the Na trapping by the structure or the surface was more important for samples such as HC_rice-husk_1400, and is achieved at lower potentials. The larger N2 SSA of this sample and the presence of important graphitic domains, as evidenced by the XRD patterns presented in Chapter V support this. However, it is complex to distinguish between the structural and surface trapping of Na ions, as these phenomena can occur at different potentials, overlapping with the SEI irreversibility.

-The atypical behaviour observed for the HC_orchard-apple_1400 samples, needs further research, as it would seem to behave as the grass derived samples, despite its low N2 SSA and high C content. As mentioned before, its structural parameters might be responsible for this tendency. inspired by [START_REF] Van Den Brink | Identifying supply risks by mapping the cobalt supply chain[END_REF], and [START_REF] Eftekhari | Sodium-ion batteries: New opportunities beyond energy storage by lithium[END_REF]. The bottom layer of the SEI, close to the electrode surface, is thought to present mainly organic degradation products such as sodium carbonates, alkyl carbonates, alkoxides, or polyethylene oxide. The top layer of the SEI, close to the electrolyte is thought to be composed mainly of inorganic species such as NaF, Na2O or NaxPF(3-x)O. A detail description of the SEI species is given in Table 20.

The SEI formation in the early cycles of the battery is of crucial importance as it prevents further electrolyte degradation and provides stability to the system.

Consequently, the SEI layer has to be insoluble and homogeneously distributed over the electrode surface, to protect it from extra electrolyte degradation.

Moreover, the SEI has to be an electronic insulator, again to prevent electrolyte degradation, but a cationic conductor, to allow the diffusion of the Na ions through the electrode.

[214]

SEI composition

The SEI is composed of an organic and an inorganic fraction. The organic fraction originates from the decomposition of the organic solvents and or the additive, while the inorganic fraction is formed from the decomposition of the anion of the salt.

The inorganic products are located close to the electrode, while the organic ones are preferentially found on the external surface, close to the electrolyte interface [START_REF] Eshetu | Impact of the electrolyte salt anion on the solid electrolyte interphase formation in sodium ion batteries[END_REF].

The electrolyte used in this work is a mixture of a Na salt, 1M NaPF6, in the organic solvents, ethylene carbonate (EC) and dimethyl carbonate (DMC) with 1.5 % of fluoroehtylene carbonate (FEC) stabilization additive [START_REF] Simone | Identification and Quantification of the Main Electrolyte Decomposition By-product in Na-Ion Batteries through FEC: Towards an Improvement of Safety and Lifetime[END_REF]. Table 20 summarizes the possible degradation products that can generate from our electrolyte that form the SEI layer. Notice that the Na salt can undergo several decomposition reactions and the products of these reactions can keep producing new species, or even interact with the organic products to form other species.

XPS reference analysis

The XPS analysis started by characterising the surface composition of pristine hard carbon powders and the composite electrodes (prepared as described in chapter III). The survey spectra for the reference pristine powders are presented in Figure 3 of Annex III. While the survey spectra for the reference composite electrodes are shown in Figure 73.

Table 20. Electrolyte degradation products. Data gathered from [211]

, [START_REF] Eshetu | Confronting the Challenges of Next-Generation Silicon Anode-Based Lithium-Ion Batteries: Role of Designer Electrolyte Additives and Polymeric Binders[END_REF]. The list of degradation products is not exhaustive and is based on the most commonly reported data for LIB's SEI formation, adapted for SIB's chemistry.

Electrolyte Anionic Salt NaPF6

Organic solvents EC -DMC

Additive FEC

Degradation products ( non-exhaustive list) 1.

2.

3.

4.

5.

6.

-Sodium alkyl carbonates (ROCO2Na) and sodium alcoxides (RONa) from linear carbonates such as DMC: 12.

13.

14.

-Sodium carbonates and ethylene oxide oligomers: polyethylene oxide (PEO) from EC: 15. 16. Based on reference study performed in the laboratory, the C 1s spectra for the pristine powders were deconvoluted using 9 peaks, associated to the different chemical carbon bonds present at the surface of the samples as shown in Figure 74a. The relative concentration of each related peak is reported in Table 21.

Pristine hard carbon powders

The C1s peaks at 284.8 and 285.5 eV are associated to C atoms in sp 2 and sp 3 hybridizations, as reported in reference samples of graphene and diamond [START_REF] Lascovich | Evaluation of the sp2/sp3 ratio in amorphous carbon structure by XPS and XAES[END_REF],

respectively. The difference in peak position between C (sp 2 ) and C (sp 3 ) was considered to be 0.7 eV from this reference study.

Both powder samples presented similar contents of C sp 2 and sp 3 . As expected, both samples surfaces are mainly composed of ~49 at% of C (sp 2 ) and ~11 at% C (sp 3 ). The concentration of C bonded to O functionalities was also very similar, except for the HC_rice-husk_1400, that presented slightly higher concentrations of C-O groups. CaCO3 was also found which was consistent with the fact that the pristine precursors have Ca in their composition and to the low volatility of Ca at 1400°C. However, carbonate species might be formed after the thermal treatment, because of air exposure. The CaCO3 content is higher for the HC_pine-forest-residues_1400 samples, which is consistent with the higher Ca content of the precursor when compared to rice husk. We noticed that the K 2p peak is visible at the higher binding energies in the C 1s spectra range. The K concentration seemed to be higher for the HC_pine-forest-residues_1400 sample, despite that pine fores residue precursor had less original K than rice husk precursor. This could indicate that K is more volatilized in the rice husk sample, or that in the HC_rice-husk_1400 it is preferentially present in the bulk of the particle and not on its

surface.

An interesting difference between the two samples is the presence of a C 1s peak at low binding energies. For the HC_pine-forest-residues_1400, this peak is thought to be related to the presence of defective carbon and or to porosity in the HC structure. For the HC_rice-husk_1400, this peak is slightly shifted towards lower binding energies and could be related to the mixed contribution of the structural defects and the presence of carbides. This would be consistent with the presence of SiC in the rice husk hard carbon, already confirmed by XRD and SEM-EDX analysis discussed in Chapter V.

Composite electrodes

The composite electrodes exhibited the presence of additional peak associated to -CH2 at 285.6 and -CF2 at 298.9 eV bond in (CH2-CF2)n. The -CH2 contribution to the experimental spectra may disturb the C(sp 3 ) relative concentration. The concentration of the -CF2 was very similar for both composite electrodes. The concentration of the O functionalities, as well as the K, CaCO3, and carbide peaks varied when compared to the pristine powders. This might result from the heterogeneity of the different samples that were analyzed, as the powders are not 100% homogenous in composition. However, this contribution should be minimal, as all the composite electrodes were prepared under the same conditions and from the same pristine powders used in the analysis. It is, therefore, logical to attribute these differences to the coating effect of the PVDF binder that might mask some surface species. In conclusion, the hard carbon electrodes presented on their extreme surface, mainly sp 2 C, O functionalities (mainly C-O), K, and CaCO3. Similar values of the afforded mention species were found for the two selected electrodes. The main difference between the HC_pine-forest-residues_1400 and the HC-rice-husk_1400, was the presence of carbides for the latter. PVDF binder was also observed in the extreme surface and is thought to be homogeneously distributed. From Figure 75, it is seen that the C1s spectra of the HC_pine-forest-residues_1400 discharged until 0.3V, differ from the reference electrode spectra presented in Figure 74c. A peak at low binding energies is observed around 282.5

eV and is denoted as Na-HCx. This peak was attributed to Na-C interactions on the surface of the hard carbon electrodes. This peak has also been reported when NaPF6 electrolyte in EC/DEC solvents were used on a hard carbon electrode [START_REF] Eshetu | Impact of the electrolyte salt anion on the solid electrolyte interphase formation in sodium ion batteries[END_REF].

The Na-C species are related to the intercalation of Na ions into the turbostratic domains or to adsorption of the Na ions by the defects or surface functional groups.

Therefore, this peak is not related to the SEI formation. The presence of the Na-HCx is clear evidence that not only the SEI formation is responsible for the capacity loss on the hard carbons. Moreover, as the depth analyses in XPS is about ~10nm, the Na-HCx related peak has been used as an indication of the thickness of the formed SEI layer. The lower the Na-HCx signal, the thicker the SEI layer [START_REF] Eshetu | Impact of the electrolyte salt anion on the solid electrolyte interphase formation in sodium ion batteries[END_REF].

The intensity and concentration (see Table 1 of Annex III) of the Na-HCx related species decrease with further DoD, and presents a minimum at 0.01V of discharge.

This might indicate that the SEI got thicker at 0.01V, before the floating step.

However, this is valid for homogeneous electrode surfaces and multilayer models of SEI formation. Herein, hard carbon electrodes are far from presenting and homogeneous surface, as can be seen from the SEM images in the preceding chapters. Moreover, it was accepted that the SEI formation followed a mosaic The concentration of the species presented before was not observed to vary significantly with increasing the DoD. The most abundant species seemed to be the alkyl ramifications (C-H SEI peak). Given their abundance, it is likely that there are long alkyl chains, present in the Na alkyl carbonates. Na alkoxides seemed to be present in the lowest concentrations.

If the evolution of the C-C/C-H HC and Na-HCx peaks is indeed related to a thicker, SEI, the fact that the concentration of the organic products in the SEI seemed to remain stable with different discharge stages, indicates that the growth of the SEI layer is probably because of the inorganic contribution. This contribution originates from the NaPF6 salt degradation. This hypothesis will be confirmed with the following discussion.

The F 1s spectra were deconvoluted using four contributions associated to NaF and NaPF6 at lower and higher binding energy, respectively, and two peaks assigned to different degradation by-products (F. D. P-1 and F. D. P-2).

The NaPF6 peak originated from the trapping of the electrolyte salt during the SEI formation or residual electrolyte pollution on the electrode. The reduction of the NaPF6 salt can involve several reactions, as presented in Table 20. These reactions depend on the cut-off voltage and the purity of the electrolyte. Given the numerous possibilities of degradation products for the NaPF6 salt, it is extremely difficult to assign with precision the F 1s spectra to each of these products. For instance, the F. D.P-1 peak could be assigned to degradation products such as PFx, while F.D.P-2 could relate to NaxPFyOz products.

It was observed from Figure 75 and Table 1 of Annex III, that NaF is the major product from the F 1s spectra, at 0.3V. With deeper discharge, NaF concentration decreased while the F.D.P-1 increased. This might indicate that the salt degradation proceeds until 0.01V producing more species than NaF. The F.D.P-2 concentration did not vary significantly with discharge, and might suggest that the formed products are stable from 0.3V. The NaPF6 concentration did not vary either with further discharge and might confirm that this signal originates from the trapped electrolyte salt during the SEI formation.

The Na 1s spectra of the cycled electrodes are also presented in Figure 75. These spectra confirmed the presence of Na-C interactions and NaF formation. A peak at high binding energies was added to group the SEI products containing Na ( Na D.

P-1), analogously to the F degradation product peaks presented before. As expected, the Na D. P-1 seemed to increase with discharge, however, its concentration is the lowest of the Na detected species.

The Na-HCx concentration obtained from the Na 1s spectra was observed to be at least two times higher than the concentration estimated from the same related peak in the C 1s spectra. That might suggest that 2 Na ions can interact with one C atom, which is a quite higher ratio compared to the Li-graphite intercalation compound LiC6. This finding might suggest that:

-The intercalation of Na ions into the turbostratic domains does not follow the same formation of intercalation compounds as in graphite.

-The adsorbed Na ions on the defects are wide more important that the intercalated ions.

-There exists an important alkaline Na clustering.

Further research is needed to corroborate these hypotheses. Here, it is only possible to conclude, that there is a contribution of the Na-C interactions to the irreversible loss of a hard carbon derived from pine forest residues.

The addition of the carbon species involved in the SEI (C-H SEI, C-O, C=O and CO3) is 15.6, 16.1, and 12.2 at% at 0.3V, 0.01V and 0.01V + floating, respectively.

The addition of the F containing species involved in the SEI (except for the trapped NaPF6) is 19.6, 16.8 and 18.2 at% for the 0.3V, 0.01 V and 0.01V + floating stages, respectively. It would seem that the SEI formed over the HC_pine-forest-residues_1400 sample is mainly composed of inorganic species, as suspected before.

At 0.3V, the SEI seemed to be rich in NaF, while at lower discharge states, NaF concentration decreases and other NaPF6 degradation products are developed, most probably related to P containing compounds.

HC_rice-husk_1400

As for the HC_pine-forest-residues_1400 cycled electrodes, the high-resolution spectra of the cycled HC_rice-husk_1400 samples presented intense signals of Na 1s, and F 1s species, as can be seen from Figure 5 of Annex III. The Si 2p signals identified before cycling were no longer observed on the high-resolution spectra of the cycled electrodes. This indicates that the SEI is also recovering the SiC whiskers that exist over the hard carbon surface.

The C 1s spectra of the HC_rice-husk_1400 electrode discharged until 0.3V is similar to the one obtained for the HC_pine-forest-residue_1400 sample at the same potential. Notice from Table 1 of Annex III, that the concentration of the Na-HCx contribution seemed to be lower for the rice husk than for the pine forest residue sample at the same potential. This is consistent with the higher C-C/C-H HC contribution in rice husk, indicating that less Na ions are loss from non-SEI related phenomena at 0.3V for the HC_rice-husk_1400 sample. However, the Na-HCx contribution strongly increases with deeper discharge, as can be seen from Figure 76 and Table 1 of Annex III. This is in contrast with the diminution of the Na-HCx contribution for the HC_pine-forest-residues_1400 sample, observed before.

Consequently, the C-C/C-H HC contribution was observed to decrease with DoD.

This trend might suggest that the SEI is thinner after the 0.01V discharge with floating, that at 0.3V. However, it should be kept in mind that the surface of this sample is even more heterogeneous than the pine forest residues sample, given the presence of the SiC whiskers. Therefore, more analysis are needed to confirm the thickness evolution of the SEI layer on the rice husk electrode.

The concentration of the C-H SEI peak seemed to decrease with discharge while the C=O and the C-O peaks slightly increased. This variation might indicate that the organic fraction of the SEI layer is not as stable as for the HC_pine-forest-residues_1400. Moreover, the concentration of the C-H SEI peak would seem to be higher for the HC_rice-husk_1400 discharged at 0.01V + floating, that the pine forest residues counterpart.

The F 1s spectra in Figure 76, also differed from that of HC_pine-forest-residues_1400 samples. At 0.3V, the NaF concentration is significantly lower than for HC_pine-forest-residues_1400 samples, discharged at the same potential. The NaF concentration increases when the rice husk electrodes are further discharge beyond 0.01V. While the concentration of the FD.P-1 and FD.P-2 decreased with further discharge. In contrast to the pine forest residues, the formation of NaF in the rice husk SEI seemed to be favour at 0.01V and not at 0.3V, while the F. D.P-2 seemed to be preferentially formed at 0.3V. Its concentration decrease with further discharge suggested that the F.D.P-2 products are not stable at lower potentials. The NaPF6 concentration also seemed to decrease with discharge, which would be consistent with a thinner SEI layer, trapping less electrolyte salt on it.

The Na 1s spectra for the HC_rice-husk_1400 cycled samples also exhibited a significant increase of the Na-HCx peak with discharge, in agreement with the C1s results. The Na to C ratio for the Na-C interactions in rice husk samples increased from 0.9 to 1.7 between 0.3V and 0.01V + floating, respectively. These ratios were lower than the ones obtained for the pine forest residues electrodes. This difference might indicate that lower interaction between the Na ions and the carbonaceous subtract are occurring when compared to the pine forest residue sample. This could be explained by the important presence of the SiC phase, which might decrease the available C to trap or adsorb Na ions. In addition, HC_rice-husk_1400 exhibited long-range graphitic domains, which are not favourable to host Na ions. The presence of these domains could also explain the lower Na to C ratio.

Interestingly, the NaF peak of the Na 1s spectra was observed to remain quite stable with discharge. An important difference between the NaF derived from the F 1s spectra and the Na 1s spectra at 0.3V is seen, which seemed to be higher for the Na 1s spectra. This might evidence that the rice husk electrodes are extremely heterogeneous, as expected from the high content of non-C phases such as SiC that are preferentially distributed on the samples.

Notice that no Na-Si interaction was observed from the Na 1s spectra. This might evidence that no Na ion insertion or adsorption is happening on the SiC whiskers, and consequently, the SiC whiskers do not contribute to the reversible capacity of the hard carbon electrodes. However, as evidence by the high-resolution spectra recorded on the rice husk cycled electrodes, the SiC whiskers were recovered by SEI products, which do consume Na ions and contribute to the irreversible loss.

This SEI is already achieved from 0.3V.

The evolution of the SEI formation through discharge for the HC_rice-husk_1400 electrode is in agreement with the gradual irreversible capacity loss from 0.7 to 0.01 + floating stages, reported in section 1. The effect of the higher surface area of the rice husk electrode is not well evidence. It is suspected that it could explain the formation of a thinner layer when compared to the pine forest residues sample. At 0.3, 0.01 and 0.01V + floating, the HC_rice-husk_1400 samples presented 12, 20, and 18 at% of total F content, respectively, and 8, 13, and 11 at% of C content, respectively. This could suggest that for the rice husk samples there is also a higher concentration of inorganic SEI species.

It is evident from the previously presented results that the SEI formation on rice husk hard carbons do not seem to follow the same mechanisms as in the pine forest residues electrodes. Figure 77 schematically summarises the SEI formation evolution for the studied samples.

Conclusions

The results presented in this chapter clearly demonstrated that at least two major phenomena are leading to the irreversible loss in hard carbon materials derived from biomass:

1. The SEI formation over the carbonaceous surface, originated from the electrolyte's degradation during the first cycles. This contribution seemed to be the main source of irreversible lost in hard carbons derived from pine forest residues precursors.

2. The lost generated by trapping and adsorption of Na ions in the turbostratic domains and defects, respectively, which seemed to be more significant for the rice husk hard carbon.

The SEI formation does not follow the same mechanisms regarding the used hard carbon electrode. For pine forest resides derived hard carbons, the SEI would seem to grow with deeper discharge, while it is thinner for the rice husk electrodes.

Moreover, the SEI was observed to be rather stable after 0.3V of discharge for the pine forest residues sample, while it was only stable after 0.01V of discharge for the rice husk samples. The difference on the hard carbon composition surface, especially the presence of SiC whiskers, as well as the higher specific surface area of rice husk, might explain the distinctive SEI formation. However, both SEI were observed to be mainly composed of inorganic species.

correlations between their composition and the final hard carbon properties and electrochemical performances.

It was confirmed that ash content and composition seem to play a major role in the hard carbon formation and therefore properties and electrochemical performances.

In general, it was observed that precursors presenting less than 3 wmf% of ash content, produced high-quality hard carbons, with reversible capacities over 280 mAh.g -1 and ICEs over 75% at 37.2 mA.g -1 . At 372 mA.g -1 , the same hard carbons exhibited reversible capacities over 250 mAh.g -1 .

The biomass precursors presenting an ash content of over 3 wmf%, delivered less performant hard carbons, presenting low C content (74 -91 wmf%), N2 SSA higher than 20 m 2 .g -1 , and mineral and graphic phases on their structure.

The biomasses presenting more than 3% of ash were mainly grasses and cereals, for which the ash fraction was mainly composed of Si and Ca compounds. Additionally, some agricultural precursors such as seeds and fruit tree branches, also presented more than 3% of ash, which were mainly composed of K and Ca compounds.

Si compounds were shown to be related to the formation of a SiC phase on the surface and the bulk of the hard carbon particles. SiC formation led to a considerable increase of the N2 SSA, and to a lower C content. Si and Ca compounds were also related to the formation of graphitic domains in the materials, of approximately 10 stacked graphene layers of more than 30 nm long, which is thought to limit the Na storage. Interestingly, biomass precursors presenting high Ca content but low Si content did not develop graphitic domains. This behaviour evidenced a possible synergetic effect of Si and Ca compounds towards catalytic graphitization of biomass precursors.

Therefore, hard carbons obtained from high ash content biomass precursors delivered lower electrochemical performances than the low-ash derived hard carbons. The former hard carbons exhibited reversible capacities between 170 and 262 mAh.g -1 at C/10, ICEs between 48 and 72 %, evidencing an important irreversible loss in these materials.

Finally, in order to understand the origin of the irreversible capacity observed before, two hard carbons were chosen to study the electrochemical phenomena at the origin of the irreversible loss at the first cycle. The hard carbons were selected according to the ash and Si content of their precursors. In this way, pine forest residues precursor was chosen as a lowash derived hard carbon, while rice husk precursor was selected as the high ash, high Si derived hard carbon.

It was observed that the irreversible loss was originated by both the SEI formation, and structural trapping of Na ions.

The SEI formation seemed to proceed in a different way for pine forest residues derived hard carbon (low ash and low Si content) than for rice husk derived hard carbon (high ash and Si content). The former electrodes presented a stable SEI around 0.3V, which became thicker with deeper discharge, while the rice husk electrodes exhibited a thinner SEI, which was stable only below 0.01V of discharge. This was related to the higher N2 SSA due to the presence of SiC in rice husk hard carbon. Moreover, the SiC in rice husk hard carbon was observed to be covered by the SEI film and to be inert towards the Na ions insertion. On the other hand, structural trapping of Na ions was observed for the studied samples and was found to be more important for the sample presenting graphitic domains, rice husk derived hard carbon.

This research work has demonstrated that the biomass composition can severely affect the final hard carbon performance. Poor performance is associated mainly to high Si contents in the biomass precursors. Ash rich in Ca or K compounds also led to low-performance hard carbons. However, the main reason for this seemed to be the lower C content, rather than the apparition of parasitic phases on the hard carbon structure of elevated N2 SSA.

Future work studies

The presented research work allows proposing some complementary fundamental studies, as well as some application-oriented perspectives.

Fundamental studies

2.1.1 Understanding of the macromolecules role in hard carbon formation

Regarding the understanding of hard carbon formation from biomass precursors, a gap is still found for the role of the macromolecules in the hard carbon formation.

To address this subject, a more detailed characterisation of the biomass composition should be performed, together with close tracking of the biochar formation.

In this study, biomass macromolecular composition was presented in terms of sugars, lignin, and extractives. However, it would be more accurate to quantify the lignin composition in terms of basic phenolic units and the detailed extractives composition. For this, more interdisciplinary studies are needed between the biochemestry and the biomass sciences, which are surprisingly dissociated despite that biomass composition intrinsically, relies on biochemical processes.

Additionally, it would be highly enriching, to properly characterised the raw mineral matter contained in the biomass precursors, and not only to quantify their elemental content. Proper identification of the inorganic compounds would allow to follow their behaviour under thermal treatment, and to clearly identify their effect on the hard carbon formation and properties.

Once a detailed composition quantification is done, the focus should be given to biochar formation of low ash precursors. Earlier stages of pyrolysis are crucial for setting the hard carbon structure. It is at these earlier stages where the macromolecules play a major role, determining the chemical reactions of biochar formation.

Besides following the thermal degradation by thermogravimetric analysis and differential scanning calorimetry techniques, more attention should be given to the solid evolution of the carbonaceous phase among the thermal treatment. Operando techniques such as SAXS could be considered for this purpose. Besides providing information regarding the local and macroscopic organisation of the solid phases, this technique might be useful to identify the formation of accessible and closed ultramicropores. Ex-situ electron tomography could also be a powerful tool to characterise the three-dimensional evolution of the carbonaceous matrix.

The former techniques could be complemented with ex-situ 13 C and 1 H nuclear magnetic resonance (NMR) to monitor the aromaticity and hybridisation of the solid, to correlate this information with the biomass composition. ex-situ XPS and TPD-MS have demonstrated to be powerful tools to characterise the surface chemistry, and should be implemented in this methodology.

This approach might not only be helpful for biomass-derived hard carbons, but could also contribute to the understanding of the hard carbon structure formation in general, whatever the initial material used.

This would allow a clearer starting point for models and simulations of the hard carbon structure. These studies are of extreme importance to fully understand the Na ions storage mechanisms and optimise the Na ions uptake.

Understanding the electrochemical phenomena in hard carbons during slow and fast cycling

Deeper electrochemical studies should be done over the hard carbons from low ash containing precursors, in order to identify the fundamental differences between these samples. In particular, it is proposed to perform and evaluate the Na diffusion kinetics (by galvanostatic intermittent titration technique -GITT), the Na ions and electrodes impedance behaviour (by electrochemical impedance spectroscopy -EIS) and the electrodes redox activity (by cyclic voltammetry -CV).

A full-cell configuration should also be considered for further negative electrode materials testing. As shown here, the metallic Na counter electrode might not give accurate results, especially at high current densities. Full-cell experiments using a selected positive material would be more accurate to evaluate the rate capability of the hard carbons materials.

Perspectives towards hard carbons industrial application

The production of hard carbon electrodes from biomass precursors seems to be far from being fully understood and optimised. The results presented in this study demonstrated that if an interdisciplinary approach is taken, the biomass-derived hard carbons properties can be properly tailored towards highly performant materials, that would lead to competitive full SIB devices. Some propositions are described below, that could help to achieve this goal.

Estimation of the hard carbon performance from its precursor composition

The results gathered in chapter V are valuable insights for a correct selection of the biomass precursors. However, they could also be used to estimate the hard carbon performance of a new biomass precursor, before a synthesis step, by producing statistical correlations of the presented data. A preliminary statistical analysis was done in order to explore this perspective. Using the Minitab software, the composition parameters of the precursors and the electrochemical performance parameters of the obtained hard carbons, were analysed. High-quality regression functions were obtained from this analysis, with high predictive potential.

Moreover, interactions between the composition parameters could be identified by using this technique. Validation of the regressions is still needed.

Reducing the cost of the process

Despite the good performances of hard carbons from low ash precursors, efforts are still needed to maximize these performances and minimise the process cost.

Options are abundant regarding both the synthesis procedures and electrode engineering.

Improvement of the synthesis conditions

Pre-treatments such as hydrothermal carbonization (HTC) of biomass might be useful towards decreasing the final pyrolysis temperature, and therefore, decreasing the process cost, as seen in chapter II. Interestingly, HTC also allows valorising wet biomasses, which are not suitable for pyrolysis treatment. A preliminary study of HTC of horse manure was done in collaboration with the Chemical department of the Umeå University and the MTC enterprise in Umeå, Sweden. It was found that HTC treatment might indeed reduce the final pyrolysis temperature requirements. However, in the case of horse manure, a purification might be needed, especially to remove sand particles, before a second thermal treatment. HTC is a promising route for hard carbon production. A whole set of experimental conditions, such as wet biomass precursors, HTC temperature, HTC residence time, HTC biomass to water ratio, and post pyrolysis final temperature, are still to be explored.

Moreover, it should be remembered that hard carbon is just one of three-stream products resulting from pyrolysis or HTC. Therefore, the recovery and valorisation of the other streams might also represent an economic advantage for he hard carbon production. Therefore, it is worth to focus some research efforts towards the development of integrated processes. Biomass pyrolytic polygeneration is an example that these approaches are attractive towards a complete biomass valorisation [START_REF] Yang | Biomass-Based Pyrolytic Polygeneration System for Bamboo Industry Waste: Evolution of the Char Structure and the Pyrolysis Mechanism[END_REF].

Improvement of the electrode's engineering

Regarding electrode engineering, two perspectives can be proposed. First, some biomasses such as woody residues or high lignin content precursors could be used to prepare binder-free-self-supported electrodes. In this way, no binder, carbon additive or current collector would be needed, representing higher energy densities and lower costs for the final full cell. Moreover, the use of toxic solvents such as Nmethyl-2-pyrrolidone -NMP is avoided by this method [START_REF] Shen | Ultra-Thick, Low-Tortuosity, and Mesoporous Wood Carbon Anode for High-Performance Sodium-Ion Batteries[END_REF].

Other option would be to reduce the biomass-derived hard carbon particle size, in order to obtain higher conductivity. In this scenario, the grinding conditions and the loading of the electrode should be studied. The missing parameters are presented in chapter V, Figure 6.

ANNEX I. complementary results for Chapter IV.

Figure 2 .

 2 Figure 2. Performance in terms of working potential as a function of the specific capacity delivered by some negative electrode materials for SIBs. TMO: transition metal oxides, TMS: transition metal sulphides. Reproduced from [5].

Figure 3 .

 3 Figure 3. Final carbonaceous materials as a function of the O/C ratio of the precursor and the pyrolysis temperature. The inset images correspond to a) optical microscopy of the microtexture of the solidified product after low pyrolysis temperature, of low O/C ratio precursors. b) Optical microscopy of the microtexture of the product of low pyrolysis temperature of high O/C ratio precursors. Adapted and reproduced from [32]. c) Soft carbon obtained from the pyrolysis of perylenetetracarboxylic dianhydride (PTCDA) at 1600°C. Adapted and reproduced from [35]. d) Hard carbon obtained from resol-type phenolic resin pre-oxidised at 1000°C and pyrolysed at 1600°C. Adapted and reproduced from [36]. e) Hard carbon derived from pinecone biomass at 1600°C. Adapted and reproduced from [37]. f) Graphitic carbon derived from anthracene at 2300°C. Adapted and reproduced from [38]. g) GL-200 Toyo Tansin commercial glassy carbon obtained at 3000°C. Adapted and reproduced from [39].

Figure 4 .

 4 Figure 4. Schematic representation of the biomass structure and composition, showing how the secondary cell wall of plants is composed of macrofibrils. The macrofibirls are at the same time composed of microfibrils, in where cellulose, hemicelluloses and lignin macromolecules are found. Illustration by Juana Molina, inspired by [53].

Figure 5 .

 5 Figure 5. a) Biochar phase evolution and b) composition variation with pyrolysis temperature until 700°C. Adapted and reproduced from [65]. c) Hard carbon structuration from biochar at temperatures above 1000°C.

Figure 6 .

 6 Figure 6. a) d002 interlayer space, b) ID/IG ratio, c) N2 SSA, d) reversible capacity and ICE of the raw argan shell and washed argan shell derived hard carbons. Data was taken from [110]. e) Thermogravimetric curves for raw argan shell and washed argan shell precursors under Ar atmosphere. Reproduced from [110].

Figure 7 .

 7 Figure 7. a) C and b) O content evolution of hard carbon with final pyrolysis temperature for different precursors: peat moss (PM) [82], banana peels (BP) [112], apricot shells (APS) [113], pectin free apple waste (PFAW) [78], and corn straw piths [95]. *O content measured by XPS.

Figure 8 .

 8 Figure 8. a) XRD patterns, and b) Raman spectra, for the series of hard carbon derived from mangosteen shell at different temperatures. c) TEM image for the mangosteen shell derived hard carbon at 800°C and d) 1600°C, pyrolysed during 2 hours. Adapted and reproduced from [114].

Figure 9 .

 9 Figure 9. a) d002 interlayer space, b) ID/IG ratios, c) N2 SSA of different hard carbons obtained from direct pyrolysis: pectin free apple waste (PFAW) [78], mangosteen shell (MGS) [114], date palm seeds (DPS), date palm pulp (DPP) [76], sugar cane bagasse (SCB) [97], corn straw piths [95], tamarin fruit shells (TS) [116] and apricot shells (APS) [113]. d) SAXS patterns for argan shell derived hard carbons at different temperatures, evidencing the presence of bigger ultramicropores at higher temperatures, adapted and reproduced from [110].

Figure 10 .

 10 Figure 10. Hard carbon electrochemical profile.

Figure 11 .

 11 Figure 11. GITT profile of sucrose derived hard carbon. The inset presents the calculated diffusivity as a function of the state of discharge. Reproduced from [42].

Figure 12 .

 12 Figure 12. a) Insertion-adsorption mechanisms for Na ions storage in hard carbon. Reproduced from [42]. b) Na ions storage model proposed by Morikawa et al., showing the influence of the final pyrolysis temperature in the mechanisms. Reproduced from [121].

Figure 13 .

 13 Figure 13. a) First reduction capacity, b) first reversible capacity, c) ICE, and d) reversible capacity at 1 A.g -1 of different hard carbons obtained from direct pyrolysis. Pectin free apple waste (PFAW) [78], mangosteen shell (MGS) [114], date palm seeds (DPS), date palm pulp (DPP) [76], sugar cane bagasse (SCB) [97], corn straw piths [95], tamarin fruit shells (TS) [116], and apricot shells (APS)[START_REF] Zhu | A porous biomass-derived anode for high-performance sodium-ion batteries[END_REF]. The legend specifies the pyrolysis gas and residence time, and the current density used during cycling.

Figure 14 .

 14 Figure 14. a) Plateau and b) slope capacities of selected hard carbons as a function of temperature. mangosteen shell (MGS) [114], date palm seeds (DPS), date palm pulp (DPP), apricot shells (APS), banana peels [112], and pinecone (PC) [37]. The legend specify the pyrolysis gas and residence time, and the current density used during cycling.

Figure 15 .

 15 Figure 15. a) Charge discharge profiles for peat-derived hard carbons prepyrolysed at different temperatures and pyrolysed at 1400°C. Galvanostatic tests were done at a current density of 50 mAh.g -1 . b) Rate performance of the different peat-derived hard carbons, pre-pyrolysed at different temperatures. The electrolyte used was 1 M NaPF6 in EC:PC (1:1). The pre-pyrolysed chars were treated with a KOH solution for activation, and further washed with HCl solution. Reproduce from [111].

Figure 16 .

 16 Figure 16. a) Apricot nutshell derived hard carbon (temperature not specified). Reproduced and adapted from [113]. b) lotus stem derived hard carbon at 1600°C. Reproduced and adapted from [128]. c) Wood derived hard carbons at 1000°C (cross-section). Reproduced and adapted from [126]. d) Peat moss derived hard carbon after air activation and pyrolysis at 1100°C. Reproduced and adapted from [82].

Figure 17 .

 17 Figure 17. SEM images for a) rape pollen grain precursor, b), c) rape pollen grain derived hydrochar, d), e) rape pollen grain hard carbon obtained at 600°C from the resulting hydrochar, f), g) directly pyrolysed rape pollen grain at 600°C hard carbon without HTC treatment. Adapted and reprinted with permission of [72]. h) Directly pyrolysed hard carbon at 900°C from pistachio shells, i) Pistachio shell hard carbon obtained at 800°C after HTC treatment of the precursor. Adapted and reprinted with permission of [139]. j) Pomelo peel derived hard carbon at 600°C, and k) 700°C, after HTC treatment of the precursor. Adapted and reproduced from [138].

  pyrolysis b. After pyrolysis c. Reduction capacity : Na uptake d. Reversible Capacity : Na release e. C content in the hard carbon quantified by elemental analysis or XPS f. DI: deionized water g. HTC: hydrothermal carbonization h. DW: distillated water * Different fitting procedures might apply in-between the reported values.

Figure 18 .

 18 Figure 18. Schematic representation of a) the home-made reactor for the first pyrolysis stage and b) the tubular furnace for high-temperature treatment. c) Thermal treatment until 450°C from biomass to biochar. c) Thermal treatment until final temperature (T final) from biochar to hard carbon. The second thermal treatment was performed using an alumina tubular furnace of Nabertherm in the laboratories of DTNM of CEA-Grenoble. Figure 18b describes the oven set-up. Ar atmosphere was used with a flow rate of approximately 55 L.h - 1 . Before the Ar flow, a vacuum was created for one hour to ensure an oxygen-free atmosphere in the oven. The thermal treatment is described in Figure 18d and consisted of a single ramp of 200°C.h -1, until a final temperature of 1000, 1200, and 1400°C for the hard carbons studied in chapter IV, and of 1400°C for the hard carbons studied in chapter V and VI. The final temperature was held for one hour, and the hard carbons were let to cool down under argon atmosphere overnight. The obtained hard carbons were weight and stored for further characterisations.

  of the obtained hard carbon particles. SEM observations are coupled to Energy Dispersive X-ray Spectroscopy (EDX), in the same instrument, allowing the analysis of the surface composition. Briefly, SEM principle consists of irradiating the sample surface with electrons at a given energy, which is determined by the operating voltage. The incident electrons will interact with a specific area of the sample surface, generating backscattered, secondary, and Auger electrons, together with X-rays, from the studied sample. The depth of the interaction will depend on the operation voltage, the atomic mass of the sample elements, and the incident angle of the electrons beam. To obtained deeper interaction zones, the operation voltage an incident angle should increase, while the atomic mass of the sample should decrease. When the sample is irradiated with a high-energy beam, X-rays are generated, that are specific for each element in the sample, and therefore, qualitative and semi-quantitative analysis of the surface composition can be done.The observations were made in a Zeiss microscope equipped with a field emission gun (LEO 1530), provided by the platform of nano-characterisation techniques (PFNC) of the CEA-Grenoble. For morphology observations, the InLENs detector was used to obtain a better define image of the surface with no shadows. The microscope was configured with an aperture size of 30 μm, an operating voltage of 5 kV and a working distance of 5 -7 mm. For EDX analysis, the Secondary Electrons detector was used, together with an aperture size of 60 μm, an operating voltage of 15 kV and a working distance of 10 mm. A cartography of the sample surface was taken for specific hard carbon samples, during 5 min of acquisition.

Figure 19 .

 19 Figure 19. Characteristic gas isotherms defined by IUPAC. Reproduced from [19].

  pores. By modelling the resulting intensity pattern, it is possible to quantify the pore size and the pore volume of both closed and open pores. A typical SAXS pattern for porous carbon is composed of three characteristic zones, as represented in Figure 20. The first zone at low Q values, related to the macroscopic surface area of the powder grains. It is a result of the scattering of well define sharp interfaces. This zone has a characteristic Q -4 intensity decay behaviour. If the intensity decay turns toward Q -3 behaviour, it is an indication that the interfaces are not well defined anymore. The second zone in intermediate values of Q, is characterised by a pseudo plateau of the intensity, denominated Guinier's domain. This region is related to the micropores or ultramicropores diffusion and corresponds to the region of interest of this work. Finally, at high Q values, the intensity drops again, which is an indication of the diffusion of the micropores surface or graphene layers.

Figure 20 .

 20 Figure 20. Schematic representation of the information comprised in the SAXSpatterns. The pattern corresponds to a hard carbon derived from beechwood biomass at 1400°C. The schema is inspired by[START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF]. The molecular schema is taken from PubChem.

Figure 21 .

 21 Figure 21. Half-cell configuration.

Figure 22 .

 22 Figure 22. Carbon and oxygen content evolution for the biochar and obtained hard carbons.

Figure 24 .

 24 Figure 24. XRD patterns for a) pine and b) beechwood biomass, biochars and hard carbons. Cellulose PDF: 00-060-1502. Graphite 2h PDF: 01-075-1621. SiO2-Quartz PDF: 00-046-1045.

Figure 25 .

 25 Figure 25. XRD patterns for a) miscanthus and b) wheat straw biomass, biochars and hard carbons. Cellulose PDF: 00-060-1502. Graphite 2h PDF: 01-075-1621. SiO2-Quartz PDF: 00-046-1045. K2CO3*1.5H2O PDF: 00-011-0655. CaCO3 PDF: 04-012-0489. SiC-3C PDF: 00-029-1129

  increase the specific surface area of the obtained hard carbons, as is further discussed in section 4.2.

Figure 28 .

 28 Figure 28. a) Un-zoomed TEM observation for the pine derived hard carbon. b) TEM image of a particle of pine derived hard carbon presenting spherical nanoparticles. c) TEM image over the SiC whiskers and spherical SiC nanoparticles for miscanthus and d) wheat straw derived hard carbons. The table under the TEM images reports the atomic elemental composition of the analysed zones for c and d.

Figure 29 .

 29 Figure 29. Raman spectra for a) pine, b) beechwood, c) miscanthus and d) wheat straw hard carbons (HC).

Figure 31 .

 31 Figure 31. SEM images for pine derived hard carbons at a) 1000, b) 1200, and c) 1400°C.

Figure 32 .

 32 Figure 32. SEM images for beechwood derived hard carbons at a) 1000, b) 1200, and c) 1400°C. The red square indicates the zoomed zone, presented on the right image.

Figure 33 .Figure 34 .

 3334 Figure 33. SEM images for miscanthus derived hard carbons at a) 1000, b) 1200, and c) 1400°C. The red square indicates the zoomed zone, presented on the right image.

44 EDX analysis of zones 1 and 2 Figure 36 .

 44236 Figure 36. N2 adsorption/desorption isotherms for a) pine, b) beechwood, c) miscanthus and d) wheat straw derived hard carbons.

Figure 37 .

 37 Figure 37. CO2 adsorption/desorption isotherms for a) pine, b) beechwood, c) miscanthus and d) wheat straw derived hard carbons.

  [START_REF] Buiel | Model of micropore closure in hard carbon prepared from sucrose[END_REF]. These phenomena give place to partially accessible or completely inaccessible pores, towards the N2 molecule. This explains the rather low values of N2 SSA and volume for the

Figure 41 .

 41 Figure 41.Quantified electrochemical parameters for the series of hard carbons. a) Reduction capacity, b) reversible capacity, c) ICE and d) average oxidation potential. Error bars correspond to the variance of the mean values for the two coin-cells, for each material.

Figure 43 .

 43 Figure 43. Stability of cycling for a) pine, b) beechwood, c) miscanthus and d) wheat straw derived hard carbons at C/10 rate.

  In this chapter, 25 biomass precursors presenting significant variations in their composition were selected to synthesise hard carbons at 1400°C and confirm the observations of the previous chapter. The obtained hard carbons were characterised in terms of composition, structure, porosity and specific surface area. The electrochemical performances of the hard carbons were also quantified at slow and fast cycling rates. The obtained physicochemical and electrochemical properties were analysed in relation to the biomass precursor composition, and correlations were proposed.

Higher C contents were

  obtained for 450°C-biochars from precursors with lower ash contents. Over 70 wmf% of C content was obtained for the 450°C-biochars of GI, GIII precursors groups and for GII-10, 16 and 17 precursors. O content in the 450°C-biochars ranged from 1 to 14 wmf%. The lowest values were observed in the GII precursors, given their higher ash concentration (O content calculated by difference). GII-15: rice husk presented the highest ash value of all the biochars, of 34 wmf%, as expected. The seeds derived biochars GIII-19, GIII-23 and GIII-25, presented the highest N concentrations of all the biochars, with values between 2 and 7 wmf%. This indicated a N concentration on the biochars derived from GIII-23 and GIII-25 precursors. It has been observed that N volatilization during slow pyrolysis under N2 flow is limited up to intermediate temperatures (~500 -700°C).

Figure 44 .

 44 Figure 44. XRD patterns for the hard carbons derived from GI. Graphite 2H PDF : 01-075-1621. SiC 3C PDF: 00-029-1129. The star denotes a non-identified phase.

  enhanced graphitization can proceed by different phenomena, such as C diffusion into the Ca compound, and by the formation of intermediate phases such as calcium carbides.

Figure 46 .

 46 Figure 46. XRD patterns for the hard carbons derived from GII. Graphite 2H PDF: 01-075-1621. SiC 3C PDF: 00-029-1129. The stars denote non-identified phases.

Figure 48 .

 48 Figure 48.Raman spectra of the selected hard carbons of a) GI, b) GII, and c) GIII groups. d) AD1/AG ratio obtained for the selected hard carbons.

  the K content in biomass might suggest that K compounds are involved in the development of a less defective carbon lattice, as suggested before.

202 Figure 50 .

 20250 Figure 50. N2 adsorption-desorption isotherms for a) GI, b) GII and c) GIII derived hard carbons. For the isotherms of GI-6: pine, GI-7 beechwood, GII-10: miscanthus and GII-11: wheat straw-1, please refer to Figure 36. The N2 adsorption-desorption isotherms for most of the samples did not exhibit a pure shape, characteristic of single microporous or mesoporous material. GI-4: SRF-poplar and GI-5: vineyard derived hard carbons seemed to present mainly mesoporous isotherms. GII-16: bagasse, GII-17: cornhusk and GII-18: switchgrass hard carbons, presented rather characteristic microporous isotherms. For the GIII derived hard carbons, only the GIII-20: sunflower-shell, GIII-23: papaya seed and

Figure 51 .

 51 Figure 51. CO2 adsorption isotherms for a) GI, b) GII and c) GIII derived hard carbons. For the isotherms of GI-6: pine, GI-7 beechwood, GII-10: miscanthus and GII-11: wheat straw-1, please refer to Figure 37.

Figure 52 .

 52 Figure 52. a) N2 SSA vs. C in biomass, b) CO2 SSA vs. C in biomass, c) N2 SSA vs. C in hard carbon, d) CO2 SSA vs. C in hard carbon, e) N2 SSA vs. Ash, f) CO2 SSA vs. Ash, g) N2 SSA vs. Si, and h) CO2 SSA vs. Si, for all the hard carbons.

Figure 53 .

 53 Figure 53. N2 and CO2 SSA as a function of a) Na, b) Mg, c) P and d) S concentration in the biomass precursors of GIII group.

Figure 56 , 4 . 2

 5642 Figure56, respectively. For each hard carbon, the first, the fifth, and the tenth cycle profiles were plotted, when available.

Figure 57 .

 57 Figure 57. 1 st Reduction capacity vs. a) C, and b) N content in the hard carbons. c) Si content in biomass precursors, d) N2 SSA and e) CO2 SSA of the hard carbons.

Figure 58 .

 58 Figure 58. Corrected reduction-oxidation profiles of the first cycle of the GII derived hard carbons, where the specific capacity is expressed per g of C in the electrode.

Figure 59 .

 59 Figure 59. Corrected 1 st Reduction capacity vs. a) C content in hard carbon, b) Si, c) ash content in biomass, and d) N2 SSA of the hard carbons.

60 .

 60 No clear relation could be found between the macromolecular composition of the precursor and the reversible capacity of the derived hard carbons, as reported in Figure12of Annex II.

Figure 60 .

 60 Figure 60. 1 st Oxidation (or reversible) capacity vs. a) ash, b) C, c) H, d) Si, e) Ca, f) K content in biomass; g) C content, h) N2 SSA, and i) CO2 SSA of the hard carbons.

Figure 61 .

 61 Figure 61. Plateau and slope capacity of the first oxidation for all the hard carbons.

  This behaviour was not observed in the previous chapter. It might suggest that the Na ions storage mechanisms of hard carbons are highly dependent on the hard carbon composition, surface and structure. The observations presented here might explain the lack of agreement in literature, concerning the storage mechanisms in hard carbons.

Figure 62 .

 62 Figure 62. Plateau Capacity vs. a) N2 SSA, c) CO2 SSA, e) Si content in biomass, and g) Ash content in biomass. Slope Capacity vs. b) N2 SSA, d) CO2 SSA, f) Si content in biomass, and h) Ash content in biomass; for all the hard carbons.

Figure 63 .

 63 Figure 63. Plateau capacity vs. a) AD1/AG ratio of the studied hard carbons, c) Ca and e) K content of the biomass precursors. Slope Capacity vs. b) AD1/AG ratio of the studied hard carbons, d) Ca, and f) K content of the biomass precursors.

Figure 64 .

 64 Figure 64. Average oxidation potential vs. a) ash, b) Si content in biomass, and c) C content in hard carbon.

Figure 65 .

 65 Figure 65. 6 th Oxidation capacity vs. a) C content in hard carbon, b) ash, and c) Si content in biomass.

Figure 66 .

 66 Figure 66. 10 th Ox. Cap / 5 th Ox. Cap. vs a) ash, b) Si content in biomass, c) C content in hard carbon, d) N2 SSA, e) CO2 SSA, and f) d002 interlayer space.

4. 3 . 1 Figure 67 .

 3167 Figure 67. Cycling stability for hard carbons derived from a) GI, b) GII, and c) GIII groups. The first 5 cycles were held at a C/10 rate, and the following 90 cycles were held at a C rate.

Figure 68 .

 68 Figure 68. Capacity fading at C as a function of a) Ash content in biomass, b) Si content in biomass, c) N2 SSA, d) CO2 SSA, and e) C content in hard carbon.

Figure 69 .

 69 Figure 69. a) Schematic representation of the metallic Na electrode plating/stripping process that lead to the formation of dendrites and dead Na. Adapted and reproduced from [209]. b) Last cycle of all the assembled half-cells as a function of the surface capacity at the C regime.

Figure 70 .

 70 Figure 70. Galvanostatic profiles for the selected hard carbon samples. Black curves (0.7V), blue curves (0.3V), green curves (0.01V) and orange curves (0.01V + floating). Red curves represent continuous cycling.

Figure 71 .

 71 Figure 71. Cumulative irreversibility generated at each cut-off potential as a function of the DoD for the selected samples.

Figure 72 .

 72 Figure 72. Schematic representation of the mosaic SEI formation. Illustrationinspired by[START_REF] Van Den Brink | Identifying supply risks by mapping the cobalt supply chain[END_REF], and[START_REF] Eftekhari | Sodium-ion batteries: New opportunities beyond energy storage by lithium[END_REF]. The bottom layer of the SEI, close to the electrode surface, is thought to present mainly organic degradation products such as sodium carbonates, alkyl carbonates, alkoxides, or polyethylene oxide. The top layer of the SEI, close to the electrolyte is thought to be composed mainly of inorganic species such as NaF, Na2O or NaxPF(3-x)O. A detail description of the SEI species is given in Table20.

-

  Semi carbonates, such as Na double alkyl carbonates Signals of C 1s and O 1s are in agreement with the composition of the hard carbons, while the F 1s signals originate from the PVDF binder. Notice that the HC_pineforest-residues_1400_electrode sample, also exhibited K 2p and Ca 2p signals, in agreement with the inorganic composition of the precursors. The apparition of these signals confirms that this hard carbon presents inorganic matter. For the HC_rice-husk_1400_electorde sample, C 1s, O 1s, F 1s, Ca 2p and K 2p signals were detected. An additional Si 2p peak associated with the presence of silicon carbide (SiC) in the material, was also detected. The latter is in agreement with the XRD and SEM-EDX observations of the HC_rice-husk_1400 samples, presented in the previous chapters. The C 1s spectra for the hard carbon pristine powders and composite electrodes are shown in Figure 74.

Figure 73 .

 73 Figure 73. Wide scan spectra for the HC_pine-forest-residues_1400 and HC_rice-husk_1400 electrodes. The right insets are the zoom of the red zones on the left.

Figure 74 .

 74 Figure 74. C1s spectra for the pristine hard carbon electrodes of a) HC_pineforest-residues, and b) HC_rice-husk_1400. C1s spectra for the composite electrode of c) HC_pine-forest-residues, and d) HC_rice-husk_1400. The black line corresponds to the experimental measurements while the red line is the result of the spectra deconvolution. The peaks use for deconvolution are also indicated in a) and c). The colour conventions used in a) and c) spectra are the same for the b and d spectra, respectively.

Figure 75 .

 75 Figure 75. Galvanostatic profiles until the different discharge stages for the samples of HC_pine-forest-residues_1400. The C 1s, F 1s and Na 1s deconvoluted spectra for the cycled electrodes. The colour conventions used in the spectra of the left are the same for the spectra of the right, respectively.

  Figure74c) are most probably no longer present in the cycled stated, as they could have reacted during the reduction stage and interact with the SEI organic products. Therefore, the O-containing groups observed in the C 1s spectra of the cycled electrodes (Figure75), were considered to originate from the SEI organic products. The C-O peak was assigned to the presence of Na alkoxides (RONa) and ethylene oxide oligomers, such as polyethylene oxide, PEO (-CH2-CH2-O-)n. The C=O peak was ascribed to the presence of Na alkyl carbonates (ROCO2Na). While the final CO3 peak was assigned to the presence of Na carbonate (Na2CO3) in the SEI. It is no longer possible to assure the presence of CaCO3, as the Ca 2p signal was very weak. Moreover, a C-H SEI peak was used to follow the alkyl ramifications (R) of polymeric or long-chain compounds of the SEI products.

Figure 76 .

 76 Figure 76. Galvanostatic profiles until the different discharge stages for the samples of HC_rice-husk_1400. The C 1s, F 1s and Na 1s deconvoluted spectra for the cycled electrodes. The colour conventions used in the spectra of the left are the same for the spectra of the right, respectively.

Figure 77 .

 77 Figure 77. Schematic representation of the SEI formation for the HC_pine-forest-residues_1400 and HC_rice-husk_1400samples. The thickness variation presented in the illustration is purely indicative and needs to be confirmed with additional techniques.

Figure 1 .Figure 2 .Figure 3 .

 123 Figure 1. Experimental XRD patterns under Cu radiation for the different types of Opal in biomass. Reproduced from [176].

Figure 2 .

 2 Figure 2. Higher resolution XRD pattern for GII-15: rice husk hard carbon. The pattern evidence the presence of the (200) and (222) peaks and their low intensity compared to the (111) family of peaks.

Figure 78 .

 78 Figure 78. d002 interlayer distance for all the obtained hard carbons plotted as a function of the macromolecular composition of the biomass precursors.

Figure 79 .

 79 Figure 79. d002 interlayer distance for all the obtained hard carbons plotted as a function of the elemental composition of the biomass precursors.

Figure 80 .

 80 Figure 80. AD1/AG ratio for all the obtained hard carbons plotted as a function of the elemental composition of the biomass precursors.The missing parameters are presented in chapter V, Figure6.

Figure 81 .Figure 82 .Figure 84 .Figure 85 .

 81828485 Figure 81. Pore size distributions from N2 isotherms for a) GI, b) GII and c) GIII obtained hard carbons.

Figure 86 .Figure 87 .Figure 88 .Figure 89 .Figure 90 .

 8687888990 Figure 86. First Reduction Capacity plotted as a function of the elemental composition of the biomass precursors.

6 Capacity 6 Capacity 6 Capacity

 666 Fading at C (mAh.g -1 .cycle -1 ) Fading at C (mAh.g -1 .cycle -1 Fading at C (mAh.g -1 .cycle -1 )

Figure 191 .

 191 Figure 191. Cycling program to quantify the irreversible capacity at different cutoff potentials. Data take from a beechwood derived hard carbon. Capacity was set to cero at each step.

Figure 2 .

 2 Figure 2. Impedance analysis over HC_walnut-shell_1400. The right figure is a magnification of the left figure at high frequencies.

Figure 4 .Figure 5 .

 45 Figure 4. Wide scan for the cycled electrodes of HC_pine-forest-residues_1400.

  Tar species being highly unstable and mostly in the gas phase

	at the pyrolysis temperature, undergo further reactions with increasing
	temperature, and can interact with other volatile species or the biochar
	surface or porosity to produce other compounds. This stage corresponds to
	the secondary reactions phase, and proceed mainly by cracking or
	recombination of the products mentioned above. The formed products
	could keep reacting, forming secondary and tertiary tars, permanent gases
	and char; although secondary char could present different physicochemical
	comprise CO, CO2, CH4, H2, HCN, NH3, or HCNCO, water vapour,
	carbonyls, alcohols, heterocyclics, sugars (levoglucosan), phenolics,
	aromatics with a single ring (benzene, toluene, xylene), and polyaromatic
	hydrocarbons (naphthalene, anthracene or pyrene).

The macromolecules in biomass begin a breakdown by the action of heat from approximately 200°C. The bonds that hold together the lignocellulosic fibres start breaking by the action of fragmentation and depolymerisation reactions. These reactions produce volatile species and unstable tar species. Volatile species -Because of fragmentation and depolymerisation reactions, the solid matrix starts re-arranging into a polycyclic aromatic structure, by means of the formation of benzene rings by intra and intermolecular reactions. The previous phenomena are commonly described as primary reactions.

Table 2

 2

	also

Table 2 .

 2 Pristine hard carbon properties for SIBs application and their evolution with high thermal treatment.

	Property	Values	Behavior	Reference Requirement
			with	for SIBs
			HTT
	Density	1 -1.6 g.cm -3

Table 3

 3 Moreover, a wide range of pyrolysis temperatures was observed, ranging from 600 to 1600°C. Both N2 and Ar gases were identified as pyrolysis atmospheres. The heating rates of pyrolysis varied from 1 to 13 °C.min -1 , while the residence times varied from 30 minutes to 6 hours.

	Regarding the electrochemical characterisations, different set-ups were employed.
	Electrolytes used varied between NaClO4, NaPF6 and NaTFSI salts, in organic
	carbonate based solvents such as Ethylene Carbonate (EC), DiEthylene Carbonate
	(DEC), DiMethyl Carbonate (DMC), or Propylene Carbonate (PC), with or without
	FluoroEthylene Carbonate (FEC) additive. For the electrode slurry, different
	binders were also used, such as PolyVinylidene DiFluoride (PVDF),

, many synthesis conditions have been explored, including pre or post-treatments other than pyrolysis.

Table 3 .

 3 Review of the reported hard carbons derived from different lignocellulosic biomass precursors.

	Ref				[82]									[134	]						[112	]				[142	]
	C e	(at	%)		76,4	84,6	87,1		93,4	78,9	86,4	89,4	93,3								82,4	87,4	92,4	83,4	88,3	92,9
	ID/IG*				1,96	1,32	1,15		1,12	2,13	1,41	1,16	1,10	>				<			1,28	1,10	1,08	1,27	1,09	1,06
	d002	(nm)	*		0,398	0,910	0,387		0,382	0,399	0,395	0,388	0,383								0,397	0,391	0,384	0,399	0,392	0,386	0,392	0,382	0,380	0,376
	SSA	N2	(m 2 .	g -1 )	55,3	45,1	24,5		20,8	369,1	271,2	196,6	92,2	1272,	0			1,0			33,0	19,3	14,5	217,3	130,8	62,1	25,5	68,6	82,8	39,0
	Rate	capability	at 1A.g -1	) (mAh.g -1	80	75	75		55	130	125	155	90	118				50			150	140	90	180	150	110	125	125	100	90
	ICE	(%)			46,1	55,0	60,7		64,3	43,9	50,1	57,5	60,1	27,0				54,0			63,0	70,0	73,0	61,0	68,0	71,0	62,8	63,0	67,7	69,2
	1 st ReV	Cap. d	(mAh.g -	1 )	230,0	219,0	325,0		275,0	267,0	280,0	332,0	294,0	181,0				110,0			206,6	253,4	256,2	205,0	261,1	263,4	262,8	314,5	430,5	359,7
	1 st ReD	Cap. c	(mAh.g -	1 )	498,9	398,2	535,4		427,7	608,2	558,9	532,0	489,2	670,4				203,7			328,0	362,0	351,0	336,0	384,0	371,0	418,6	499,1	635,9	520,0
	Current	density	) (mA.g -1		50									200							50					30
	Binder				PVDF									PVDF							PVDF					PVDF
	Electrolyte				1 M NaClO4	in EC:DEC								1 M NaClO4	in EC:PC						1 M NaClO4	in EC:DEC				1 M NaClO4	in EC:DEC
	HC	yield	(%)		45				25
	Final	Pyr. T	(°C)		600	900	1100		1400	600	900	1100	1400	700				700			800	1100	1400	800	1100	1400	800	1000	1200	1400
	Pyr.	Cond.			5°C/min	under Ar								Pyrolysis	-2h	under N2					Pyrolysis	-5°C/min	5h under	Ar		Pyrolysis	-2h	under N2
	Post-	treatment b			KOH and HCl	wash				KOH and HCl	wash + 3h air	activation at	300°C								KOH and HCl	wash		KOH and HCl	wash + 3h air	activation at	300°C	wash with	diluted HCl	and drying
	Pre-	treatment a												drying +	H3PO4	impregnation	for 24 h	drying and	NO	impregnation	DI f wash +	drying				DI wash and	drying
	Argument	for	utilization		abundance	-	macroscopi	c structure	-low	cellulose	content			abundance	-low cost -	accessibilit	y				abundance	in USA				abundance	in	Southeast	Asia -	waste
	Precursor	-origin			Peat moss	-Canada								Pomelo	peels -	China					Banana	peels -	Canada			Pomelo	peels -	China

Table 4

 4

gathers all the biomass precursor generalities.

Table 4 .

 4 Description of the biomass precursors used in this study.

	Precursor	Type	Description	Origin
	ashwood	wood -hardwood forestry harvesting a	France
	beechwood	wood -hardwood	forestry harvesting	France
	orchard apple	wood -hardwood	AR * -landscape	North Macedonia
			harvesting b
	pine (pinus pinaster)	wood -softwood	forestry harvesting	France
	pine forest residues	wood -softwood	primary residues c	Sweden
	scot pine bark	wood -softwood	secondary residues d	Sweden
	Short Rotation	wood -hardwood	forestry harvesting	France
	Coppice willow
	Short Rotation	wood -hardwood	forestry harvesting	France
	Forestry poplar

*AR: Agricultural residue. a wood products harvested from forest lands or tree plantations. b wood products harvested out of the forest lands. c residues from cultivation and harvesting. d residues from wood processing. e long-term crops.

Table 8 .

 8 Composition of the biomass precursors.

	Measurement	Unit	Pine	Beechwood	Miscanthus	Wheat straw
			Elemental Analysis
	C		51.3 ± 0.51	49.1 ± 0.49	48.3 ± 0.48	45.2 ± 0.45
	H		6.2 ± 0.12	6 ± 0.12	6 ± 0.12	5.7 ± 0.11
	N O a	wmf%	0.3 ± 0.05 40.9 ± 2.04	0.16 ± 0.20 43.8 ± 2.19	0.3 ± 0.05 42.6 ± 2.13	0.8 ± 0.12 39.9 ± 2
	Cl		0.03 ± 0.0045	<0.002	0.05 ± 0.0075	0.2 ± 0.03
	S		0.05 ± 0.0075	0.1 ± 0.015	0.05 ± 0.0075	0.14 ± 0.021
	K		590 ± 30	1123 ± 56.15	3364 ± 168	8823 ± 441
	Ca		2128 ± 106	1496 ± 75	2539 ± 127	4082 ± 204
	Si		2074 ± 104	773 ± 39	7009 ± 350	25226 ± 1261
	P Na	mg .Kg -1	75 ± 4 57 ± 3	141 ± 7 43 ± 2	291 ± 15 56 ± 3	1248 ± 62 326 ± 16
	Mg		597 ± 30	269 ± 13	329 ± 16	735 ± 37
	Fe		288 ± 14	107 ± 5	91 ± 5	297 ± 15
		Immediate and Macromolecular Analysis b
	Moisture (±4%)	%wt	7.3	8.7	6.55	6.6
	Ash (±5%)		1.3	0.8	2.8	8.3
	Extractives		8.4	1.85	8.6	15.7
	Cellulose	wmf%	36.7	44.4	45.7	33.8
	Hemicelluloses		26.1	27.3	22.8	21.7
	Lignin		27.5	26.51	20.2	20.5
			Sugar Monomers Analysis
	Glucose	% of	70.6	63.9	69.9	63.7
	Xylose	total	8.2	29.1	26.1	29.7
	Mannose	sugars	16.9	3.0	0.5	0.8
	Galactose	(normal	2.8	2.0	0.7	1.6
	Arabinose	_ised)	1.6	2.0	2.9	4.2
	Acetyl Groups	wmf%	1.7	8.3	2.6	1.7

a Calculated by difference: C+H+N+O+S+Ash = 100 wmf%. b Cellulose+Hemicellulose+Lignin+Extractives+Ash ≈ 100 wmf%.

Table 9 .

 9 Biochar composition and yield. Hard carbon yield for the 1400°C materials is also reported.

	450°C-	C	H	N	Ash	O*	O** Biochar	Biochar	HC
	Biochar							yield	Yield	Yield***
	Units				wmf%				wmaf%	wmf%
	Pine	78.4 ±	4.3 ±	0.21	4.4 ±	12.7 ±	17.3	29.3 ±	26.3 ±	18.7 ±
		0.78	0.09	±	0.22	0.63	±	1.47	1.32	1.31
				0.03			0.02
	Beechwood	78.8 ±	4.1 ±	0.40	2.6 ±	14.0 ±	17.3	30.5 ±	28.3 ±	22.5 ±
		0.79	0.08	±	0.13	0.7	± 0.2	1.53	1.41	1.58
				0.06
	Miscanthus 74.8 ±	3.8 ±	0.4 ±	8.8 ±	12.2 ±	16.1	31.7 ±	24.5 ±	21.3 ±
		0.75	0.08	0.06	0.44	0.61	± 0.2	1.58	1.23	1.49
	Wheat	63.5 ±	3.4 ±	1.2 ±	23.2	8.6	16.1	35.8 ±	24.1 ±	21.5 ±
	Straw	0.64	0.07	0.18	±	±0.43	± 0.7	1.79	1.20	1.51
					1.16

*By difference. ** Measured. ***Hard carbon yield at 1400°C.

Table 10 .

 10 Chemical composition by EDX analysis of pine and beechwood hard carbons obtained at 1400°C.

	Chemical Element in overall spectra	C, at% O, at% Si, at% Ca, at% K, at%
	HC_pine_1400	87,8	11,9	0,13	0,04	0,12
	HC_beechwood_1400	85,0	14.3	0,11	0,26	0,30

Table 11 .

 11 Summary of the hard carbon properties evolution with thermal treatment between 1000 and 1400°C.

		Pine	Beechwood	Miscanthus	Wheat Straw-1
	Increasing Temperature from 1000 to 1400°C
	C content
	d002	-
	AD1/AG
	N2 SSA
	CO2 SSA
	Non-accessible			-	-
	porosity
		-	-	Spheres (SiO2) +	Spheres (SiO2) +
	Morphology			whiskers formation	whiskers
				(SiC)	formation (SiC)
	Reduction Capacity	-	-	-
	Oxidation Capacity			-
	Plateau/slope
	Capacity
	Average oxidation
	potential

Table 12 .

 12 Classification of the biomass precursors.

	Criteria	Precursor	Code in the text
		Group I
	Type:	orchard apple	GI.1
	Woody related precursors,	pine forest residues	GI.2
	and woody residues.	SRC-Willow	GI.3
	Properties:	SRF-Poplar	GI.4
	-low ash content	vineyard	GI.5
	-high lignin content	pine	GI.6
		beechwood	GI.7
		ashwood	GI.8
		scot pine bark	GI.9
		Group II
	Type:	miscanthus	GII.10
	Grass and cereals, and some	wheat Straw-1	GII.11
	related agricultural residues	wheat Straw-2	GII.12
	Properties:	reed canary grass	GII.13
	-Moderate to high silicon	triticale	GII.14
	content	rice husk	GII.15
	-low lignin content	bagasse	GII.16
		corn husk	GII.17
		switchgrass	GII.18
		Group III

Type: Very diverse: seeds, shells, cob and husk (Only agricultural residues)

		grape seed cake	GIII.19
		sunflower seed shell	GIII.20
		corncob	GIII.21
	Properties:	walnut shell	GIII.22
	-heterogeneous	papaya seed	GIII.23
		coconut husk	GIII.24
		tamarind seed	GIII.25

Table 13 .

 13 Elemental composition of the biomass precursors.

*by difference. **0.2% of error. Measurement standards are presented in chapter III.

Table 14 .

 14 Sugar monomers and macromolecules concentrations. Macromolecular composition of the biomass precursorsLignin, cellulose, hemicelluloses, and extractives contents of all the precursors are reported in Table14. Macromolecules composition of the biomass precursors varied significantly between the samples. GI and GIII groups contained precursors with the highest lignin content ranging from 25 to 63 wmf%. High lignin content is characteristic of woods, and of some agricultural residues such as fruit seeds, shells and husk, as reported by Vassilev et al.[START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF]. Consequently, the highest lignin contents of over 40% were found for GI-9: scots pine bark, GIII-19: grape seed cake and GIII-22: walnut shell. With the available quantification methods, it is not possible to assess the lignin composition of each biomass. However, lignin in woods likely presented a different composition than lignin in the walnut shell or in the grape seed cake precursor. Lignin content in GII precursors was low when compared to the other precursors groups, and ranged between 13 and 25 wmf%,

	Extractives		16.6	18.9	8.0	4.7	10.5	8.4	1.9	10.0	15.2	8.6	15.7	7.8	6.4	25.5	8.8	3.0	9.5	8.5	7.6	10.7	6.9	6.1		27.5
	Hemicelluloses	wmf%	22.2	27.9	22.6	24.1	19.4	27.3	25.3	20.8	18.3	22.6	23.8	24.7	25.7	31.3	22.2	35.9	39.8	36.7	21.9	26.4	34.8	30.6		17.3
	Cellulose		29.1	22.3	44.3	46.8	32.9	38.5	41.3	36.6	21.9	45.3	37.5	38.1	39.5	20.6	36.2	43.9	36.9	31.3	7.8	37.0	38.4	21.1		26.4
	Glucose Xylose Mannose Galactose Arabinose Lignin	% of total sugars (normalised)	59.1 25.0 3.7 5.3 6.8 24.0	58.8 10.7 14.4 8.7 7.5 27.5	69.1 21.8 2.8 3.0 3.4 25.1	72.1 22.3 3.7 0.9 0.9 22.0	64.6 26.2 2.5 3.8 2.9 27.3	70.6 8.2 16.9 2.8 1.6 27.5	63.9 29.1 3.0 2.0 2.0 26.5	69.2 24.9 2.0 1.1 2.7 26.3	64.3 7.9 9.9 7.1 10.9 40.7	69.9 26.1 0.5 0.7 2.9 20.2	63.7 29.7 0.8 1.6 4.2 20.5	61.0 32.5 0.3 1.6 4.6 22.4	61.0 29.9 0.4 3.1 5.6 24.7	39.9 46.7 0.0 5.7 7.8 16.4	64.3 31.3 0.0 1.3 3.0 18.7	62.6 34.0 0.0 0.5 3.0 19.5	51.8 43.0 0.0 0.9 4.2 12.8	49.3 45.9 0.0 3.7 1.2 20.3	33.6 48.2 7.4 5.2 5.6 62.7	59.5 30.7 1.0 2.2 6.6 25.9	52.4 39.4 0.0 2.5 5.7 15.9	46.5 50.0 0.0 2.2 1.4 43.9	unquantified	65.2 29.1 0.0 1.4 4.3 25.9	unquantified
		Units	GI-1	GI-2	GI-3	GI-4	GI-5	GI-6	GI-7	GI-8	GI-9	GII-10	GII-11	GII-12	GII-13	GII-14	GII-15	GII-16	GII-17	GII-18	GIII-19	GIII-20	GIII-21	GIII-22	GIII-23	GIII-24	GIII-25

  2 and GI-9 precursors, with values of 17, 19 and 15 wmf%, respectively. In grasses and cereal species of GII group, the extractives quantification can also include soluble sugars, and starch. The particular high amount of extractives in GII-14: triticale of 26 wmf% is mainly related to starch extraction, as previously reported by Da

Silva et al.

[START_REF] Da | Characterisation of the Most Representative Agricultural and Forestry Biomasses in France for Gasification[END_REF]

, while the other biomasses in this group presented extractives content from 3 to 16 wmf%. It is worth mentioning that the two wheat straws used in this study presented very different extractives values, whereas they presented very similar content of lignin, cellulose, hemicellulose and ash. GII-11: wheat straw-1 presented the double amount of extractives than GII-12: wheat straw-2, 16 to 8 wmf%, respectively. This could suggest that for GII-11, a higher fraction of starch or soluble sugars are recovered from the sample.

Table 15 .

 15 Elemental composition and yield of biochar and hard carbons. *by difference. ** Calculated.

	Hard	Carbon	yield		± 7%	18.5	18.6	19.9	16.9	15.5	18.7	22.5	19.0	21.4	21.3	21.5
	O +	ashes*			4.3	4.4	3.5	0.3	1.8	1.6	2.5	2.7	7.2	6.2	15.4
		S		wmf%	± 15%	0.47	0.16	0.13	0.17	0.14	0.06	0.10	0.34	0.12	0.12	0.27
		N			± 15%	0.10	0.20	0.00	3.30	0.10	0.13	0.13	0.04	0.20	0.16	0.50
		H			± 2%	0.8	0.2	0.5	0.8	0.3	0.3	0.2	0.2	0.3	0.3	0.3
		C			± 1%	94.3	95.0	95.9	95.4	97.7	98.0	97.0	96.7	92.2	93.2	83.5
	Biochar	yield	wmaf%	± 5%	27.3	26.3	29.8	31.7	26.0	26.3	28.3	27.7	35.9	24.6	24.1
	Biochar	yield		± 5%	31.1	28.0	29.7	27.4	29.6	29.3	30.5	31.9	36.4	31.7	35.8
		Ash**				12.2	7.9	7.7	16.8	8.8	4.4	2.6	8.8	7.7	8.8	23.2
		O*				12.0	11.7	12.3	5.9	13.8	12.6	14.0	10.8	12.9	12.1	8.6
		S		wmf%	± 15%	0.07	0.04	0.06	0.05	0.04	0.02	0.08	0.03	0.05	0.03	0.12
		N			± 15%	1.4	0.9	0.8	0.4	1.1	0.2	0.4	0.5	0.5	0.4	1.2
		H			± 2%	3.9	4.1	3.8	3.9	4.3	4.3	4.1	3.9	4.3	3.8	3.4
		C			± 1%	70.4	75.4	75.3	73.0	72.0	78.4	78.8	76.0	74.6	74.8	63.5
					Units	Error	GI-1	GI-2	GI-3	GI-4	GI-5	GI-6	GI-7	GI-8	GI-9	GII-10	GII-11

Table 16 .

 16 N2 and CO2 SSAs and porous volumes of the hard carbon series.

		N2 SSA	CO2 SSA	N2 Total Volume	CO2 Total microporous volume
	Units		m 2 .g -1		cm 3 .g -1
	Error	<1%	<5%	Non-reported
	GI-1	5.0	23.0	0.01	0.01
	GI-2	5.0	12.0	0.01	0.01
	GI-3	3.0	19.0	0.01	0.01
	GI-4	11.0	<1	0.02	-
	GI-5	<5	33.0	0.11	0.01
	GI-6	<5	66.0	0.01	0.01
	GI-7	<5	35.0	0.01	0.01
	GI-8	6.7	5.0	0.01	0.01
	GI-9	7.0	22.0	<0.01	0.01
	GII-10	12.0	29.0	0.02	0.01
	GII-11	42.0	39.0	0.09	0.02
	GII-12	73.0	85.0	0.11	0.04
	GII-13	52.0	70.0	0.06	0.04
	GII-14	50.0	199.0	0.07	0.09
	GII-15	97.0	85.0	0.15	0.04
	GII-16	19.0	338.0	0.02	0.14
	GII-17	14.0	33.0	0.02	0.02
	GII-18	16.0	23.0	0.03	0.01
	GIII-19	7.0	8.0	0.01	<0.01
	GIII-20	6.0	7.0	0.01	<0.01
	GIII-21	8.0	12.0	0.01	0.01
	GIII-22	-	18.0	-	0.01
	GIII-23	10.0	8.0	0.03	<0.01
	GIII-24	14.0	60.0	0.03	0.04
	GIII-25	2.0	32.0	<0.01	0.01

Table 17 .

 17 Average oxidation potentials for reference hard carbons reported in literature.

	Precursor	Pyrolysis Temperature	Average Oxidation potential	Reference
	Commercial Cellulose	1400	0.30	[204]
	Carbotron*	Not reported	0.31	[28], [29]

Table 18 .

 18 Summary of observed correlations steps. A schematic representation of the cycling procedure is given in Figure1of Annex III. To ensure that the hard carbon electrode was completely wet by the electrolyte, 4 hours of rest in an open circuit voltage (OCV)

	Input/Output	1 st	Ox.	Ox. P at	Ox. Cap.	Plateau	Slope	C
		Red.	Cap. at	C/10	At C	Cap. At	Cap. at	Cycling
		Cap	C/10			C/10	C/10
	Lignin	//	//	//	//	//	//	//
	Cellulose	//	//	//	//	//	//	//

Table 19 .

 19 Summary of the main physicochemical parameters of the selected hard carbons and their percentage of irreversible lost per cut-off potential.

	Irreversible capacity per cut-
	off voltage

Table 21 .

 21 Concentration of each chemical specie derived from the spectra deconvolution shown in Figure74.

			C
		C sp 2	sp 3a	C-O	C=O	OCO	K	-CO3	Carbides	Other	-CF2 b
							K 3/2
	Binding energy (eV)	284,2	284,9	286	287.4	288.7	291.7 K 1/2	289.8	283.2	283.4	290.6
							293.7
							(at% ± 0.1)
	HC_pine-forest-
	residues_1400_	49,3	11,4	6,4	4,4	3,5	9,7	9,4	//	5,9	-
	powder
	HC_rice-
	husk_1400_	52,7	11,8	8,2	4,3	3,5	4,9	6,2	8,4	//	-
	powder
	HC_pine-forest-
	resdues_1400_	44,0	17,8	8,0	3,9	6,1	3,7	5,0	//	6,0	6,6
	electrode
	HC_rice-
	husk_1400 _	49,4	15,8	9,0	4,3	2,5	4,2	5,6	3,2	//	6,1
	electrode
	a for the composite reference electrode, this peak includes the contribution of the CH2 signal from
	the PVDF binder

b -CF2 * is only quantify for the electrodes, and is originated by the PVDF contribution.

Table 1 .

 1 Electrochemical performance parameters for C/10 cycling. The reported values correspond to the mean of the two coin-cells results. STD stands for the standard deviation of the observations.

	STD		0.00	0.00	0.00	0.01	0.00	0.00	0.00	0.00	0.00	0.01	0.02	0.01	0.01	0.00	0.01	0.00	0.04	0.01	0.00	0.00	0.01	0.00	0.00	0.00
	5 th Av	Ox Pot.	V	0.2	0.2	0.2	0.2	0.2	0.3	0.3	0.2	0.2	0.3	0.4	0.3	0.3	0.3	0.4	0.3	0.3	0.3	0.2	0.2	0.3	0.3	0.2	0.3
	STD		1.48	4.23	2.18	0.52	0.80	5.33	3.31	0.43	0.00	11.63	16.05	1.91	4.42	1.08	1.18	0.58	3.46	0.14	2.62	4.22	0.47	3.84	2.80	3.04
	5 th Ox	Cap	mAh.g -1	222.9	297.0	291.4	274.5	271.3	314.4	302.4	268.7	297.6	251.2	208.5	232.6	243.1	237.9	171.2	283.9	278.0	259.4	262.1	296.6	291.8	310.7	256.3	291.3
	STD		1.75	3.46	1.46	0.24	0.06	1.70	0.47	1.18	4.06	5.83	2.04	5.51	4.79	0.42	1.78	0.78	1.02	0.84	1.15	4.41	0.55	3.93	4.35	3.81
	1 st Ox	Cap	mAh.g -1	225.9	299.1	291.4	273.6	265.3	314.7	298.1	260.3	302.1	254.6	200.3	231.8	243.8	241.5	170.3	283.5	279.9	262.1	263.3	297.9	293.1	310.6	255.8	291.0
	STD		2.95	3.40	4.50	0.86	2.90	11.44	17.02	7.81	9.21	0.27	6.43	6.80	5.68	0.62	0.22	0.22	1.00	11.26	1.20	7.65	5.73	4.71	5.26	5.11
	1st Red	Cap	mAh.g -1	322.2	363.9	356.5	365.2	359.0	375.6	352.6	297.3	377.3	343.6	334.7	383.1	352.4	379.3	356.4	374.0	369.7	361.6	332.9	376.8	375.6	380.4	359.1	381.2
	STD		0.29	1.57	2.27	1.62	0.32	3.69	5.95	3.52	2.75	1.86	0.19	0.95	0.75	0.01	0.24	0.02	0.14	2.97	0.20	3.36	0.93	0.03	0.05	0.15
	5 th CE	%	99.5	99.2	104.3	98.5	99.2	104.3	107.2	116.8	98.4	102.4	101.1	97.7	100.3	97.8	98.4	99.8	99.2	98.5	99.6	98.6	98.3	99.7	99.1	99.3
	STD		0.10	1.72	1.44	0.11	0.58	3.01	4.22	1.90	3.03	1.64	1.76	0.36	0.59	0.01	0.53	0.17	0.48	2.49	0.06	2.78	1.34	0.02	0.17	0.03
	ICE	%	70.1	82.2	81.7	74.9	73.9	83.8	84.7	87.6	80.1	74.1	59.9	60.5	69.2	63.7	47.8	75.8	75.7	72.5	79.1	79.1	78.0	81.6	71.2	76.4
			Units	GI-1	GI-2	GI-3	GI-4	GI-5	GI-6	GI-7	GI-8	GI-9	GII-10	GII-11	GII-12	GII-13	GII-14	GII-15	GII-16	GII-17	GII-18	GIII-19	GIII-20	GIII-21	GIII-22	GIII-23	GIII-24

Table 2 .

 2 Electrochemical performance parameters for C cycling. The reported values correspond to the mean of the two coincells results. STD stands for the standard deviation of the observations.

	STD		0.01	0.02	0.01	0.02	0.02				0.01			0.03	0.02	0.05	0.05	0.00	0.05	0.01	0.02	0.01	0.01	0.01	0.01
	10 th Av Ox	Pot	V	0.3	0.4	0.3	0.3	0.3				0.3			0.4	0.3	0.5	0.5	0.3	0.4	0.4	0.3	0.3	0.3	0.4	0.3
	STD		1.11	20.72	0.21	1.04	1.67				1.47			2.91	5.49	8.03	2.64	0.76	10.07	0.35	6.50	3.45	0.94	2.85	3.54
	10 th Ox	Cap	mAh.g -1	200.0	226.2	273.9	264.9	262.7				286.9			221.2	227.9	154.3	157.5	270.1	249.5	191.0	234.0	281.0	278.5	273.1	241.1
	STD		1.42	9.68	0.37	1.23	1.63				3.26			5.46	5.34	0.20	2.59	0.23	13.38	8.01	9.02	2.48	0.97	1.38	4.02
	6 th Ox	Cap	mAh.g -1	199.4	268.6	274.8	265.4	250.9				287.7			221.1	231.0	200.6	155.1	272.5	247.3	225.8	230.4	280.9	279.7	278.9	242.8
	STD		0.91	8.37	2.37	0.10	1.91				0.15			6.38	5.03	0.24	0.59	1.27	10.06	8.74	5.17	5.11	0.89	2.04	4.10
	6 th Red	Cap	mAh.g -1	216.3	281.0	282.9	279.2	258.3				300.2			234.4	241.5	216.0	167.2	281.1	261.1	239.6	246.4	293.7	294.9	289.8	256.7
	STD		0.86	0.30	0.59	0.26	0.09				0.99			0.31	0.14	0.32	1.40	0.36	0.11	1.29	1.65	1.09	0.09	0.13	0.06
	10 th CE	%	101.7	99.9	100.8	99.0	101.1				99.6			99.1	100.1	99.2	100.3	100.9	100.2	99.5	99.3	99.3	99.5	100.3	99.8
	STD		1.04	0.60	0.68	0.41	0.09				1.04			0.23	0.22	0.01	1.22	0.52	1.47	0.09	1.70	0.82	0.61	0.20	0.06
	6 th CE	%	92.2	95.6	97.1	95.1	97.1				95.8			94.3	95.7	92.9	92.7	96.9	94.7	94.3	93.5	95.6	94.8	96.2	94.6
			Units	GI-1	GI-2	GI-3	GI-4	GI-5	GI-6	GI-7	GI-8	GI-9	GII-10	GII-11	GII-12	GII-13	GII-14	GII-15	GII-16	GII-17	GII-18	GIII-19	GIII-20	GIII-21	GIII-22	GIII-23

Table 322 .

 322 Calculated electrochemical performance parameters for C cycling. The reported values correspond to the mean of the two coin-cells results. STD stands for the standard deviation of the observations.

Table 1 .

 1 Deconvolution results and concentration of the C 1s, F 1s and Na 1s spectra for the pine forest residues and riche husk cycled hard carbons.

	Total							99.5			99.6			99.7	91.01
		Na	P.D.-1	1071.5-	1072	1.5-2		2.3			4.8			4.8	2.3
	Na 1s	NaF	1070.3-	1071	1.3-2		7.7			8.3			6.6	15.8
		Na in	HC	1069.3-	1069.6	1.7-2.1		24.6			18			26.7	8.3
	F 1s	F D. NaPF6 P.-1 F D. P.-2	683.4-686.5-684.5-	683.6 687.1 685.2	1.6-1.7 1.96-2.4 1.5-2.5		4.6 2.2 2.7			7.0 3.1 4.2			6.8 1.7 3.4	6 4.3 5.0
		C-O NaF	286.5-682.2-	287 682.3	1.59 1.5-1.6	at% (±0.1)	3.4 12.4			2.6 5.6			2.1 7.9	2.3 4.3
		CO3	289.78	1.84		1.5			1.5			0.6	1.2
	C 1s	C-H in C=O SEI	285.36 288.2-288.4	1.59 2.09		7.0 3.7			8.4 3.6			6.6 3	1.1 3
		C-C/C-	H HC	284.3-	284.4	1.30		15.8			25.3			20	28.5
		Na-	HCx	282.2-	282.9	1.5-1.8		11.7			7.3			9.4	9
			Binding energy	(eV)	FWHM (eV)	HC_pine-forest-	residues-	_1400_0.3V	HC_pine-forest-	residues-	_1400_0.01V	HC_pine-forest-	residues-	_1400_0.01V +	Floating	HC_rice-	husk_1400_0.3V

and ICEs as low as 30% were reported. These results evidenced that the high values of reduction capacity are related to irreversible Na storage mechanisms, such as an extended formation of the SEI over the elevated SSA.

IHE Delft Institute for Water Education, Department of Environmental Engineering and Water Technology, Wesvest 7, 2611 AX Delft, The Netherlands.

International Centre for Diffraction Data.

Reduction and oxidation terms will be used from now on, instead of discharge and charge terms, respectively, unless otherwise indicated, to avoid terminology confusion.

Calculated as the ratio between the difference in the carbon (or oxygen) content of each biochar sample and the carbon (or oxygen) content of the biomass sample, over the carbon (or oxygen) content of the biomass sample.

), and pine hard carbon showed neither a lower reversible capacity, nor a significant increase of its reversible capacity with temperature. Therefore, for

The overpotential refers to the experimental potential at which an electrode undergoes oxidation/reduction, which normally differs from the theoretical thermodynamic potential.

Abbreviation List

BET

Brunauer-Emmett-Teller Model CO2 SSA and pore volume can also be explained by the coalescence of the micropores.

Is it worth noting that a wide range of CO2 SSA values was observed for the different hard carbons obtained at 1000°C. The HC_miscanthus_1000 sample presented the highest CO2 SSA of all the samples, with 438 m 2 .g -1 . This elevated value was a result of the combination of the high holocelluloses content of miscanthus and to a lesser extent, to the intrinsic activation triggered by K compounds. A high content in holocelluloses will produce higher amounts of volatiles, which can lead to microporous materials [START_REF] Tripathi | Effect of process parameters on production of biochar from biomass waste through pyrolysis: A review[END_REF].

HC_pine_1000 exhibited again higher CO2 SSA values than HC_beechwood_1000 with values of 403 and 338 m 2 .g -1 , respectively. This difference might confirm that pine developed higher microporosity than beechwood because of the different arrangement of the macromolecules in the cell wall of the raw biomass.

HC_wheatstraw_1000 presented the lowest CO2 SSA value of 257 m 2 .g -1 . This lower value with respect to the HC_miscanthus_1000 could be explained by the lowest holocelluloses content of wheat straw with respect to miscanthus, of 55.5 wmf% vs 68.5 wmf%, respectively. Moreover, the lowest CO2 SSA of HC_wheatstraw_1000 would suggest that the porosity generated by an intrinsic activation of K compounds did not produce microporosity at 1000°C.

At the highest temperature of 1400°C, the CO2 SSA values are higher than the N2 SSA values for the pine, beechwood and miscanthus derived hard carbons, thus suggesting that their structure at this temperature presented a large quantity of small inaccessible ultramicropores. HC_wheat-straw_1400 exhibited rather similar values of N2 and CO2 SSAs, while the N2 pore volume was higher than the CO2 microporous volume (0.085 cm 3 .g -1 for N2 and 0.016 cm 3 .g -1 for CO2). These values suggested that HC_wheat-straw_1400 was the only sample presenting a significant amount of accessible pores, when compared to the other hard carbons obtained at the same temperature.

The pore size distributions (PSD) for all the samples are presented in Figures 2 and3 of Annex I. The average pore size did not change significantly with the

Inorganic composition and ash content

The main inorganic elements present in the 25 precursors were quantified and are reported in Table 13, as well as their ash content.

The ash content was in general, lower for the GI precursors, ranging from 1 to 5 wmf%. Low ash content is expected in wood species [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF], [START_REF] Da | Characterisation of the Most Representative Agricultural and Forestry Biomasses in France for Gasification[END_REF]. However, bark in woody biomasses is known to present higher ash contents [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF]. GII group gathers precursors with the highest content of ash, with GII-15: rice husk having 15 wmf%, followed by the 8 wmf% content in the wheat straw studied in the precedent chapter. Agricultural residues from grass or cereals are expected to have high values of ash [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF], [START_REF] Da | Characterisation of the Most Representative Agricultural and Forestry Biomasses in France for Gasification[END_REF]. However, some of the residues of GII group presented ash values very close of the GI precursors, such as GII-16: bagasse and GII-17:

cornhusk, with only 1 and 2 wmf%, respectively. GIII group presented ash content from 1 to 7 wmf%. As for GII, some of the agricultural residues presented very low ash values such as in the case of GIII-22: walnut shell, and GIII-21: corncob, presenting only 1 and 2 wmf%, respectively.

A higher ash fraction in the precursor was observed to yield in lower C contents in the final hard carbon, therefore, lower conductivity, and potentially, lower Na uptake. The influence of the ash content in biochar and hard carbon formation needs to take into account the inorganic fraction composition that forms the ash, which is discussed below.

Each group of precursors seemed to have one major inorganic element. For GI precursors, the most abundant inorganic element was Ca, ranging from 1496 to 13000 mg.Kg -1 . Inorganic fraction in woody precursors is expected to contain mainly Ca compounds [START_REF] Vassilev | An overview of the organic and inorganic phase composition of biomass[END_REF], [START_REF] Vassilev | Ash contents and ash-forming elements of biomass and their significance for solid biofuel combustion[END_REF], which is confirmed by our observations. For GI-2:

pine forest residues, Si and Ca contents were very close, 3936 mg.Kg -1 of Si vs. 3927 mg.Kg -1 of Ca, and probably not significantly different. K seemed to be the next most abundant inorganic element in GI group.

In GII group, the predominant inorganic element was Si, which varied from 3494 to 63955 mg.Kg -1 , except for GII-13: reed canary grass, presenting a higher amount of K than of Si, with 14705 mg.Kg -1 of K vs. 11380 mg.Kg -1 of Si. The second most Galactose and arabinose contents were lower than xylose and glucose contents for all the precursors, but higher than those of mannose for some of the precursors.

Galactose ranged from 0.5 to 9 wmf% for all the precursors and was higher than mannose for the landscape woods, and all the GII and the GIII precursors, with the exception of GIII-19. Arabinose ranged from 0.9 to 10 wmf% and was, overall higher than galactose for the GII and GIII precursors.

Where n is the final cycle until which the half-cell was able to keep cycling, if the 95 th cycle was not reached.

The electrochemical parameters described before are reported in Table 1, Table 2 and Table 3 of Annex II, for all the hard carbons.

All the previous values were reported as the mean value of the two half-cells measurements with their respective standard deviation (STD). An error bar of 10 mAh.g -1 was assigned for the capacity values in the subsequent analysis, calculated from the maximum capacity variation between all the half-cells.

General observations

At the first cycle, all the obtained hard carbons exhibited the typical reductionoxidation profile as seen from Figure 54 Thus, the oxidation capacity quantifies the remaining Na ions that are actually capable of keeping cycling, and it is also referred to as the reversible capacity of the electrode.

The formation of the SEI layer is normally completed and stabilised after the first 5 to 10 cycles. The green profile in Figure 54, Figure 55, and Figure 56, corresponds to the 5 th cycle of the electrodes at a C/10 regime. Notice that after five cycles, the electrodes did not seem to lose any more capacity between reduction and oxidation, as there is no more hysteresis between the curves. This behaviour confirms the stabilization of the SEI in the electrodes.

Analysis of the C rate cycling

Increasing the rate during cycling implies that the hard carbon electrodes need to accommodate and release the Na ions faster during the reduction and the oxidation steps. Consequently, the mechanisms related to the Na ions storage in the hard carbon may not have time to be completed to the same extent as for a slower rate of C/10. This hypothesis is confirmed by the decrease of the reduction and the oxidation capacity of the materials observed in Figure 54, Figure 55, and Figure 56 and reported in Table 2 of Annex II.

Notice that the increase of the cycling rate induced an increase of the polarization between the reduction and the oxidation profiles of the electrodes. This can be seen from the red profile of the tenth cycle in Figure 54, Figure 55, and Figure 56, for all the hard carbons. This increase of polarization arises from the higher resistance of the system at higher current rates, decreasing the delivered capacity.

The capacity decay was also seen in the decrease of the CE from the fifth to the sixth cycle. The CE of all the hard carbons was above 97 % at the fifth cycle at C/10 rate, however, for the sixth cycle it decreased, taking values from 92,2 to 97,1%.

Additionally, at higher current rates, the overpotential 9

DERIVED FROM BIOMASS PRECURS ORS

This chapter aims for the understanding of the phenomena at the origin of the electrochemical irreversibility of Na ions insertion/de-insertion in hard carbons derived from biomass. In order to identify these phenomena, selected hard carbon samples were used to perform iterative cycling until different cut-off potentials and therefore, to quantify the irreversible loss at each potential. Furthermore, an ex-situ XPS analysis was performed over the surface of two cycled electrodes until different cut-off potentials, to study its composition, and correlate it to the irreversible loss.

Irreversible loss as a function of the depth of discharge

We evidenced in the previous chapter that in between the 25 hard carbon materials evaluated, the Initial Coulumbic Efficiency (ICE) varied from 40% to more than 80%. Hard carbons exhibiting higher C contents, lower accessible surface area, no evident graphitisation and no evident formation of inert phases such as SiC, presented the highest ICEs.

The irreversible loss is commonly attributed to the formation of the Solid Electrolyte Interface (SEI) during the first reductions of Na-ions (herein, discharges) in the hard carbon electrode. The SEI is normally more important for larger accessible surface area materials, leading to lower ICEs [START_REF] Dou | Hard carbons for sodium-ion batteries: Structure, analysis, sustainability, and electrochemistry[END_REF]. However, it is also possible that Na-ions remain trapped inside the hard carbon structure, surface and defects sites after discharge [START_REF] Lotfabad | Origin of non-SEI related coulombic efficiency loss in carbons tested against Na and Li[END_REF].

To quantify the capacity that the hard carbons lost at given cut-off potentials of the first cycle, half-cells containing selected hard carbon electrodes were assembled and cycled under a particular program. Each half-cell underwent an iterative reduction until cut-off potentials of 0.7, 0.3, and 0.01 V. A fourth stage was added, consisting on a floating at 0.01 V. After each cut-off potential, the half-cells underwent oxidation until 1.5 V. A relaxation of 30 minutes was added after each During the floating step, at a constant potential, the reduction of the electrolyte species is not favoured, and the SEI formation is thought to be stable. Therefore, we suggest that the irreversible loss produced during the floating stage, was rather related to structural trapping or adsorption on the defects of the Na ions. The trapping of Na ions at low potentials has also been observed by Lotfabad and coworkers [START_REF] Lotfabad | Origin of non-SEI related coulombic efficiency loss in carbons tested against Na and Li[END_REF], using ex situ XRD and XPS measurements.

The discussed observations strongly support the fact that the SEI formation is not the only mechanism at the origin of the irreversible loss on hard carbons. The structural trapping apparently promoted at low potentials, is also responsible for the lower ICE. While adsorption of the Na ions in the defects could also contribute to the irreversible loss, it is not possible to identify its effect by galvanostatic tests, as it probably occurs at the same potentials as the SEI formation. From this point of view, we performed ex-situ XPS to go deeper in the understanding of this phenomenon.

ex-situ XPS analysis of selected cycled hard carbon electrodes.

XPS analysis was performed to probe the surface chemical structure change (depth analyses of XPS is about 10 nm) of HC_pine-fores-residues_1400 and HC_rice-husk_1400 samples during discharge. The latter samples exhibited the fastest and the slowest irreversible formation, respectively. Hence, the XPS pave the way to probe the SEI composition at different cut-off voltages, corresponding to different discharge states. XPS analyses could also provide insights regarding the structural trapping of Na ions, as the Na-C interactions can be also identified through a careful analysis of the C and the Na 1s core level spectra.

SEI formation model

Before presenting the XPS analysis, it is important to discuss the existing insights regarding the formation of the SEI layer. Studies concerning the SEI formation on SIBs are very limited, and are based on the knowledge gained from the Li-ion batteries (LIBs). It is thought that the SEI formation process in SIBs is analogous to those on LIBs. However, the different chemistry of the Na and Li anions, may lead to different specific phenomena in the SEI formation. Eshetu et al. [START_REF] Eshetu | Impact of the electrolyte salt anion on the solid electrolyte interphase formation in sodium ion batteries[END_REF] highlighted in particular that the higher redox potential, the lower Lewis acidity, the lower desolvation energy, the larger ionic radius, the larger coordination shell, and the higher water reactivity of Na with respect to Li anion, should be considered. Komaba et al. [START_REF] Komaba | Electrochemical Na Insertion and Solid Electrolyte Interphase for Hard-Carbon Electrodes and Application to Na-Ion Batteries[END_REF] confirmed that the composition of the SEI differed in the proportion of the inorganic and organic species between the Na and the Li studied systems.

However, as no model has been proposed for the SEI formation in SIBs, here, the accepted model for LIBs will be used as a reference.

The SEI is the result of the instability of the electrolyte components at a given potential. Typically, electrolyte solutions consist of a Na salt in polar aprotic organic solvents, such as alkyl carbonates [START_REF] Moshkovich | Investigation of the Electrochemical Windows of Aprotic Alkali Metal (Li, Na, K) Salt Solutions[END_REF]. These solutions have a stability window, which is normally between 0.5V and 4.5V vs. Na + /Na, depending of the electrode material and the electrolyte composition and concentration [START_REF] Carboni | Analysis of the Solid Electrolyte Interphase on Hard Carbon Electrodes in Sodium-Ion Batteries[END_REF], [START_REF] Muñoz-Márquez | Na-Ion Batteries for Large Scale Applications: A Review on Anode Materials and Solid Electrolyte Interphase Formation[END_REF], [START_REF] Bhide | Electrochemical stability of non-aqueous electrolytes for sodium-ion batteries and their compatibility with Na 0.7 CoO 2[END_REF]. Therefore, when the system undergoes reduction outside this window, the electrolyte will decompose, forming new products that constitute the so-called SEI layer. Initially, it was thought that the decomposition products deposited over the electrode surface, forming a single uniform layer [START_REF] Peled | The Electrochemical Behavior of Alkali and Alkaline Earth Metals in Nonaqueous Battery Systems-The Solid Electrolyte Interphase Model[END_REF]. Then, it was proposed that the SEI layer is composed of two thin layers, one compact near the electrode and another porous near the electrolyte [START_REF] Peled | Film forming reaction at the lithium/electrolyte interface[END_REF]. Peled et al. [START_REF] Peled | Advanced Model for Solid Electrolyte Interphase Electrodes in Liquid and Polymer Electrolytes[END_REF] have demonstrated that the electrolyte degradation products are more favourable to precipitate, forming a mosaic aggregation or microphases, as depicted in Figure 72. This is known as the mosaic SEI model and will be used for the analysis presented here.

ex-situ analysis

HC_pine-forest-residues_1400 and HC_rice-husk_1400 electrodes were used to prepare half-cells. Three half-cells were prepared per sample that were discharged until different cut-off voltages each. The chosen cut-off voltages were of 0.3 V, 0.01 V and 0.01 V with further floating. After cycling, the half-cells were disassembled in an glove box, washed with dimethyl carbonate (DMC) and transferred to the XPS device into a transfer case, to avoid oxidation of the surface of the electrodes (details are described in Chapter III). The galvanostatic cycling program and the XPS spectra for the C 1s, F 1s, and Na 1s are presented in Figure 75 and Figure 76, for the electrodes of HC-pine-forest-residues_1400, and HC_rice_husk_1400, respectively. The electrodes were analysed after oxidation of the half-cell, at the point that is indicated in Figure 75 and Figure 76 with the red arrow. The atomic concentration of all the deconvoluted species is gathered in Table 1 of Annex III.

Commonly the organic products on the SEI are investigated from the C 1s spectra, while the inorganic products are analysed from the F 1s, Na 1s and P 2p spectra.

P was found to be in very low concentrations (under 0.5 at%) therefore, the resolution of the spectra did not allow further interpretation. Only F and Na spectra were used for the analysis of the inorganic fraction of the SEI.

HC_pine-forest-residues_1400 sample

The survey spectra for the HC_pine-forest-residues_1400 cycled electrodes are reported in Figure 4 of Annex III. All the cycled electrodes presented intense signals of Na 1s, and F 1s, as well as less intense signals for C 1s, O 1s and P 2p.

The K and Ca signals are barely detectable. These observations confirmed that there had been a modification of the surface composition after discharge.

Moreover, the XPS observation showed that the SEI recovers the SiC whiskers of the rice husk electrodes. Additionally, it was confirmed that there is no insertion or adsorption of the Na ions into the SiC whiskers.

GENERAL CONCLUSION AND FUTURE WORK OUTLOOKS 1. General Conclusion

The main objective of this research work was to identify the role of biomass composition on hard carbon properties and performance in Na-ion batteries. To achieve this objective, three axes were followed.

First, the final pyrolysis temperature was studied between 1000 and 1400°C on four-selected biomass precursors in order to identify the main parameters affecting the hard carbon formation.

It was observed that the final pyrolysis temperature had a stronger effect on the composition, morphology and structure of the hard carbons derived from miscanthus and wheat straw, than from pine and beechwood. This behaviour was related to the lower ash content of the latter precursors, which led to a direct hard carbon formation, with no parasitic phenomena observed. For the miscanthus and wheat straw derived hard carbons, activation was detected at 1000°C, while graphitisation and SiC formation were detected at 1400°C. These phenomena were related to the high content of K, Ca and Si elements in the biomass precursors.

Given the high C purity, turbostratic domains concentration and low N2 SSA, pine and beechwood derived hard carbons exhibited competitive electrochemical performances, with reversible capacities between 282 and 315 mAh.g -1 at 37.2 mA.g -1, between 1000 and 1400°C. The ICE of these materials was between 79 and 85%. On the other side, miscanthus and wheat straw derived hard carbons showed lower electrochemical performances at all temperatures, with reversible capacities between 189 and 256 mAh.g -1 and ICEs of 60 to 74%. Because of their lower C content, the presence of SiC and graphitic phases and the higher N2 SSA.

To confirm the previous observations, hard carbons were further synthesised from 25 different biomass precursors at 1400°C to analyse the ANNEX II. Complementary results for Chapter V.