Laura C Loaiza

Elodie Salager

Nicolas Louvain

N Dupré

C Davoisne

L Madec

L Monconduit

Complex mechanism of lithiation and sodiation in Siloxene and Germanane: Promising electrodes for batteries. In preparation.

behind an improved performance and the failure mechanism. Thus, in this work we have studied the electrochemical lithiation mechanism of the Ge and the Si0.5Ge0.5 alloy, some of the lithiated intermediates were identified and the impact of the cycling conditions (morphology, composition, C/rate, electrode formulation) was assessed. Particular attention was given to the role of the solid solution over the improved performance. In addition, the electrochemical mechanism of the layered compounds was explored in an attempt to identify the intermediates and establish the differences with their bulk counterparts.

Finally, the results presented in this thesis set ground for future studies, particularly regarding the use of layered materials derived from Zintl phases for battery applications, which is a novel topic in the literature. The possibilities for finding new electrode materials within the Zintl family is vast, several other elements present layered phases (Sn, Sb, Bi, Ar, P), all of them with potential applications in batteries and in other fields like optoelectronics or sensors. As previously mentioned, the understanding of the processes taking place during cycling is vital to really identify the performance of a given material. The results presented here set the bases for these type of studies in layered Si-and Ge-based materials and complement the existing literature for the electrochemical mechanism of Si, Ge and Si1-xGex alloys. All this knowledge is crucial to effectively design new electrode materials to satisfy the increasing demand in energy storage devices.

Table of Contents

General Introduction

Nowadays, we are facing a climate crisis with the effects of global warming becoming more and more evident. Most of the the green-house effect gases that cause the global warming are originated from the human activity, like agriculture, transport and energy. These two last ones have been mainly based in the consumption of fossil fuels, a limited source, responsible for the release of enormous quantities of CO2 and other green-house effect gases into the atmosphere. In order to reduce these emissions, there must be a change in the way the energy is produced and stored. The increased awareness on the climate change, has led to the establishment of new policies for the energetic transition from fossil fuels to alternative energies, comprising the use of solar cells, windmills, geothermal, among others. These sources of energies are renewable and can be adopted at a worldwide scale. Nonetheless their energy production is intermittent. Therefore, a device of energy storage is necessary, and batteries play a key role in this context. They are used in stationary applications to supply the demand of energy when the production is low (peak shaving). For this specific application, the devices are not constrained by size, and several technologies (NaS, redox flow and other Li-ion based chemistries) with different specific energy densities and cycle life have been implemented at relatively low cost.

On the other hand, the transportation sector is expected to move from fossil fuel based to electric in the upcoming years. Here batteries are among the most promising candidates to participate in this substitution and they have started to be implemented in different transportation media, like trucks, buses or personal cars. The market of electric cars is led by lithium-ion batteries (LIB), with chemistries frequently based on Nickel-Manganese-Cobalt oxides (NMC), Nickel-Cobalt-Aluminum oxide (NCA), lithium iron phosphate (LFP), among others, coupled with a graphite anode. [START_REF] Tsiropoulos | Li-ion batteries for mobility and stationary storage applications[END_REF] Other lithium-ion chemistries or even sodium-ion batteries (NIB) are being considered also.

The number of devices that require an energy storage device is increasing exponentially, in order to satisfy this demand one single technology will not be enough. For example, if all the cars that are in circulation are transformed into LIB-based electric ones, the sources for producing the batteries would be limited. Several different technologies must be implemented considering each specific application. Such technologies might include fuel cells, hydrogen storage and batteries. Among these last ones, other systems besides LIB must be explored, such as sodium, potassium, calcium or aluminum-based batteries as well as redox flow systems. Additionally, studies on the current technologies in order to improve their efficiency, cycle life and address the failure mechanisms are being performed. The quest is open then, for the search of new electrode materials for different systems, Li, Na, K, etc. that provide higher energy densities with longer cycle life.

In this context, silicon and germanium appear as good candidates, due to their high theoretical capacities. [START_REF] Tsiropoulos | Li-ion batteries for mobility and stationary storage applications[END_REF] Each element presents its advantages, silicon is abundant on the earth crust and is relatively cheap, germanium is more expensive but has an improved Li ion diffusion and electric conductivity compared to Si. Although, prior to the implementation of these materials at an industrial level, several challenges must be addressed. The high delivered capacities come at the expense of a volume expansion and contraction upon alkali insertion and deinsertion. These volume changes in the Si and Ge particles cause; i) particle pulverization, isolation from the conductive network and detachment from the current collector, ii) the cracks expose new surface to the electrolyte, causing an excessive and uncontrolled formation of SEI layer, iii) the significant SEI layer increases the resistance of the electrode and irreversibly trapping the Li ions available. [START_REF] Feng | [END_REF][3][4] At last, the cell experiences capacity fade. The volume expansion is an inherent process of the alloying reaction and cannot be avoided, therefore the strategies to avoid the failure of the cells have been reported in the literature. These works include the improvement of the contact between particles by the introduction of binder that can alleviate the volume expansion and maintain the structural stability. The reducing of the particle size to the nanoscale and the design of specific architectures that reduces the mechanical stress along a specific orientation and avoids the cracking. [5][6][7] The implementation of special electrolyte formulations that stabilize the SEI layer. [7,8] The design of effective networks made of composites type Si(Ge)-graphite/graphene. The use of alloys that benefit from the superior desired properties of two or more elements. The use of 2D or layered structures that provide sufficient interlayer space to accommodate the volume changes while providing fast transport kinetics. [9] In this work, two of these approaches are considered, the study of the Si1-xGex alloys and the use of a layered morphology. In the first one, the formation of the Si1-xGex solid solution improves the capacity retention properties and the electronic conductivity. In the second one, the use of a layered Si-and Ge-based material (Siloxene and germanane), derived from the CaSi2 and CaGe2 Zintl phases buffers the volume changes and improves the kinetics of the system. For the layered siloxene and germanane, the electrochemical performance has been evaluated with Li, Na and K. Additionally, given the high versatility of these phases, the siloxene interlayer distance was expanded by the intercalation of organic molecules of various carbon chain lengths and its influence over the electrochemical performance was evaluated. Lastly, considering the synergy between Si and Ge, the layered Si1-xGex, the siligane, was prepared and tested with Li and Na.

Nonetheless, finding a material and studying its electrochemical behavior is only one of the aspects.

The fundamental study of their electrochemical mechanism is crucial to understand the reasons

Thesis outline

The aim of this thesis is to propose two different approaches to address the challenges associated with the Si and Ge electrodes; by the implementation of a solid solution type Si1-xGex that improves the capacity retention and the electronic conductivity, and by the use of a layered morphology that buffers the volume variations. Additionally, the electrochemical mechanism for Ge, Si1-xGex and the layered siloxene and germanane was studied, with the objective of shed light to the different processes behind the improved performance of these materials. These main aspects are discussed as follows:

Chapter I is dedicated to the state of the art. The different challenges concerning the use of Si and Ge as anode materials are presented as well as the strategies to address them. Following, the Zintl family is introduced by briefly reviewing the Zintl-Klemm concept. The different Zintl phases involved in this work are presented, LixM, NaxM, CaM2 with M=Si,Ge. Finally, the Xnes phases, siloxene and germanane (derived from CaM2) and their modification are covered.

Chapter II introduces the alloying electrochemical lithiation mechanism of Si based on the reviews in the literature. Following, the alloying electrochemical lithiation mechanism of Ge and Si0.5Ge0.5 is studied, certain lithiated intermediates are identified and the influence of the cycling conditions over the electrochemical behavior is considered. The sodiation and potassiation mechanism of Si and Ge are not contemplated in this study due to the limited experimental information available.

Chapter III is devoted to the synthesis and characterization of the layered siloxene and germanane, derived from the CaSi2 and CaGe2 Zintl phases. Their electrochemical performance with Li, Na and K is explored. The modification of siloxene by the intercalation of organic molecules and its impact over the electrochemistry is evaluated. Finally, the layered Si1-xGex are prepared from the Ca(Si1-xGex)2 phases, the material is characterized and its electrochemical behavior with Li and Na is studied.

Chapter IV is dedicated to the exploration of the electrochemical lithiation mechanism of the layered siloxene and germanane. The lithiated intermediates are studied and the results are compared with the information obtained in Chapter II for the lithiation of Si and Ge.

This manuscript ends by the Conclusion that summarizes the main findings of this research and gives some perspectives for future works.

Chapter I. State of the Art 1 Li, Na and K-ion batteries 1.1 Fundamentals of batteries

Lithium ion batteries are the first choice of energy storage devices for portable applications and have entered the market of electric vehicules and grid storage. A combination of small atomic weigth, low reduction potential and a monovalent charge, makes Li a suitable candidate for batteries with very few elements that can outperform it in terms of gravimetric capacity, nevertheless when converting to volumetric capacity other elements such as Be, Ca, Mg or Al have higher values. (Table I.1). [START_REF] Larcher | The Science and Practice of Lithium Therapy[END_REF][START_REF] Winter | [END_REF][12] The concept of Lithium battery appeared in the late 1960s, where metallic Li was used as anode with a non-aqueous electrolyte resulting in a primary battery with high energy density. Following, in 1970 Stanley Whittingham developed the first rechargeable Li-ion battery, by using an intercalation material like TaS2 and TiS2, which can reversibly host Li + ions, and Li-metal as an anode. However, the stripping and re-plating of Li is inhomogenous and derives in dendritic, mossy or granular deposits that may cause shortcircuit in the battery. Additionally, Li is thermodynamical unstable towards the electrolyte and passivation products that consume Li are continuously formed on the surface. Consequently, due to safety concerns the commercialization of this technology was not sucessful. [START_REF] Reddy | Handbook of Batteries[END_REF][START_REF] Tarascon | Electrochemical Energy Storage[END_REF][START_REF] Placke | [END_REF] Several strategies have been adopted to overcome these problems, one of the first ones was proposed by Prof. Michel Armand, who suggested the replacement of the liquid electrolyte by a conductive polymeric membrane. Nowadays, this approach has been pushed forward and Li-metal is being used in All-Solid-State batteries. [START_REF] Tarascon | Electrochemical Energy Storage[END_REF] The second option, was the replacement of Li metal by a compound that can reversibly host Li, implying that the cathode must be the source of Li. In 1980 John B. Goodenough, discovered a new class of cathode materials, the LiCoO2 that can serve as a source of Li with a high redox potential. Later on, Akiro Yoshino combined Goodenough's cathode with a carbon anode, a material that can reversibly host Li and has a potential close to the one of Li metal. Note that one of the main differences between Li-metal and Li-ion battery is that the last one is assembled in a discharge state, therefore it requires a process of formation of an effective SEI layer that is usually carried out by the manufacturer. [START_REF] Placke | [END_REF] These discoveries are the fundamentals of the Li-ion technology that still being used nowadays, and their impact is so high that the Nobel Prize in Chemistry for 2019 has been awarded to John B. Goodenough, Stanley Whittingham and Akira Yoshino.

On the other hand, sodium can be a possible contender for LIB, given the lower Na metal cost, its higher abundance on earth's crust and precisely its higher equilibrium potential that eliminates the need of Cu current collector in the anode side, reducing the cost of battery manufacturing. [16] The volumetric and specific capacity of sodium-ion batteries (NIB) are much smaller than for lithium-ion batteries (LIB) due to a bigger and heavier Na + ion and an equilibrium voltage that is 0.3 V higher than the operational potential in Li/Li + (E°(Li + /Li)=-3.04 V and E°(Na + /Na)=-2.71 V). [17,18] In addition, the Na + ion suffers from sluggish diffusion and excessive volume expansion. The early research in Na metal based batteries was focused in high-temperature systems, such as Na/S or Na/NiCl2, which had high energy densities but low power and a complicated implementation. [19] Na metal is not used as the negative electrode at room temperature, mainly due to safety concerns. [18] Potassium has a similar abundance in the earth crust as Na, its equilibrium potential is 0.11 V (E°(K/K + ) =-2.92 V) slightly higher with respect Li/Li + and it is possible to use some of the already established materials for LIB, like graphite, which is able to store K-ions. [20] Furthermore, the Na + and K + ions have lower Lewis acidity compared to Li + , meaning that their energy for desolvation is lower.

The larger alkali ions have weaker interactions with the solvents and the anions, property that could be beneficial for high power applications. [18,21] Table I.1. Comparison critical parameters Li, Na and K-ion. Adapted from ref. [22]

Parameter

Li + Na + K + Ionic structure Ionic (Shannon) radii (Å) 0.76

Negative Electrode Materials

Since the development of the LIB, NIB and KIB, several types of anode materials have been proposed.

-Graphite: it is the most common anode used for LIB with 370 mAh/g corresponding to the formation of LiC6. [21] The storage Na + in graphite is unsuitable due to the lack of sufficient host sites and the unfavorable energetics, the capacity is limited to ~35 mAh/g. [16] Except if a glyme-based electrolyte is used, in this case the solvent co-intercalation will increase the storage capacity. [17,18] On the contrary,

Graphite is able to store up to 1/8 K-ions (KC8) corresponding to a theoretical capacity of 278 mAh/g. [20] -Hard Carbon. In contrast to graphite all Li, Na and K can be inserted. This material is of particular interest for NIB, since reversible capacities as high as 400 mAh/g can be obtained, several studies have been performed on the relationship between the source and synthesis methods of carbon and the morphology and degree of disorder which has a direct impact on the electrochemical performance. [17,18,21] -MXenes. Are a family of 2D materials with the general formula Mn+1XnTx where M=transition metal (Sc, Ti, Zr, Hf, V, Nb, Ta, Cr, etc), X=carbon or nitrogen and T=surface terminations (-OH, H, -F, ..). They have alternated layers of M and X, leaving the outermost M layers with terminations T that allow tuning the properties of the material. Alkali ions like Li + , Na + and K + can be reversibly intercalated in the interlayer space, [22,23] with capacities of 225, 250 and 1200 mAh/g for the lithiation of Ti2CTx (T=-OH, -O, -F), Nb2CTx and MoS2/Ti3C2Tx-Mxene@C, respectively. [25] For the sodiation, capacities of 130 and 181 mAh/g for Ti2CTx and Nb2CTx [26] are observed while for the postassiation Ti3CNTx exhibited 90 mAh/g and Ti3C2Tx and Nb2CTx 30 mAh/g. [27] -Layered van der Waals Solids/Transition metal dichalcogenides. These compounds have the general formula MX2 (M=W, Mo, Hf, V, Ta, Re, etc; X= S, Se, Te). They feature strong covalent bond toward in-plane and weak van der Walls interactions towards out-of-plane, due to these weak electrostatic interactions between the layers they are easily exfoliated, similarly to graphite. Their rather large interlayer distances provide place for accommodating bulky cations like Na + or K + . [22,24] -Alloys/Intermetallics. Group 14 and 15 elements (Si, Ge, Sn, P, Sb, Bi) have been investigated as anodes for LIB, NIB, KIB (Table I.2). Tin is electrochemically active with all three alkalis, forming Li22Sn5, Na15Sn4 and KSn corresponding to a theoretical capacity of 990, 847 and 226 mAh/g, respectively. [20,28,29] The obtained experimental capacities are lower (eg. 640, 655, 180 mAh/g for Li, Na and K, respectively [21] ). Phosphorus forms Li3P, Na3P and K3-xP all with a theoretical capacity of 2596 mAh/g. [16,20,22] Antimony has a theoretical capacity of 660 mAh/g for Li (Li3Sb), Na (Na3Sb) and K (K3Sb). [28] Silicon displays a capacity of 4200 mAh/g for Li22Si5 and 3579 mAh/g for Li15Si4. [START_REF] Feng | [END_REF]3] Theoretical studies have predicted the reaction of Si with Na and K with theoretical capacities of 954 mAh/g for NaSi [30] and 995 mAh/g for KSi. [20] In both cases the reaction suffers from poor kinetics and experimentally the c-Si has been considered to be inactive or at least partially inactive towards Na and K, however a-Si appears to be a potential candidate. [20,[31][32][33][START_REF] Sangster | Phase Diagr[END_REF][35] Germanium, reacts with Li to form Li22Ge5 and Li15Ge4 with capacities of 1623 [36] and 1384 mAh/g, [37] respectively. The sodiation presents a theoretical capacity of 369 mAh/g, experimental values around 350 mAh/g corresponding to the formation of NaGe have been found. [29] The potassium insertion in Ge has a theoretical capacity of 369 mAh/g. [38] All the elements in this group suffer volume variations upon alloying reaction, in general with increasing alkali content and cation size the highest is the volume expansion. [18,20,21,28] Table I.2. Theoretical capacities of lithiation, sodiation and potassiation of group 14 and 15 elements and their approximated volume expansion after alloying reaction.

Material

Alkali Stoichiometry final product Theoretical Capacity (mAh/g)

Approximated

Volume change (%)

Sn [20] Li Sb [20] Li Li3Sb 660 147 Na Na3Sb 660 390 K K3Sb 660 407

Si [START_REF] Feng | [END_REF]3,30,39,40] Ge [29,[36][37][38]41]

Electrolytes

The role of the electrolyte is to provide efficient ion conduction pathway while being electronically insulator. A good electrolyte should fulfill the following requirements; i) high dielectric constant and low viscosity, ii) stability in a wide voltage and temperature range and towards other battery components (separator, current collector, packaging materials), iii) it should promote the formation of an effective SEI layer that prevents further electrolyte decomposition and prolongs the cycle life. It is hard to find a single solvent that meets all these requirements, thus most of the electrolytes adopt a mixture of two or more solvents. Common solvents are carbonates, ethylene carbonate (EC), diethyl carbonate (DEC) or dimethyl carbonate (DMC). [START_REF] Feng | [END_REF]16,42] Ethers are used as well due to their low viscosity and high ionic conductivity, but they are prone to decompose at low potentials. [42] Regarding the salts for Li, the LiPF6, Li(BOB), LiTFSI and LiFSI are used. The LiPF6 is moisture sensitive and has low thermal stability (70°), thus its applications are limited by safety issues, although its formed SEI layer posses the desired stability and ionic conductivity. LiFSI offers a higher solubility, anionic conductivity and a thermal stability up to 180°C, but it passivates the aluminum current collector. [42] In analogy with Li, for Na the electrolyte salts include NaClO4 and NaPF6, although NaClO4 is potentially explosive. [17] For potassium, due to the high reactivity of the K the choice of salts and solvents can drastically affect the cycling performance. Indeed, KFSI containing electrolites form a more stable SEI layer than KPF6, ensuring a better cyclability. In KPF6 the hydrolysis is absent, and the reduction is restricted, so the SEI is mainly formed from organic species (C-C/C=C, R-O-K, RCOOK) and K2CO3 that constantly dissolve in the electrolyte. KFSI on the contrary is easier to decompose and forms KF, which plays an important role in the stabilization of the SEI layer. [43] 2 Electrochemical Mechanism

Intercalation (Graphite and MXenes)

A typical example of an intercalation material is graphite, who can reversibly host Li + , K + , Cs + cations.

The highest stoichimetry that Li can achieve is LiC6 while the other alkalis form MC8. The process takes place through a series of staged graphite intercalation compounds, described in the literature by two models (Figure I.1); the Rüdorf and Hofmann and the Daumus-Hérold. The first one proposes a sequential filling up of alternating graphene interlayer spaces with no structural distortions induced within the individual graphene sheets, while in the second one, the graphene layers are flexible and deform around the Li ions. [45] An intercalation mechanism is often assumed for the MXenes phases as well. For lithiation the intercalation of a bilayer of Li ions in the MXene layers. For sodium, the intercalation occurs electrochemically and the Mxene layers are pillared with Na + ions, in some cases the Na-ions stay trapped in these pillars. In the case of MXenes the ions can be intercalated in the surface or the edges of the layers, with their solvation shell. Once the solvation shell is lost, the ions can diffuse in the bulk of the layer. For sodium mono and bilayer can be formed in the interlayer space. [26] Although given the huge structural variety of MXenes, all the known mechanism can be found, including conversion, insertion/extraction and plating /stripping. Intercalation pseudocapacitance occurs through bulk redox reactions with ultrafast ion diffusion. A complete exfoliated layer will show totally different electrochemical mechanism compared to the bulk material. [46] Since the ion storage is mainly present in the interlayer space a good use and rationalization of this space is necessary to increase the capacity.

Alloying

Reaction mechanism

The alloying reaction between group 14 or 15 (M) elements from and alkali (Li, Na, K) (Equation 1)

in most cases delivers higher theoretical capacities compared to the conventional carbon-based anodes. (Figure I.1). Nonetheless, it is accompanied by a huge volume expansion and particle pulverization that is detrimental for the electrode.

𝑀𝑀 + 𝑥𝑥𝐴𝐴 + + 𝑥𝑥𝑒𝑒 -↔ 𝐴𝐴 𝑥𝑥 𝑀𝑀 where M=metal and A= Li, Na, K (1)

Figure I.1 Top, Schematic of the (A) Rüdorf and (B)

Daumas-Hérold models for the Li intercalation into graphite. [45] Bottom, Theoretical capacities of lithiation, sodiation and potassiation of C, Si, Ge, Sn. Adapted from ref. [20,47] For both silicon and germanium, the lithiation leads to the formation of amorphous lithiated a-LixM (M=Si, Ge) intermediates, for voltages below 50 mV a crystalline Li15M4 phase appears. The delithiation proceeds with the amorphization of this crystalline Li15M4 phase and at the end of the process an amorphous M network is obtained. Both Li15Si4 and Li15Ge4 phases are very similar, they crystallize in the same space group I-43d and their lattice parameters differ in as little as 0.9%. [48] Despite the similarities between Si and Ge their lithiation phase formation sequences, kinetics and mechanical responses are different.

Lithiation

For Silicon the lithiation of the crystalline phase happens in an anisotropic way and it is governed by the movement a two-phase boundary between the inner crystalline Si core and the outer amorphous LixSi alloys. Note that this is only applicable for the first cycle, after which the silicon remains amorphous. [49] The lithiation of a-Si is isotropic and leads as well to the formation of a-LixSi intermediates. [48,50] At the end of discharge, the spontaneous crystallization of Li15Si4 is not common in other alloy anodes, in which crystalline phases are attained via a two-phase reaction that implies nucleation and growth processes. For instance, as the x value in a-LixSi approaches 3.75, the electronic structure of Li3.75Si becomes very similar to the Li15Si4, the phase transformation is possibly driven by the similarities of the electronic structures between the amorphous and the crystalline phases at high Li contents. This process only takes place electrochemically and no composition fluctuations are observed during the crystallization. The formation energies of Li15Si4 must be then lower than the amorphous counterparts and other crystalline phases with similar composition like Li12Si7 and Li13Si4.Thus the lithiation mechanism of Si does not follow the thermodynamically stable phases in the phase diagram. [47,50,51] In the case of germanium compared to silicon the Li diffusion is about 400 times faster [52] and the electronic conductivity is two orders of magnitude higher. [36] Germanium experiences isotropic volume expansion for the crystalline and the amorphous phase, which results in a lower particle cracking improving the capacity retention compare to Si. [53] In parallel with Si, the lithiation happens from the surface of the particles to the core, with the cracks nucleating in the particle edges and propagating inwards. [53,54] During its lithiation the Ge-Ge bonds gradually break to form a-Li2.26Ge and later on a-Li3.5Ge (like in the crystalline Li7Ge2). Both phases are constituted of Ge dumbbells and isolated Ge atoms, since these motifs are present in several LixGe phases with similar structures and formation energies the transformation might not be linear and several other intermediates could be involved. As the lithiation proceeds, the number of isolated Ge atoms increases and the crystallization of c-Li15Ge4 is energetically favoured. [36,55] Indeed, for Ge the nature of the formed phases depends on particle size and morphology and cycling conditions like C/rate or electrode formulation, these factors affect the kinetics of phase formation. [36,56,57] These processes will be discussed in depth in Chapter I.

Sodiation.

The sodiation of silicon has been less investigated, whether crystalline Si is electrochemically inactive or at least partially is subject of debate, one can find reports supporting both statements. This reaction is considered to be impeded by the poor Na kinetics and the positive binding energy during the sodiation of c-Si. [30,31,58] Some authors have reported the sodiation of nano-sized c-Si, which after the first sodiation is transformed in a-Si, the following cycles take place in the amorphous domain. Micronsized c-Si does not experience a sodiation process. [32,33] On the contrary, the sodiation of amorphous Si is possible. According to the Na-Si phase diagram the most sodiated phase has a composition 1:1, leading to a theoretical capacity of 954 mAh/g. [30] Recently a theoretical report showed that the maximum Na storage in Si is 0.76 Na, corresponding to a theoretical capacity of 725 mAh/g. [59] Experimentally it is not clear how many atoms of Na can be stored in Si, neither the achievable capacities or which dimensional scale or form (nano, micron, amorphous, crystalline) is required. [35] Reports in the literature have showed the feasibility of sodiation of a-Si, yet with severe transport limitations that can be partially improved by electrodes nanodesigning. [30,31] In analogy with Si, the insertion of sodium in c-Ge is almost negligible. The major obstacle lies in the high diffusion barrier for Na in Ge lattice, due to its larger size the activation energy for migration between the interstitial sites in the lattice is higher (1.5 eV for Na compared to 0.51 eV for Li). [60] Still, c-Ge can be amorphised after the first discharge/charge cycle while the sodiation continues in the amorphous phase. [60] Amorphous Ge is active towards sodiation, [38,61] the phase diagram indicates NaGe as the most sodiated phase, but recent reports have shown the insertion of more than one sodium in the Ge lattice, corresponding to Na1.6Ge. [60,62] In addition, GeOx has shown some activity vs Na, with the formation of Na2O that serves as a buffer to accommodate the volume changes during cycling. [63,64]

Potassiation

The knowledge regarding alloying compounds with potassium is limited. It has been theoretically

predicted that Si undergoes this reaction to form KSi with 955 mAh/g. This process is expected to happen through a series of alloys, like K12Si17, K7Si46, K8Si46 and KSi. Crystalline silicon is believed to be inactive towards potassiation. [20] Yet, theoretical predictions indicate that amorphous Si can store up to 1.1 K-ion leading to a theoretical capacity of 1049 mAh/g, the K-ion can diffuse rapidly in the a-Si network, due to the weak electrostatic K-Si attraction and the formation of isolated Sin clusters during the potassiation. [40] Experimentally, few work has been performed, Sultana et al. [20] reported the electrochemical inactivity of Si with K, whereas Komaba et al. [21] presented in a preliminary study of the potassiation of a-Si with a capacity of 510 mAh/g. In the case of germanium, the crystalline phase shows no activity while the amorphous counterpart is expected to form KGe with a theoretical capacity of 369 mAh/g. Experimentally, a nanoporous Ge electrode has delivered a capacity of 120 mAh/g. [38]

Challenges and limitation in Si and Ge anodes

Whether the Si or Ge alloying reaction with alkali is anisotropic or isotropic they lead to volume expansion and particle pulverization. The lithiation of Silicon happens preferentially along the <110> direction with a swelling of ≈200% while the <111> direction experiences only 10%. In contrast, germanium has a more flexible lattice that undergoes facile atom rearrangements during lithiation. As a consequence, the volume expansion happens both in the axial and radial direction, resulting in a more uniform and faster lithiation and in a reduction of particle cracking. [48,65] For the sodiation and potassiation, given the inconsistencies with the maximum known values for alkali insertion in Si and Ge (NaSi, [66] Na0.76Si, [59] NaGe, [61] Na1.6Ge [41] ), the information regarding the nature and degree of volume expansion is limited, though it is known to be related with the quantity and size of inserted alkali. [18,20,21,28] Indeed, this volume expansion is one of the main challenges for the implementation of Si or Ge in commercial applications. The full lithiated states can expand up to 300%, compromising the stability of the electrode morphology by the formation of cracks and change in the connections with the conductive network. After delithiation, the particle contracts and it is unable to return to its pristine state. With the constant expansion/contraction the particles suffer from cracking and pulverization that derive in a fracture/disconnection from the conductive network. [START_REF] Feng | [END_REF] The pulverized or disconnected particles cannot participate in the electrochemical reaction leading to capacity fading. [8] Another limitation is the low initial coulombic efficiency due to the consumption of Li + by the SEI formation during the first cycle, which contributes to the irreversible capacity and depletion of the Li reservoir in the cathode (for a full cell). Moreover, the trapping of Li increases upon cycling due to the continuous formation of fresh surfaces in the alloying/dealloying process. To summarize the silicon/germanium electrode failure is due to: [START_REF] Feng | [END_REF][3][4] i) Large volume changes in bulky films and particles lead to high internal stress causing pulverization of the particles.

ii) Pulverization leads to detachment from the current collector and conductive network, causing isolation of the material, meaning that they no longer contribute to the energy delivered by the electrode.

iii) Unstable and uncontrollable growth of the SEI layer that increases the resistance of the electrode and irreversibly consumes Li. The expansion and contraction is continuously exposing fresh particle surface for the formation of new SEI layer.

Since the volume change cannot be avoided, the key solution is the improvement of the cyclability by preventing the loss of electrical contact between the electrode and the current collector and by stabilizing the SEI layer. Strategies to achieve this goal are the reducing of the operating voltage at the expense of the capacity and the use of composites and nanostructures.

Nanostructures

Limiting the particle size has been proven to be a good strategy to improve cyclability. Indeed, both silicon and germanium have a critical particle size after which the cracks appear. For crystalline silicon this value is limited to 150 nm [48] while for the a-Si and c-Ge due to their isotropic volume expansion these values are limited to 870 [50,67] and 620 nm, [54] respectively. The use of nanostructures appears then as an efficient approach to improve the cyclability since i) their smaller size buffers the volume expansion and their unique architecture minimizes the stress along a specific orientation and avoids cracking or pulverization, [START_REF] Feng | [END_REF]7] ii) their high surface area maximizes electrochemical reactions in the electrode but also increases the amount of parasitic reactions, [5] their interconnected electrolyte-filled network enables fast ion transport and their structure offers a directional path for fast electron transport. [6] Each type of architecture contributes in a different way to the electrochemical performance but has its own limitations (Table I.3), from here the interest of combining them in the electrode design. In the following segment we will revise the characteristics of the most common nanostructures.

0D Structures-Nanoparticles(NP):

Their smaller size grants a lower associated volume expansion, higher specific surface area and tolerance to tensile stress and improved transport pathways for electrons and ions along the binders or conductive additives. All of these conditions minimize the electrode degradation and improve the cycling performance. [4,8,49,68] Nevertheless, the cycling stability is restricted by the limited space to accommodate volume changes, undesirable reactions and excessive formation of SEI in their surfaces. [68] In addition, with repeated cycling the NP tend to merge into bigger particles, leading to a decrease in the surface to volume ratio and eventual peeling from the current collector. [50,69] 1D structures: They distribute the volume expansion along the radial direction, in order to minimize particle cracking and the abundant empty space between the structures serves as a buffer for volume changes. Along the length direction they provide efficient electron transport and intimate contact with the current collector, allowing rapid charge transfer. However their synthesis methods are complicated and expensive, hindering their implementation at large scale. [3,8,9,49,70] Examples of these type of structures are the nanowires (NW) (Figure I.2), which can be grown directly in the current collector eliminating the need for conductive additive or binders. Unfortunately, the NW still can be broken and expose surface to the formation of new SEI layer that will thicken with repetitive cycling limiting the Li + transfer. [4] Hence, there is a critical diameter (<300 nm) in order to obtain the best performance in terms of coulombic efficiency and cycle life. [47,69] The nanowires can be added as active material in a composite containing the binder and the conductive additive. [START_REF] Feng | [END_REF] The nanotubes (NT) compared to NW and nanofibers usually bear much higher specific surface area when they possess a hollow structure. This last one helps to buffer the volume expansion, allowing it to be released inward and outward. [START_REF] Feng | [END_REF] Since the outer and inner parts are exposed to the electrolyte, the diffusion paths for the ions are shorter, but an excessive formation of SEI layer is promoted. [9] The length, outer and inner diameters and the wall thickness are important parameters that determine the electrochemical performance of the NT. [69] 2D structures: In order to avoid the particle rearrangement due to volume expansion the initial film should have enough void space (interparticle space) to accommodate the changes. [3] The film thickness has an impact on the electrochemical performance and in general thinner films perform better due to a better conductivity due to a shorter distance to the current collector. [START_REF] Feng | [END_REF] For these factors the thin films have limited mass loading and thickness and are only foreseen for microbatteries. [8,9,69] Note that films with thicknesses lower than 2.5 µm do not form Li15Si4. [69] Another 2D structures are the nanosheets (NS), due to their low dimensionality they provide fast kinetics for ion transport and buffer the strain generated by the expansion/contraction process. Their large surface area allows an effective contact between the electrolyte and the electrode. [9] Micron-sized Si (Ge). Despite the advantages of nanostructures to buffer the volumetric expansion, they depict complex synthesis methods, low mass loading and low coulombic efficiency. Hence it is practical to engineer micrometer silicon with characteristics of the nanosized one. [3] With this approach, the tap density can be increased and as a consequence the volumetric energy. [49] This strategy can be implemented either by the introduction of nanopores/nanograins or by the incorporation of nano/micro-meter Si(Ge) in a micrometer host such as graphite or other carbon frameworks. [START_REF] Feng | [END_REF]70] The pores diminish particle pulverization, serve as buffer for volume changes and facilitate the electrolyte diffusion in the structure; while the composites containing carbon benefit from a better structural stability, enhanced conductivity and stabilization of the SEI layer. [START_REF] Feng | [END_REF]8,68] Prelithiation: The high surface area of nanoparticles decreases the coulombic efficiency during the first cycle due to irreversible consumption of Li in the formation of the SEI layer, this process depletes the Li reservoirs in the cathode (for a full cell). With the prelithiation an artificial SEI layer is created that is activated upon contact with the electrolyte, reducing the consumption of Li during the first cycle. [49] This process can be performed in three ways; electrochemically by contacting lithium foil with electrolyte-wetted anodes, mechanically by mixing the Si(Ge) with lithium metal powder and chemically by the synthesis of LixSi (LixGe) nanoparticles that are mixed with the rest of the slurry. [3,70] Composites: Since it is challenging to tackle all of the Si(Ge) drawbacks by solely morphology control, the design of effective networks is vital for this purpose. The Si(Ge)/core carbon shell (Figure I.2 a), is an interesting structure, in which the void space between the Si(Ge) core and the carbon shell, accommodates the volume changes without deterioration while the SEI layer is stabilized in the surface of the shell. A porous but rigid shell allows to keep the integrity of the conductive network without particle disconnection and provides an active Li + diffusion path and efficient electron transport. [3,7,8,49] Tough, in some cases the volume expansion of the core affects the shell. Hence, the following factors must be considered for an effective design; i) the size of the core (less than 100 nm, to avoid excessive volume expansion), ii) the material of shell should be conductive and electrochemically stable, iii) the synthesis conditions should lead to an uniform, mechanically stable and conductive shell. [49] Ideally the shell should be thick enough to provide good mechanical and chemical stability but thin enough to avoid high ionic or electronic resistance. A similar structure is the Yolk-Shell (Figure I.2 b) in which the void space is engineered specifically to buffer volume expansion. [49,50,69] The continued cycling leads to an uncontrolled growth of SEI layer in the core/shell and nanowires while the yolk/shell-type structures can stabilize a uniform SEI layer. Adapted from ref. [3,68] Silicon-graphite composites: This composite presents a remarkable improvement compared to graphite with good cycling stability and low cost of fabrication. The synergy between Si and graphite increases the capacity and improves the electrical conductivity and stabilizes the SEI layer. Note that the ratio between both components is critical in the electrode design to ensure the benefits from both materials. [START_REF] Feng | [END_REF] Silicon-graphene: In this composite the Si particles are deposited in between the layers of graphene which can deform easily to accommodate the volume expansion (Figure I.3). The graphene layers that are far away from Si particles can reconstruct the graphite network due to van der Waals forces. [4] This method presents good mechanical stability and prevents the Si particles to be disconnected from the conducting network; it allows as well the preparation of free standing electrodes. [START_REF] Feng | [END_REF]4,68,69] However, since NP tend to aggregate, selecting the ideal matrix that can provide flexibility, high surface area and efficient electron and ion transport is important. [68] Figure I.3. Schematic of a Silicon/Graphene composite. Adapted from ref. [71] Metal alloys: the use of metal alloys is another alternative to benefit from the desired properties of two or more elements. A very interesting composite is the Si-Ge alloy. Since Ge exhibits a series of advantages over Si, such as faster Li + diffusion, higher electronic conductivity and lower volumetric expansion; the enrichment with Ge affects the rate capability and allows a better capacity retention. [72,73] Sn has also demonstrated to contribute to improve the capacity retention of a Si electrode by a synergetic effect between both elements. [74] It is worth to note that most of these approaches are implemented as well for NIB, but despite the similarities between Li and Na, certain functions used for LIB may not work for NIB, because of the different kinetic limitations between the two systems. [6] Certain examples include the c-Si core/ a-Si shell composite, in which the c-Si core improves the electronic conductivity whereas the sodium storage occurs in the a-Si shell. [35] A porous Si-sponge, in which the synergy between the porous and the amorphous character helps to alleviate the volume expansion and improve Na kinetics, allowing cycling at high C/rates. [30] a-Si Nanomembranes with enhanced Na + transport kinetics is able to store Na through a pseudocapacitive/bulk mechanism. [58] A silicon/graphene and silicon graphite composite alleviates the Na transport limitation issues and improves the electronic conductivity. [75,76] Ge nanocolumns improve Na kinetics due to their short diffusion length. [77] a-Ge NP supported in Ni nanopyramids with enhanced cycling stability and rate capability, due to enough space between each pyramid to buffer the volume changes and the intimate contact between Ni and Ge that shorten the diffusion paths. [60] The use of a-Ge thin film or porous c-Ge NP in which all the sodiation products are amorphous. [60,78,79] Finally the Li activation of Ge NWs, induces amorphization and porosity in the structure, reducing the barrier for nucleation of the NaxGe phase and accelerates the solid state Na diffusion. [80] Indeed a-Ge NW can store up to 1.6 Na per Ge (Na1.6Ge) due to the high concentration of defects in a-Ge. [41] Table I.3. Electrochemical performance of some Si anodes for LIB and NIB. Adapted from ref. [START_REF] Feng | [END_REF]8,35,70] Type of Si(

Electrolyte formulation and additives

In addition to the electrode architecture and composition, other features must be considered to improve the performance of Si (Ge) anodes. One of them is the electrolyte formulation since it is related to the formation of the SEI layer that is expected to inhibit further side electrochemical reactions on the surface of the electrodes and to guarantee a good cycle life via effective ion transport. [8] However, the alloying reaction promotes a continuous and excessive SEI formation due to its facile detachment from the particles surface during the expansion/contraction process and the particle cracking that exposes new surface area. This uncontrolled growth of SEI layer impedes the transport of ions and increases the resistance of the electrode. [8] In general the stability of the SEI layer depends on the choice of salts, solvents and additives. [4] The SEI layer derived from carbonate solvents is composed of polymeric and oligomeric organic units and it is semipermeable to the electrolyte components, meaning that it is continuously decomposed with little passivation effect. [START_REF] Feng | [END_REF]4] Hence, an efficient approach to stabilize this SEI layer is the use of additives, which most of the times are designed to be consumed during the initial formation the SEI layer, but can serve also to improve the electrochemical stability window of the electrolyte, or to enhance the safety (flame retardants). [42] Common additives include vinyl carbonate (VC), fluoroethylene carbonate (FEC) or Lithium bis(oxalate)borate (LiBOB). [START_REF] Feng | [END_REF]9] Both FEC and VC reduce before the carbonates providing an initial protective layer on the particle surface. VC decomposes into polycarbonates at about 1 V vs Li/Li + , the presence of double bonds in its structure rend it is more susceptible to reduction. [START_REF] Feng | [END_REF]81] FEC decomposes at 1.3 V vs Li/Li + into a passivating layer that is much more compact resulting in a lower mass transfer resistance and improved capacity retention. [START_REF] Feng | [END_REF]18,42] FEC decomposes into polyenes and a high concentration of Li2O which prevents the diffusion of HF through the SEI. [42] It promotes the formation of Li-F and Si-F (Ge-F) bonds that are much stronger and give more stability to the SEI layer, compared to Si-C (Ge-C) or Si-O (Ge-O) bonds that are constantly decomposing and reforming if there is no additive. However, if FEC is used in excessive quantities, the formation of HF and LiF is boosted, causing cathode dissolution, consumption of Li reservoirs and increase of the irreversible capacity. [START_REF] Feng | [END_REF]3,81] It is crucial then to optimize the additive content and to use mixtures (VC+FEC) with synergetic effects that impact the electrode performance. [81] For Sodium and potassium, given their different reactivity with the electrolyte solvents, the composition of the SEI layer differs from their Li analogue. For Na a lower quantity of oxygenated or organic species is found, being mainly composed of Na2CO3 and alkyl carbonates ROCO2. Since the Lewis acidity of Na is lower than Li, the coulombic interaction of Na + ions with negatively charged species or lone pairs is weaker, as a result the solubility of the electrolyte decomposition products is higher for NIB. [18,[82][83][84] Therefore the use of additives is crucial to avoid further SEI dissolution and electrolyte decomposition. [18,85] For potassium, certain electrolytes are suitable for the anode but not for the cathode, due to composition of the SEI layer on top of the K surface (for half cells). For most of the anodes the KFSI salt has been found to be the most stable. The carbonate based solvents allow reversible cycling whereas the glymes-based present systematic failure. [44] Whether the additives improve the performance or not is not clear yet. [20] In any case, in order to avoid undesired reactions of K with the electrolyte, full cell test must be implemented. [44]

Binder

The role of the binder is to keep all the particles together, attach them to the current collector and alleviate the stress induced during volume changes maintaining the mechanical stability of the whole electrode. Polyvinyldene fluoride (PVDF) is one of the most common binders for different technologies, but due to the lack of polar functional groups it only offers weak van der Waals interactions that are not sufficient to maintain the integrity of the electrode. [START_REF] Feng | [END_REF]4,9] On the contrary other binders such

Carboxymethyl cellulose (CMC), poly acrylic acid (PAA) or Sodium alginate contain polar groups that interact strongly with the hydroxide groups on the particle surface via hydrogen or covalent bonding, improving the adhesion. [86,87] These binders are soluble in water, meaning that the slurry can be prepared in aqueous media, which is more environmentally friendly compared to PVDF processing in N-Methyl-2-pyrrolidone (NMP). [START_REF] Feng | [END_REF]4,9,69] Nonetheless, during the volume expansion process the integrity of the binder can be compromised. The polymer chains can slide, isolating the active material particles from the conductive network, a chain crosslinking can address this issue. [9] In this sense, the use of self-healing binders is of great interest since they can repair spontaneously the mechanically destruction of the electrode, thanks to the re-association of hydrogen bonds. [3] Another concept is the use of conductive binders that ensure the electrical connectivity of all the particles even after the particle pulverization due to volume changes (Figure I.4). [50] Finally, in order to improve the contact between the current collector and the active material to decrease the electrical resistance and improve the cycle life, a porous or rough treatment by sandpaper and chemical or electrochemical etching of the current collector is suggested. [9] For further reading about the challenges and strategies for the implementation Si and Ge anodes reader can be referred to [START_REF] Feng | [END_REF]3,9,49,50,70] .

Figure I.4. a)

Schematic of a Si electrode with non-conductive binder (e.g. PVDF) and conductive additive (e.g. Carbon black) and with conductive polymer with dual functionality (conductor and binder). Adapted from ref. [88] b) Capacity vs cycle number for a Silicon anode with different binders. [89] 4 Zintl Phases

The Zintl-Klemm Formalism

The Zintl phases are named after Eduard Zintl who in 1929 studied a family of compounds in which one component is a highly electropositive material and the other is an element located in the right side of the so-called the Zintl boarder; today known as semimetals, metametals or metalloids. [90] (Figure

I.5)
In Zintl phases the valence electrons are formally transferred from the electropositive to the more electronegative component, the anions do not achieve an electronic octet rule as isolated species but may rather bond to each other, forming cationic and anionic frameworks. The versatility of the Zintl phases arises from the characteristics and interactions between these frameworks, resulting in an enormous variety of crystal structures, ranging from cluster compounds to extended solids, in which certain physical and chemical properties can be tuned. [91,[START_REF] Sevov | Intermetallic compounds: Principles and Practice[END_REF] Zintl (red) and to Klemm (blue). [START_REF] Nesper | [END_REF] Zintl defined some rules to distinguish the Zintl phases from intermetallics. A Zintl phase should; [91] i) Contain an alkali, alkaline-earth metal or p-element that is a metal, semimetal or small bandgap semiconductor.

Figure I.5. Periodic Table showing the Zintl border according to

ii) Have electrochemically balanced or closed-shell compounds, meaning that the number of electrons provided by the constituting elements is equal to the number of electrons needed to form a covalent bond.

iii) Have a well-defined relationship between the chemical (geometrical) and the electronic structure. iv) Be semiconductors (band gap < 2 eV) or poor conductors. v) Show diamagnetism or weak temperature dependent paramagnetism (Pauli paramagnetism). vi) Be brittle.

Since a vast number of compounds that do not follow all the Zintl rules, during the years the definition of Zintl phases has changed and adapted. Wilhelm Klemm extended the Zintl concept through the pseudoatom approach in which he studied isoelectronic Zintl anions. This is, after receiving the electrons from the electropositive metal, the electronegative will behave like the isoelectronic neutral atom at its right. [91] In 1980 Von Schnering extended the Zintl-Klemm concept from an 8N rule to a 18 e -one, not only to allow the study of transition metal compounds but also to prove that in some phases both concepts can work simultaneously. [90] Additional extension of the Zintl Klemm formalism have been proposed by Miller, Mooser and Pearson, [94] Corbett, [95,96] Van der Lugt [97] and Nesper [90,[START_REF] Nesper | [END_REF]98] and Belin. [99] Since the discovery of the Zintl phases, several compounds have been successfully synthesized and characterized.

Structural diversity of Zintl phases

The huge structural variety of Zintl phases is determined by the size and interactions of cations and anions. In other words, different structures varying from 0D clusters, 1D wires to 2D layers can be obtained from the same anion by changing the stoichiometry or the cation size. For example, in layered CaSi2 each Si is bonded to other three Si atoms, giving 4 + 3 valence electrons and the octet rule is complete by the electron donated from the Ca. This structure is analogous to the 3D bulk Si, but all the bonds along the [111] direction have been replaced by Ca 2+ cations. If the stoichiometry is changed to CaSi, then each Si is bonded to only two Si atoms, forming zig-zag wires, here each Si atom has 4+2 valence electrons and the octet rule is competed by 2 electrons from the Ca 2+ . If the Ca 2+ is replaced by a monovalent cation, like in KSi, NaSi, each Si atom is bonded to other three Si atoms forming isolated [Si4] 4-tetrahedra, this gives 4+3 valence electrons, and the octet rule is completed by one electron from the cation (Figure I.6). [START_REF] Arguilla | Electronic and Magnetic Materials from Two-dimensional Honeycomb Tin Lattices[END_REF] This tetrahedral unit (X4) 4-similar to P4 and As4 and can be found in alloys of Si, Ge, Sn, Pb (M) with Na, K, Rb, Cs but not for Li. This effect is related to the size of the cation, as the size increases the metal atoms are more separated one from each other, reducing their interactions, compared to a LixM mixture, and in order to decrease their mutual distances and bond energies, they prefer to cluster together. [97]

Figure I.6. Structural models for CaSi2 (layers), CaSi (wires) and NaSi (isolated tetrahedra).

Other type of Zintl phases are the cage-like clathrates, composed of tetravalent group 14 (Si, Ge, Sn and their mixtures) and trivalent group 13 elements as host structure. The guest atom (electropositive metal) is in the cavities of the network and balances the electron deficiency by the formal charge transfer. The most common types of clathrate are Type I with the formula A8E46 and Type II with A24E136. [START_REF] Vegas | Inorganic 3D structures[END_REF] Another type of structures are the allotropes, composed of different arrangements of the same element that determine the specific properties of the material, examples are the allo-Ge and allo-Si. The first one is obtained from oxidation of Li7Ge12, a phase with two dimensional polyanionic [Ge12] 7-separated by Li atoms. The second one from Li3NaSi6 which has polyanionic layers of [Si6] 4- separated by Li and Na atoms. In both cases the alkali is extracted while preserving the anionic structure. [START_REF] Zeilinger | Synthesis , Characterization and Reactivity of Lithium-Containing Silicides , Germanides and Borosilicides[END_REF] These structural versatility opens a field of study, since the material properties can be changed in function of the desired application (optoelectronics, spintronic, magnetic, catalysis, etc.). [START_REF] Arguilla | Electronic and Magnetic Materials from Two-dimensional Honeycomb Tin Lattices[END_REF]

Binary Zintl phases

LixM (M=Si, Ge): Lithium based Zintl phases present different characteristics compared to the heavier alkali metal homologues, even with the same stoichiometry they exhibit completely different structures. The main reason is the Li high electron affinity, strong polarizing power and smaller size, which condensate the Zintl anions around the Li + . [START_REF] Nesper | [END_REF] For example, the LiM (1:1) consists of infinite [X -] linkages while the Na, K, Rb, Cs analogues ideally form X4 4-tetrahedra. Interestingly, the lithiated compounds with group 14 elements are mostly isotypic, meaning that Si can be replaced by Ge up to some extent. [START_REF] Nesper | [END_REF] Most of the Li-Si and Li-Ge phases do not follow the Zintl-Klemm concept. Both Si and Ge are not electronegative enough to completely strip all the Li atoms of their electrons and the charge transfer is not complete. At low lithium contents the compounds behave like Zintl phases but as the Li ratio increases the standard electron counting rules fails. [START_REF] Nesper | [END_REF][START_REF] Chevrier | [END_REF] Examples of compounds that do not follow the electron counting rules are almost all lithium silicides and germanides, Li12Si7, Li13Si4, Li21Si5, Li12Ge7, Li9Ge4, Li13Ge4, Li14Ge4 and Li21Ge5 except the Li7Ge12, LiGe and LiSi. The electronic structure of all these compounds can be understood if the concept of an additional electronic state corresponding to a cage orbital is used. A cage orbital corresponds to metal (Li) electronic state being delocalized uniformly over other Li atoms that envelops certain Zintl ions. [START_REF] Nesper | [END_REF][START_REF] Chevrier | [END_REF] Note that the composition of several silicides and germanides have been revised in the literature after the first time they were reported, some of them are not included in the phase diagram, since there is still controversy about their true composition. [START_REF] Zeilinger | Synthesis , Characterization and Reactivity of Lithium-Containing Silicides , Germanides and Borosilicides[END_REF] In the next segment we will revise the properties of some of the most common LixM phases

(Figure I.7)
LiM. Among the LiSi series several metastable phases with same composition and pressure dependent atomic arrangements can be found. Examples are the Li tetrahedral in a three-fold coordinated Si network, layered structure with two dimensional (nontetrahedrally) four-fold coordinated Si with Li intercalated in the sheets, the buckled Si sheets with eight and four-rings and the Li intercalated silicene (Si sheets with six-member honeycombs). [104] Some of the previously listed structures are isostructural to LiGe. [91,105] At lower pressures three-dimensional three-fold silicon/germanium networks prevail, at high densities both Si and Ge exhibit sheets. [104] Li12M7. These compounds contain planar five membered rings and star-trigonal planar M4. They have a cage orbital that helps to stabilize the structure in a planar configuration. [START_REF] Nesper | [END_REF][START_REF] Chevrier | [END_REF]106,107] Li7M3/Li9M4.The Li7M3 and Li9Ge4 phases have double bonded X2 dumbbells and electrons in cage orbitals distributed over the surrounding Li atoms. [36,[START_REF] Nesper | [END_REF][START_REF] Chevrier | [END_REF] Li13M4/ Li15M4. In the Li13Si4 half of the atoms are in Si dumbbells and the other half are isolated atoms surrounded by Li atoms. The Li13Ge4 crystallizes in a different type of structure with a higher ratio of isolated Ge atoms per dumbbells. [START_REF] Nesper | [END_REF]104,108] The Li15M4 phases are isostructural for Si and Ge and are composed of only isolated M atoms surrounded by Li atoms. They do not satisfy the electron counting rules and are electron deficient. [START_REF] Nesper | [END_REF]104,109,110] These phases can be obtained electrochemically or by mechanical synthesis, Li15Si4 is metastable and decomposes into other crystalline phases at ≈170°C, while Li15Ge4 is thermodynamically stable and melts at 720°C. [104,108,110] Li21M5/Li22M5. Both phases are composed of isolated M atom surrounded by Li atoms. The Li22Si5 and Li21Si5 have two and one extra electron respectively, these electrons are distributed in cage orbitals. Since the first report of the Li22Si5 [111] there has been controversy about its true composition, which was later revised to Li21Si5. [111,112] It was determined that both phases are temperature dependent and have similar XRD patterns; with Li21Si5 being stable at room temperature whereas Li22Si5 only at 415°C. [START_REF] Chevrier | [END_REF]104,112] Li17M4. There has been great controversy regarding the composition of the Li-rich phases, as a result some of them do not figure in the Li-Si or Li-Ge phase diagrams. The Li17M4 structure is closely related to the Li21M5 only differing in the occupation of one fourfold special position, in both Li17Si4 and Li17Ge4 the positional parameters are almost identical. Some authors claim that the Li17M4 is the real composition of Li21M5. [108,113] As observed, there are many Lithium silicides and germanides that have the similar composition and formation energies, in a LIB the system does not necessarily have the time to equilibrate thermodynamically. In general, it has been observed that upon lithiation both Si and Ge exhibit dumbbells and as the lithium concentration increases these are broken into isolated atoms. The dumbbells break at lower concentrations for the Li-Ge system than in Li-Si. Since these dumbbells have been observed for ground-state and metastable phases, it is likely to find them in LIB. [104] NaxM (M=Si,Ge). The NaSi and NaGe are monoclinic and contain isolated [Si4] 4-or [Ge4] 4-clusters surrounded by four alkali metal cations, each Na + donates one electron to the [X4] 4-unit (Figure I.8). [114] The Si -and Ge -are pseudoelements of group 15 since they are isoelectronic with As and P and adopt the same tetrahedral units like in P4 and As4. [START_REF] Nesper | [END_REF]114] In contrast LiSi has a different structure formed from interpenetrating sheets and does not contain isolated clusters. [105]

CaxMy (M=Si, Ge

). The group 12, 13 and 14 elements can form extended covalently bonded networks of different dimensionality depending on their coordination. The AX2-type Zintl phase with A= electropositive divalent element (Ca, Sr, Ba, Eu) and X=group 12, 13 or 14 element stabilizes in structures like orthorhombic KHg2, hexagonal CaIn2, Hexagonal AlB2, cubic MgCu2, hexagonal MgZn2 or MgNi2. [91] These phases typically form 0D clusters, 1D wires or 2D sheets of the main element held by the cation; the dimensionalities can be tuned by varying the stoichiometry or the size ratio of the anion and the cation. The layered Zintl phases are very sensitive to their stacking sequences and the nature of the surrounding layers, in this sense understanding the interactions of neighboring atoms can lead to exotic forms of the material. [115] For instance, many layered materials have polymorphs with different stacking sequences that exhibit different properties. [90] Group 14 elements are characterized by the formation of 2D compounds (AlB2-type), like graphene, consisting of π-bonded honeycomb lattices of carbon atoms. Si and Ge, due to their larger atomic sizes do not form this type of π-bonding, their atomic bonding distances are longer thus each atom would preferentially bond to another ligand than to form a π-bond with its neighbor. This issue leads to the formation of puckered [Si -]n and [Ge -]n layers separated by monolayers of the alkali cation, like in CaSi2 and CaGe2 (Figure I.9). [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] These basic units are isoelectronic and structural analogous to black phosphorous, grey arsenic or bismuth. [START_REF] Nesper | [END_REF] One of the most impressive features of these 2D phases is their ability to undergo topotactic alkali deintercalation to form another type of van der Waals materials, the sp 3 -hybridized group 14 graphene/graphane analogues. [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF]

Figure I.9. Structural representations of CaSi2 (left) and CaGe2 (right).

Ternary Zintl phases

Liy(Si1-xGex), Ca(Si1-xGex)2. Some Zintl phases can be stabilized by the introduction of a second anion if a common array of cations is supplied. These phases are known as the double salts, and originally appeared by the introduction of oxygen impurities in the lattice. The stabilization of binary compound is done by different phase formation and it is only possible if the distances between cation-anion are not too small. In fact, the presence of the cation plays an important role, in many cases the binary phase diagram for the electronegative (semi)metals pair does not predict the existence of any compound in the solid state (e.g. Zn-Sn, Al-Ge, Al-Sn), nevertheless the introduction of an active phase or cation stabilized new phases (e.g. Na-Zn-Sn, Sr-Al-Ge, Ae-Al-Sn (Ae=alkaline earth metal)). As the combination of electronegative elements approaches the border line between the triel elements (B, Al, Ga, In, Tl) and tetrel elements (C, Si, Ge, Sn, Pb) the resulting structures begin to deviate from the Zintl-Klemm formalism. [91] It has been demonstrated that certain ternary systems behave differently depending on their electronic structure changing the chemical bonding; in some cases, the system prefers to adopt a two-phase configuration rather than an electronic and structural distortion. [91] Since Si and Ge are miscible, the existence of Si1-xGex compounds with Li and Ca is possible.

Zintl phases in batteries

In the previous segment it was possible to have a brief picture of the immense structural variety of Zintl phases, hence their use in batteries is coherent. Different Si and Ge Zintl phases have been used to prepare alternative structural organizations by chemical extraction of the alkali. E.g. the delithiation of c-Li15Si4 that leads to an a-Si and the delithiation of Li12Si7, Li7Si3 and Li13Si4 that generates layered Si. [START_REF] Zhao | [END_REF] Amorphous Si nanoparticles can be prepared as well from the NaSi Zintl phase. Similarly, Ge nanoparticles can be obtained from Na12Ge17. [118,119] The previously mentioned phases showed superior Li storage compared to their bulk counterparts. [START_REF] Zhao | [END_REF][118][119] Another example is the preparation of allo-Si or allo-Ge from Li3NaSi6 and Li7Ge12, respectively; [START_REF] Zeilinger | Synthesis , Characterization and Reactivity of Lithium-Containing Silicides , Germanides and Borosilicides[END_REF] and the layered Si and Ge obtained from CaSi2 and CaGe2, respectively. These last ones will be studied in detail in the following chapters.

Xenes, Xanes (X=Si,Ge)

The Xenes refer to the single element layered graphene analogues made from group 14 and group 15 elements. Their ligand-functionalized derivatives are known as Xanes.

Silicene/Siloxene

As observed in the previous sections, silicon can show an impressive versatility in terms of structural variety and specific properties. Likewise carbon, it has the ability to form layered structures; like the buckled sheet based on the bulk Si (111) structure with honeycomb lattice and sp 3 hybridization and the weakly corrugated graphene analogue with mixed sp 2 -sp 3 hybridization. [120] In order to understand the formation of these structures it is convenient to review the similarities between carbon and silicon, listed in the following Table I.4. [120,121] As observed in the previous table Si and C are very alike, indeed the silicon equivalent of graphene, the silicene (Figure I.10 c), was predicted more than 50 years ago. However, this Si monolayer is not stable and cannot be found in nature. Despite graphene which has a sp 2 hybridization leading to planar structures, in silicene due to the bigger size of Si, a mixed sip 2 /sp 3 hybridization is favored resulting in a buckled structure while the planar one is energetically non-favored. [122][START_REF] Spencer | Silicene: Structure, Properties and Applications[END_REF][START_REF] Molle | [END_REF] Consequently, the synthesis of free-standing silicene has not been achieved, although it has been grown on metallic substrates. The lifetime of this material is very short and oxidizes rapidly outside the ultra-high vacuum. [START_REF] Spencer | Silicene: Structure, Properties and Applications[END_REF][125][START_REF] Cahangirov | Introduction to the Physics of Silicene and other 2D Materials[END_REF][START_REF] Houssa | [END_REF] Silicene is of interest due to its unique physical and chemical properties, strikingly different from the bulk ones, modifiable by chemical doping. Because of the 2D nature, charge carriers are highly mobile.

This property is of particular interest for applications like transistors, photo-detection, biomedical imaging, optoelectronic devices, quantum information technology remote sensing, and optical communications, among others. [24,[START_REF] Spencer | Silicene: Structure, Properties and Applications[END_REF] Silicene can be functionalized; the hydrogen terminated silicene (polysilane) has higher stability and it is a free standing sheet (Figure I.10 b). [START_REF] Matsuda | Basics and Families of Monatomic Layers[END_REF] The adsorption and desorption of hydrogen is reversible and can be used for controlled hydrogen storage. [START_REF] Spencer | Silicene: Structure, Properties and Applications[END_REF] The chemisorption and intercalation of species like Li, Na, Mg in the layers has been theoretically studied. [START_REF] Houssa | [END_REF][START_REF] Jose | [END_REF] Indeed, Li can be reversibly adsorbed in both sides of a Silicene layer with no structure degradation. This process takes place with a lower Li diffusion barrier compared to bulk Si. The improved kinetics and the absence of volume expansion settle silicene as a potential anode for LIB. [130] An extended study of this topic will be presented in the next chapters.

In the practice the layers of Silicene can be stabilized through terminations of hydrogen or hydroxide, forming the polysilane (Si6H6 the analogue of graphane) and siloxene (Si6H3OH3)2 (Figure I.10). These phases were synthesized for the first time by Wöhler in 1863 [131] who studied the effects of acid in CaSi2 (Equation 2). This last one has planes composed of Si6 honeycomb rings with Ca intercalated in between, and by an acid treatment at low temperatures a yellow-green powder was obtained. [132] Subsequently, in 1921 Kautsky [133] modified the synthesis and obtained a gray-green compound called "Siloxene" (Figure I.10 a). From the experimental data he suggested a structure constituted of planes of Si6 honeycomb rings connected by oxygen bridges with each Si bonded to a hydrogen above or below the plane. Since the oxygen insertion in the Si planes causes stress and eventually structure degradation, no crystallinity evidence was found.

3CaSi2 + 6HCl + 3H2O  Si6H3(OH)3+ 3CaCl2+3H2 (g) (2)
It was only until 1979 when Weiss et al. [134] synthesized a crystalline Wöhler siloxene and performed the first structural characterizations. They proposed a structure of Si planes terminated by H or OH and no O inserted into the Si-Si bond system. They claimed that in the presence of moisture, the siloxene undergoes hydrolysis of Si-H to Si-OH followed by a condensation to Si-O-Si. These results were later confirmed by a powder and single crystal XRD characterization by Weber et al. [135] Posteriorly, Dahn et al. [136,137] attributed presence of Si-O-Si groups to the siloxene oxidation into a glassy SiO2 by exposure to air and moisture, this crosslinking is intensified by thermal annealing >300°C leading to decomposition into SiO2. Finally, Yamanaka et al. [138] synthesized layered polysilane (Si6H6) by topotactic reaction of CaSi2 at -30°C with no hydrogen evolution, reactions carried out at higher temperatures resulted in oxidized siloxenes. The polysilane has a considerable amount of highly reactive dandling bonds, thus it is extremely reactive and complicated to handle. [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] Since the siloxene synthesis is different from conventional bulk Si, for a long time its research and understanding have remained in the organosilicon domain, being considered as a silicon-based polymer. This assumption was reinforced by the fact that oxygen serves as a link between two silicon atoms, allowing polymerization of Si-clusters to macroscopic structures. [139] As a matter of fact, this Si-O bonding isolated Si chains connected via Si-O-Si. [139]

Figure I.11. Structural model of the different siloxene modifications. a) corrugated silicon layers with three Si-Si bonds and alternating -H or -OH substituents (Wöhler siloxene).b), two-dimensional Si sheets made of Si6 rings linked by oxygen bridges (Kautsky siloxene). c) one dimensional Si-Si chains interconnected by Si-O bonds.

Diverse modifications have been performed on the siloxene and polysilane, here, the synthesis conditions play an important role in the specific reaction products. [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] The following segment presents a brief revision of these particular synthesis (Figure I.12).

Figure I.12. Synthetic routes for the preparation of organic modified siloxene (Germanane).

Organic modification. As observed in the figure, it is possible to introduce/intercalate different organic groups in the siloxene and polysilane layers. For example, unsaturated organic groups like alkenes and alkynes assisted by a catalyzer. These organic moieties do not only intercalate but also react with Si to form Si-C linkages, originating the characteristic IR vibrations of organic molecules (Si-CH2 and CH) and a decrease in the Si-H signal intensity, indicators of a covalent attachment with the silicon. In the process the original siloxene/polysilane backbone is oxidized. [120,[START_REF] Spencer | Silicene: Structure, Properties and Applications[END_REF]132] The Si-C bond can be inserted as well via the Grignard reagent (PhMgBr). This method allows the synthesis of phenyl-modified and oxygen-free silicon nanosheets that are more stable towards oxidation and hydrolysis. The presence of Si-Ph bonds is determined through IR and 1 H NMR by the identification of the aromatic C=C bond and the different hydrogen environments of the phenyl group, respectively. The introduction of organic aromatic groups broadens the light absorption ranges and has applications in photovoltaics. [120,[START_REF] Spencer | Silicene: Structure, Properties and Applications[END_REF]132,140] Finally, the use of a HCl solution in methanol, butanol, C12H25, Benzyl alcohol or CH2COOMe forms the corresponding alkoxide-terminated organosiloxenes (Si6H5OR, where R= methanol, ethanol, butanol, C12H25, benzyl alcohol or CH2COOMe). [START_REF] Spencer | Silicene: Structure, Properties and Applications[END_REF]132] This process allows as well the production of carbon coated sheets by a non-oxidative pyrolysis, a better dispersion in organic solvents by the introduction of Si-C or Si-N bonds and a tuning of the band gap depending on the interlayer distance. [24] While c-Si has an indirect band gap and shows extremely weak luminescence, the Si sheets and chain polymer have a direct band gap resulting in a strong visible and near-UV photoluminescence. [141] Amination. The Si center experiences a facile nucleophilic attack by N due to the Si δ+ -H δ-polarization. This way by reacting Siloxene or polysilane with N-Decylamine or N-Butylamine the amine group is inserted on the surface. The Si-H bonds are transformed into Si-NR, subsequently by reaction with another Si-H into Si-NR-Si, originating a mixed composition of Si-NH-R and Si-NR-Si bonds. The IR confirmed the presence of the Si-N-Si vibrations and the absence of Si-H, N-H, Si-O ones. The obtained nanosheets are oxygen-free. [120,132] Exfoliation. Several methods can be employed to achieve the exfoliation of the siloxene and polysilane layers. One of them is the Mg doping of CaSi2. This action weakens the electrostatic interaction between Ca 2+ and the [Sin] -layers, by reducing the Si charge, producing CaSi1.85Mg0.15. The Ca is deintercalated with polyamine hydrochloride (PAHCl) to obtain nanosheets. The exfoliation occurred preferentially in the sections where Mg atoms were present and the obtained nanosheets are oxidized. [120,132,142] Also, the CaSi2 can be doped with K, the produced silicon nanosheets have lower degree of oxidation and a high degree of ordering. [120] Another method for exfoliation is the dispersion of siloxene/polysilane in a solution of sodium dodecyl sulfate (SDS) surfactant. Silicon nanosheets with lateral dimensions of <200 nm are obtained after several days of reaction. In this case the cleavage occurred thorough the particle rather than on a layer by layer basis starting from the surface. The electron diffraction showed the presence of crystalline Siloxene, although the eventual oxidation of the layers cannot be prevented. [132,143] In addition, silicon nanoflakes can be prepared via a solid-state reaction, by heating a mixture of CaSi2 and NiCl2 at 350°C under Ar atmosphere to obtain the silicon nanoflakes and a mixture of CaCl2 and NiSi2. The CaCl2 is easily removed with dimethylformamide (DMF), but the NiSi2 remains. If the reaction is performed with TaCl5, the TaCln and CaCl2 by-products are removed by washing with ethanol. The resulting silicon nanoflakes are in agreement with Kautsky siloxene, the sheets are bridged by Ca atoms providing additional structural stability. [122,[START_REF] Spencer | Silicene: Structure, Properties and Applications[END_REF]144] Lastly, the interlayer distance can be controlled until complete exfoliation by the intercalation of organic functional groups of different sizes, the bulkier the group the larger the interlayer distance and the weaker the interlayer adhesion. The complete exfoliation is assisted by sonication, this step determines the size and quality of the nanosheets. [24,120] Germanene/Germanane The prediction and successful synthesis of silicene and siloxene opened the quest to find the graphene germanium analogue, the germanene (Figure I.13). This material, similarly to silicene, consists of a buckled Ge honeycombs monolayer, its free-standing form does not exist in nature.

Although, it has been successfully deposited on metallic substrates like Au (111), Ag (111), Cu (111) by molecular beam epitaxy in ultra-high vacuum, but it decomposes upon contact with air. [145] The choice of substrate is crucial as interactions with the deposited Ge atoms can take place, eg. in Ag (111) Ge atoms prefer to stabilize in Ag2Ge, [146] and the degree of buckling can be influenced. Then, a substrate that interacts weakly with germanene must be used to obtain free-standing monolayers. [START_REF] Spencer | Silicene: Structure, Properties and Applications[END_REF][125][START_REF] Cahangirov | Introduction to the Physics of Silicene and other 2D Materials[END_REF][START_REF] Houssa | [END_REF] Due to its 2D nature the germanene has electron mobility 5 times higher than bulk Ge and the ability to undergo topological functionalization that produces a series of electronic states like semi-metallic, semiconducting, superconducting and trivial insulating. Applications include low-energy electronics, piezo-magnetism and thermoelectricity. Its band gap can be tuned by applying an electric field perpendicular direction to the sheets, property of great interest for transistor applications. A monolayer of hydrogen terminated germanene has already been used as a transistor. [START_REF] Matsuda | Basics and Families of Monatomic Layers[END_REF] Unfortunately, most of these properties are hindering by its instability at ambient conditions. [START_REF] Molle | [END_REF] The functionalization of germanene is an effective strategy to tailor physical and chemical properties (band gap, thermal stability, etc.), depending on the nature of the ligand. Functionalization includes fluorination, hydrogenation, organic group termination, foreign element doping. [145] Fluorination. The F diffusion in the CaGe2 lattice segregates the crystals into Ge and CaF2 while maintaining their layered structure. With this technique depending on the F concentration monolayer, bilayer and trilayer germanene are stabilized. [147] Hydrogenation. The hydrogenation to form the germanane (analogue of graphane) with a puckered structure (like siloxene) with a mix sp 2 -sp 3 hybridization was first reported in 2000. [145,148] Similarly to CaSi2, CaGe2 undergoes a topotactic deintercalation of Ca at -30°C in concentrated HCl to form the germanane. (Equation 3). [START_REF] Molle | [END_REF] The formation of germanane is accompanied by a change in color from the gray-metallic in CaGe2 to red in germanane. [148] As observed for siloxene, the germanane stabilizes in the polymorphs that are inherent to the parent CaGe2 compound, the tr-6 (R-3m) trigonalrhombohedral with a 6-fold stacking sequence and the 2H structure (P63mc) hexagonal structure with 2-fold stacking sequence. [149] nCaGe2 + 2nHCl  (GeH)2n + nCaCl2

(3)

Contrary to siloxene, the formation of germanane does not lead to a hydrolysis reactions even at temperatures >0°C. [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] Therefore, the product is a hydrogen-terminated Ge honeycomb layers (isomorphic to (111) [141] Other way, prolonged moisture exposure of germanane eventually leads to the formation of Ge-OH bonds, forming the Germoxene Ge6OH3H3. [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF]150] Germanane has been reported to be stable up to 75°C above which amorphization and dehydrogenation occur. [151] The germanane layers can be exfoliated into nanosheets (GeNS), by mechanical processes, like scotch tape or sonication. [151,152] These GeNS have buckled honeycomb structures identical to Ge (111) but are easily oxidized. This issue can be addressed by the introduction of covalent bonds improving the chemical and thermal stability. [153] Organic group termination. The formation of (GeCH3)n (methyl terminated germanane) has been reported by reacting CaGe2 with CH3I/H2O. The prepared GeCH3 had more -H than -CH3 terminations and oxidized germanane (GeOx). The crystals obtained by this method (≈1mm and <100 µm thick) share the same honeycomb lattice as in CaGe2 and the thermal stability is increased to 250°C (GeH=150°C). [154] A non-oxidative synthesis route that produces homogeneously covered GeCH3 to avoid local variations in the electronic structure, can be performed by limiting the hydrogen source (water) and using CH3I in CH3N. This GeCH3 has an increased ratio -CH3/-H with no oxidation, it is air stable and has a greater thermal stability (380°C). [155] This organic modification can be extended to other functional groups; by reacting CaGe2 with -CH2OCH3, -CH2CH=CH2, -CH3, -CH2CH3, CH3I/CH3CN, etc., in HCl resulting in GeRxH1-x. [156] The ligand size and electron withdrawing/donating properties have a strong influence on the final product; smaller ligands show almost full replacement of Ge-H bonds by Ge-R while bigger ones lead to partial hydrogenation of the framework. The band gaps and the Raman shifts are dictated by the relative Ge-H/Ge-R ratio in a similar fashion to the Vegard's Law. Larger and more electronegative ligands expand the germanane framework and lower the energy of the Ge-R interactions, reducing the band gap. [156,157] The methyl terminated germanane shows as well enhanced visible-light driven photocatalytic properties. [157] Foreign element doping. The introduction of foreign elements in the germanane structure is done by doping the CaGe2 precursor with group 13, 14 or 15 elements that remain in the backbone after topotactic deintercalation. For instance, Ge can be partially substituted by As (1.1%), Ga (2.3%), [158] P [159] and Sn. The Ge0.91Sn0.09H0.91(OH)0.091 [160] and Si1-xGexH1-y(OH)y [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] alloys have been reported, the dopant content is determined by the maximum quantity of Sn (Si) that can be incorporated in the precursor Zintl phase Ca(Ge1-xSnx)2, Ca(Si1-xGex)2. [160] In both cases the Ge is terminated with H while the Sn (Si) with OH. Unfortunately these phases decompose rapidly at ambient conditions. [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF]158,160] Alkali adsorption/intercalation. The adsorption of more than one Li in germanene is favored in both sides of the layer. Calculations predict the adsorption of a total of 8 Li atoms, 4 above and 4 below a single plane, with a good electron mobility and a Li diffusion barrier between 0.21-0.76 eV which is lower than silicene (1.2 eV). Based on a single and bilayer germanene the theoretical capacity are 369 and 276 mAh/g, respectively. [161] The adsorption of Na has been theoretically predicted as well in germanene and germanane. [162] The exfoliated germanane has been used as anode in LIB, with a reversible capacity of 1108 mAh/g. The advantage of the use of a 2D material in batteries is that the bulk diffusion can be eliminated, improving the rate capability. [152] 5 Discussion and aim of the work

In the previous sections we were able to have a brief introduction on the use of alloying materials like silicon and germanium as anodes for LIB. Both deliver considerably higher capacities compared to the commercially available graphite anode in LIB. Their investigation and use in the NIB and KIB technologies is in its very beginning. However, their application at a large scale is hindered by a series of challenges derived from the alloying reaction. The incorporation of alkali in Si and Ge causes volumetric expansion that can attain values as big as 300% (for Li21Si5) upon alkali deinsertion the particles in most of the cases are unable to recover their initial state. This process of volume expansion and contraction is inherent of Si (Ge), therefore it cannot be avoided, the most suitable strategies are to mitigate their overall impact over the electrode. Additional issues caused by the volume expansion included the constant exposure of fresh electrode surface to the formation of new SEI layer, the particle cracking with posterior pulverization and disconnection from the current collector due to constant volume changes. Several strategies have been placed to address these problems, among them one can find the use of nanostructures specifically designed to host the volume expansion while maintaining a good structural stability, and to offer a better conductive network that will ensure the participation of all particles to the electrochemical reaction, increase the rate capability and extend the life cycle. Other strategies may include the use of an adequate binder to ensure the integral stability of the electrode and its adherence to the current collector or the use of an effective electrolyte formulation that will promote the formation of a good SEI layer preventing further electrolyte decomposition. Our approach to address the associated problems with volume expansion is the use of a silicon-germanium alloy in which one can benefit from the specific properties of each material, this will be discussed in the Chapter II, and the structural approach in which instead of a bulk material a layered one is used, this will be revised in the chapter III and its electrochemical mechanism will be discussed in chapter IV. The goal of this project is to understand how differently silicon and germanium work as a bulk and as a layered material. This knowledge is crucial for unveiling their reaction mechanisms during cycling and use them as alternatives for addressing the problems associated with the alloying reaction of group 14 elements

Chapter II. Si and Ge based alloys 1 Introduction

Silicon and germanium have both been at the very centre of electronics development and are now becoming of huge interest for the next generation of negative electrode materials particularly for Liion batteries. Both elements undergo alloying reactions with lithium while delivering capacities up to ten times higher than conventional graphite-based electrodes (372 mAh/g). [163] For instance, silicon in its highest lithiated states delivers 4200 mAh/g for Li22Si5 and 3579 mAh/g for Li15Si4, [164][165][START_REF] Touidjine | Optimisation de l'électrode négative à base de silicium pour les batteries lithiumion[END_REF] while germanium 1623 and 1384 mAh/g for Li22Ge5 [36] and Li15Ge4, [164] respectively.

The vast specific capacity of silicon comes at the expense of a huge volumetric expansion that eventually leads to a capacity fading. Germanium suffers from similar limitations, and compared to silicon, is less abundant and therefore more expensive, thus its commercialization potential is limited.

Nevertheless, it presents many advantages over silicon as an electrode: i) it allows very fast Li + diffusion, (ca 400 times faster than in silicon) [52,[START_REF] Graetz | [END_REF] and exhibits an electronic conductivity two orders of magnitude larger. [36] ii) Lithiated Ge alloys have a slightly lower volumes than the corresponding Si analogues, hence their associated volumetric expansion is smaller. [168] iii) Ge experiences an isotropic volume expansion that permits a facile stress relaxation; while this process in Si is anisotropic and promotes the formation of cracks in the electrode. [53] As reviewed in the Chapter I, the electrochemical alloying reaction of silicon and germanium presents a series of challenges and several strategies have been adopted to face them. In this context, it is of great interest to evaluate the synergetic effect of Si1-xGex alloys in order to extract the benefits from each element. Previous works indicate that the specific capacity is increased compared to a Gerich electrode while the capacity retention is improved compared to a Si-rich electrode, [73] but the exact interest of these type of electrodes will depend on factors like specific capacity, C-rates, cost, etc.

The present chapter is devoted to the alloying mechanism of Si, Ge and Si1-xGex alloys. Initially, the silicon system is described based on the different reports in the literature. Following, a combination of in situ and operando characterization techniques is adopted to study the germanium and Si1-xGex alloys systems. For Ge, it was possible to identify how the cycling conditions impact the different lithiated intermediates during cycling while for Si1-xGex, the Si0.5Ge0.5 composition was studied in depth tracing a relationship between its electrochemical mechanism and the improved cycling performance. Note that the mechanism for sodiation and potassiation are not revised in this chapter due to the very limited information available in the literature, for more details the reader can refer to Chapter I section 2, where the mechanism is briefly highlighted based on the state of the art.

Sample preparation and characterization techniques details

This section describes the experimental details employed for the study of the Ge and Si0.5Ge0.5

systems, for additional details the reader can refer to the Annex I section.

Synthesis. The Si0.5Ge0.

Experimental details Operando Techniques

X-ray diffraction (XRD):

The synthesized phases were analyzed by ex situ XRD (Panalytical Empyrean diffractometer, Cu-Ka1/Ka2 radiation, theta-theta configuration) with a silicon sample holder covered with a Kapton foil. The measured 2θ range was 10°-60°, during 1 hour. The in situ XRD experiments were conducted using a cell with a beryllium window. [169] The electrolyte and the cycling rates used were the same as for the electrochemical measurements. A Whatman glass fiber membrane was used as the separator and the counter electrode was lithium. The recorded spectra were processed with the FullProf software and the cell parameters were refined.

Raman spectroscopy:

The synthesized phases were ex situ analyzed at room temperature in a Dilor RT30 spectrometer with a HeNe laser source, a wavelength of 633 nm, in a 180° backscattering geometry, and scanned from 200 to 600 cm -1 . The laser was focused with a 50X optical microscope objective. The recording time was 10 minutes. For the operando analysis a self-supported film was used. The in situ cell described in ref. [170] was assembled using a lithium foil as counter electrode, a

Whatman glass fiber separator wetted with LP30 electrolyte (1 M LiPF6 in EC: DMC (1:1)) and cycled at C/5. A drilled copper grid was used as the current collector to enable the laser light to reach the electrode. The measurements were carried out at room temperature using a confocal Dilor LabRAM spectrometer, equipped with a HeNe laser source, a 600 grooves/mm grating, and a Peltier-cooled CCD detector. A backscattering geometry with a 50X optical microscope objective, under an excitation laser line of 647 nm was used. The spectra were recorded by accumulation of two spectra of 10 minutes each with a resolution of 1 cm -1 . Additional experimental details can be found in the Annex I.

X-ray absorption spectroscopy (XAS):

The in situ electrochemical cell dedicated to XAS was assembled using the same electrochemical conditions described above for in situ Raman experiment. [170] XAS measurements at the Ge K-edge were performed in transmission mode at the ROCK beamline [171] of synchrotron SOLEIL (France). A Si (111) channel-cut quick-XAS monochromator with an energy resolution of 1 eV at 30 keV was used. The intensity of the monochromatic X-ray beam was measured by three consecutive ionisation detectors. The in situ electrochemical cell, [172] cycled at C/8, was placed between the first and the second ionisation chambers. Successive spectra were collected at a rate of 2 Hz and averaged out over periods of 20 minutes. The energy calibration was established with simultaneous absorption measurements on a Ge metal foil placed between the second and the third ionisation chamber. The first inflection point of the XAS pattern of Ge (11103 eV)

was used for the energy calibration. The absolute energy reproducibility of the measured spectra was ±0.05 eV.

The operando EXAFS (Extended X-ray Absorption Fine Structure) spectra, collected up to k = 16 Å -1 , were first globally analyzed with Principal Component Analysis (PCA) [START_REF] Massart | Handbook of Chemometrics and Qualimetrics: Part A; 1st Editio[END_REF] using Matlab in order to determine the number of independent components contributing to the whole series of collected spectra during electrochemical cycling. The number of principal components was then used as the basis for Multivariate Curve Resolution-Alternating Least Squares (MCR-ALS) [START_REF] Ruckebusch | [END_REF]175] analysis. The reconstructed spectral components as well as the spectrum of pristine Ge and Si0.5Ge0.5 were fitted using the Demeter software package. [176] Fourier transforms of EXAFS oscillations with different k weights were carried out in k-range from 0.3 to 1.3 nm -1 . Fitting was performed in R-range from 0.1 to 0.4 nm using k 1 , k 2 and k 3 weights. EXAFS amplitudes and phase-shifts were calculated by FEFF7 starting from the calculated lattice parameters of Ge and Li15Ge4 for Si0.5Ge0.5 and Ge, LiGe, Li7Ge2, Li15Ge4 and Li22Ge5 for Ge. Interatomic distances (R) and the Debye-Waller factors (σ 2 ) were calculated for all paths included in the fits. Further experimental details are reported in the Annex I. and dried under vacuum. Thereafter, the samples were packed in air-tight rotors, inside the glove box and spun at the magic angle with a spinning frequency of 62.5 kHz. All spectra were acquired using a

Hahn echo (32 μs) with a radio-frequency field of 190 kHz and referenced to a 1 M LiCl solution (at 0 ppm). The repetition time was chosen to ensure full relaxation of the magnetization. The spectra were processed without any line broadening and fitted using the dmfit software. [177] The operando experiments were performed on the same spectrometer with a Bruker diffusion probe equipped with a 10 mm saddle coil tuned to 77.7 MHz. The Kel-f electrochemical cell design has been published elsewhere. [178] No spinning can be applied to this design. A self-supported Si0.5Ge0.5 electrode of 5 mm diameter was used. The electrolyte was 1 M LiPF6 in EC:PC:DMC with 1% VC and 5% FEC, a Whatman glass fiber membrane was used as the separator, and the cell was cycled at C/5 vs. Li/Li + .

The NMR and the electrochemistry were synchronized using TTL pulses. Transients were acquired following a 90° pulse excitation (RF field of 14.3 kHz), with a repetition time of 10 s during the whole measurement (86h30). 176 transients were added for each spectrum shown in the following (average over 30 minutes). The spectra were processed using an exponential line broadening of 50 Hz.

Scanning electron microscopy -Energy dispersive X-ray analysis (SEM/EDX):

Morphology and composition of the synthetized samples as well as of the electrodes were checked by SEM (Hitachi S4800) operating at 20 kV and EDX (Hitachi S4500).

3 Alloying Mechanism of Si, Ge and Si 1-x Ge x alloys

Review on the electrochemical mechanism of silicon

The electrochemical behavior of silicon has been widely studied in the literature; the present section comprises a review of the different works performed on Si that are relevant for the comprehension of the following sections.

The galvanostatic cycling of c-Si (Figure II.1-a), is characterized by a rapid decrease in the potential until reaching a plateau between 120 and 80 mV, then there is a change in slope close to the end of the discharge where the potential rapidly decreases. [179] During charge the processes take place at higher voltages, the potential increases rapidly until 0.4 V followed by a plateau, after which the slope increases upwardly suggesting a one-phase region. For the second discharge at least two different plateaus are observed. [179] According to the literature the plateau during the first discharge is associated with the amorphization of the c-Si into a-LixSi intermediates, with the crystallization of c-Li15Si4 at the end of discharge. The delithiation of Li15Si4 is characterized by a plateau at 0.45 V, corresponding to a two-phase region where c-Li15Si4 is converted into a-LixSi. Further delithiation occurs as a solid solution in which a-LixSi is delithiated to form a-Si. This last one can be affected by a series of factors like structure or residues of c-Si in the electrode. The second and subsequent lithiation/delithiation of a-Si are characterized by gently sloping profile indicating that the process occurs in the amorphous phase, then the energy differences between all the formed phases are smaller, the lithiation will occur preferentially on the previously lithiated regions. [180] These processes can be evidenced in the derivative curve (Figure II.1-b), where the peaks marked as A and B correspond to the amorphization of c-Si into a-LixSi and the peak C corresponds to the crystallization of Li15Si4. [179] If the cutoff potential is set to 50-100 mV, the formation of c-Li15Si4 is avoided and the peak C is absent. The delithiation is dependent on the state of lithiation and the cutoff potential, if c-Li15Si4 is formed, the delithiation will be characterized by a sharp peak (F) corresponding to the amorphization of c-Li15Si4 into a-LixSi; the magnitude and the sharpness of the peak are measure of the presence of Li15Si4 at full lithiation. If no c-Li15Si4 is formed during the discharge, the delithiation will be characterized by two broad peaks (D and E). The nature of these processes is not yet well understood and could be highly dependent on other factors such as dopants or structure of the Si. [50,180,181] The suppression of the formation of Li15Si4 has been observed at high C/rates due to a kinetic limitation to achieve full lithiation, translated into a decrease in the capacity [5] and for films lower than 500 nm. [51] Nevertheless, this last one is subject of debate, according to Obrovac [182] during the first 20 cycles of lithiation/delithiation of a Si thin film the profile of the derivative curve is similar to a-Si, but after continuous cycling the sharp peak near 0.42 V typical of the formation of Li15Si4 appears. Indeed is the formation of c-Li15Si4 believed to be the responsible for the delamination of the Si films after certain cycles. [182] Figure II.1. a) Galvanostatic curve of Si cycled vs Li (dashed line corresponding to the 2 nd cycle),3 b) Cyclic voltammetry of Si vs Li with 0 V and 50 mV cut-off voltage, [50] c) Galvanostatic curve of nano-and bulk-Si vs Li and their respective derivative curves. [183] For nano-silicon, during the first discharge, the galvanostatic profile is the same as in bulk silicon.

The main difference appears upon charge, where instead of a rapid increase in the voltage to 0.4 V followed by a plateau there is a gradual increase in the voltage. In the second discharge two plateaus appear in nano Si at 0.25 V and 0.10 V, indicating a different reaction compared to the first discharge. [183] Regarding the lithiation of amorphous silicon, initially it occurs at a higher potential compared to c-Si, but this potential decreases as the lithiation proceeds. In order to keep the lithiation in the amorphous region the cutoff voltage must be set before the crystallization of c-Li15Si4. [184]

Electrochemical Mechanism insight

The understanding of the behavior of Si in a battery is crucial, but the phase diagram based on thermal processes does not necessarily represent the processes taking place in a battery. A typical phase diagram includes LiSi, Li12Si7, Li13Si4, Li1Si4 and Li22Si5, all of them with less formation energy for the crystalline phases than for the amorphous one. However, during lithiation no crystalline phases are formed despite their lower Gibbs free energy. The LixSi products are always amorphous with composition 0<x<3.75, at x=3.75 the crystallization of a phase takes place through a spontaneous process without long-distance atomic diffusion. For a long time this crystalline phase was presumed to be Li22Si5/Li21Si5 [111,112] but in 2004 Obrovac [185] established it as Li15Si4 corresponding to the most lithiated phase formed electrochemically at room temperature while Li22Si5/Li21Si5 can only be obtained for very slow rate cycling at T°>100°C. [186,187] The Li15Si4 has lower formation energy than other alloys with the same Li content and it is metastable, decomposing rapidly into other lithiated compounds such as Li13Si4, Li7Si3 or other weakly crystallized phases. [50,186] Figure II.2. In situ XRD for a self-supported Si electrode cycled vs Li at a rate of C/5. 4 The in situ and ex situ XRD analysis of the electrochemical lithiation mechanism of Si (Figure

II.2),
indicates a gradual decrease in the diffraction peaks of c-Si upon lithiation, corresponding to the amorphization of the crystalline network. Following the system shows no diffraction peaks indicating the presence of amorphous phases, until 100 mV where a phase corresponding to Li15Si4 starts to crystallize, and below 50 mV is transformed into a non-stoichiometric phase c-Li15+δSi4. [181] All these processes are in line with the previous description of the plateau and peaks in the galvanostatic and derivative curves for the first discharge. During charge, the diffraction peaks associated with c-Li15Si4 gradually disappear, in agreement with the plateau and peak at around 0.4 V in the galvanostatic and derivative curve, respectively, indicators of the amorphization of c-Li15Si4. At the end of charge an amorphous pattern is recorded. [180] Note that if the cutoff voltage is set to 100 mV no XRD pattern associated with Li15Si4 appears. [180] The presence of amorphous silicon, influences the potentials at which the processes take place during the second discharge; the formation a-LixSi intermediates happen at a higher potential compared to the first discharge, while the crystallization of c-Li15Si4 happens at a slightly lower potential. [179,188] Given the amorphous nature of the different lithiated intermediates in the Li-Si system, alternative studies such as NMR, XAS, Raman, among others, must be implemented in order to study the shortrange order interactions and relate them with the electrochemistry phenomena. For instance Grey et al. [104,109,189,190] have devoted great part of their studies to the identification of these intermediates, indeed the lithiation happened to be more complicated than previously thought and dependent on several factors. By studying the 7 Li NMR, it is possible to identify different lithiated intermediates. As discussed previously in Chapter I, the lithiation starts by the Si surface, the presence of Li in the Si framework eventually results in the breaking of the Si-Si bonds, with the loss of all the long range ordering in c-Si (as indicated by the PDF measurements) and the formation of small Si clusters. Given the large associated energy to break the crystalline Si network, once the lithiation has started and the small clusters are formed (stars, rings and dumbbells), it is kinetically easier to continue breaking these clusters than the original Si framework. Both processes compete and prevent the system to reach equilibrium. Due to the inhomogeneity of the processes no single type of cluster is present at this stage, hindering the crystallization of any phase as it will require bond breakage and rearrangement involving high activation energies. The breakage of the small clusters continues until the formation of mainly isolated Si ions, from which the c-Li15Si4 nucleates. The preference in the nucleation of c-Li15Si4 over Li22Si5 ( a phase composed also of isolated Si ions), lies in the ability of the first one to accommodate defects such as the dumbbells that are presumed to remain at this stage. [181] Additionally the atomic arrangements of a-Li3.75Si and c-Li15Si4 are very similar, favoring the crystallization of this last one and no other phase. [51] These phenomena originate characteristic Li environments that are related with the degree of lithiation. This is, as the Li/Si ratio increases the chemical shift moves to lower frequencies and the shielding around the Li ions increases. i) signals between 16-22 ppm are attributed to small Si clusters like in Li12Si7 (planar Si5 rings and Si4 stars), ii) signals between 12-14 ppm are attributed to Si dumbbells like in Li7Si3. iii) signals between 8-10 ppm described lithium ions interacting with both dumbbells and isolated Si 4-ions like in Li13Si4. iv) signals between 3-6 ppm describe Li interacting exclusively with isolated Si ions like in Li15Si4. v) signals between -1 to 3 ppm are attributed to Li species in the SEI layer. [110,181,191,192] During the first discharge ( discussed above. As the potential decreases the peak at 18-16 ppm increases in intensity indicating the formation of more small clusters, which later are broken into isolated ions in agreement with the shift of this peak to lower frequencies and the decrease of its intensity. At 50 mV both small clusters and isolated Si ions are present and eventually the resonance at 6-3 ppm (isolated Si ions) dominates coinciding with the crystallization of c-Li15Si4. At the same time the shift corresponding to isolated ions moves to higher frequencies, indicating that the remaining a-LixSi phase contains fewer isolated Si ions. This suggests that c-Li15Si4 nucleates inhomogeneously from regions with high contents of isolated Si ions, leaving behind an a-LixSi phase with higher amounts of small clusters. Below 30 mV a resonance at -10 ppm is observed only in the in situ experiments, the negative shift suggests the presence of a more shielded environment. This resonance grows at the expense of the peaks with positive shifting, its intensity is enhanced by holding the voltages at 0 V but it rapidly disappears if the cell is relaxed.

The XRD findings at this stage indicate a diffraction pattern corresponding only to c-Li15Si4, meaning that there is no change in the macrostructure. Thus, since Li15Si4 is an electron deficient phase with high degree of defects and electron deficient, it is possible to accommodate extra Li to form c-Li15+δSi4, since for the system it is more favored to further lithiate c-Li15Si4 than to convert the residual a-LixSi into c-Li15Si4 originating the negative shift in the NMR experiment. [110,181,191,192] The fact that this metastable phase disappears rapidly upon charge or relaxation is associated with a process of "selfdischarge" due to a reaction with the electrolyte. The fact that this phase is only observed for the in situ experiments is possibly related to the different cycling conditions, defects, particle size, etc. For instance an electrochemically formed c-Li15+δSi4 phase from a pristine Si electrode presents a different 7 Li NMR shift compared to the one obtained from mechanically synthesized c-Li15Si4 (discharged to 0 V). [110,181] Finally the formation of c-Li15Si4 is incomplete, in some cases there are residual unbroken Si-Si pairs that limits the maximum of Li atoms that can be bonded to Si, explaining the inability of the system to go towards higher x values, x=4.2 or 4.4 like in Li21Si5 and Li22Si5. [193] During charge (Figure II.3), the delithiation occurs via formation of small Si clusters that serves as nucleation sites for bigger ones from which the amorphous network is built. This process involves the addition of another Si ion to an existing cluster and is favored over the nucleation of several small clusters which will require Si ions to migrate and recombine. Note that the kinetics of the reaction is also involved and affects the composition of the formed intermediated phases. [181] These phenomena originate a gradual loss of the NMR signal, both resonances at 18-16 and 6 ppm (Li in Si-Si clusters and Li in isolated Si ions, respectively) loose intensity, and above 440 mV only a signal between 5-6 ppm is preserved. As the potential increases this signal shifts to lower frequencies, originating a signal between 2-3 ppm ascribed to a Li close to a bigger Si cluster, where little transfer of electron density between the Si clusters and the Li + ions have taken place, meaning that the Li ions are still in a diamagnetic environment with a high partial charge. [110,181,191] The second cycle happens in the amorphous domain, this matrix is more accessible to the Li ions compared to the crystalline one, resulting in a more complete lithiation at the end of discharge. At first the transformation of a-Si into a-Li2.0Si a phase with extended Si networks and large Si-Si clusters is observed between 300-200 mV, originating an increase in intensity of the peak at 18-16 ppm (Figure II.3). At 100 mV the Si-Si bonds are broken to form dumbbells and eventually isolated Si ions (composition a-Li3.5Si). At 50 mV the crystallization of c-Li15Si4 from a-LixSi takes place, at this point the peak at 5 ppm dominates. Finally at 30 mV an overlithiation of the c-Li15+δSi4 is observed. [110,181,191,192] If the cutoff voltage is set to 85 mV, at this stage there are more clusters than isolated Si ions. Hence, upon delithiation, these clusters serve as nucleation sites for the reconstruction of the amorphous network by the addition of isolated ions and/or the fusion of several small clusters. This is translated in a decrease in the signal of the 18-16 ppm peak after 300 mV. This peak shifts to lower frequencies close to 10 ppm at around 450 mV, indicating a depletion of small clusters due to a preferential delithiation while the isolated Si ions still present in the system (like in Li13Si4). Beyond this point the 10 ppm resonance disappears and a new one appears at 3 ppm nearly at the end of charge. [181] Figure II.3. Ex Situ 7 Li NMR MAS of batteries stopped at a selected potential during a) first discharge, b) first charge and c) second discharge. [181]

Germanium

The relatively high cost of metallic germanium (Ge) as a lithium-ion battery negative electrode material is more than counterbalanced by its high capacity, high lithium diffusivity, and electronic conductivity. Using a unique and highly complementary set of operando characterization techniques, we propose a complete mechanism for the reversible lithiation of Ge. The electrochemical mechanism is characterized by: i) the formation of amorphous a-LiGe as the first intermediate during the lithiation of c-Ge independently of the charge/discharge rate, followed by Li7Ge3, and ii) the observation of c-Li15Ge4 at low potential, for moderate rates while at low rates more lithiated phase is preferred. The complementarity of the data obtained from XAS, Raman spectroscopies, and XRD, all in operando mode, was crucial in order to understand the complex mechanism based on reversible formation of the various crystalline and amorphous phases.

Electrochemical lithiation Mechanism

A perspective from literature

Most reports in the scientific literature are devoted to the improving of the Ge electrode performance, [168] particularly regarding the optimisation of the volume expansion by using nanostructuring [52,[START_REF] Graetz | [END_REF][194][195][196] or composites, [197][198][199][200] while fewer information is focused on the basic lithiation mechanism. [36,57,[201][202][203][204] Other studies demonstrate that the Ge particle size and morphology have a strong impact on the lithiation mechanism and stabilized phases. [36,55,57,104,[201][202][203][204][205] Indeed, many different amorphous and crystalline phases can be obtained upon lithiation (Table II.1). The results are, however, difficult to rationalize; even if it seems clear that crystalline Li15Ge4 and Li15+δGe4 are the end-products under electrochemical conditions, there is still no consensus on the different intermediated phases. Moreover, the presence of such intermediate phases depends strongly on the current density, revealing the importance of kinetics in the electrochemical mechanism.

Table II.1. Literature review of the chemical species formed during lithiation of germanium electrodes

Author (ref.) Type of electrode C/Rate Proposed species during lithiation Identification Technique

Yoon et al. [205] Micrometric c-Ge and carbon-coated composite prepared by ball milling and pyrolysis

100 mA/g c-Li9Ge4 c-Li7Ge2 c-Li15Ge4 c-Li22Ge5
Ex situ XRD Baggetto et al. [201,202] Evaporated/sputtered a-Ge thin film (50 nm thickness) C/20 a-LiGe a-Li7Ge2 c-Li15Ge4

In situ XRD In situ XAS Graetz et al. [START_REF] Graetz | [END_REF] Evaporated/ballistically deposited a-Ge thin film

C/4 c-LiGe c-Li7Ge2 c-Li15Ge4 c-Li11Ge6 c-Li9Ge4 c-Li22Ge5 a-LixGe

Ex situ XRD c-Ge upon delithiation

Lim et al. [203] Micrometric c-Ge C/15

a-Ge a-Li9Ge4 a-LiGe a-Li7Ge2 a-LixGe c-Li15Ge4
Operando XRD and XAS

Lim et al. [57,204] Micrometric c-Ge Jung et al. [36] Micrometric c-Ge

C/50 c-Li7Ge3 a-Li7Ge2 a-Li13Ge5 c-Li15Ge4 c-Li15+δGe4
Ex situ XRD In situ/Ex situ Solid state NMR Pair Distribution Function

Loaiza et al. [56] Micrometric c-Ge

C/5-C/8 a-LiGe a-Li7Ge3 c-Li15Ge4 Operando XRD, XAS and Raman Spectroscopy. C/14 a-LixGe c-Li17Ge4
Tang et al. [55] Ge Nanorods embedded in multiwall carbon nanotubes C/5 a-Li2.26Ge a-Li3.5Ge c-Li15Ge4 c-Li15+δGe4

In situ XRD In situ Solid State NMR Morris et al. [104] ------

LiGe Li7Ge3 Li13Ge5 Li8Ge3 Li15Ge4 Li17Ge4

Density

Functional Theory calculations

In analogy with Si, the information obtained by XRD is very limited (mainly sensible to materials with long-range order) and other characterization techniques are implemented, for example, X-ray absorption spectroscopy (XAS), which is widely used for heterogeous and amorphous materials with no long-range order but with a well-defined coordination around the central atom. The XAS spectra can be divided into two parts, the XANES (X-ray Absorption Near Edge Structure) and EXAFS (Near Edge X-ray Absorption Fine Structure), the first one is sensitive to the oxidations state, bonding environment and local geometry around the central atom, while the second one allows to determine the interatomic distances, coordination number and type of ligands. Table II. 1 shows some of the different lithiated intermediates and the respective characterization techniques used to identify them. It is worth to note that Graetz et al. [START_REF] Graetz | [END_REF] have identified some of the lithiated intermediates by ex-situ XRD, this point is particularly surprising since most of the intermediate phases are reported to be amorphous, hence, the presence of crystalline phases in this study might be related with the sample treatment before the XRD analysis, which given the metastable nature of the involved phases could induce to their crystallization. The XAS technique has been implemented commonly for the identification of the intermediates, and phases such as a-LiGe, a-Li7Ge2, a-Li7Ge3 and a-Li9Ge4 have been identified, these phases present very similar properties and their interpretation from the results could variate depending on the used method. It is concluded as well that one or more of the previous phases coexist at the same potential, thus the lithiation process is highly inhomogenous. [56,201,203,204] The solid state NMR is another powerful technique to identify the different lithiated intermediates which in analogy to the findings for silicon have a characteristic signal at different frequencies. In this sense, signals between 20-25 ppm are attributed to phases with Ge dumbbells like in Li7Ge3 and Li9Ge4, with the formation of Li7Ge3 being more energetically favored. Signals between 12-14 ppm are attributed lithium ions interacting with both dumbbells and isolated Si 4-ions like in Li7Ge2 and Li13Ge4. Signals between 0-10 ppm describe Li interacting exclusively with isolated Si ions like in Li15Ge4, while the ones between -24 to -21 ppm featured a process of overlithiation of the Li15Ge4 into Li15+δGe4. This process of overlithiation is identified merely by means of NMR spectroscopy, where it is possible to observe the difference of the chemical shift of the phase due to a new Li environment that is more shielded.

Other characterization techniques did show any difference in the results at this voltage. Indeed the Li15Ge4 is an electron deficient phase (does not satisfy the electron counting rules of the Zintl-Klemm concept, as observed in Chapter I, section 4.2) and can accommodate extra charges without any change in the crystal structure, this is a purely electrochemical driven reaction that terminates once the cell is allowed to rest at OCV. [36,55]

New insights by Operando spectroscopy and diffraction

The electrochemical response for the Ge electrode at different rates delivers a practical capacity of 1370 and 1351 mAhg -1 at C/5 and C/14, respectively (Figure II.4). For both C/rates, the galvanostatic curve presents four main processes during the lithiation. From 1.0 to 0.7 V the continuous drop in potential is attributed to partial reduction of the electrolyte leading to the formation of the solid electrolyte interphase (SEI) at the surface of the electrode. [80,[START_REF] Graetz | [END_REF]202] From ≈380-250 mV and ≈250-150 mV, two pseudo-plateaus indicate the coexistence of two phases, and are followed by a drop of the potential (after ca. 200 mV), corresponding to a one-phase region. Upon delithiation, the slope increases abruptly until approximately 350 mV (independently of the rate used), followed by the presence of two pseudo-plateaus at 420-450 and 590 mV for C/14 and C/5 respectively. The second lithiation is slightly different, the first pseudo-plateau (at 410 mV) presents a steeper slope, whereas the second (at 250 mV) is more marked. Both processes, however, occur at higher potentials, suggesting the electrochemical path followed during the second discharge is possibly different. All the processes in the galvanostatic curves are well defined in the derivative curves, except for the intensities of the peak at 590 mV during charge and the broadening of the peak at 380 mV during the second discharge indicative of a change in the nature of the Ge network compared to the pristine sample.

Figure II.4. Galvanostatic discharge-charge curves of the Ge electrode at a) C/5 and b) C/14 rate (second discharge in dashed lines) and their corresponding derivative curves (c, d).

The lithiation of germanium is isotropic; this difference compared to Si lies in the properties of the crystallographic planes of unlithiated Si (Ge) crystals that adjoin the amorphous product. This is, the orientation dependence of interfacial mobility at the sharp boundary of the two phases (pristine crystal and amorphous lithiated intermediates) is expected to govern the lithiation anisotropy in Si, rather than the long-range transport. While Si has a preferential lithiation the (110) planes, Ge does not experience such phenomena and lithiates evenly in all the directions without the formation of cracks during lithiation. These factors influence greatly the final electrochemical performance, an in general an isotropic expansion is favoured in order to extend the lifetime of the battery. [50] Li15Ge4. [36,104,206] Indeed crystalline Li15Ge4 coexists with Li7Ge2 or Li7Ge3. In contrast at C/14, after the full amorphization of c-Ge, there is no growth of any new peaks during the insertion of about 2.7 Li, only three broad peak centred at about 23.3°, 24.5° and 40.5° appear during the end of the discharge, and no crystalline Li15Ge4 is observed. Again it is difficult to unambiguously associate these features to Li7Ge3, Li7Ge2 or Li17Ge4, whereas the finally achieved insertion of 4 Li is more in agreement with either Li7Ge2 or Li17Ge4. [36,104,113] Figure II.7 shows the diffraction pattern for the discharged product of this operando XRD. The observed Bragg reflections are likely attributable to Li17Ge4 (main peak at 40.5°), in analogy with the reports of Iwamura et al. [207] and Zeilinger. [START_REF] Zeilinger | Synthesis , Characterization and Reactivity of Lithium-Containing Silicides , Germanides and Borosilicides[END_REF] Upon delithiation at the C/5 rate, the Li15Ge4 phase rapidly disappears and only the humps centred at 24.5° and 41° are still visible (Li7Ge2 or Li7Ge3) with a shift to higher angles, corresponding to the volume contraction upon Li extraction. For the C/14 rate the broad peaks disappear and are replaced by the hump at 24.5°, which continuously shifts to 26° at the end of charge. In both cases the humps eventually disappear and give place to a completely flat diffraction pattern, confirming the formation of a-Ge and/or low lithiated a-LixGe. [36] Figure Most substantial differences in the XRD evolution between the two C-rates are observed for the second discharge/charge cycle. For the C/5 rate, the second discharge path that resembles the first one, but with a better crystallization of Li15Ge4 at the expense of the humps centred at 2θ= 24.5° and 41°. Indeed, the presence of a-Ge or low lithiated a-LixGe rather than pristine bulk c-Ge may contribute to better kinetics of lithiation, facilitating the formation of Li15Ge4. In stark contrast, for the C/14 rate, no Li15Ge4 is formed at the end of the second discharge; reason why the two broad peaks centred at 23.3° and 24.5°, together with the reflection at 40.5° (with an increased intensity), are more likely to be a highly lithiated phase such as Li17Ge4 (in agreement with x≈4.1) rather than Li7Ge2. [104,113] The possible formation of such highly lithiated phases has been previously observed by Yoon et al., [205] Iwamura et al., [207] Zeilinger, [START_REF] Zeilinger | Synthesis , Characterization and Reactivity of Lithium-Containing Silicides , Germanides and Borosilicides[END_REF] Goward et al. [113] and predicted by Morris et al. [104] for very low potentials, close to 0.10 V. The presence of a-Ge cycled and the optimised electrochemical interfaces along with the low rate cycling render such very highly lithiated phases accessible.

For both rates the Bragg peaks of the crystalline phases that grow during the discharge process initially slightly shift to lower angles (Figure II.6-b) during the discharge and then towards higher angles during the beginning of charge, in agreement with an increasing/decreasing of the crystallographic cell parameters during lithiation/delithiation. [36,72] This structural evolution might be related to the existence of solubility domains. Such domains have already been explained as an overlithiation process transferring c-Li15Ge4 into Li15+δGe4, as observed by in situ solid state NMR by both Jung et al. [36] and Tang et al. [55] conditions, in agreement with the literature. [194] As the lithiation progress, this peak decreases continuously until its complete extinction at 240 mV, in line with the depletion of c-Ge (also observed by XRD), through the first biphasic process leading to an amorphous a-LixGe intermediate. Shortly after (at ∼200 mV), as the potential keeps decreasing, a new weak peak that appears at around 220 cm -1 , grows and shifts to lower frequencies, until reaching 208 cm -1 at the end of discharge (∼ 20 mV). During delithiation (Figure II.8-b), two broad signals emerge at ∼200 (∼500 mV) and 278 cm -1 (∼750 mV). The first one losses intensity and disappears at ∼1 V, while the second one gradually shifts to ∼300 cm -1 , reaching a maximum intensity at ∼1 V. The peak at 220 cm -1 that appears towards the end of the discharge pairs with the spectrum of the Li15Ge4 reference (Figure II.8-c, D). Its shift towards lower frequencies suggests the evolution of the freshly formed Li15Ge4 to a new amorphous LixGe-like states which could possibly correspond to Li7Ge3 (or Li7Ge2), proposed to compete with Li15Ge4 at the end of discharge, based on the XRD analysis. The feature might, however, also correspond to the evolution of Li15Ge4 to Li15+δGe4, as observed by Jung et al. [36] and Tang et al. [55] by in situ solid state NMR and already discussed above (vide supra). In order to complement the latter propositions, a Ge electrode discharged down to 0.01 V and held there for 40 hours (Figure II.8-c, C), in order to allow a better crystallization of the lithiated Ge phase(s) was measured. The resulting spectrum shows several features matching those of the Li15Ge4 reference (Figure .8-c, D), and also resembles the spectra of the analogous LixSi phases. [208] After the crystallization process, however, the main peak is well centred at 222 cm -1 , corresponding to the presence of "normal" Li15Ge4. Finally, during delithiation, the presence of a final peak at ∼300 cm -1 indicates the reformation of Ge-Ge bonds, while its broad shape and slightly shifted position compared to pristine c-Ge agrees well with the formation of a-Ge during the charge process, as observed in the previous XRD measurements.

Figure II.8. a) Galvanostatic discharge-charge curves at C/5 rate for the Ge electrode, b) evolution of the active Raman modes during the galvanostatic cycling and c) Raman spectra: (A, B) from the operando experiment at E=0.20 and 0.06 V, respectively, (C) in situ of the Ge electrode after discharging and 40 hours of potentiostatic step at 0.01 V, (D) ex situ of the Li15Ge4 reference electrode prepared by ball milling.

Further complementary information about the electrochemical mechanism of Ge during lithiation was gathered by operando XAS, a site selective and local probe technique ideal for the study of amorphous phases. In total, 104 spectra were collected during the first discharge at C/8 rate (a rate chosen as optimal in terms of low rate cycling without exceeding the beam time allocated) and are shown in Figure II.9. Previous XAS studies have mainly centred on interpreting the signal corresponding to the first peak of the Fourier transform (FT) of their extended fine structure (EXAFS) in terms of contributions of Ge-Ge and Ge-Li shells. [57,201,203,204] Here, we have proposed a different analysis approach based on chemometric tools that take into account the whole evolution of the XAS spectra during the lithiation. This approach was used in order to extract the maximum information from this new whole series of operando XAS spectra, especially on the intermediate phases. First, the number of orthogonal "pure" components was obtained by Principal Component Analysis (PCA); then, by Multivariate Curve Resolution-Alternating Least Squares (MC-ALS) method, we were able to reconstruct these "pure" components. In fact, this statistical method allows not only to reduce the dimension's problem but also to reveal the hidden intermediated species.

Figure II.9. XAS spectra collected during the first discharge at C/8 for a Ge electrode and reconstructed components from the analysis of the whole series: a) XANES and b) EXAFS regions of the experimental spectra (the evolution of the spectra from the beginning to the end of the discharge is represented by spectra going from dark blue to green); c) XANES and d) EXAFS regions of the reconstructed components

The PCA (Figure II.10) strongly indicates the existence of 4 dominant components expressing more than 99.9% of the variance of the whole series of spectra. The inspection of their shape and of the evolution of their scores strengthen this assumption. These four components were then used in the MCR-ALS analysis, which provides "pure" spectral components expressing, by their linear combination, the main modifications of the XAS spectra. The evolution of their respective relative concentrations along the series of spectra is shown in By observing the shape of the absorption edge of the four components, we find that the evolution of the XANES spectra through the first discharge cannot simply be expressed as a linear combination of the XANES spectra of the pristine and the fully lithiated material, but at least two new XANES spectra are needed. Since the shape of the XANES spectra is directly related to band structure, this means that at least two additional phases are formed during the first discharge. The analysis of the EXAFS portions of the MCR-ALS components show the fit in both k and R spaces (Figure II.12, Table II.

2).

The fitting of the first component, representing pristine c-Ge, is rather trivial, and is performed up to R = 5 using the first three distance shells of the diamond structure of Ge. [209] While the bond distances coincide perfectly with those expected from the structure, rather high values of mean-square relative fluctuations (σ 2 ) are observed for the second and the third shell, indicating a somewhat disordered structure or a partial amorphization of the pristine c-Ge, in line with the ball milling activation pre-treatment before preparation of the electrodes. The FT of the second component shows the same three peaks as for the first component. However, the peaks are now globally less intense and have different intensity ratios, suggesting this component does not represent a pure compound, but rather a combination between some remaining c-Ge and at least one lithiated phase. A fit was attempted using a two-phase approach, dividing the total intensity of the spectrum between the three-shell contribution of c-Ge and the two-shell contribution of LiGe, where the different Ge-Ge and Li-Ge bonds distances were averaged out. While the LiGe fraction dominates, more than 2/3 of the total component intensity corresponds to the three-shell of c-Ge, like the pristine material, but with clearly lower σ 2 ; indicating an increase in the crystallinity. Taking into account the general amorphization during this part of the discharge, an increase in crystallinity is rather unexpected, but this apparent incoherence can be explained by the fraction of more a-Ge, contained in the pristine sample, reacting first and thereby increasing the relative concentration of better crystallised c-Ge reacting at a later stage.

The FT of component 3, which replaces component 2 about halfway through the discharge, shows only a main contribution corresponding to the first coordination shell of Ge, which can be shared between a Ge-Ge and a Ge-Li shell. It is interesting to note the shifting of the main peak of FT towards higher R distances compared to the first two components, associated to an important change of the unit cell volume. These data could not be fitted by using the typical Ge-Ge and Ge-Li shells intensity of Li7Ge2, whereas a better fit was obtained by increasing the Ge-Ge contribution, indicating Li7Ge3 as an alternative. [104,210] Therefore, we used coordination shells simulated for Li9Ge4, a phase with a similar composition and numbers of Ge and Li neighbours. Both phases contain the same type of Ge-Ge dumbbell-like structures and taking into account plausible longer Ge-Ge distances in Li7Ge3, the fit confirms this component to be due to Li7Ge3.

Finally, component 4 could be easily fitted with three shells: two Ge-Li shells describing the first coordination shell of Ge, and a Ge-Ge shell at longer distances indicating a somewhat better mid-range order. These shells correspond perfectly with the bond distances observed in Li15Ge4, in line with the formation of this phase with a relatively good crystallinity at the end of the discharge.

Discussion and proposed electrochemical mechanism

By correlating all the information provided by operando XRD, Raman spectroscopy, XAS and the electrochemical results, we can follow the relation of the electrochemical mechanism of Ge upon lithiation at different C/rates:

At relatively higher C/rates (C/5 and C/8), the lithiation of c-Ge leads to a phase amorphization, followed by the formation of amorphous/poorly crystallized intermediates; LiGe as first lower lithiated state and Li7Ge3 as second one. Subsequently, at the end of discharge there is a competitive growth between these amorphous lithiated states and c-Li15Ge4. The mechanism is reversed during delithiation with the eventual formation of a-Ge. Hence, the lithiation of Ge occurs by the successive biphasic processes and can be summarized as follows:

Discharge: c-Ge  LiGe  Li7Ge3  c-Li15Ge4 Charge: c-Li15Ge4  LiGe  Li7Ge3  a-Ge
At lower C/rates (C/14), the lithiation of c-Ge also leads to amorphous intermediates, but with the formation of Li17Ge4 at the end of the first discharge and the absence of c-Li15Ge4. Upon delithiation, a-Ge is present. The mechanism can be described as:

Discharge: c-Ge  a-LixGe  c-Li17Ge4 Charge: c-Li17Ge4  a-LixGe  a-Ge
The second discharge/charge cycles in both cases is characterized by the enhanced crystallinity of the different phases/states due to the presence of a-Ge and optimized electrochemical interphases.

On the other hand, the end of the discharge/beginning of charge for both C/rates is characterized by a solid solution process with cell parameters modification that is related to the existence of solubility domains and to a process of accommodation of extra charges in the defects of the structure (overlithiation).

The formation of the Li17M4 phase has been controversial, and has its origins in the discrepancy between the formation of Li22M5 and Li21M5, which originally was thought to be Li22Si5 due to the similarities in the characteristics of both space groups, however the formation of Li22Si5 will originate unreasonable short interatomic distances, thus the stabilization as the Li21Si5 is the most reasonable.

In fact, these both phases differ in the positions occupied by 4 Li atoms, which for Li21Si5 cannot be occupied. Later on, Nazar et al. [113] , determined that some of these positions are indeed occupied and the formula corresponds more to Li21+5/16M5 or approximately Li17M4. The occupation of these special sites varies depending on the element in group 14, with an increased occupancy as the series progress, and the interatomic distances have reasonable values. The bigger size of Ge, Sn and Pb allow the occupation of extra tetrahedral sites in the unit cell stabilizing structures of the type Li17M4, silicon on the contrary due to its smaller size does not follow this tendency and stabilizes in Li21Si5. [108,113] However, Zeilinger et al. [108] claim that the given the similarities between Li17Ge4 and Li17Si4, this last one could be the real structure for Li21Si5.

Summary and conclusion

In summary, we have investigated the electrochemical mechanism of a Ge electrode during lithiation/delithiation by operando XRD, Raman spectroscopy and XAS. The results indicate that the lithiated phases/states formed during lithiation/delithiation are kinetically dependant on the C/rate.

Meaning that relatively higher C/rates promote the growing of c-Li15Ge4 at the end of discharge after the successive formation of amorphous/poorly crystallized low lithiated LixGe (x=7/3 or 7/2) states whereas, low C/rates are characterized by the absence of c-Li15Ge4 and the presence Li17Ge4. In both cases during delithiation, the mechanism is reversed but the crystalline Ge network is not rebuilt and an amorphous product is obtained (as confirmed by the Raman and XAS results). Therefore, the cycling rate and the morphology of the pristine electrodes play an important role in the kinetics and the participating phases in the electrochemical mechanism; underlining that the differences in the literature arise from the material and the conditions used for cycling.

Finally, this work illustrates the high interest of performing long-range and short-range sensitive experiments in operando conditions and brings new essential information to understand the puzzling mechanism of alloys-based electrode materials like Ge in order to improve their performances.

Si0.5Ge0.5 Solid Solution

GexSi1-x alloys have demonstrated synergetic effects as LIB anodes; due to the silicon high lithium storage capacity and the superior germanium electronic and Li ionic conductivity and capacity retention. Previous studies primarily focused on intricate nanostructured alloys with high costs of production, here we studied the simpler Si0.5Ge0.5 alloy as a composite electrode. The electrochemical mechanism is explored by a combination of in situ and operando techniques such as XRD, XAS, Raman spectroscopy and 7 Li solid state NMR, all providing unique and complementary information about phase transformations during cycling. It was possible to observe the amorphization of c-Si0.5Ge0.5 upon lithiation and the crystallization of a new phase at the end of the discharge. This last one presents an evolution of the refined cell parameters related to an overlithiation process. The crystallinity of Si0.5Ge0.5 was not restored upon delithiation and an amorphous phase was obtained. The results have been combined with a series of ex situ experiments to better comprehend the mechanism.

The results from this section have been adapted from the publication:

Previous reports and background

Various formulations of Si1-xGex alloys with promising LIB anode performance have been reported, [73,164,[211][212][START_REF] Phan | Etude et développement de couches minces de type Si1-xGex pour une utilisation comme électrode négative dans des microbatteries Li-ion[END_REF] with most of the works performed on complex nanostructures resulting from high-cost synthesis. Abel et al. [164] found that Si1-xGex nano-columnar thin films have increasing electronic conductivity and specific capacity with increasing germanium and silicon content, respectively; while a high Si content is related with a lower capacity retention. Ge et al. [211] investigated the performance of Si56Ge44 nano-particles with better capacity retention compared to both Si-rich and Ge-rich nanoparticles, originated from a structural evolution during lithiation that decreased the formation of additional SEI layer. Phan [START_REF] Phan | Etude et développement de couches minces de type Si1-xGex pour une utilisation comme électrode négative dans des microbatteries Li-ion[END_REF] studied a thin film of Si0.7Ge0.3 deposited by cathodic pulverization with a capacity retention value in-between pure Ge and Si.

The studies of bulk electrodes are relatively scarce; Hashimoto et al. [START_REF] Hashimoto | [END_REF] synthesized Li4.4GexSi1-x alloys by means of ball milling and used them as negative electrodes in all-solid-state batteries.

Particularly, their Li4.4Ge0.67Si0.33 alloy presented a specific capacity of 1900 mAh/g with good chargedischarge reversibility. More recently, Duveau et al. [73] studied the Si0.9Ge0.1 composition as composite electrodes and demonstrated that the addition of Ge improves the performance in terms of cyclability as compared to a pure Si electrode. Their in situ XRD study showed formation of Li15(Six/Gey)4 to take place at the end of the discharge, accompanied by a structural evolution.

Synthesis and characterization

Synthesis and Characterization of the Si0.5Ge0.5 and Li15M4 (M=Si, Ge) phases.

The Si0.5Ge0.5 phase and the Li15M4 references were synthesized following the procedure described in section 2 of the present chapter. phase was observed, therefore the synthesis approach for sample B leads to a phase with higher purity.

The characterization by Raman spectroscopy is guided by the signatures of SiXGe1-X alloys which are well known in the literature. [215,216] The Si0.5Ge0.5 phase (Figure II.14-top) undergoes partial amorphization during the ball milling process leading to a red shift of the Si-Si signal from 520 cm -1 to 480 cm -1 . Similarly, the Ge-Ge signal showing a band at 300 cm -1 red-shifts to 291 cm -1 . The formation of Si0.5Ge0.5 alloy can be evidenced by a band at ca. 400 cm -1 assigned to the Si-Ge bond. [208,[215][216][217] For the lithiated references (Figure II.14-bottom) it was possible to identify a peak at 217 cm -1 for Li15Ge4, while, for the other phases, only the signal coming from the quartz window could be recorded. 16 shows the 7 Li MAS NMR experiments were performed on mechanically synthesized Li15Si4 and Li15Ge4. Despite being structurally related, their NMR signatures differ. [53] For Li15Si4 two sharp components at 10 ppm and 3 ppm, and a broad one at 16 ppm are observed. These findings are in relatively good agreement with the literature that reports shifts between 6 and 8 ppm (see section 3.1 of the present chapter). [110,192] Note that spinning of the sample at 62.5 kHz increases the temperature to 60-70°C which might change slightly the values of the shift. The other contributions are most probably due to decomposition products of this metastable phase, which is more remarked at high temperatures and could include Li13Si4 phases. [218] For Li15Ge4 there are three narrow contributions at -7, 3 and 7 ppm with a broad one at 4.6 ppm.

The component at -7 ppm is assigned to Li15Ge4, based on literature. [36,55] The remaining peaks probably come from partial decomposition of Li15Ge4 into compounds like Li7Ge3 (Ge-Ge dumbbells) or Li7Ge2 (isolated Ge 4-atoms), as observed by Jung et al. [36] on aged Li15Ge4 samples. This sample was indeed measured two months after synthesis. 7 Li MAS NMR spectra for samples A (Li+Si0.5Ge0.5), B (Li+Si+Ge), C (Si0.5Ge0.5 discharged and allowed to rest at OCV) and D (Si0.5Ge0.5 discharged and held at 0 V). The experimental spectra (blue) are overlaid with the fit (dotted red). Individual components of the fit are shown below each spectrum.

In order to investigate the impact of aging on the sample, freshly electrochemically synthesized Li15(Si0.5Ge0.5)4 was analyzed. For this purpose the cells were discharged down to 0 V. One cell was allowed to rest at open circuit voltage mode (OCV) (sample C) while another was held at 0 V (sample D).

The spectra (Figure II.17) of both samples contain a contribution at -10 ppm and confirm the assignment of Li15(Si0.5Ge0.5)4. The signal at 0 ppm is associated with the SEI and the remaining electrolyte. [36,55,110,191,192] For sample C, the main contributions are centered at -11 ppm (72%) and -12 ppm (14%), with an additional peak at 10 ppm (9%). The peaks at -11 and -12 ppm are assigned to Li15(Si0.5Ge0.5)4, with slight differences in the overlithiation or disorder, consequence of cycling at low voltage. [36,55,110,191,192] The peak at 10 ppm is most probably issued from amorphous Si/Ge phases produced by decomposition during the OCV. For sample D, a peak with two contributions at -4 ppm (54%) and -10 ppm (45%) is recorded. The observation of a signal -4 ppm is indicative of the particularly high reactivity of Li15(Si0.5Ge0.5)4; we observed this parasitic phase even for freshly prepared samples.

In situ studies are therefore essential to catch the reactivity of Si0.5Ge0.5 during (de)lithiation.

Electrochemical mechanism study and structural evolution upon cycling

The Si0.5Ge0.5 and pure-ball-milled Si and Ge electrodes were cycled at C/14 (Figure II.18). The capacity in the first discharge was 3709 mAhg -1 , 2026 mAhg -1 and 1351 mAhg -1 for Si, Si0.5Ge0.5 and Ge, respectively, in good agreement with the literature. [183,191,198,[219][220][221] For the three systems part of this capacity is assigned to the irreversibly formation of the SEI layer and decomposition of the electrolyte. [52,197,217,222] The delithiation and second lithiation processes happen at higher voltages compared to the first discharge. [193] In the literature the Si and Ge electrochemistry has been extensively studied; for Si (Figure II.18-a) the derivative curve shows a small broad peak during the first discharge around 250 mV (A) attributed to lithiation of c-Si into a-LixSi intermediates. At 120 mV there is a broad peak (B) with a small sharp shoulder at 50 mV, coinciding with the crystallization of Li15Si4 at low potentials. Upon charge, the derivative curve shows a sharp peak at 450 mV (C), characteristic of the delithiation of c-Li15Si4. [179,180,182,183,187,223] For Ge (Figure II.18-c), the derivative curve shows the presence of two peaks (A and B) during the first discharge centered at 380 (A) and 250 mV(B), assigned to the formation of a-LixGe and the crystallization of Li15Ge4 below 150 mV. [START_REF] Graetz | [END_REF]198,201] Upon charge, the processes of delithiation of Li15Ge4 and formation of a-Ge (C and D) can be detected at 430 and 590 mV, respectively. [START_REF] Graetz | [END_REF]198,202] Note that additional details of the galvanostatic cycling of Si and Ge can be found in sections 3.1 and 3.2 of the present chapter.

For the Si0.5Ge0.5 composite (Figure II.18-c) the voltage drops to ≈280 mV in the galvanostatic curve, followed by a gently sloping profile until ≈100 mV, after which the potential drops until the end of discharge. The derivative curve displays one broad peak (B) around 200 mV with a small shoulder at 100 mV. Based on the trends for Si and Ge the broad peak could be assigned to the amorphization of c-Si0.5Ge0.5 to form a-Lix(Si/Ge) and to the crystallization of Li15M4 below 100 mV. Upon charging, the potential rises continuously until 480 mV, followed by a plateau and at the end of charge an increase in the slope indicates the presence of a one-phase region. These features correspond to the amorphization of the c-Li15M4 phase and the presence of a sharp peak at 470 mV in the derivative curve is a slight indicator of this biphasic process. The profile for the second discharge is different, with two pseudo-plateaus (two peaks in the derivative curve) at ≈300 and 100 mV. The absence of multiple peaks coming from pure Ge and Si confirms that the observed electrochemical behavior corresponds to the Si0.5Ge0.5 solid solution and not to an overlap of Si and Ge. [73,211] the formation of other amorphous lithiated phases. [179] The cell parameter for the XRD pattern collected at the end of discharge (xLi= 4.21) was calculated (at 10.745(2) Å) in the I-43d space group and it is an intermediate value between those of Li15Si4 (10.6546 Å-ICSD 43235) and Li15Ge4 (10.72(5) Å-ICSD 237945). [73,164,211] An EDX mapping indicates that Si and Ge are well-dispersed at this stage (Figure II.21). Thus, the Li15X4 phase at the end of discharge indeed corresponds to Li15(Si0.5Ge0.5)4. Our results are consistent with the reports for Si [180,[START_REF] Phan | Etude et développement de couches minces de type Si1-xGex pour une utilisation comme électrode négative dans des microbatteries Li-ion[END_REF]224] and Ge. [START_REF] Graetz | [END_REF]198,203] When the cell is charged, the intensity of the c-Li15(Si0.5Ge0.5)4 phase peaks gradually decreases until complete amorphization and the same humps observed during the first discharge reappear. At the end of charge, the material remains amorphous, in agreement with the reports for pure Si and Ge. [36,192,198,203] The second discharge is consistent with the first one, and despite a reduced Li insertion, the Li15(Si0.5Ge0.5)4 peaks present a higher intensity. Possibly, the configuration of the in situ electrochemical cell introduced a limitation in the lithiation/delithiation of the material. [73] . Just before the end of discharge the refined cell parameters start to decrease from 10.772(5) Å (xLi=4.07) to 10.745(2) Å (xLi=4.12). Upon charging, the decreasing tendency continues from 10.732 (8) The crystalline to amorphous transitions in the Si0.5Ge0.5 electrode limits the information that can be extracted from XRD, in this context 7 Li NMR brings valuable complementary information. The in situ 7 Li NMR experiment was performed for a self-supported Si0.5Ge0.5 electrode, cycled at a rate of C/5, using a cell especially designed for in situ solid-state NMR [178] (Figure II.24-28). The cell was discharged, followed by potentiostatic step at 0 V and an open-circuit (OCV) period and finally charged. Subsequently, the procedure was repeated. All along the experiment, a peak around 0 ppm could be observed and associated with the electrolyte, slightly lithiated phases and the formation of the SEI layer. [36,109,110] An additional signal registered at 250 ppm is assigned to metallic Li from the negative electrode.

Figure II.24.

In situ 7 Li NMR of a Si0.5Ge0.5 self-supported film cycled vs Li/Li + at the rate of C/5. a) Contour plot of the 7 Li NMR spectra as a function of time. b) stacked plot of the 7 Li NMR spectra with increasing time. c) voltage and current as a function of time.

In the first discharge (Figure II.24, 25), the evolution of the spectrum follows the same general trends as pure silicon [110,192] and pure germanium. [36,55] A small and broad peak at 20 ppm appears when the voltage reaches 200 mV. The intensity of this peak increases and it progressively shifts towards lower frequency to reach 10 ppm at 100 mV. Based on previous reports for Si [110,192] and Ge [36,55] this peak is assigned to a lithiated amorphous phase (a-Lix(Si0.5Ge0.5)) in which the SiGe network is not completely broken (Si-Si, Si-Ge and Ge-Ge small clusters and dumbbells). The slow shift towards lower ppm values is assigned to the further breakdown of the network, resulting in fewer dumbbells and more isolated Si and Ge ions. Interestingly, the voltages at which the spectrum evolves are different from Si [110,192] and Ge. [36,55] The peak at 20 ppm appears below 200 mV, between the typical voltages of silicon (100 mV) and germanium (300 mV). It reaches 10 ppm at 100 mV, lower voltage than germanium (150 mV), while the shift was already 6 ppm at 100 mV in silicon. These observations confirm the specific behavior of the Si0.5Ge0.5 solid solution, different from the simple sum of the Si and Ge.

When the potential reaches 95 mV a second peak appears with a negative shift (-17 ppm) (Figure II.25). Note that this peak is broad and covers the -10 ppm region assigned to Li15(Si0.5Ge0.5)4 from the ex situ 7 Li NMR experiments. It is therefore ascribed to the crystallization of c-Li15(Si0.5Ge0.5)4 and c-Li15+δ(Si0.5Ge0.5)4 (overlithiated phase) based on the reports for germanium [36,55] and silicon. [110,192] As the voltage decreases, this peak increases at the expense of the 10 ppm peak. The later disappears at the end of discharge, which indicates a good conversion of the amorphous lithiated phases to the c-Li15(Si0.5Ge0.5)4 and its overlithiated c-Li15+δ(Si0.5Ge0.5)4 phase. According to the literature it is kinetically more favored at really low voltages to overlithiate Li15M4 (M=Ge, Si) into Li15+δM4 than to form new Li15M4, as this last one implies the breaking of more Si-Si bonds to form sufficient isolated Si 4-atoms to allow the crystallization of c-Li15M4 [36,110,192] During this time, the peak at decomposition/evolution into other phases not detected by NMR. This process is accompanied by an increase in the potential of the cell and resembles the observations of Key et al. [110] for Si and Tang et al. [55] for Ge. The former study suggested that the disappearance of the peak with negative shift could be related to a reaction of Li15Si4 with the electrolyte and subsequent reformation of small Si-Si clusters. The latter identified that upon removal of the electrical contact right after the end of discharge, the peak at -24 ppm (Li15+δGe4) quickly disappeared and hence the formation of Li15 + δGe4 is an electrochemically driven reaction. Here the -17 ppm peak also disappears during the open-circuit step with no intensity increase in the other regions of the spectrum. This observation will be the subject of further investigation in the future. 7 Li NMR of a Si0.5Ge0.5 self-supported film cycled vs Li/Li + at the rate of C/5. shifts towards lower frequencies (-20 ppm at the end of the potentiostatic step). This behavior could be related with the amorphous nature of the Si0.5Ge0.5 network formed after the first cycle. The lithiation of all the a-Si0.5Ge0.5 is faster, hence the already formed Li15(Si0.5Ge0.5)4 is not overlithiated unless the voltage is held at 0 V. [36,110,192] Finally, near 250 ppm, the reversibility of the Li plating can be observed. The cell was then allowed to rest and like in the OCV period during the first cycle, the peak at -17 ppm broadens, shifts towards more negative frequencies and ultimately completely disappears.

Figure II.26. Extracted slices during the first open circuit (OCV) step (Left) and first charge (right) of the in situ

Upon charging (

Figure II.27.

Extracted slices during the second discharge (left) and the potentiostatic step at 0 V (right) of the in situ 7 Li NMR of a Si0.5Ge0.5 self-supported film cycled vs Li/Li + at the rate of C/5.

We believe that the structural evolution for c-Li15(Si0.5Ge0.5)4, found at the end of discharge/beginning of charge by means of in situ XRD, is closely related to the process of overlithiation of c-Li15+δ(Si0.5Ge0.5)4. The fact that the in situ XRD patterns recorded at the end of discharge corresponded only to c-Li15(Si0.5Ge0.5)4, with no other contribution from other phases, allows us to assume that the Li15+δ(Si0.5Ge0.5)4 is structurally related with Li15(Si0.5Ge0.5)4. In fact, the defects in Li15(Si0.5Ge0.5)4 could play an important role in accommodating the additional charges during the process of overlithiation without a structural variation. [225] In addition, the refined cell parameters for that phase, differ slightly at the end of the first and second discharge (a=10.745(2) and a=10.725(2) Å, respectively). This difference could be related with an incomplete overlithiation (smaller peak at -17 ppm) during the second cycle. The amorphous character of the Si0.5Ge0.5 phase in the second cycle may play a role since it will impose a difference in the energy required for the formation of the lithiated phases. The possible Ge-Li bonds formed during this part of the lithiation are very weak in intensity compared to both Ge-Ge and Ge-Si, and therefore difficult to observe (at least during the first part of the lithiation when the starting material dominates the overall EXAFS signal). The decrease in the signal can be completed by adding a Ge-Li shell (calculated from the crystal structure of Li15Ge4). Similar results were obtained by Lim et al. [203] in an in situ XAS study of c-Ge electrodes, who also attributed the decrease in the intensity of the Ge-Ge shells to the lithiation of Ge. Additionally, there is a slight but significant decrease of the Ge-Ge bonds compared to the Ge-Si bonds, in agreement with a preferential lithiation reaction of Ge rather than Si at the beginning of the discharge.

During the second part of the discharge the third component replaces the second, leading to a stronger decrease of the amplitude in line with a further depletion of the Ge from the reacting compound and with an increasing formation of lithiated Ge species. Finally, the third component is replaced by the fourth, which represents the final discharged product. In this spectrum, the contributions of Ge-Ge and Ge-Si have completely disappeared and are replaced by an oscillation at low k values that can be fitted with the first two Ge-Li shells of the Li15Ge4. In this case, the fitting region is restricted to 2.5 < k < 8. The nature of the reconstructed components is not straightforward to explain. Both PCA and MCR-ALS analyses are developed and more adapted to the analysis of bilinear systems (pure biphasic reactions for electrochemistry). For solid solutions the components that are created by MCR-ALS are fictive and allow us to follow gradual monophasic evolutions as sequences of several biphasic steps. This could be our case; the progressive replacement of one component by the next can be interpreted in terms of gradual evolution of the system. to the formation of a conductive phase. [224] Upon charging, both Ge-Ge and Si-Ge bands reappear, the peaks are though broader (due to the amorphous character of the material at this stage) and slightly shifted (due to a decrease in the particle size). Therefore, the Raman spectra confirms the re-formation of the Si-Ge bonds in Si0.5Ge0.5 alloy after delithiation, in agreement with the EDX mapping made on a charged Si0.

Summary and conclusion

The Si0.5Ge0.5 composition was successfully prepared by means of mechanical synthesis. The refined Si0.5Ge0.5 cell parameter was in between those ones of Si and Ge. The solid solution is Si0.5Ge0.5 is confirmed by Raman spectroscopy which showed the Si-Ge bond contribution at 400 cm -1 with additional resonances for the Ge-Ge and Si-Si bonds at 291 and 480 cm -1 , respectively. EXAFS analysis suggests the formation of subnanosized clusters of Ge at the very local scale, but a rather homogenous composition close to the nominal one in the medium distance range.

The electrochemical mechanism of the new negative electrode Si0.5Ge0.5 was studied by a combination of complementary technics XRD, XAS, Solid State NMR and Raman spectroscopy, which allowed to have a complete view of the structural modifications induced by the lithiation/delithiation. It was possible to observe the Si0.5Ge0.5 amorphization at the beginning of discharge. Further lithiation induces the formation of a-Lix(Si/Ge) intermediates and the crystallization of Li15(Si0.5Ge0.5)4 at the end of the discharge. At really low voltages a reversible process of overlithiation and formation of Li15+δ(Si0.5Ge0.5)4 was identified and related with a structural evolution of Li15(Si0.5Ge0.5)4. Upon charging, c-Li15(Si0.5Ge0.5)4 was transformed into a-Lix(Si/Ge) intermediates and at the end of charge an amorphous phase assigned to a-SixGey was recovered. Thereby, it was demonstrated that Si and Ge are conjointly active in the phase along the cycling, upon discharge with the formation of a ternary Li15(Si0.5Ge0.5)4 phase (with a step of overlithiation) and upon charge with the rebuilding of the a-Si-Ge phase. This process is undoubtedly behind the enhanced performance of Si0.5Ge0.5 compared to a physical mixture of Si and Ge. [73]

Conclusion

The Si, Ge and their solid solution Si0.5Ge0.5 exhibit very high theoretical capacities for lithiation due to their alloying reaction. Their behavior towards sodiation or potassiation have been less studied and at the present it is believed that only certain morphologies are accessible for this reaction. In any case, the alloying reaction is associated with a volume expansion/contraction of the particles that leads to pulverization and excessive formation of the SEI layer. As reviewed in Chapter I, several strategies have been implemented to address these challenges. In this Chapter the electrochemical mechanism of lithiation has been studied. Upon reaction with Li, the Si, Ge and Si0.5Ge0.5 undergo amorphization and form a series of lithiated intermediates, only at the end of discharge a crystalline phase corresponding to c-Li15X4 (X=Si, Ge or Si0.5Ge0.5) appears. The delithiation process with the amorphization of the Li15X4 phase into amorphous lithiated phases and at the end of the discharge the original crystalline network is not recovered and an amorphous product is obtained. The subsequent cycles take place in the amorphous domain with the crystallization of the c-Li15X4 phase at the end of discharge. As observed, a great part of the processes during lithiation implicate amorphous phases, thus, a combination of Operando characterization techniques (XRD, XAS, NMR; Raman Spectroscopy) was implemented in this study. For the lithiation of Ge, a-LiGe and a-Li7Ge3 were identified as intermediates. In addition for this compound the cycling conditions such as C/rate, electrode formulation and particle morphology can directly impact the type of intermediates, for instance, when cycled at C/14 the more lithiated c-Li17Ge4 phase is stabilized over the Li15Ge4. Following, the Si0.5Ge0.5 solid solution was evaluated, this phase is of particular interest since it profits from the high reversible capacities from Si and the good capacity retention properties of Ge. The study of the electrochemical mechanism indicate the presence of a solid solution all along the cycling with no segregation into Si or Ge, at least two lithiated intermediates are involved in the reaction and the depletion of the Ge-Ge bonds is slightly faster than the Si-Si or Si-Ge. All these factors might be behind the enhanced performance of the solid solution. Finally, a phenomenon of overlithiation for the Si, Ge and Si0.5Ge0.5 was identified in the c-Li15M4 phase formed at the end of discharge. The Li15M4 is an electron deficient Zintl phase, hence it is possible to accommodate extra Li + without any structural degradation. In fact at this stage, it is energetically easier for the system to overlithiate the Li15M4 phase into Li15+δM4 than to form new Li15M4, as this last one would involve a series of M-M bond breaking to form isolated M 4-atoms. In conclusion, the fundamental study of the processes taking place during the cycling is crucial to understand the failure mechanism and extract the best potential of a given electrode. This work demonstrated the importance of the Operando characterization techniques in the identification of the real processes taking place during cycling and set the basis for a better understanding of the Si and Ge anode materials.

Introduction

The alloying reaction between 14 th group elements and alkali can be detrimental for Li, Na or K-ion batteries due to the large associated volume expansion that leads to a rapid capacity fading, as previously observed in Chapter II. In order to overcome this issue, their associated layered phases are promising anodes that enable alkali insertion with low volume expansion and high alkali-ion diffusivity.

Silicon is a crucial material for different technologies such as integrated circuits, photovoltaics, optoelectronics and most recently, batteries. It is one of the most promising anodes for lithium-, sodium-and potassium-ion batteries due to its high theoretical capacity, 3590 mAh/g for Li15Si4, [226] 954 mAh/g for NaSi [227] and 995 mAh/g for KSi. [40] Nevertheless, its practical application is hindered by a series of obstacles, such extreme volume expansion, resulting in a rapid capacity fade. And for Na and K, the slow kinetics and larger ionic radius restrict the sodiation/potassiation of c-Si. Germanium is also a suitable candidate but its application remains hindered by its high cost. The theoretical capacities are 1384 mAh/g for Li15Ge4 [36] and 369 mAh/g for NaGe [41] and KGe [38] . It presents a series of advantages over the Si, such as an improved Li + diffusivity (almost 4 times faster) and electronic conductivity (2 orders of magnitude bigger), meaning that high power applications are accessible. The Chapter II demonstrated the possibility to extract the best characteristics of both elements and use this synergy for improving the electrochemical performance; while Chapter I showed that the dimensionality of a given material plays an important role in its properties. Indeed, a layered structure is expected to be beneficial for buffering the volume expansion, while providing improved electronic conductivity and ion diffusivity. For the group 14 the dimensional diversity is based in the nature of the bonding between atoms, allowing some Si and Ge with layered structures based on a sp 3 bonding with the fourth bond completed by -H or -OH, such in siloxene or germanane. [228] These two last ones present several differences compared to their bulk counterparts, for instance siloxene and germanane have a direct band gap and exhibit a strong photoluminescence, which is directly related with the presence of -OH and -H groups. [141] Hence the nature of the ligands has a repercussion in the physical and chemical properties. This characteristic is absent in most of the Van der Waals compounds due to their neutrality, while for siloxene and germanane it can be easily achieved by a modification of the synthesis conditions. [154] This versatility of siloxene and germanane sets them as promising candidates for battery applications, not only due to their layered structure and improved electronic conductivity and ion diffusivity but also by their feasibility to be modified to meet certain requirements; properties that undoubtedly would be beneficial for improving the electrochemical performance of Si and Gebased anodes in batteries.

In this work we have synthesized the layered siloxene and germanane from the layered CaSi2 and CaGe2. The compounds have been characterized and their electrochemical performance in Li, Na and K half-cells has been evaluated. The ligand modification has been proposed for siloxene, evaluating its impact over the electrochemical performance. Finally, in an analogy with the Si1-xGex solid solution presented in Chapter II, the solid solution HSi-GeH (siligane) has been prepared in different composition ranges from the layered Ca(Si1-xGex)2 in order to identify a possible synergy between both elements.

Sample preparation and characterization techniques details

This section describes the experimental details employed for the study of the layered silicon-and germanium-based phases, for additional details the reader can refer to the Annex I section. Characterization. The Powder XRD patterns were recorded in a Bruker D8 Advanced diffractometer with Cu radiation (λ1 = 1.54056 Å, λ2 = 1.54439 Å), using either an airtight sample holder with an X-ray transparent dome-like cap or an in situ cell with Be window described in ref [169] . The Xenes

CaSi2, CaGe2 and

CaSi2

The CaSi2 is a Zintl phase (for further details see Chapter I, section 3) in which each Si atom is bonded to other 3 Si and completes the octet rule by accepting 1e -from Ca 2+ , the result is a 2D Si network of corrugated Si6 rings layers separated by Ca atoms, isoelectronic and isostructural to the As/Sb/Bi analogues. This layered structure is related to bucked (111) ). [144,229,230] These last ones, are the most thermodynamically stable and differ by their stacking sequences, AABBCC and ABC for the R6 and 3R, respectively. The 6R polymorph is synthesized by annealing at 800°C stoichiometric amounts of Ca and Si in a sealed tubed, whereas the 3R is obtained by annealing the 6R-CaSi2 either between 200-700°C under hydrogen pressure (30 bars) [230] or at 800°C in a evacuated fused-silica ampule, [231] although both polymorphs can be stabilized in the same sample. [230] The role of hydrogen in the stabilization of the 3R polymorphs is still unknown but possibly related to the incorporation of hydrogen defects and/or the incorporation of hydrogen in the lattice. [230,231] ). [229,232] Note that the peak intensity calculation presents some ambiguity due to a strong preferential stacking of the layers and turbostractic disorder, as observed for other layered materials. [135,[START_REF] Drits | X-Ray Diffraction by Disordered Lamellar Structures[END_REF] The Raman spectra (Figure III.2-b) for the CaSi2 shows an intense peak at 493 cm -1 , along with less intense ones at 127, 222, 283, 410 and 905 cm -1 . According to the literature, the first order Transversal Optical Phonon (TO) of a 2D silicon-based material appears at ≈500 cm -1 (E2g Symmetry). [START_REF] Fuchs | [END_REF] This mode is attributed to the Si-Si vibration in the planes, which in comparison with bulk Si (520 cm -1 ) is shifted towards lower frequencies. The peaks at 127, 222, 283

and 410 cm -1 correspond to a combination of the Transverse Acoustic Phonon (TA) and Longitudinal

Optical phonon (LO), while the one at 905 cm -1 indicates a second order Raman scattering involving two TO phonons. [START_REF] Fuchs | [END_REF][START_REF] Zhang | Raman Spectroscopy and its Application in Nanostructures[END_REF][START_REF] Yu | [END_REF][237][238][239] Note that since the 6R and 3R polymorphs share the same space group and the Si-Si and Ca-Ca bond distances are very similar, their Raman modes are almost the same. [231] Lastly, the SEM picture (

Siloxene

Synthesis and characterization

The layered silicon structures, polysilane (Si6H6) and siloxene (Si6H3(OH)3), can be obtained by facile soft-chemical synthesis. As previously mentioned, the CaSi2 has a Si -corrugated layered structure interconnected by Ca 2+ ions. Upon reaction with concentrated HCl the Ca 2+ is topotactically deintercalated while the integrity of the Si layers is preserved, producing the Siloxene. [134,135] This reaction was first reported by Wöhler [131] and Kautsky [133] while Weiss et al. [134] deintercalation, with an erosion and increased separation of the layers. [240] The XRD pattern (Figure III.5-a) shows broad peaks centered at 14, 27.7 and 47° accompanied with sharp reflections at 28 and 47°. The former ones are ascribed to the siloxene, while the later ones to c-Si, originated during the synthesis through the formation of Si-Si bonds, as reported for other Zintl phases like KSi and NaSi. [135] The diffraction peaks present a broadening due to a decrease in the crystallite size or dislocations and stacking faults. [241] The Le Bail refinement provided a cell parameter a=3.93(2) Å which is close to the CaSi2 value (3.8575(2) Å), indicating the preservation of the corrugated layers. The c cell parameter, has slightly increased from 5.11(1) Å (c/6) in CaSi2 to 6.092(5) Å in siloxene due to a larger interlayer separation, consequence of the replacement of Ca 2+ by two -H/ -OH (up and down the plane) or the intercalation of HCl/H2O. [134,135] The Siloxene presents a turbostratic disorder, with some of the layers with a staggered alignment and some others randomly stacked, though as in CaSi2 the 6R stacking sequence predominates. [135,137,242] The electron diffraction (Figure III.5-b), performed under cryogenic conditions to avoid sample amorphization with the electron beam, showed a slight preferential stacking over the (001) plane and an interatomic distance d= 3.35(1) Å, while the d value calculated from the Bragg's equation (001) was 6.31 Å, both in agreement with the literature (d=3.2-3.3 Å, from Bragg's equation 6.5 Å). [240,243] Three different interpretations of the siloxene structure are reported in the literature (Chapter I, section 3); [134,135,[START_REF] Stüger | Silicon Chemistry: From the Atom to Extended Systems[END_REF] Si-OH and condensation to Si-O-Si, enhanced by light and UV irradiation. [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF]134] For the chain structure few experimental evidences have been found. [135,[START_REF] Stutzmann | [END_REF] Based on the experimental results our siloxene is more conform to the Wöhler siloxene. accompanied by bands at 288, 373, 636, 727, 903 and 2115 cm -1 . In analogy with CaSi2, the Raman modes at 497, 288 and 903 cm -1 are assigned to the E2g vibration, the 2TA and the 2TO of the Si-Si in the Si layers, respectively. The mode at 497 cm -1 is accompanied by a shoulder, attributed to the presence of residual c-Si. The bands at 635, 735 and 2115 cm -1 indicate the appearance of Si-H vibrations after Ca deintercalation. [START_REF] Fuchs | [END_REF][246][247][248] The Infrared spectrum in III.1). The peak at 518 cm -1 is assigned to the Si-Si mode in Si plane; that becomes active due to a dipole moment induced by the difference in electronegativity of the Si ligands (H and OH). [START_REF] Fuchs | [END_REF] The Si-H vibrations are centered at 643 and 2110 cm -1 , the Si-OH, H-Si-O and H-SiO3 at 875, 895 and 2250 cm -1 , respectively. [START_REF] Fuchs | [END_REF]249,250] The -OH bending mode is characterized by a narrow line at 1628 cm -1 while the -OH groups that participate in a hydrogen bond are defined by a broad band at 3400 cm -1 . [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] The peak at 465 cm -1 corresponds to the Si-O-Si bending accompanied by the Si-O-Si symmetric and asymmetric stretching at 809 and 1060 cm -1 , respectively. [START_REF] Fuchs | [END_REF] The appearance of the Si-O-Si modes in our siloxene indicates the presence of SiOx amorphous phases or an oxygen crosslinking (oxidation) between the planes. [240] This oxidation occurs gradually and without any change in the stoichiometry as the oxygen from the -OH is incorporated in the Si-planes in the form of Si-O-Si bridges, forming isolated Si6 rings (Kautsky siloxene). As the oxidation degree rises, the Si-O-Si band becomes broader and splits into two bands. [243]

Table III.1. Assignment of infrared modes in Siloxene

ω (cm -1 )
Assignment 465 Si-O-Si stretching [START_REF] Fuchs | [END_REF] 518 Si-Si stretching in the Planes [START_REF] Fuchs | [END_REF]249] 643 Si-H bending [START_REF] Fuchs | [END_REF]249] 809 Si-O-Si symmetric stretching

875/895

Si-OH/H-Si-O [250] 1060 Si-O-Si asymmetric stretching [START_REF] Fuchs | [END_REF] 1628 -OH bending [START_REF] Fuchs | [END_REF]249] 2110 Si-H stretching [START_REF] Fuchs | [END_REF]249] 2250 H-SiO3 stretching [START_REF] Fuchs | [END_REF]249] 3200-3400 Si-OH bonded [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF]249] 3615 -OH free [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] The XPS spectra for siloxene is presented in Si 4+ from SiO2 possibly arising from decomposition products or glassy oxides in the siloxene. These findings are consistent with the literature in which the peak at ≈99eV is attributed to Si 0 , the one at ≈103 eV to SiO2 (Si 4+ ) and those ones between 99 and 103 eV to Si 1+ , Si 2+ and Si 3+ . [251][252][253][254] Figure and Cl and Ca 1 and 0.4%, respectively of the atomic percentage.

In order to determine the thermal stability of siloxene TGA and DSC analysis were performed

(Figure III.8).
The DSC shows an exothermic peak at 320°C while the TGA a first mass loss at around 100°C corresponding to the loss of surface water molecules followed by a mass gain approximately after 250°C assigned to the formation of sub-oxides with diverse oxygen ratio. This oxidation causes the loss of the diffraction peaks and favors the formation of defects due to Si-H bond breakage. [24,[START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF]139] Even though the process of oxidation is exacerbated with high temperature or an oxidizing environment, a simple exposure of siloxene to ambient conditions is enough to start the degradation, hence the best strategy to avoid it is the replacement of the -OH ligand by oxygen-free groups. [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] Figure III.9. In situ temperature-controlled XRD for siloxene, zoom (right) of the zone containing the siloxene and c-Si Bragg reflections.

Electrochemical behavior of siloxene in Li, Na and K-ion half-cells

The potential use of siloxene as anode material for lithium, sodium and potassium ion batteries has been evaluated; the galvanostatic profiles are shown in Figure III.10. The discharge capacities for the first cycle were 3958, 760 and 790 mAh/g for Li, Na and K, respectively. Unfortunately, a substantial amount of this capacity is irreversible and the second discharge delivered capacities of 2300, 311 and 296 mAh/g for Li, Na and K, respectively. Note that the conductive additive contributes with a reversible capacity of ≈270, 120 and 150 mAh/gC45 for Li, Na and K, respectively (Figure Annex III-1); though this approximation is not ideal and the real capacity could be lower for the composite electrode.

For the three systems the galvanostatic curves can be divided into three regions; i) a pseudo-plateau at 1.2 V for Li and Na and 0.8 V for K (only observed for the first discharge), ii) a sloping profile from 0.7-0.5 V and iii) a steady decrease in the potential for voltages <0.5 V. Likewise, the processes during charge can be categorized into a rapid increment of the potential until ≈0.3 V, followed by a region with a gentle slope, after which the potential increases gradually until the end of charge. Certain differences are identified; and for Li, the region with a different slope upon charge is more prolonged, constituting almost a pseudo-plateau while in the second discharge a plateau at ≈0.5 V appears and prevails during the following cycles.

Figure III.10. Galvanostatic curves and their respective derivative curves for the first cycle and second cycle for siloxene vs a), b) Li, c), d) Na and e), f) K.

All the aforementioned processes are displayed in the differential curves; in the first cycle, for Li (Figure III.10-b) it is possible to observe three reduction peaks at 0.37, 0.16 and 0.06 V and an oxidation one at 0.3 V followed by a very broad signal. The second cycle presents the reduction peaks at 0.45 and 0.20 V with a shoulder at 0.16 V, and two oxidation ones at 0.3 and 0.43 V, in line with the plateau observed in the galvanostatic curve. For Na (Figure III.10-d), the reduction presents peaks at 1.10 and 0.09 V, and a very sharp one at 0.03 V, whereas the oxidation at 0.1 V and a broad signal at 0.6 V. In the second cycle, only a peak at 0.02 V appears. The oxidation displays a peak at 0.01 V followed by a hump with a maximum at 0.6 V. For K (Figure III.10-f), two reduction peaks appear at 0.8 and 0.04 V, and an oxidation one at 0.46 V. During the second cycle, only one reduction peak is observed at 0.03 V, while for oxidation two peaks take place at 0.02 and 0.50 V. These preliminary tests demonstrate the possible reversible potassiation of siloxene; however, its electrochemical behavior and cyclability depends on several factors; particularly the electrolyte formulation, as it will be discussed later. [44] These observations can be compared with the lithiation/sodiation of other silicon-based materials (see Chapter I, section 2 and Chapter II, section 3.1), in which the first discharge presents a substantial amount of irreversible capacity, due to the irreversibly trapping of Li ions in the SEI layer. In siloxene this process is probably enhanced by the presence of -OH and -H groups, as they can react with the electrolyte for form additional decomposition products. The plateau at 1.2 V for Li and Na corresponds to the reduction of the FEC additive; as evidenced, when no FEC is used (discussed after in this chapter).

For K since electrolyte formulation did not contain any FEC, the plateau at 0.80 V during the first discharge, must be attributed to a different phenomenon than the FEC reduction.

The cycling of siloxene vs Li presents different features from the alloying reaction with Si (Figure III.11-a). For instance, the amorphization of c-Si and crystallization of c-Li15Si4 during the first discharge, displayed in the galvanostatic curve by two pseudo-plateaus at 0.3 and 0.18 V and in the derivative by two reduction peaks at 0.17 and 0.005 V and a sharp oxidation one at 0.45 V, which are not evidenced in the siloxene curves. [179,180,182,187,223,255] Though, the traces of c-Si in siloxene (as observed by XRD and Raman spectroscopy) and the carbon additive may also originate certain peaks in the derivative curves. The prolonged sloping profile for the electrochemical reaction of siloxene in the galvanostatic curve, accompanied by the reduction peaks at 0.37, 0.16 and 0.006 V and an oxidation ones 0.3 V, are in line with the literature and could be indicators of a different electrochemical mechanism or different lithiated intermediates, this approach would be explored in the next chapter. [144,246,251,[256][257][258] The improved capacity retention for the lithiation of siloxene compared to a pure Si electrode (Figure III.11-b), is believed to be related with the layered morphology of siloxene; which buffers the volumetric changes during lithiation and shortens the Li diffusion distances, [240,259] as predicted by the theoretical calculations in the literature. [179,180,182,187,223,255] Additionally, the oxygen present in siloxene can be transformed into a Si/Li2O/LiSixOy composite at the nanoscale that can help to accommodate the volume expansion and stabilize the SEI layer. [240,253] However, when the oxygen content reaches a critical point the system suffers from sluggish diffusion. [253]

Figure III.11. a) Derivative curves for the first and second cycle of bulk Si and Siloxene, b) capacity retention of siloxene and bulk Si cycled vs Li under the same conditions at a rate of C/20.

The sodiation of Si has been less studied and, as reviewed in Chapter I, it is feasible in a-Si and a series of nanomaterials and composites and very limited in c-Si. It is worth to note that no work has been performed in the sodiation of siloxene or any other layered silicon-based material. According to the literature, the galvanostatic curve for the sodiation of Si presents a sloping voltage region with a plateau at 0.5 and 0.1 V; whereas the desodiation proceeds with a gradual increase in the voltage until the end of charge. Both processes are visualized in the derivative curve by two reduction peaks at 0.5 and 0.04 V and an oxidation one at 0.08 V. [32,33,260] Additionally, the sodiation of the conductive additive is characterized by a galvanostatic curve with a sloping voltage region at V >0.1 and a plateau at 0.05 V, corresponding to a reduction peak at 0.8 and 0.02 V and an oxidation one at 0.1 V. [33,261] Considering these information, some of the features in the siloxene galvanostatic and derivative curves are in line with the signature of the conductive carbon (as it contributes with 120 mAh/g to the total capacity) while it is more complex to identify a peak related to the reaction of sodium with silicon.

Likewise for the sodiation, the potassiation of silicon is only possible in a-Si and very limited in c-Si, yet it has been reported with a very limited capacity (<100 mAh/g) [20] probably mainly originated from a capacitive process. The corresponding galvanostatic curve is characterized by a continuous sloping profile for discharge and charge. Other reports concerning composites claim the presence of two plateaus at 0.25 and 0.10 V [21] during reduction and one at 0.30 V during oxidation. The striking differences between both profiles highlights the influences of the cycling conditions in the electrochemical behavior. The siloxene presents a different profile as well and a section of this chapter will be devoted to this particular system. Note that all the following values refer to the second cycle. For lithium, sodium and potassium, 1845, 252 and 220 mAh/g are delivered after 58, 77 and 47 cycles, corresponding to capacity retention of 79, 81 and 74 %, respectively. The coulombic efficiency (CE) increases from 51 to 91%, 37 to 90% and 34 to 82% from the first to the second cycle for Li, Na and K, respectively. For Li and Na, CE keeps an average value of ≈96 % while for K of 87%. Note that for K, after the 40 th cycle there is an increase in the capacity, the origin for this behavior has not been identified, although it coincides with an increase in the temperature of the cycling room (±5°C). If the anomalous cycles are not taken into account, then the siloxene delivers 186 mAh/g vs K/K + after 40 cycles with a capacity retention of 63%. The siloxene cycled vs Li exhibits excellent rate capability properties, at high C/rates (1C, 2C, 5C) the electrode still It can be observed that with exception of the cell at C/5, all the capacity values are relatively similar regardless the cycling rate, with no specific increase in the polarization at higher rates (Table III.4).

Perhaps, the cell at C/5 presented atypical behavior since it was prepared from another batch of electrodes, a duplicate of this cell showed similar results, reinforcing this assumption. In spite these fluctuations, all the cells presented a good reproducibility and their galvanostatic curves have the same profile as previously described (vida supra). Their corresponding capacity retention is shown in and 166 cycles corresponding to a capacity retention of 68, 70, 60 and 42% respectively. These results demonstrate that siloxene presents higher capacities compared to a conventional graphite electrode and an improved rate capability and capacity retention compared to bulk Si. The obtained capacities for our siloxene are also superior to certain publications in the literature for siloxene-type materials. [262,263] Unfortunately, the significant irreversibility remains a challenge, particularly the -OH and -H groups in the siloxene planes may contribute to an excessive formation of SEI layer. In order to decrease the irreversible capacity in the first cycle, different electrolyte formulations were tested. Initially, the influence of FEC on the cycling was evaluated. Indeed, as observed in When 1% FEC is used there is an improvement in the performance with higher capacities values for the first and subsequent cycles, whereas, for 5% the overpotential and the hysteresis increase for both Li and Na. For Li the first cycle present similar discharge capacities for 1 and 5%, but for the second one the values are lower when 5% FEC is used, indicating a higher irreversibility of the system. For Na, the irreversibility during the first discharge is higher for 5% (1150 mAh/g) compared to 1% (750 mAh/g) while for the second cycle both capacity values are relatively similar (≈260 mAh/g). Therefore, even though the addition of FEC is necessary, substantial amounts contribute to increase the irreversibility, probably due to an excessive formation of impedant SEI layer that is already exacerbated by the presence of -OH and -H groups in the siloxene. Consequently, all the following electrochemical measurements were performed with 1% of FEC.

Figure III.17. Galvanostatic curve of siloxene cycled at C/20 vs Na using 1 M LiPF6 in a) 1:1 EC/DEC 1% FEC, c) 1:1 EC/PC 1% FEC and e) Tetraglyme 1% FEC and vs Li using 1 M NaPF6 in b) 1:1 EC/DEC 1% FEC, d) 1:1 EC/PC 1% FEC and f) Tetraglyme 1% FEC.

Subsequently, the influence of the electrolyte solvents in the electrochemical performance was evaluated. For this purpose different formulations were prepared for Li and Na, note that in all cases the salt (1 M LiPF6 or NaPF6) and the FEC content (1%) was kept constant while the solvent mixture was changed. The prepared electrolyte formulations were 1 M LiPF6 (NaPF6) in 1:1 EC/DEC 1%FEC, 1 M LiPF6 (NaPF6) in 1:1 EC/PC 1%FEC and 1 M LiPF6 (NaPF6) in tetraglyme with 1% FEC, their respective c,e) the capacities for the first discharge are 3958, 3577, 3319 and 3409 mAh/g for EC/DMC; EC/DEC, EC/PC and tetraglyme, respectively, while the second discharge delivered 2300, 2165, 2002 and 2302 mAh/g, respectively. The galvanostatic profile for discharge/charge for EC/DEC and EC/PC present relatively the same behavior as EC/DMC, except for the first discharge where the processes take place at lower voltages, particularly for EC/PC. On the contrary for tetraglyme the first discharge happens at a higher voltage and the voltage drop from 1.6 to 0.5 V is more pronounced compared to the other electrolytes.

Additionally, there is a remarked capacity fading accompanied by a gradual loss of the discharge plateau at 0.4 V and a steady increase in the overall delithiation potentials as the cycling proceeds. For

Electrolyte formulation: the particular case of Potassium

The previous segment demonstrated that the solvent directly influences the electrochemistry of the cell and particularly for the potassium system the electrolyte formulation is a critical issue, as the capacities can change drastically from one electrolyte to another one, due to their very specific reactivity towards K metal. In order to determine the most suitable cycling conditions for the K-system, three different salts were evaluated KPF6, KTFSI and KFSI with two solvent mixtures EC/DMC and EC/DEC and the role of the FEC addition. Figure III. 20 shows the galvanostatic curves for the siloxene cycled vs. K using 1 M KPF6, KTFSI and KFSI in 1:1 EC/DMC with 1% FEC at C/20. For the first discharge, the delivered capacities are 708, 513 and 834 mAh/g and for the second one 123, 80 and 151 mAh/g for KPF6, KTFSI and KFSI, respectively. The subsequent cycles are characterized by a remarked capacity loss and by the 5 th cycle <50 mAh/g are obtained. Concerning the profile of the galvanostatic curve, in the three cases a plateau at ≈1.5 V is observed for the first discharge followed by a gradual decrease in the voltage until the end of discharge. The subsequent charge/discharge present an abrupt change in the voltage that is more pronounced as the cycling proceeds. Except for KFSI where a change in the slope takes place at about 0.4 V, though this region gradually disappears. Consequently, the siloxene does not cycle correctly with K under these conditions.

Figure III.20. Galvanostatic curve for siloxene cycled vs K at C/20 using a) 1 M KPF6 in 1:1 EC/DMC with 1% FEC, b) 1 M KTFSI in 1:1 EC/DMC with 1%FEC and c) 1 M KFSI in 1:1 EC/DMC with 1% FEC as electrolyte.

Four other different electrolyte formulations were prepared, 0.8 M of KPF6 in 1:1 EC/DMC, 0.8 M of KPF6 in 1:1 EC/DEC, 0.8 M of KFSI in 1:1 EC/DMC and 0.8 M of KFSI in 1:1 EC/DEC, the 0.8 M was chosen based on previous reports. [44] In this case no FEC was added to the electrolyte and KTFSI was excluded This issue can be related to an excessive and continuous consumption of the K 0 in the counter electrode by the reduction of the FEC and formation of the SEI layer or other decomposition compounds, reflected in the plateau at 1.5 V in the galvanostatic curve. Additionally, the potassium salts are highly reactive and more soluble in the electrolyte, meaning that there is a constant formation and dissolution of the electrolyte decomposition products. For instance, in KPF6 there is no hydrolysis, and the reduction is restricted, the SEI is mainly formed from organic species (C-C/C=C, R-O-K, RCOOK) and K2CO3 that constantly dissolve in the electrolyte. On the contrary the KFSI decomposes easier and forms KF, which plays an important role in the stabilization of the SEI layer due to its insulating nature, allowing the cells to attain higher capacity values (Figure III.21). [43] The salts have also a different stability towards K 0 , in a very simple experiment a piece of K metal was submerged in a vial containing 0.8 M KPF6 in EC/DEC and another one with 0.8 M of KFSI in EC/DEC electrolyte. For the KPF6 in EC/DEC the K metal immediately reacts with the electrolyte forming gas and the colorless solution turns yellow after a few minutes, meaning that the electrolyte chemical decomposition starts even before the cycling. On the contrary the KFSI in EC/DEC did not show gassing or color change, indicating a higher stability, reflected in a better cyclability. Finally, the solvent mixture also is important, and the use of EC/DEC improved the delivered capacities compared to EC/DMC for both KPF6 and KFSI. Hence, the choice of salt and solvent impacts the type of decomposition products formed which at the same time are related with the cyclability of the cell. Nevertheless, due to the high reactivity of K 0 counter electrode, a full cell would be the ideal conditions for the study of the K system.

Siloxene modification

The siloxene is a very versatile compound and as observed in Chapter I a simple variation in the synthesis conditions has a direct impact in the physical and chemical properties, such as interlayer distance, band gap or functionalization with different gorups. A very interesting modification is the intercalation of functional groups to tune the band gap and the interlayer spacing until a partial or complete exfoliation of the siloxene layers. In this section the synthesis of siloxene has been performed in the presence of organic alcohols with different carbon chain length, the obtained products were characterized and the influence of the modification over the electrochemical performance was explored. In order to outrule the changes induced by the synthesis of different batches, the modified synthesis were performed twice.

Synthesis and characterization

The modified siloxene was synthesized by dissolving CaSi2 in solution of 1.25 M HCl in R-OH with R= C2H5 (ethanol), C3H8 (propanol), C4H10 (butanol), C5H12 (pentanol) and C10H22 (decanol). The reaction temperature was -20°C (as for conventional siloxene) except for the decanol solution (melting point 6.4°C) which was performed at room temperature. The obtained products were separated by centrifugation, washed with ethanol and acetone and dried under vacuum for 30 min. An annealing at 200-300°C under an Ar flow for 6 hours was performed on the siloxene derived from ethanol, pentanol and decanol. The exfoliation of the planes was attempted by placing the already synthesized siloxene in a solution of 11 mg of dodecyl sulfate (C12H25SO4Na) in 20 mL of distilled water. The mixture was magnetically stirred for 3 weeks at room temperature under air conditions, following it was sonicated for 7 hours divided in intervals of 1 hour each to avoid sample heating. The powder was recovered by centrifugation, washed with water and ethanol and dried under vacuum for 30 min. All the products were characterized, and their electrochemical performance was studied in coin cells following the same protocol for cell testing and electrode preparation as for siloxene (see section 2 of the present chapter).

The synthesis delivered a series of greenish products, with slight variations in tones depending on the used alcohol. The color is in line with the previous siloxene synthesis, except for the exfoliated one.

In this case, the greenish color in pristine siloxene vanishes upon mixing with the dodecyl sulfate solution and after a few days the solution turns white. This change in color is probably related to an oxidation process caused by the aqueous conditions for a prolonged period of time. The siloxene treated with ethanol, propanol, butanol, pentanol and decanol are named as Eth-Siloxene, Prop-Siloxene, But-Siloxene, Pent-Siloxene and Dec-siloxene, respectively.

The XRD of the modified siloxene (Figure III.23) feature Bragg reflections that are found in the pristine siloxene (blue) and the c-Si by-product. Particularly, the (001) reflection shows an interesting progression, with a decrease in the peak intensity and a shift towards lower angles with increasing carbon chain length until its disappearance for the dec-siloxene and the exfoliated siloxene. This shift is associated with an increase of the interlayer distance, as the carbon chain in of the intercalated alcohol becomes longer and is an indicator that the ligand is oriented parallel to the c-axis. The d distance for the modified siloxene was calculated from the Bragg equation (Table III.5), and as the (001) reflection shifts towards lower angles from 14.02° to 8.87°, the d value increases from 6.31 Å in the pristine siloxene to 9.96 Å in pent-siloxene, respectively. [264] These values are not available for dec-siloxene and exfoliated siloxene due to the absence of this reflection. Similarly to pristine siloxene, the modified one presents traces of c-Si, which were not present in the CaSi2 precursor, thus they might be a by-product of the synthesis process. This reflection is particularly intense for the dec-siloxene, which was the only synthesis performed at room temperature. Then, the fraction of c-Si could be related with the synthesis temperature, specifically low contents are promoted by negative temperatures, although this was not quantified from the XDR pattern. The diffractogram for the exfoliated siloxene is slightly different and a huge hump appears at around 23°. This change is likely originated from an oxidation and/or formation of amorphous phases or a loss in the short range order caused by an increase of the interlayer distance and exfoliation or a decrease in the particle size due to the sonication process.

Figure III.23. XRD patterns for (bottom to top) the pristine siloxene and the modified siloxene with ethanol (eth-siloxene), propanol (prop-siloxene), butanol (but-siloxene), pentanol (pent-siloxene), decanol (dec-siloxene)and exfoliated siloxene.

The morphology of the modified siloxene was analyzed by SEM (Figure III.24). The different phases presented degradation upon exposure to the electron beam. The particles have a different amount of staked layers and different lateral sizes are found. Certain particles present signs of cleavage in the edges of the layers, but there is no direct relationship with the carbon chain length. [264] For the exfoliated siloxene, completely exfoliated and staking of several layers are observed, meaning that the exfoliation process is not complete. The particles appear to be smaller, here probably the intercalation of the dodecyl sulfate group increases the interlayer distance and the exfoliation is promoted by the sonication, which also breaks the layer into particles of smaller lateral size.

Following, the Raman spectroscopy was performed (Figure III.25). All the spectra resemble the one of pristine siloxene with a very sharp peak at around 495 cm -1 and a shoulder at 518 cm -1 , note that this shoulder is not always present. Additional Raman bands at 377, 637, 734 and 2120 cm -1 show the presence of the Si-H bonds. The band at 495 cm -1 is assigned to the Si-Si bonds in a plane, in analogy with the previous interpretations, and indicates the preservation of the Si planes after synthesis. The Raman shift of this mode is associated with the size of the organic ligand that impacts directly the bond length and the frequency of vibration. The shoulder/band at 518 cm -1 is assigned to the presence of c-Si as a by-product (in analogy to the XRD); which changes in intensity for some phases, indicating probably a variation of the c-Si content.

Indeed, samples like prop-siloxene and but-siloxene have a slight vestige of this shoulder and a small peak intensity for c-Si in the XRD pattern, while for the dec-siloxene both the Raman and XRD peaks associated with c-Si have a high intensity. On the other hand, the presence of the Si-H Raman bands could indicate a merely intercalation of the organic molecule without chemical bonding or a partial degree of substitution. Particularly for but-siloxene, no Si-H bands are visible, then either the substitution was complete or there was a decomposition. Note that this sample underwent spontaneous combustion upon exposure to air. The exfoliated siloxene has a relatively different Raman spectrum, there is a shift of the 495 cm -1 band to 489 cm -1 and an intense peak at 518 cm -1 . This shift could be associated with an oxidation or with the exfoliation process that reduces the number of layers redshifting the Raman bands. Additional non-attributed bands are observed for all the samples; in ethsiloxene, prop-siloxene, pent-siloxene and dec-siloxene at 2900 cm -1 , and for exfoliated siloxene at approximately 1300-1400 cm -1 are characteristic of the asymmetric and symmetric bending of (C-H) CH2 and (C-H) CH3 while the second group centered between 2700-3000 cm -1 are associated to the asymmetric and symmetric stretching of (C-H) CH2 and (C-H) CH3. The C-O vibration region is located in a similar frequency as the Si-O-Si, complicating their differentiation. The -OH vibrations in the 3000-3600 cm -1 region present a weak intensity for the modified siloxene, these bands are expected to have a strong signal due to -OH present in the organic alcohol, siloxene ending groups and trapped water molecules in the interlayer space. The absence of these bands for but-siloxene is probably related to an optimal removal of trapped water molecules. The presence of all these vibrations could indicate a chemical bonding of the organic molecule with the Si plane, though the Si-CH3 vibration centered at 1250±10 cm -1 is hardly observed. [140] Once again, the behavior for the exfoliated siloxene is different and displays none of the C-H Table III.6. Band assignment for the Infrared spectra of the modified siloxene [259,265] ω (cm -1 ) Assignment ω (cm -1 ) Assignment 465 Si-O-Si stretching [START_REF] Fuchs | [END_REF] 1628 -OH bending [START_REF] Fuchs | [END_REF]249] 518 Si-Si stretching in the Planes [START_REF] Fuchs | [END_REF]249] 2110 Si-H stretching [START_REF] Fuchs | [END_REF]249] 643 Si-H bending [START_REF] Fuchs | [END_REF]249] 2250 H-SiO3 stretching [START_REF] Fuchs | [END_REF]249] 809 Si-O-Si symmetric stretching 2869 (C-H) CH2 symmetric stretching

875/895

Si-OH/H-Si-O [250] 2881 (C-H) CH3 symmetric stretching 1060 Si-O-Si asymmetric stretching [START_REF] Fuchs | [END_REF] 2937 (C-H) CH2 asymmetric stretching 1035-1085 C-O stretching 2966 (C-H) CH3 asymmetric stretching 1380 (CH) CH3 symmetric bending 3200-3400 Si-OH bonded [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF]249] 1465 (CH) CH3 asymmetric bending, (CH) CH2 symmetric bending 3615 -OH free [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] In summary, the relation between the XRD and the Raman spectroscopy allowed to stablish a possible parallel between the presence of the 518 cm -1 Raman band and the diffraction peaks for c-Si, as their peak intensities seem to be directly related. Nevertheless, contrary to XRD, the Raman spectroscopy provides information of a local environment, thus both parameters cannot be directly linked. A comparison of the Raman and Infrared spectroscopy, allowed to suggest a partial intercalation of the organic molecule in the layers of siloxene possibly with a chemical bonding, given the presence of both C-H and Si-H vibration bands. The -OH IR signal has decreased in intensity but the Si-OH and H-Si-O bands between 870-890 cm -1 are still present, supporting the previous assumption. In certain cases the intercalation of the organic molecules can be incomplete, this is related with the interactions between the different -H that generates repulsion at a certain substitution degree, hence the introduction of new ligands becomes energetically demanding. This phenomenon is more remarked for ligands of bigger size. [156] The introduction of covalently bonded (Si-C) and/or oxygen-free ligands in siloxene, promotes the exfoliation, improves the stability towards oxidation and hydrolysis; though, degradation via external factors such as UV light still occurs and increases the solubility in organic solvents. [140,259,264,266,267] For the exfoliated siloxene, both the Raman and the IR indicate an oxidation due to the oxygen incorporation in the planes with gradual loss of the planar configuration. However, the presence of oxygen is necessary to promote the exfoliation, although once exfoliated, the sheets tend to re-aggregate. [266]

Electrochemical performance of modified Siloxene in Li-and Na-ion half cells

The modification of siloxene with electron-rich and electron-poor substituents can control the electronic states to achieve a desired property, opening the door for a variety of applications such as semiconductors or chromic and luminescence materials. [267] In this section, the applications for Li-and Na-ion batteries were tested. Given the layered character and the expanded interlayer distance of the modified siloxene it is expected that the insertion of the alkali ions would be facilitated. Hence in an exploratory study the electrochemical performance was tested for the eth-siloxene, pent-siloxene and dec-siloxene. The electrodes were prepared using the same protocol as for pristine siloxene (see with a constant decrease/increase in the potential and; for the first discharge a plateau at 1.2 V, assigned to the reduction of FEC additive and for the charge a pseudo plateau between 0.25-0.5 V.

Interestingly, the plateau observed at around 0.4 V after the first cycle for the pristine siloxene, is absent in the modified siloxene. Similarly, the sodiation resembles the one of pristine siloxene with a constant sloping profile for discharge/charge, the plateau at 1.2 V during the first discharge and the region with different slope between 0.5-0.8 V during charge, only visible for eth-siloxene, but that gradually disappears.

The capacity retention properties for the modified siloxene are presented in Figure III.28, note that due to the low capacities of the pent-siloxene and dec-siloxene vs Na, the cells were not chosen for long-term cycling. Despite the low reversible capacities little capacity fading is observed for at least 30 cycles. Then the lithiation of eth-siloxene, pent-siloxene and dec-siloxene, respectively, retains 562, 581 and 886 mAh/g after 30, 30 and 21 cycles; corresponding to a capacity retention of 67, 75 and 87%

with respect to the second cycle. For the sodiation of the eth-siloxene 113 mAh/g are delivered after

89 cycles with a capacity retention 78%. These capacities values are far from the pristine siloxene, considering that the conductive additive contributes with ≈270 mAh/g for Li, then the modified siloxene provides about 350 mAh/g for eth-siloxene and pent-siloxene and 630 mAh/g for dec-siloxene, meaning that a great fraction of the reversible capacity comes from the carbon. For Na, the conductive additive contributes with ≈120 mAh/g, which is almost the electrode total capacity.

Although it is hard to differentiate the specific contribution of both modified siloxene and the conductive additive to the total capacity. Whether the modification is beneficial for the electrochemical performance or not is hard to infer as more information is necessary, but until this point no improvement is observed. It appears that the intercalated molecule imposes a steric effect that hinders the insertion of the alkali ion. [26] Therefore, the lower reversible capacity might be originated from a capacitive behavior with the alkali storage close to the edges of the layers, while the bulk remains hardly accessible due to the presence of the organic molecule. Additionally, the-OH and -H groups in the organic molecule could interact differently with the electrolyte changing the composition of the SEI layer. The larger reversible capacity presented in the dec-siloxene could be due to an efficient interlayer distance that compensates the steric effect of the organic molecule and grants the access of the alkali ion to the bulk; or to an increased fraction of c-Si by-product (as observed by XRD and Raman spectroscopy) that remains electrochemically active. Assuming this last one to be correct, the c-Si fraction will be active towards the lithiation and inactive to the sodiation (as described in Chapter I), which is the case. Therefore, the extra capacity for the dec-siloxene might come from the c-Si fraction. Finally, considering that the steric effect that hinders the alkali intercalation and the absence of the 0.4 V plateau after the second cycle for the modified siloxene, one can interpret this plateau as an inherent process to the lithiation of pristine siloxene. This topic will be developed in depth in the next chapter.

In an attempt to diminish this steric effect, a thermal treatment was performed on the modified samples to remove the excess of volatile molecules, and weakly bonded -OH /-H groups. For this purpose the eth-siloxene, pent-siloxene and dec-siloxene were annealed at 250°C for 6 hours in an argon flow, to further observe the effect of the temperature, the eth-siloxene was also annealed at 300°C. For comparison, Sa reference of pristine siloxene was annealed under the same conditions (250 and 300°C). These temperatures were chosen based on the in situ temperature-controlled XRD, in which siloxene is stable until ≈250°C. Figure III. 29 shows the XRD pattern for the annealed modified siloxene and as evidenced the profile of the diffraction peaks has evolved. Some of the Bragg reflection from siloxene remain while a big hump appears for 2θ<30°. These humps might be related with the oxygen incorporation in the Si planes or the formation of sub-oxides and the possible loss of the layered arrangement, particularly for the sampes annealed at temperatures <300°C. [253] These results indicate an improvement of the reversible lithiation capacities for the annealed modified siloxene from 840 to 1002 and 848 mAh/g for the eth-siloxene 250°C and 300°C, respectively, and from 750 to 986 mAh/g for pent-siloxene 250°C. For Na no significant improvement is observed for the annealed samples. Previously it was suggested that the reversible capacities might be originated from a capacitive storage in the edges of the layers and not at the bulk. Perhaps the thermal treatment removes some of the trapped molecules facilitating the access of the alkali ions to the bulk. This last one could be behind the manifestation of the plateau at 0.4 V for the second discharge.

Unfortunately, the thermal treatment is not enough to achieve the reversible capacities of pristine siloxene, probably some molecules still trapped and produce steric effect. A full removal could be attained at higher temperatures, but as observed in the in situ temperature-controlled XRD, siloxene is stable until 250°C and the annealing at 300°C (as in eth-siloxene 300°C) did not show improvement in the electrochemistry. Therefore, the thermal treatment must be done under certain conditions that grant the removal of the organic molecule while without degrading the siloxene structure. Additional strategies to effectively expand the interlayer distance and its impact over the electrochemistry must be considered.

Germanium-based layered phases

The previous section put in evidence the advantages of a layered material for the electrochemical performance; the germanium analogue of siloxene, the germanane (GeH)n, is obtained from the layered CaGe2 in which the Ge -Zintl ion allows the topotactic deintercalation of Ca while preserving the Ge backbone. [268] This germanium-based layered material was recently discovered and is less known compared to the siloxene. In this study, the alkali-ion storage properties of germanane were evaluated, the results are very promising with 2150, 495 and 205 mAh/g of reversible capacity for Li, Na and K respectively. As expected from the study performed in Chapter II, the germanium-based phases present an improved capacity retention, electron mobility (3900 cm 2 /V.s, ≈3 X Si) and Li-ion diffusivity, allowing high rate applications. [152] These properties are translated into the germanane phases, for instance the electron mobility is predicted to be 18000 cm 2 /V.s. [269] The layered morphology of germanane buffers the volume expansion, increases the surface area and lowers the energy barrier for diffusion. [152] Consequently, its utilization as an anode for Li, Na and K-ion batteries is very promising. The applications of the germanane are not limited to batteries, the characterization analysis performed in this work could set also the bases for other usages such supercapacitors, optoelectronics or catalysis. The results from this section have been adapted from the publication: Loaiza, L. C., Monconduit L., Seznec V. Towards germanium layered materials as superior negative electrodes for Li-, Na-and K-ion batteries, Batteries & Supercaps 2020, 3, 1-11.

CaGe2

The CaGe2 is a Zintl phase with a layered structure consisting of Ge6 rings interconnected to form planes, each Ge atom is bonded to other 3 Ge atoms and completes the octet rule by accepting one electron from the Ca 2+ . Likewise CaSi2, the CaGe2 phase presents several polymorphs (Figure III.32).

The β-CaGe2 or 6R-CaGe2, that crystallizes in the rhombohedral CaSi2-type (space group R-3m), known for several decades as the only thermodynamically stable polymorph. Recent studies have revealed the existence of the 2H or α-CaGe2 that crystallizes in a hexagonal cell (space group P63mc), its formation is related to the synthesis conditions and is a metastable phase stabilized by the presence of impurities. [270] The 3R-CaGe2 has a rhombohedral unit cell (space group R-3m) and is the most stable polymorph at low temperature, while the 6R stabilizes above 673 K. [270] The aforementioned polymorphs mainly differ by their stacking sequences, the layers are organized in a staggered fashion, and follow the ABAB sequence for 2H-CaGe2, ABC for 3R-CaGe2 and AA'BB'CC' for 6R-CaGe2. [115,271] The stacking compatibility allows the formation of other metastable phases, meaning that the 4H polymorph can be easily formed from the 3R-CaGe2 by the addition of an A' stacking to the matrix. [115,270,271] Figure III.32. Crystal structures of the different CaGe2 polymorphs. [270] As observed in the Figure III.32, the stacking originates different interlayer distances that are related to the CaGe2 stability, indeed a large interlayer distance leads to a rapid decomposition upon contact with moisture, according to the reaction: [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] CaGe2 + H2O  Ca(OH)2(GeH)2

(2)

This decomposition product contains layers of germanane (GeH)n separated by Ca(OH)2 monolayers with no destructive oxidation of the Ge backbone. The stacking sequence is preserved, although the c parameter is increased due to the bigger size of Ca(OH)2. [150,271] [148,229,270,271] Likewise CaSi2 the CaGe2 presents a turbostratic disorder with a preferential orientation of the planes, impeding an adequate refinement. [271] The Raman spectra for CaGe2 (Figure III.33-b), presented two A1g modes at 300 and 202 cm -1 and three Eg modes at 244, 235 and 125 cm -1 . In layered hexagonal materials such as the 3R-and 6R-CaGe2 the Eg modes correspond to the in-plane lattice vibrations while the A1g to the out-of-plane. [229] In certain occasions some modes can be activated depending on the polarization of the incident light while some others are independent of this variable. The SEM pictures for the CaGe2 showed in

Germanane

Synthesis and characterization

The germanane was obtained from the CaGe2, by dissolution in concentrated HCl at -20°C, according to Equation 3. Contrary to siloxene, the formation of germanane does not undergo a hydrolysis reaction even at temperatures >0°C, yet the low temperatures increase the crystallinity. [148] Consequently germanane mainly forms Ge-H bonds and no Ge-OH, following the formula Ge6H6, which contrary to polysilane (Si6H6) is stable in air. [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF] The resulting powder is reddish and exhibits photoluminescence in the infrared. [148] 3CaGe2 + 6HCl + 3H2O  Ge6H6+ 3CaCl2+3H2(g)

The preferential bonding of Ge-H over Ge-OH reveals the different oxidation behavior of Si and Ge sheets. This difference can be understood in terms of the binding energies, the Si-O bond (8.0 eV) is significantly stronger than the Ge-O (6.6 eV) and the Si-H bond (3.0 eV) is slightly weaker than the Ge-H bond (3.2 eV). The absence of -OH bonds improves the germanane stability. [272] Nevertheless, the prolonged exposure to moisture causes substitution of Ge-H bonds by Ge-OH, forming the Germoxene Ge6OH3H3. This compound has much less structural stability and the introduction of -OH groups predominately influences the band gap. [START_REF] Brandt | Silicon Chemistry: From the Atom to Extended Systems[END_REF]268] Figure III.34 reveals the preservation of the layered morphology after the topotactic deintercalation of Ca in CaGe2. The particle size distribution is inhomogeneous, various lateral sizes and stacks of different layers are observed. Indeed, the TEM images show the presence of several staggered layers as well as voids on the particle surface, probably caused by the H2 released during the synthesis. In some cases, a partial layer exfoliation can be observed.

An essay to exfoliate the germanane was performed by sample sonication prior the recovery of the powder by centrifugation. The resulting suspension was stable for a prolonged period of time, but the recovery of the germanane particles was difficult. According to the literature the exfoliated particles re-aggregate upon drying and show a decrease in the lateral size due to process of sonication. [152] Undoubtedly this approach must be considered for future experiments, particularly the influence of the exfoliated sheets over the electrochemical performance. inherent turbostratic disorder that limits the quality of the refinement, although the obtained cell parameters are in line with the literature. [148,149,268] The c lattice constant expands from 10.32 (7) Å in 3R-CaGe2 to 11.876 (7) Å in 3R-GeH and from 30.714(3) Å in 6R-CaGe2 to 34.6(1) Å in 6R-GeH. The a cell parameter experiences a slight decrease from 4.008(2) Å in CaGe2 to 3.922(4) Å in germanane for the 3R polymorph, while for the 6R there is a slight increase from 3.999(2) Å in CaGe2 to 4.027(9) Å in germanane, in agreement with the literature. [115,151,268] The expansion in the c parameter from CaGe2 to germanane, is caused by the replacement of Ca 2+ by two -H between each layer. [151] The interlayer distance d, calculated from the Bragg Equation, is 6.14 Å which is slightly lower than the one of siloxene 6.31 Å. This decrease is related with the preferential bonding of germanane with -H and the almost absent -OH groups. [148,268] The electron diffraction (Figure III.35-b) for germanane presents a high crystallinity and contrary to siloxene, it is stable with the electron beam, thus the experiment can be performed under normal conditions. The pattern shows a preferential orientation of the planes and can be indexed in a hexagonal cell assuming a [011] zone axis.

Several polymorphs are reported for germanane: the 1T (P-3m1-one layer per trigonal unit cell), 2H

(P63mc-two layers per hexagonal cell), 3R (R-3m three layers per rhombohedral cell), 4H (P-3m1 four layers per hexagonal cell) and 6R (R-3m six layers per rhombohedral cell). They all present a similar value for the a cell parameter, while the c varies according to the number of layers per unit cell. [115,268] Each polymorph present its stacking sequence, the 3R has ABC, the 4H A'ABC and the 6R A'AB'BC'C. [270] In analogy with CaGe2 the interlayer distance is directly related to the staking sequence in the polymorph, the larger the distance is the highest is the reactivity. [149,268,270] Finally, several polymorphs can coexist in the same sample, thus in some cases their identification can be complex. [270] Figure it is composed of only one band centered at 302 cm -1 corresponding to the in-plane E2 mode of the Ge-Ge bond. This signal is blue-shifted compared to the E2 Ge-Ge signal for c-Ge centered at 297 cm -1 . [151] A similar band is observed in CaGe2, corresponding to an out-of-plane A1g vibration at 300 cm -1 . These results are in line with the literature, although the A1 mode should be present at at 2000 cm -1 (Ge-H stretching) and 570 and 507 cm -1 (Ge-H wagging) preferentially oriented parallel to the c axis. [148] Other weak vibration modes are found at 770 and 850 cm -1 , assigned to the bending of Ge-H2, probably located at the edges of the layers or originated from Ge vacancies. [151,272] The band at and is completed at 175°C. The dehydrogenation occurs between 200-250°C and at 300-350°C the residual Cl from the synthesis is desorbed. Due to a stronger Si-H bonding compared with Ge-H, the dehydrogenation temperature for germanane is lower than the one of siloxene (400°C). [148,151,269] The thermal stability of germanane can be improved by replacing the Ge-H bonds by the less reactive Ge-CH3 ones. [154] 4.2.2 Electrochemical behavior of germanane in Li-, Na-and K-ion half cells

Figure III.38. Galvanostatic curves and their respective derivative curves for the first cycle and second cycle for germanane vs a), b) Li, c, d) Na and e), f) K.

The electrochemical behavior of germanane was evaluated by galvanostatic cycling in Li, Na and K half-cells. The electrodes were prepared following the indications in section 2 of the present chapter.

The electrolyte was 1 M LiFP6 (NaPF6) in 1:1 EC/DMC + 1% FEC, while 0.8 M KPF6 (KFSI) in 1:1 EC/DEC were used for potassium. Figure III. 38 shows the galvanostatic curves and their respective derivative curves. The discharge capacities for the first cycle are 3732, 982 and 485 mAh/g and for the second one 2150, 495 and 205 mAh/g for Li, Na and K, respectively. Note that the conductive additive contributes to the capacity with ≈270 mAh/g for Li, ≈120 mAh/g for Na and ≈150 mAh/g for K (Figure Annex III-1). The lithiation capacities are lower compared to siloxene, as expected from the inferior theoretical capacity of Ge compared to Si (3579 mAh/g for Si and 1384 mAh/g for Ge [37] ). Surprisingly, the sodiation of germanane deliver higher capacities than for siloxene, as expected from the enhanced electrochemical reactivity of Na towards Ge. This capacity also exceeds the theoretical capacity for NaGe (396 mAh/g), according to the literature more than one Na can be inserted in the Ge matrix. [60,62] For potassium, the optimal cycling conditions must be established before assuming an improved cyclability.

The galvanostatic cycling of germanane takes place at higher voltages compared to siloxene. The first discharge for Li and Na is characterized by a plateau at ≈1.5 V and 1.3 V, respectively, assigned to the reduction of the FEC additive. Following, for Li there are two regions with different sloping profiles, the first one at around 0.6-0.3 V and the second one from 0.3-0 V. For the first charge and subsequent ones, two pseudo-plateaus are observed from 0.3-0.5 V and 0.5-1.0 V. The second discharge is slightly different from the first one, and a plateau is observed at 0.65 V, followed by a region with a sloping profile prolonged until 0.15 V. These features are reflected in the derivative curve (Figure III.38-b) by a three reduction peaks at 1.2, 0.5 and 0.14 V during the first discharge, while the second one presented reduction peaks at 0.65 (plateau in the galvanostatic curve), 0.32 and 0.16 V. For the first and second oxidation, peaks at 0.4 and 0.95 V are observed, in line with the two-pseudo plateaus in the galvanostatic curve.

For sodium, in the first discharge after the plateau at 1.3 V, the potential decreases with a constant slope. For the first charge and subsequent ones, a pseudo-plateau is observed between 0.5-0.7 V, which gradually disappears as the cycling proceeds. The second discharge presents a constant sloping profile until 0.2 V followed by a steady decrease in the potential. All these processes are displayed in the derivative curve (Figure III. 38-d), in the first discharge by reduction peaks at 1.3 and 0.04 V, accompanied by two other small signal at 0.6 and 0.7 V. For the first charge, one oxidation peak is observed at 0.7 V, in line with the pseudo-plateau in the galvanostatic curve. The second cycle, presents a double reduction peak at 0.08 and 0.04 V with a shoulder at 0.2 V, and an oxidation peak at 0.7 V with a shoulder at 0.4 V.

For potassium, the first discharge shows a very prolonged plateau at ≈1.0 V, followed by a gradual decrease in the voltage. Since this system has no FEC in the electrolyte formulation, this plateau must be attributed to another phenomenon than the reduction of FEC. The first and subsequent charge steps present a small region with a different slope approximately from 0.45-0.55 V that gradually vanishes with the cycling. The second discharge has a constant decrease in the voltage. The derivative curve (Figure III.38-f) showed a reduction peak at 1.0 V, in line with the plateau in the galvanostatic curve, and two other peaks at 0.1 and 0.018 V during the first discharge. For the second discharge two reduction peaks are observed at 0.31 and 0.01 V. The first and second charge presents only one oxidation peak at 0.5 V.

The observed phenomena for the lithiation of germanane are in line with the only report available in the literature. [152] The galvanostatic curve presents a sloping profile. The derivative curve displays a broad reduction peak at 0.3 V and two oxidation ones at 0.5 and 0.9 V. Given the lack of information regarding the electrochemical mechanism of germanane, the source of these peaks remains unclear.

The reported reversible capacities for germanane are in the range of 1108 mAh/g while for this work 2000 mAh/g were obtained at the same rate. [152] Regarding the capacity retention properties ( where capacities in the range of 1100, 1000 and 900 mAh/g were delivered at C, 2C and 5C, respectively. This excellent rate capability can be attributed to the higher electron mobility in germanane compared to bulk Ge. [272] The sodiation of germanane is less accessible at high C/rates and 220, 170 and 120 mAh/g are obtained at C, 2C and 5C, respectively. At these rates, probably the capacity mainly comes from a capacitive behavior, similarly to siloxene, as the Na ions do not have sufficient time to access the bulk of the layers due to their sluggish diffusion; or from the conductive additive (120 mAh/gC45). Based on this information, the electrochemical behavior of germanane vs Li resembles to the one at C/10 described above, no significant changes, as polarization, are observed except for a size reduction of the plateau at 0.6 V during the second discharge for C and 2C. The capacities for the first discharge were 2525 and 2383 mAh/g and for the second one 1560 and 1523 mAh/g for C and 2C, respectively. A priori there is no major decrease in the capacity for C and 2C during the first cycles. For instance, after 25 cycles 1130 mAh/g are delivered at 1C and 2C, corresponding to a capacity retention (with respect to the second cycle) of 72 and 74%, respectively.

For the cell at 2C after 58 cycles 990 mAh/g are obtained with a capacity retention of 65%. These values are higher compared to the reports in the literature, 530 and 265 mAh/g for 1C and 2C. [152] Thus, the germanane is very promising for high rate applications, for comparison siloxene presents a capacity retention of 52% after 58 cycles at 2C. This improvement in germanane could be related with the almost 400 faster Li ion diffusion [START_REF] Graetz | [END_REF] and 2 orders of magnitude higher electronic conductivity [36] of Ge compared with Si. for the first discharge there is one plateau at 0.45 V, a pseudo plateau at about 0.2 V, reflected in the derivative curve by two reduction peaks at 0.43 and 0.14 V with a shoulder at 0.3 V. For the first and subsequent charge, there are two regions with a different slope between 0.3-0.5 V and 0.5-1.0 V, this last one gradually disappears with the cycling; only one oxidation peak is observed in the derivative curve at 0.43 V, accompanied by a shoulder for the second charge. The second discharge presents a plateau at 0.50 V and for the subsequent discharges the plateau appears at a higher potential 0.6 V.

The derivative curve showed two reduction peaks at 0.48 and 0.14 V with a shoulder at 0.3 V. It is possible to observe then certain differences in the electrochemical behavior of germanane and bulk Ge; for the first discharge the germanane does not present the plateau and the reduction peak at 0.45 V, neither the shoulder at 0.3 V observed for bulk-Ge, while the peak at 0.14 V is present in both.

For the charge, both materials share the oxidation peak at 0.43 V (region with different slope in the galvanostatic curve) and germanane presents an extra oxidation peak at 0.94 V (second pseudo plateau in the galvanostatic curve). These features are slight indicators of a different lithiation process, as revised in Chapter II, the Ge is very sensitive to the cycling conditions. Probably due to the germanane layered character, the lithiated intermediates or even the electrochemical mechanism might be different. The germanane presented higher capacity values compared to bulk Ge, for both Li and Na (Figure III.42). As expected, the germanium presents an excellent capacity retention, property that is transferred to the germanane. This property along with the high reversible capacity renders germanane an excellent choice as anode material for Li-and Na-ion batteries. and 485 mAh/g for KPF6 and KFSI, while the second discharge delivered respectively 202 and 205 mAh/g. For KPF6 EC/DEC, the galvanostatic curve presents a sloping profile during the discharge/charge cycles. For KFSI EC/DEC a plateau at 1.0 V is observed during the first discharge and the subsequent charge/discharge cycles present a sloping profile. The cells cycled with KPF6 EC/DEC presented more irreversibility and capacity fading compared to KFSI; and after 33 cycles only 56 mAh/g are delivered for KPF6 in contrast to 150 mAh/g for KFSI. Therefore, as observed for siloxene, the formulation KFSI in EC/DEC exhibits the best performance in terms of reversible capacity and capacity retention, though tests in full cell configuration are necessary to assess the real performance of germanane in KIB.

Summary and conclusion

In summary, the germanium-based layered phase, the germanane, was obtained from the layered CaGe2. After topotactic deintercalation of Ca 2+ , the CaGe2 transfers the same crystal symmetry to the germanane and two polymorphs were identified by XRD, the 3R (R-3m, 3 layers per rhombohedral unit cell) and 6R (R-3m, 6 layers per rhombohedral unit cell). The SEM and TEM analysis showed a layered morphology for both CaGe2 and germanane. The Raman and infrared characterization indicate a preferential bonding of Ge with -H and almost no formation of Ge-OH. The presence of Ge-O-Ge bonds was identified and ascribed to a possible oxygen incorporation in the Ge planes (in analogy with siloxene) or the formation of amorphous GeOx as by-product of the reaction. The thermal stability of germanane was estimated to ≈250°C, yet the amorphization starts at lower temperatures. The electrochemical performance of germanane was evaluated in Li, Na and K half-cells. The reversible lithiation capacities are in between 1700-2000 mAh/g while the sodiation between 400-500 mAh/g. These results are very promising and higher compared to bulk-Ge cycled under the same conditions. Particularly, the sodiation is higher compared to siloxene, proving the feasibility of Na insertion into layered group 14-derivated phases. The germanane displayed good performance at high rates and at 1C and 2C capacities in the range of 1100-1200 mAh/g are obtained, values that are superior to the previous reports in the literature. [152] For the potassiation, the results are more ambiguous, the germanane seems to deliver god reversible capacities when the KFSI EC/DEC formulation is used, however, tests in full cell configuration are necessary.

Finally, the layered morphology of germanane influences the electrochemical behavior by buffering or mitigating the volume expansion during cycling. The decrease in the irreversible capacity, remains a challenge, and understanding of the role of the -H groups in the first discharge and formation of the SEI layer is crucial at this stage. Additionally, due to its structure the lithiated/sodiation intermediates might be different and even another electrochemical mechanism could take place. For instance, the theoretical studies predict the absorption of Li and Na in the germanane layers with lower Li diffusion barriers (0.21-0.76 eV) than in siloxene (1.2 eV). [161,162] This work opens the quest for developing all the potential from the germanane as anode for LIB, NIB and KIB.

Si 1-x Ge x -based layered phases

The Chapter II introduced the benefits of the Si1-xGex alloys for LIB. Such phases profit from the high reversible capacity from bulk Si and the good capacity retention properties of bulk Ge. The previous sections demonstrated that the use of layered siloxene and germanene allowed an improvement of the cyclability of both Si and Ge-based electrodes. Hence, in the following section the feasibility for the formation of the layered Si1-xGex analogues was explored, their synthesis, characterization and electrochemical behavior in Li and Na half-cells are presented.

Ca(Si1-xGex)2

Synthesis and characterization

The synthesis of the layered Si1-xGex, begins as well from the layered Ca(Si1-xGex)2. Few reports regarding this phase can be found in the literature, and the synthesis methods comprise epitaxial alloying of Si and Ge to form the Ca(Si1-xGex)2. [273] Since Si and Ge are completely miscible, their layered analogues are expected to be soluble. For the synthesis of Ca(Si1-xGex)2, initially three different Si1-xGex compositions were prepared, Si0.9Ge0.1, Si0.5Ge0.5 and Si0.1Ge0.9. The Si1-xGex solid solutions were The Si0.1Ge0.9, Si0.5Ge0.5 and Si0.9Ge0.1 present a main diffraction peak at 2θ=27.22°, 27.81° and 28.17°, respectively. This peak corresponds to the (111) reflection of the cubic unit cell in which both Si and Ge crystallize (space group F-3dm). The 2θ peak position is related to the fraction of the major component in the phase, this is for Si0.1Ge0.9 the peak centered at 27.22°, nearby the elementary Ge at 27.3°, while the Si0.9Ge0.1 has a peak at 28.17°, close to the elementary Si at 28.5°. The Si0.5Ge0.5 composition presents an intermediate peak between those of Si and Ge, implying equal fractions of both components. For this phase, there is a small peak shoulder at ≈27°, suggesting that a fraction of Ge has not been incorporated into the solid solution. The shift towards lower angles of the (111) diffraction peak is also associated with an expansion of the cell parameter with increasing Ge content and vice versa, following the Vegard's Law. The Raman spectra shows the presence of three main bands centered at approximately 293, 395 and 500 cm -1 , assigned to the Ge-Ge, Si-Ge and Si-Si bonds, respectively. The Ge-Ge band shifts from 297 to 293 cm -1 with respect to bulk Ge and the Si-Si from 520 to 490-500 cm -1 with respect to bulk Si. The band intensity is proportional to the quantity of a specific bond, and for the Si0.1Ge0.9, the Ge-Ge band is stronger, while for the Si0.9Ge0.1 the Si-Si band is more intense. For the Si0.5Ge0.5, the Ge-Ge band presents higher intensity, probably due to the unreacted Ge fraction, as observed by XRD. Finally, the presence of the Si-Ge bond indicates the formation of a solid solution between Si and Ge. [72] Once the formation of the solid solution was confirmed, the synthesis process continued by the incorporation of Ca and the formation of the layered phase. For this purpose, the already synthesized Si1-xGex solid solutions were mixed in stoichiometric amounts with Ca by ball milling, followed by a thermal treatment at 800°C in a sealed tube under argon, in analogy CaGe2 and CaSi2 synthesis. Due to similarities between the layered Zintl phases of group 14, their elements can be partially substituted and it is expected that they remain in the backbone after topotactic deintercalation of Ca. [158] The XRD for Ca(Si1-xGex)2 is shown in Figure III.45. The diffractogram of Ca(Si0.1Ge0.9)2 resembles to CaGe2, while the one for Ca(Si0.9Ge0.1)2 to CaSi2. The Ca(Si0.5Ge0.5)2 presents peaks from both CaSi2 and CaGe2, but it is not a linear combination of both, thus it is hard to stablish it as an intermediate composition.

Interestingly, the (110) reflection shifts to lower angles as the fraction of Ge increases, corresponding to an expansion of the unit cell. No shift is observed in the (002) plane, meaning that probably the incorporation of Ge happens in the a and b plans while along c no dramatic change occurs. As observed by Vogg et al. [273] for epitaxy grown Ca(Si1-xGex)2, the a parameter increases with increasing Ge content, following the Vegard's law as for the Si1-xGex solid solution. Then, the c lattice constant must be independent of the Ge content since CaSi2 and CaGe2 present almost the same values. The XRD measurements did not show segregation into CaSi2 and CaGe2, showing that both are completely miscible. In order to evaluate the morphology and composition, SEM and EDX analysis were performed for the Ca(Si0.5Ge0.5)2 composition (Figure III.47). The SEM shows a layered morphology, with a very inhomogeneous particle size distribution and stacks of several layers. The EDX analysis indicates the presence of 23.41% of Ge, 3.86% of Si and 11.34% of Ca with traces Cl, remaining from the synthesis process. The EDX map showed areas that are richer in Ge than in Si; the Ca is present along with the Si and Ge, while the Cl is randomly distributed. The fact that the Ge fraction is higher than the Si, probably is related to a solubility range, meaning that only certain compositions of Ca(Si1-xGe)2 are thermodynamically stable. Additionally, there could be a phase segregation and different compositions could coexist. The Ca has a uniform distribution in the sample and perhaps all the phases correspond to Ca(Si1-xGex)y. For this sample the different tested areas provided similar compositions.

It is worth to note that the Ca(Si1-xGex)2 phases decompose into CaSiGeOH upon exposure to air. The decomposition implies amorphization and can be tracked by naked eye due to a change in color related to the x content (greenish for 0≤x≤0.5, yellow for x<0.7 and reddish for x=1). In this decomposition reaction the Ca ions are not topotactically deintercalated but rather they react with the moisture to form Ca(OH)2 monolayers that are stacked in between the siligane (HSiGeH) layers. Due to the presence of Ca(OH)2 the c lattice constant is almost double (65Å) than the one of 6R-CaSi2 (30.68 Å) or 6R-CaGe2 (30.71 Å). [150,273] Figure III.47. EDX analysis for the Ca(Si0.5Ge0.5)2 composition

Siliganes

Synthesis and characterization

The synthesis of the layered siligane have been performed following the same protocol as for siloxene and germanane, the Ca(Si0.1Ge0.9)2, Ca(Si0.5Ge0.5)2 and Ca(Si0.9Ge0.1)2 phases were dissolved in concentrated HCl at -20°C. The powders were recovered by centrifugation, washed with distilled water and acetone and dried under vacuum for 30 min. The color of the products was related to the composition, the Si-rich was green, the 1:1 Si:Ge was brownish and the Ge-rich was dark brown. These phases are also known as siliganes, and from now they will be referred as siligane Si0.1Ge0.9, siligane Si0.5Ge0.5 and siligane Si0.9Ge0.1 for the phases coming respectively from Ca(Si0.1Ge0.9)2, Ca(Si0.5Ge0.5)2 and Ca(Si0.9Ge0.1)2. The XRD pattern of the siliganes (Figure III.48) show slightly differences compared to siloxene and germanene. Interestingly, the crystallinity of the intermediate phases seems to increase compared to siloxene and germanane, the origin of this behavior is unknown. The siligane Si0.9Ge0.1, shares most of the reflections with the siloxene, such as the (001) at 14°, the (100) at 27.25°, the (011) at 30° and the (110) at 47.4°, this phase presents a very intense peak at 28.4° assigned to the presence of c-Si. The (011) reflection has increased in intensity with respect to the pristine siloxene. Other peaks that are present also in siloxene are observed in the siligane Si0.9Ge0.1 at 22.5, 39.7, 44.6, 46.25, 49.6, 56° with an increased intensity. For the siligane Si0.5Ge0.5 certain peaks that were present in siloxene are absent, for instance the (001) reflection is not observed, instead there is a peak at 18.67°, the (100) reflection is absent as well and only the very intense peak from c-Si (28.4°) is noticed at these angles. The reflections (011) at 30° and (110) at 47.4° presented an increased intensity compared to siloxene. Other small peaks are observed at 22.25, 39.7, 44.6, 46.25 and 57°similarly to Siligane Si0.9Ge0.1. For the Siligane Si0.1Ge0.9 some of the diffraction peaks are missing as well; the (002) in germanane, the (001) and ( 100) in siloxene and the c-Si. Instead another peak at 22.7°, the (011) reflection at 30° and the peaks at 39.7° and 46.25° appear in siligane Si0.1Ge0.9 and the other siliganes with a high intensity while in germanane and siloxene they are absent or very small. No peak corresponding to the reflection (110) at 47.4° in siloxene appears for siligane Si0.1Ge0.9, but at this angle there are two reflections (110) and ( 112) from germanane. Additional peaks at 35.7° (also in germanane) and 57.4°.

As observed the siliganes present certain characteristics proper from the germanane and siloxene.

Table III. 7 shows the different peaks for the siloxene, germanane and siliganes with their respective hkl reflections; as the fraction of Ge increases additional reflections proper from germanane stars to appear while those ones from siloxene gradually disappear. The increased intensity for some diffraction peaks in the siliganes compared to siloxene and germanane is in line with the work of Luo et al [149] who simulated the diffraction pattern of germanane and found that due to the inherent disorder only certain peaks appear in the experimental XRD pattern. This, accompanied with the possible existence of different polymorphs (non-identified) could favor the diffraction of certain orientation planes increasing their intensity. On the other hand, the studies suggest that the siligane keeps the same lattice constants as the siloxene and germanane, with an increase in the values as the fraction of Ge increases, following the Vegard's law. [228,274] The structure of siligane might consist of buckled sheets in which Si and Ge are alternatively placed in two sub lattices and the maximum Ge fraction, is determined by the maximum stoichiometry of Ge that can be incorporated into the Zintl precursors. [160,274] Nevertheless, until now no deep structural characterization has been done in these phases. The Raman spectra in Figure III.49 the presence of three main bands, the Ge-Ge at 285-295 cm -1 , the Si-Ge at 390-395 cm -1 and the Si-Si at 490-505 cm -1 , in analogy to the Si1-xGex and Ca(Si1-xGex)2. Note that these bands are in the same region but they do not present the same shift, particularly for the siligane Si0.1Ge0.9. These differences in the Raman spectra are expected due to variations in the bond length, caused by the formation of the solid solution, size and ratio of ligands (-H/-OH) and the disorder in the sample. [156] Note that Ge will preferentially bond with -H, while Si with both -H and -OH. And that the ratio between -OH/-H groups influences the vibrational properties and lattice constant in a similar way to the Vegard's law for solid solution. [156,228] The Raman shift is dictated also by the bond length that changes with Ge incorporation in the lattice. Indeed Arguilla et al. [160] observed a shift of the Raman modes to lower values as the fraction of Sn is increased in the layered Ge1-xSnxH1-x(OH)x solid solution. proper from siloxene; while the siligane Si0.1Ge0.9 had a small shoulder in the band close to 885 cm -1 (H-Si-O) possibly due to the presence of Ge-O-Ge (830 cm -1 ). No additional bands are found for the Siligane Si0.5Ge0.5. Thus, the siliganes exhibit features that are more in line with siloxene than with germanane. However, due to the similarities in their vibration frequencies their differentiation is complex for certain bands, like Si-O-Si and Ge-O-Ge. There is a remarked presence of the Si-O/H-Si-O bonding over the Si-H, indicating probably a higher degree of oxidation and the preferential bonding of Si with -OH over Ge-OH. The previous works for the siliganes from epitaxy grown Ca(Si1-xGex)2 show an increase in the Si-O-Si bands when the fraction of Si is increased. [150] Figure III.50. Infrared spectra for siligane Si0.1Ge0.9, Si0.5Ge0.5 and Si0.9Ge0.1.

The morphology was analyzed by SEM, the profile of the siligane Si0.1Ge0.9 (Figure III.51) seems different compared to the others siliganes and the siloxene or germanane. There is a partial preservation of the layered morphology but some sort of flakes and spherical particles are observed.

This sample is extremely sensitive to the electron beam, and it was not possible to perform the EDX composition analysis. The previous results combined with the EDX analysis of the Ca(Si0.5Ge0.5)2 demonstrate the inhomogeneity of these phases. Perhaps there are solubility limits and only certain compositions are thermodynamically stable, as observed for the Ge/Sn analogues. [160] Based on the studies for Ge/Sn [160] and Ge/Ga(As) [158] phases, the replacement of Ge atoms might be limited and constrained to certain crystallographic positions. The introduction of a bigger atom could generate repulsion in the structure and the creation of different neighboring environments. Indeed, in certain cases the system prefers to adopt a two-or several phase configuration rather than any electronic or structural distortion. [91] This theory could explain the differences in the XRD and Raman spectrum for the siligane Si0.1Ge0.9, meaning that only a certain fraction of Si can be incorporated in the lattice without losing the phase homogeneity. Indeed for the Sn/Ge phases only 9% of Sn can be inserted in the Ge lattice. [160]

Electrochemical cycling of the siliganes in Li-and Na-ion half cells

Following the electrochemical performance of the siligane Si0.9Ge0.1 and Si0.5Ge0.5 vs Li and Na was tested (Figure III.54), since the cycling conditions for K were not optimized, the system was not studied in this case. For the first discharge the lithiation capacities are 3181 and 2400 mAh/g and the sodiation 668 and 480 mAh/g for Si0.9Ge0.1 and Si0.5Ge0.5, respectively, while the second discharge delivered 1100 and 1130 mAh/g for Li and 238 and 170 mAh/g for Na. As expected the large presence of Si in Si0.9Ge0.1 contributes to increase the reversible capacity for Li and surprisingly for Na. The galvanostatic curve for the lithiation of both siliganes is similar to the one of siloxene. The first discharge is characterized by a plateau at 1.2 V, attributed to the FEC reduction, followed by a sloping profile that occurs at lower voltage (0.16 V) for Si0.5Ge0.5 compared to Si0.9Ge0.1 (0.4 V). Upon charge for Si0.9Ge0.1, there is a region of different slope between 0.25-0.5 V, while for the Si0.5Ge0.5 a plateau is observed at ≈0.35 V. In the second discharge, for Si0.9Ge0.1 there is a plateau at 0.4 V that is absent for Si0.5Ge0.5. For both compositions the second charge resembles the first one with no remarkable change. The galvanostatic curve for sodiation also resembles the one of siloxene. For both siliganes the first discharge has a sloping profile with plateau at 1.2 V. During the first charge, for Si0.9Ge0.1 there is a region with different slope from 0.37-0.65 V while for Si0.5Ge0.5 the potential increases constantly. For the second and subsequent discharge/charge cycles both siliganes present a sloping profile. The phenomena in the galvanostatic curve and derivative can be indicators of the presence of a solid solution, taking in consideration the reports for the lithiation of the Si1-xGex alloys. [72,73] For instance for the solid solution Si0.9Ge0.1 the derivative curve for the first cycle shows a double reduction peak 0.10 V and two oxidation ones at 0.3 and 0.45 V, in line with the results for the siligane Si0.9Ge0.1. [73] Although, the profile for the second discharge is different and the solid solution Si0.9Ge0.1 presented two reduction peaks at 0.18 and 0.28 V whereas the siligane Si0.9Ge0.1 two reduction peaks at 0.32 and 0.01 V. For the solid solution Si0.5Ge0.5 the first cycle has one reduction peak at 0.2 V with a small shoulder at 0.1 V and one oxidation one at 0.45 V, while the second discharge two reduction peaks at 0.30 and 0.1 V. On the contrary, the siligane Si0.5Ge0.5 presents different reduction peaks but the oxidation profile is almost the same as for the solid solution. This information suggests the presence of a solid solution, for instance a single signal is observed for all siliganes rather than separated phenomena from siloxene and germanane, and consequently there should not be segregation. Previously it was observed that the cycling of the Si1-xGex came at the expense of a capacity loss but the capacity retention was enhanced. Considering the previous studies that show an excellent electrochemical performance of germanene and siloxene, other factors must be taken into account in order to establish the real benefit of a solid solution.

The 2D Si-Ge phases have been studied theoretically and, the siligene (2D Si-Ge) and siligane (HGe-SiH) are predicted to be more stable than the silicene and germanene. [275] The siligane is particularly interesting for post-LIB applications, with expected reversible capacities of 532, 1064 and 532 mAh/g for Li (Li2SiGe), Na (Na4SiGe) and K (K2SiGe), respectively. Both sides of siligene can absorb alkali ions, however an excessive quantity causes detrimental distortion of the SiGe backbone, then the maximum ion incorporation is as follows Li2SiGe, Na4SiGe and K2SiGe. [274] These predictions are very promising and suggest that after an exhaustive study of the layered Ca(Si1-xGex)2 Zintl phases and their derived siliganes, excellent electrochemical performances can be assessed. Thus, this work sets the bases for the experimental study of the layered solid solution of Si and Ge.

Summary and conclusion

The siliganes were synthesized in three composition ranges Si0.9Ge0.1, Si0.5Ge0.5 and Si0.1Ge0.9. In analogy with CaSi2 and CaGe2 the synthesis begins by the preparation of their corresponding layered Ca(Si1-xGex)2. An XRD analysis of these last ones revealed a change in the a and b lattice parameters due to the Ge incorporation in the lattice, manifested in an increase in the unit cell size and a shift of the (110) Bragg peak towards lower angles. The Raman spectra showed the presence of the Ge-Ge, Si-Ge and Si-Si bands. The SEM pictures indicated a layered morphology, while the EDX revealed a difference in the composition, possibly there is a solubility range for the Ca(Si1-xGex)2 phases and only certain compositions are stabilized.

After the Ca deintercalation and formation of the corresponding siligane, the XRD pattern showed some changes, for the Si0.5Ge0.5 and Si0.1Ge0.9 the Bragg diffraction peaks proper from siloxene and germanane have disappear, while for the Si0.9Ge0.1 the pattern resembles the one of siloxene. Certain diffraction peaks have increased in intensity, respect to siloxene and germanane, probably due to a preferential diffraction of a specific plane caused by the disorder or the formation of different polymorphs in the phase. The Raman spectra showed the presence of the Ge-Ge, Si-Ge and Si-Si, although the bands are shifted due to the variation in the bond length, caused by the formation of the solid solution, the size of the -H/-OH ligands and the inherent disorder in the phases. The infrared spectroscopy, featured the bands associated with Si-OH, H-SiO3 and Si-O-Si/Ge-O-Ge, suggesting an oxidation and a preferential formation of Si-OH bonds over Si-H and Ge-H. The SEM showed a layered compound for Si0.9Ge0.1 and Si0.5Ge0.5 while the morphology of Si0.1Ge0.9 included layered and flake-like particles. The EDX analysis for the Si0.5Ge0.5 revealed a composition that is richer in silicon than in Ge (contrary to Ca(Si0.5Ge0.5)2) while the one for siligane Si0.9Ge0.1 was in agreement with the expected composition. These results support the hypothesis of the existence of solubility ranges and the stabilization of only certain compositions. Additionally, the Ge replacement in the lattice is limited and possibly constrained to certain crystallographic positions. In this case, the synthesis conditions might impact the structure of the final product, as observed for the modified siloxene. Therefore, a more rigorous structural study must be performed.

Finally, the electrochemical behavior of the siligane Si0.9Ge0.1 and Si0.5Ge0.5 was evaluated vs Li and Na. The results showed lower capacities compared to the pristine siloxene and germanane. A comparison of the derivative curve with those ones for the Si1-xGex alloys, suggest the presence of a solid solution without segregation into germanane or siloxene. Although, as suggested for the Ca(Si1-xGex)2 phases only certain solubility ranges might be thermodynamically stable and more than one composition can coexist in the same sample. A priori the siligane did not show any improvement in the electrochemical performance, but the theoretical predictions expect 532, 1064 and 532 mAh/g for Li, Na and K, respectively. [274] Hence, in order to access these capacities a detailed characterization of the siliganes is necessary. The information to gather should include among others, the identification of the polymorphs and the atomic positions for Si and Ge; the role of the preferential bonding of Ge with -H and of Si with -H and -OH and their reactivity towards the electrolyte, the different intermediates formed during cycling and the degree of accessibility of the alkali ions to the layers of siligane. This work is expected to set a base for the exploration and use of the siligane phases for battery and several other applications.

Conclusion

A comprehensive view of the silicon-and germanium-based layered Zintl phases was presented in this Chapter. The CaSi2 and CaGe2 phases present a layered structure formed of puckered Si6(Ge6) rings interconnected to form planes, several stacking sequences for the layers are possible, deriving in a whole set of polymorphs. The presence of the Si -(Ge -) Zintl ion allows the topotactic deintercalation of Ca and the preservation of the Si (Ge) backbone, forming the so-called siloxene (Si6OH3H3) and germanane (GeH)n. The Ca extraction produces an electronic instability in the layered structure that is compensated by the formation of Si-H and Si-OH bonds for the siloxene and Ge-H for germanane.

Additionally, a variation in the synthesis media permits the replacement of -H and -OH bonds by different organic molecules and functional groups, the degree of substitution is related to the size of the ligand. These modifications allow to tune specific properties in siloxene and germanane, such as band gap, thermal stability, interlayer distance, among others. In this work the expansion of the interlayer distance and the exfoliation of the siloxene layers was achieved by the intercalation of organic carbons of different carbon chain length. The electrochemical behavior of siloxene and germanane was tested for Li, Na and K. The results are very promising, as the layered structure serves as a buffer for the volume expansion of the associated electrochemical reaction, while enhancing the electronic conductivity and ion diffusivity. These properties derived in an excellent capacity retention, offering a real improvement compared to bulk Si or bulk Ge. The tuning and expansion of the interlayer distance in siloxene, did not fulfill the expectations, perhaps the presence of the organic molecule in between the planes imposes a steric effect that hinders the alkali-ion access to the bulk. Therefore an effective method for removal of the organic molecule while preserving the structural integrity of siloxene is necessary. Finally, in an analogy with the previous Chapter for the Si1-xGex solid solution, in which the synergy between Si and Ge provided a superior capacity and capacity retention, the layered Si1-xGex solid solution, the siliganes, derived from Ca(Si1-xGex)2 were synthesized. In this case despite the formation of the solid solution, the composition is inhomogeneous, possibly due to solubility domains and electronic repulsion caused by the presence of a foreign element that forces the system to stabilize in more than one phase or generates a structural change/distortion. The electrochemical behavior of such phases was not as satisfactory as pristine siloxene and germanane, but the theoretical predictions indicate that an excellent performance can be achieved. Hence a more exhaustive structural study of the siliganes is necessary to access their potential. In conclusion, this work demonstrate the high versatility of the layered Si-and Ge-based Zintl phases and sets a precedent for their use in batteries, yet the applications are not limited to this domain and the results presented here can serve as base for other studies.

Introduction

The layered siloxene and germanane have proved to be very versatile phases that can be easily tuned for certain applications. Their use as anode materials for LIB, NIB and KIB is very promising and the obtained performances are improved compared to bulk Si and bulk Ge. As observed in Chapter II, the electrochemical profiles of siloxene and germanane appear to slightly differ from their bulk counterparts. Considering that both Si and Ge are very sensitive to the cycling conditions (morphology, particle size, electrode formulation, etc.), the lithiated phases formed during cycling might not always be the same. A literature research reveals various publications regarding the electrochemical performance of siloxene and germanane-type compounds, but no work is devoted to the identification of their lithiated compounds. Consequently, in this chapter, a series of characterization techniques including XRD, Raman and IR spectroscopy, TEM, EELS, XPS and NMR have been combined in order to study the electrochemical mechanism of siloxene and germanane. The obtained results have been compared with the previous knowledge of the alloying mechanism of bulk Si and Ge, presented in Chapter II. The similarities and differences between both systems have been highlighted. This work is exploratory and set the bases for future and more comprehensive studies. The ultimate goal would be to understand how differently Si and Ge works as a bulk and a layered material, and identify their advantages and limitations in order to improve their performance as anodes for LIB, NIB and KIB.

Experimental details

This section describes the main parameters for the characterization of the lithiated intermediates in siloxene, for additional details please refer to the Annex I section.

Self-supported electrode. In order to rule out the contribution of the Cu current collector to the different characterizations a self-supported electrode was prepared. The mass ratio was 1:1:1 of active material (siloxene/germanane), C45 conductive additive and CMC binder (CMC-DS = 0.9, Mw = 700 000 Aldrich). The slurry was prepared in distilled water with 0.1% of Triton X dispersant, casted on a Mylar foil using a Doctor Blade (100 µm thickness), dried under ambient conditions for one day and vacuum dried at 70°C overnight.

In situ XRD. The diffractograms were recorded in a Bruker D8 Advanced diffractometer with Cu radiation (λ1 = 1.54056 Å, λ2 = 1.54439 Å), using a specifically designed cell with Be window. [169] The cell was assembled with a self-supported electrode, using 1 M LiPF6 in EC/DMC with 1% FEC, a Whatman glass fiber separator (GF/ D, 675 μm) and Li metal as a counter electrode. In order to ensure a complete stabilization of the intermediates, the cycling was performed at C/40 (C/n = 1Li /n h) versus Li/Li + . The diffraction patterns were recorded at a set potential, in a 2θ range of 8-40° during one hour.

In Situ Raman Spectroscopy. The specific cell with quartz window [170] was assembled following the same protocol as for the in situ XRD and cycled at C/10. For the experiments performed with Na, NaPF6 in EC/DMC with 1%FEC was used as electrolyte and Na metal as counter electrode. The Raman spectra were recorded in a Dilor RT30 spectrometer with HeNe laser source operating at a wavelength of 633 nm with a 180° backscattering geometry. The measurements were performed at a set potential with a total recording time of one hour.

Cyclic Voltammetry. The measurements for the siloxene and germanane were performed in Swagelok cells, using 1 M LiPF6 in EC/DMC +1%FEC, a Whatman glass fiber separator and Li metal as counter electrode. A VMP3 Biologic device was used, the voltage window was 0.0-2.0 V with a scan rate from 0.1-100 mV/s.

For the ex situ characterizations, self-supported siloxene electrodes were cycled in Swagelok cells, with 1 M LiPF6 (NaPF6) EC/DMC + 1%FEC, a Whatman glass fiber separator and Li (Na) metal as a counter electrode. The cells were cycled at a rate of C/20 (C/n = 1Li (Na)/n h) versus Li/Li + or Na/Na + , using a VMP3 or MPG2 (Biologic) device. The electrodes were recovered inside an Ar-filled glove box, washed with DMC and dried under vacuum.

The cycled electrodes were characterized by Raman spectroscopy using a DXR 2 Raman Microscope spectrometer (Thermo-Fisher Scientific, 532-nm excitation). In order to avoid contact with air, the samples were placed in a glass lid and sealed with Kapton foil. The Fourier Transformed Infrared spectroscopy was performed using KBr methods and a cell with NaCl window to avoid sample contact with air. The morphology and composition were analyzed by SEM (SEM-environmental FEI Quanta 200 FEG) and EDX (X INCA Oxford). The TEM, STEM and EELS were performed in a TEM-FEI Tecnai F20 S-TWIN, operating at 200 kV fitted with a Gatan Image Filter Tridiem in post column. The electron diffraction patterns were obtained with Selected Area Electron Diffraction (SAED). The Electron Energy Loss Spectroscopy (EELS) was performed with 1-2 eV energy resolution, a dispersion of 0.2 eV/ch, a conversion and a collection angle of 5.8 and 2.2 mrad, respectively. The energy loss near edge structure (ELNES) acquisitions were performed in STEM mode in defocus to limit the electron beam interaction.

All the energy losses presented an error of ±0.2 eV. The Energy Dispersive X-Ray (EDX) was acquired in order to determine the local chemical composition. The air sensitive samples were prepared in an Ar filled glove box and transferred to the TEM under an Ar flow. [276] The XPS was performed using a Escalab 250 Xi spectrometer with a monochromatized Al Kα radiation (hν =1486.6 eV). The electrodes were placed on a sample holder using an insulating uPVC tape (ref. 655 3M) and then transferred to an Ar-filled glove box directly connected to the spectrometer. The analysis were performed with the standard charge compensation mode and an elliptic 325 x 650 µm X-ray beam spot. Core spectra were recorded using a 20 eV constant pass energy with a 0.15 eV step size and short time iterative scans.

The 6 Li and 1 H NMR were performed with a Bruker Avance-500 spectrometer using a Bruker MAS probe with a 2.5 mm diameter zirconia rotor. Spinning frequencies up to 25 kHz were used. 6 Li NMR spectra were acquired with a single pulse sequence and a recycle time of 30 s.

Exploration of the electrochemical mechanism of siloxene and germanane

The galvanostatic cycling of siloxene and germanane (Chapter III) suggested a somewhat different profile compared to bulk Si and bulk Ge. As observed in Chapter I and II both Si and Ge are particularly sensitive to the cycling conditions and it is possible that the layered morphology induces some changes in the phenomena that take place. This section presents a combination of different characterization techniques employed to identify the intermediated phases during the electrochemical cycling of siloxene and germanane. This information combined with the previous knowledge from the alloying mechanism of bulk Si and bulk Ge permit to establish the differences and similarities between them and the impact of the layered morphology over the system. Note that most of the experiments were done for the lithiation of siloxene since more information is available for the interpretation of the results.

The analysis of the system begins by the in situ XRD diffraction, in order to identify the presence of the crystalline phase Li15X4 (X= Si, Ge) characteristic of the alloying reaction of Li with Si or Ge. Nonetheless, the crystalline Li15Si4 formed at the end of discharge for bulk Si, is absent in siloxene.

Instead a second hump appears at around 22.5° with no trace of any other crystalline phase. Note that this experiment was performed at a C/rate of C/40, in order to allow sufficient time for the different phases to stabilize. These observations are consistent with Gao et al. [240] who after lithiation identified the XRD peaks proper from siloxene with an increased broadness and a decreased intensity, accompanied by two broad humps at 15 and 22.5°. Thus the periodicity of siloxene must have been transformed into an amorphous-type phase. No crystalline LixSi phase was found. for the CaSi2, siloxene and discharged and charged product for Na and Li; the same basic structure of Si-Si in Si-planes defined by an intense peak centered between 492-507 cm -1 and the second order TA and TO at 283 and 900 cm -1 , can be observed in all of them. Note that upon Ca deintercalation from CaSi2 to form siloxene, new bands at 635, 735 and 2115 cm -1 appear, ascribed to the Si-H vibrations.

These bands disappear upon electrochemical (de)sodiation/(de)lithiation but the Si planes vibration modes are preserved, possibly indicating no structural degradation with cycling. Additional bands emerge at 1330 and 1600 cm -1 , corresponding to the disordered (D) A1g breathing mode of defects/vibrations and the graphitized (G) E2g vibration of the C-C plane from the conductive carbon. [277] Interestingly, the main Si-Si planes Raman band position seems to be related with the composition; for CaSi2 it is centered at 493 cm -1 and at 497 cm -1 once the Ca is replaced by -H and -OH in siloxene.

Upon sodiation/lithiation it shifts to 504 and 505 cm -1 ; and after desodiation/delithiation to 494 and 492 cm -1 , respectively. This shift could be related with a change in the interlayer separation and in the bond nature between the Si-planes and the substituent group. Then, upon electrochemical The analysis of the cycled electrodes by Raman spectroscopy was performed for germanane also (Figure IV.4). As observed in Chapter III the spectra between CaGe2 and germanane are different; the CaGe2 presents multiple bands while the germanane only one peak centered at 302 cm -1 . Upon lithiation/sodiation, a peak is observed at 290 cm -1 , accompanied by the D and G bands from the carbon additive at 1330 and 1600 cm -1 , as observed for siloxene. Upon delithiation/desodiation no drastic change in the spectra is observed and the main peak is centered at 290 cm -1 , contrary to siloxene, this peak does not present a redshift. Therein, the peak at 290 cm -1 has been assigned to the Ge-Ge bonds in the plane that is redshifted compared to pristine germanane. The redshift suggests a change in the interlayer distance, which is also related to the presence of mainly Ge-H bonds and few/no Ge-OH. The germanane has a preferential bonding with -H, in contrast with siloxene which bonds with both -OH and -H. [268] The redshift of this band upon lithiation has been reported in the literature. [152] The authors found a broad band at 269 cm -1 , attributed to the amorphization of germanane, caused by the oxidation and dehydrogenation. A zoom in the region from 400-600 cm -1 , showed a decrease in the intensity for the 495 cm -1 band as the lithiation proceeds, although no shift was observed. On the other hand, it is possible to track the Li insertion in the conductive additive; as the Li content increases the D (1330 cm -1 ) and G (1600 cm -1 ) band increase in broadness until merging into a single band. Their intensity starts to decrease and completely disappear at the end of discharge, meanwhile the fluorescence intensifies. According to the reports in the literature, the lithiation of the conductive additive is characterized by a decrease in intensity of the D band and a shift towards lower wavenumbers, while the G band generally splits into two bands (E2g2(i) 1575 cm -1 and E2g2(b) 1580 cm -1 ) and shifts to higher wavenumbers. As the content of Li increases and the conductive LiC6 phase is formed, the optical skin depth is reduced resulting in a lower Raman scattering. [45,[START_REF] Baddour-Hadjean | AIP Conf. Proc. 2010[END_REF][START_REF] Zou | [END_REF][280] accompanied by a less intense one at 495 cm -1 and the D (1330 cm -1 ) and G (1600 cm -1 ) bands from the conductive additive, are observed. In analogy for the lithiation, the signal from the quartz window at 464 cm -1 impedes the proper assignment of the bands from siloxene. The intensity of the 518 cm -1 peak increases during the OCV period until reaching a maximum at 1.5 V. Subsequently it decreases and at the end of discharge this peak is barely noticeable. This reduction in intensity can be related with the formation of the SEI layer that hinders the Raman scattering, as previously observed in other works. [32] An evolution of the conductive carbon bands is observed as well, contrary to the lithiation, these bands do not disappear at the end of discharge and instead for voltages <0.08 V the broadness increases. At some stages of sodiation the intensity ratio between the D and G band is different, likely due to a better organization of the graphite layers or to an increased interlayer distance. [281,282] Even though the in situ Raman spectroscopy offered interesting results for the lithiation/sodiation of the conductive additive, it did not provided conclusive information regarding the shift of the Si-Si band in siloxene.

The high noise/signal ratio, the presence of the electrolyte that reflects back some of the incident light and the thick SEI layer on the particle surface hinder the proper observation of the phenomena. Thus, the conditions for performing the experiment must be improved in order to get valuable information. for the expected Li15Si4 to be formed at the end of discharge for bulk-Si does not correspond to the spectrum of the lithiated siloxene. Other possible phases that might be stabilized at this stage are Li22Si5 and the intermediate phase Li12Si7, also very different from the lithiated siloxene. Note that the Raman spectra of Li15Si4 and Li22Si5 resemble, both phases are composed of isolated Si 4+ ions surrounded by Li + , whereas the Li12Si7 contains planar five membered rings and star-trigonal planar M4. On the contrary the spectrum of LiSi resembles the one of lithiated siloxene, except that the main peak for LiSi is centered at 511 cm -1 while for the lithiated siloxene at 505 cm -1 . [105] According to the literature, [104] the LiSi has several layered polymorphs, and one of them presents the same planes composed of Si6 rings as in siloxene. The theoretical capacity for LiSi is 954 mAh/g [283] which is lower than the capacity values obtained for siloxene, meaning that possibly the structure of the lithiated siloxene is somewhat similar to LiSi but not exactly the same. Consequently, given the absence of the c-Li15Si4 in the system, the intermediates and ending members formed during the lithiation of siloxene might be different to those of conventional bulk Si. [243] who observed a preservation of the layered morphology of siloxene after cycling. Interestingly, the LiF is found as an assembling of multiple nanocrystals while the areas which the sharpness of the layers is preserved consist of an amorphous compound. [284] These observations reinforce the theory that the formation of the passivation layer is responsible for the modification of the edges in the layers. cases, the double contribution is still present but at 62.3 and 67.5 eV in line with a lithium-carbonate type Li2CO3 or methyl lithium carbonate (LiC2O3H3). [285] Additionally, a more complex spectrum with multiple contributions type lithium oxide, lithium carbonate and lithium fluoride was also observed. In general the Li2O-type and Li2CO3-type compound is found in the regions where the edge of the layers is sharp. [284] An additional contribution present in all the spectra is possibly related to the Li in the siloxene layer. Unfortunately due to the penetration depth the EELS and the significant thickness of the SEI layer, an analysis of the bulk of the layers was not possible. Thus, the determination of the Li in the SEI layer and the one inserted inside the siloxene is rather hard to accomplish. It is worth to note that the decomposition layer is sensible to the electron beam and can be degraded after the EELS analysis. the characteristic signal from Si 4+ . [240,246] For the lithiated siloxene an evolution of the Si L-edge shape is observed, and the shoulder at low energy that accompanied the peak at 110 eV disappears and is replaced by a shoulder at high energy. This change in profile could be an indicator of the presence of Li in the siloxene layers, although it could have other origin like a decomposition product type-LiySiOx.

The analysis showed as well the presence of areas without lithium, meaning that certain particles do not lithiated.

The germanane was also analyzed by TEM ( In order to further investigate this decomposition layer and gather information about the presence of Li inside the siloxene layers, X-Ray Photoelectron Spectroscopy (XPS) was performed. Figure IV.17

shows the spectra for the fresh siloxene electrode and the pristine siloxene. Noticeable changes are observed in the profile after the electrode preparation. The siloxene oxidizes during this process and the Si 0 peak intensity is reduced and shifts to lower energies. This peak shifting is related to the ability of Si to form various sub-oxides via bonding with different number of oxygen atoms. [286] These variations induced by the electrode aqueous preparation have been reported for other Si-based electrodes using CMC and water as a solvent. [286] considered, then the Si 3+ peak appears at a lower energy for the fresh electrode. The peak at Si 3+ is most probably coming from the formation of Si2O3 suboxides and their disproportionation reaction [252] while the presence of Si 4+ is an indicator of an oxidation process which is reported to take place from the surface to the core. [263] The Si-O binding energies for the fresh electrode are slightly different compared to bulk SiO2 (104 eV), suggesting the presence of an additional compound rather than merely SiO2. [287] The type of oxide that is formed is highly influenced by the surface activity and it is likely to be an oxygen doped silicon oxide. [239] It is believed that the oxygen content is beneficial for the electrode, since it helps to buffer the volume variations and to keep the structural integrity.

However, an optimal oxygen content is required in order to avoid problems associated with the SiOx matrix, such as excessive formation of Li2O and Li-silicates that lead to a sluggish diffusion. [263] In all cases the Si 0 and Si + peaks have decreased in intensity, meaning that the oxidation process has also caused a loss in the fraction of the Si-H bonds. [243,250,251] This evolution of the spectra is not unexpected considering that the electrode is prepared using water as a solvent. Table IV.2 summarizes the peak position and atomic concentration for the fresh electrode. The XPS displays also peaks coming from the CMC and the carbon conductive additive. to the first one, the Si 0 is not observed, but a new peak appears at high energy (≈105 eV). This peak is likely originated from the LixSi phases. Nonetheless, in order to have an accurate analysis, the surface must be sputtered with Ar + to partially remove the thick SEI layer. Different reports in the literature state that the lithiation process in siloxene leads to a decrease in Si 0 intensity and an increase in the Si-O peaks, the ratio between the peaks depends on the morphology and surface reactivity of the sample. The degree of lithiation/delithiation can affect also the peak position and intensity. [288] The formation of the SEI layer can alter the intensity of the Si 0 and Si-O peaks, due to the interactions between Si and the decomposition products. [286,289] The presence of the Si 3+ peak and its shift to lower energies is related to the Si2O3 sub-oxides in the electrode and its disproportionation and transformation into Li2SO3/LixSiOy/Li2Si2O5. [289] The Si + possibly is originated from LixSiOy or Si-OH. Lastly, the components of the SEI layer can be listed as hydrocarbons/alkyl groups (285 eV), C atoms in a twooxygen environment like Li-esters, carboxylates or oxalates (286 eV) and Li-alkylcarbonates/Li2CO3

(289 eV), LixSiOy (531 eV) and Li-alkyl carbonates (-CO) or LiOH/-C=O (533 eV). [254,289] In order to investigate the origin of the new peak at high energy for the lithiated siloxene, the XPS spectrum of the Li15Si4 reference synthesized by ball milling was taken (Figure IV.19-rigth). The results indicate the presence of Si 0 , Si + , Si 3+ and Si 4+ , and an additional peak 98.95 eV, ascribed to a LixSi phase.

No peak at ≈105 eV was found, thus it might come from other specific Si species. This peak is in line with the one observed for the second lithiation of siloxene. Note that Li15Si4 is a metastable phase and decomposition products can be present, hence the surface has been sputtered to properly access the Si signal. Table IV.3 presents the peak assignment for the Li15Si4. The exploration of the different intermediate and ending phases during the lithiation of siloxene continues with the solid state NMR spectroscopy. Intially, the experiments were performed with 7 Li, but the resolution was not optimal. Hence, all of the following experiments were performed with 6 Li.

Despite its less abundance and longer relaxation times, the homonuclear dipole-dipole interactions in 6 Li are weaker and the quadrupolar broadening is smaller, thus under MAS conditions most of the anisotropic interactions are removed and the resolution of 6 Li MAS NMR is higher compared to 7

Li. [191,290] Figure IV. 20 shows the MAS 6 Li NMR spectra for the lithiated and delithiated self-supported siloxene anode, compared to the LiSi and Li15Si4 references synthesized by ball milling. For the lithiated/delithiated siloxene spectra the presence of two main Li environments can be identified; the peak at approximately 0.7 ppm is attributed to the SEI layer while the one at 4. [110,181,191,192,291] A deconvolution of this peak suggests an additional contribution at 9 ppm (16%), and as reviewed in Chapter II, NMR signals centered at these frequencies are characteristic of Li surrounded by both isolated Si 4-ions and small Si clusters (dumbbells), like in Li13Si4. The presence of this phase has been previously identified as a decomposition product of the metastable Li15Si4. [291] The resonance at 2 ppm (6%) is ascribed to a nonidentified LixSi phase. For the LiSi, the resonance at 19 ppm is assigned to the presence of Li atoms in Si clusters, like Li12Si7 (composed of Si5 rings and Si4 stars) or other poor lithiated phases, [110,181,191] the signal at 7 ppm is likely originated by the LiSi environment while the one at 2 ppm corresponds to a non-identified phase. A fit with only these three resonances is not satisfactory and at least two other signals must be included, at 4.5 ppm and 11.8 ppm. The former, can be ascribed to the presence of Si 4- isolated ions (Li15Si4) while the latter is in agreement with the reported 7 Li NMR for LiSi [286] and can be related as well with a Li in an environment composed of Si dumbbells and isolated Si 4-ions (Li13Si4). [191,291] The LiSi phase has a particular structure for a monosilicide, while NaSi and KSi are composed of isolated Si 4-ions, the LiSi presents different polymorphs, including flat Si sheets composed of Si4 rings, buckled Si sheets comprising four-and eight-member rings, Si sheets formed of interconnected Si6 rings like in silicene, among others. [104] The different polymorphs can be stabilized depending on the synthesis conditions, thus, various Li environments are expected for the LiSi phase.

In addition, the reported NMR studies for the LixSi phases indicate that as the Li content increases, the Si network breaks into smaller clusters until forming Si 4-isolated ions, the Li vibration frequencies decreases proportionally with the size of the cluster. Additionally, Li species with longer bond distances and their appearance at potentials below 0.18 V is in line with the galvanostatic curve in which most of the capacity is produced below 0.2 V. Once the delithiation starts, at 0.4 V the 4.8 ppm peak disappears and leaves place to a resonance at 2.2 ppm. This peak eventually fades and at the end of charge only the resonance at 0.8 ppm is observed. For the second discharge, the region at which the plateau appears in the galvanostatic curve was explored; before the plateau at 0.45 V, the spectrum is composed of a resonance at around 1 ppm that includes also the signal from the SEI layer (0.8 ppm).

As the cycling continues, the plateau in the galvanostatic appears, interestingly, the resonance at 1 ppm shifts to lower frequencies and at 0.43 V only one peak centered at 0.8 ppm is observed. The same profile is observed at 0.25 V, at this stage the plateau has already been surpassed and the potential decrease has a more sloping profile. When the potential reaches 0.10 V the peak shifts towards high frequencies and a broad signal appears at around 2.2 ppm. At the end of the second discharge, two peaks at 4.8 and 0.8 ppm have grown at the expense of the peak at 2.2 ppm. For the second charge, at 0.6 V, only a peak at 0.8 ppm is observed and remains until the end of charge.

Figure IV.23. 6 Li NMR spectra of a self-supported siloxene electrode cycled at different depth of discharge/charge.

The formation of the lithiated intermediates in siloxene is slightly different from the lithiation of bulk-Si (see Chapter II). In this last one, as the lithiation proceeds, the crystalline Si network is broken into small clusters, followed by dumbbells until the formation of isolated Si 4-ions. The breaking of the network is reflected in the NMR spectra by a continuous shift of the peaks to lower frequencies as the clusters become smaller and the amount of Li atoms surrounding the Si increases. In other words, Li environments with big Si clusters originates signal between 16-22 ppm, small Si clusters or dimmers between 12-14 ppm, small clusters and isolated Si 4-ions between 8-10 ppm and only isolated Si 4- between 3-6 ppm. [110,181] The lithiation of siloxene follows the opposite behavior, the peaks appear at low frequency and shift towards high frequencies as the lithiation proceeds. In this case no resonances are found at high frequencies, meaning that no big cluster or low coordinated Li are involved. The formation of the 2 ppm peak is intriguing; it could be related to some Li on the surface, Li next to a big cluster or with high coordination number. Indeed, the theoretical studies suggest that the most thermodynamically stable coordination site for Li in silicene is the hollow center of the Si6 ring, meaning that Li would have a coordination number of six. In addition, a Li ion in which the electron transfer has not been performed and stills in a diamagnetic environment could produce a shift close to 0 ppm. [181,283,292] Upon delithiation the peak at 4.8 ppm shifts to 1 ppm at 0.4 V and to 0.7 ppm after 0.6 V, indicating that most of the delithiation is performed at low voltages, which is consistent with the oxidation peak at ≈0.42 V in the derivative curve. The second discharge presents a particular profile, initially the plateau in the galvanostatic curve at ≈0.43 V was thought to be originated from an inherent reaction between the Li and the siloxene structure. The NMR results indicate only one peak between 1-0.7 ppm in this region. Therefore, this plateau is likely to be related with processes in the surface of siloxene, in line with the production of most of the capacity below 0.2 V. In analogy, the peak at 2 ppm, suggest a possible surface reaction. The resonance at 2 ppm is present in all cases. Hence, as stated before it could be ascribed either to a surface reaction between Li and siloxene, such as absorption of Li ions in the edges of the layers or to the presence of Li in the hollow sites of the Si6 rings (Li close to a big cluster and with high coordination number). The resonance at 4.8 ppm (end of discharge) could be related with Li15Si4 but not necessarily and considering that the lithiation of siloxene does not follow the conventional pathway for bulk Si, other type of phases with similar Li environment can coexist at that stage. Note that the siloxene presents an impurity of c-Si, thus a certain content of Li15Si4 is expected. Finally a comparison of the normalized NMR spectra with respect to the host material for the lithiated and non-lithiated siloxene indicates a mass difference of 18%. Therefore, the fact that Li is inside or not in the matrix is not negligible, in contrast to common intercalation materials, like transition metal oxides or phosphates where the change in mass is in the range of 5-7%. No process of overlithiation as previously observed for bulk Si was observed for siloxene. To shed light in the puzzling lithiation of siloxene 1 H NMR was performed. Figure IV. 25 shows the discharge/charge for the first and second cycle compared with the pristine siloxene, the fresh electrode and the CMC. For the siloxene, a main resonance is observed at 2.2 ppm accompanied by a weak signal at 0.5 ppm, the first one is assigned to the characteristic protons from siloxene while the second one to surface protons. The CMC presents a broad peak centered at 2.7 ppm and the fresh electrode at 2 ppm, which includes the contributions from both CMC and siloxene, the determination of their relative quantities is not possible due to the similarities between both peaks. The lithitiated and delithiated siloxene electrodes present also a very broad peak centered at 2.2 ppm (protons from siloxene) accompanied by a peak at 0.5 ppm (surface protons), the other signals correspond to protons in the rotor or/and in the probe. An evolution of the peak at 0.5 ppm is observed with the cycling and the peak intensity increases regardless the potential values. This phenomenon could indicate a continuous growth of the SEI layer, particularly during the first cycle, where the intensity between discharge and charge diverge more. This difference can also be associated with the degree of lithiation attained by the cell (as discussed above), in which the fraction of reversible LixSi species was higher and the amount of SEI layer was lower. Indeed, the variation in intensity of the 0.5 ppm peak for the discharge/charge in the second cycle is not large. An additional peak appears at -0.3 ppm possibly due to protons on the surface of the carbon conductive additive. Yet, according to the literature the protons with high mobility present shifts with negative frequencies, while the ones with high positive frequencies (10-20 ppm) correspond to protons with Knight shift or strong hydrogen binding. [26] A broad peak between 0-5 ppm is observed and evolves with the degree of lithiation whereas the signal between 5-6 ppm remains constant during the cycling. For the first discharge two main peaks are observed at 2.7 and -0.3 ppm, during the lithiation the intensities of both peaks vary and at the end of discharge the peak at 2.7 ppm presents a higher intensity and has shifted towards lower frequencies. Upon delithiation, a well-defined peak at 2.2 ppm is observed at 0.4 V accompanied by the signal at -0.3 ppm. At the end of charge only a very broad signal is observed, this signal most probably contains both contributions at 2.2 and -0.3 ppm. During the second discharge, no significant change is observed until 0.43 V, voltage at which the plateau takes place in the galvanostatic curve. At this point the -0.3 ppm peak decreased in intensity and the peak at 2.2 ppm is more defined. Interestingly, at 0.25 V (after the plateau) a well-defined peak is observed at -0.3 ppm while the one at 2.7 ppm has decreased in intensity. Both peaks continue to grow and at 0.10 V two clear peaks are observed. At the end of discharge, the peak at 2.2 ppm has shifted towards higher frequencies while the one at -0.3 ppm has decreased in intensity. Upon delithiation both peaks start to merge and lose resolution and the end of charge, a very broad signal accompanied by the peak at -0.3 ppm is observed.

Figure IV.27. Evolution of the protons in the siloxene for the first two discharge/charge cycles

As observed the different intensities of the peaks vary during the lithiation/delithiation processes.

Figure IV.27 illustrates the evolution of the peak intensity. As previously described the signal at 2.2 ppm is assigned to protons in the siloxene, the one at 2.7 ppm to the CMC, while the one at 5 ppm could correspond to a variation in the CMC polymer during the electrode preparation process. In overall the total amount of protons augments with the cycling, probably due an increase of protons in the SEI layer upon its formation. In analogy with previous discussions the first discharge presents a lower quantity of protons, since this cell was atypical and achieved higher reversible capacities with lower formation of SEI. For the first charge the number of total protons has increased, the second discharge follows the same tendency, except for the sample at 0.43 V (plateau in the galvanostatic curve) where the protons are less. The decrease in the total protons continues during the second charge, but suddenly increases at the end of charge. It is complicated then, to establish a correlation between the quantity of protons and the redox mechanism since there is not a clear tendency in the results. Surprisingly, the signal at 2.7 ppm corresponding to the CMC evolves during the cycling with a tendency to increase. This behavior is rather unusual since the CMC should remain constant as it does not participate in the electrochemical reaction, the same trend is found for the peak at 5 ppm. Then, the presence of another proton with a similar chemical shift to CMC is suggested. The 2.2 ppm signal, assigned to the protons in the siloxene, can be found in almost all of the spectra, although the amount of protons tends to diminish during the discharge, in agreement with the removal of the Si-H and Si-OH bonds observed by Raman and IR spectroscopy for the cycled samples. This process appears partially reversible, and upon charge the total protons increases for the 2.2 ppm signal (the sharp signal at 2.2 ppm is observed in the NMR spectra at 0.4 V). For the second discharge/charge, the protons are prone to decrease, except at 0.10 V where they increase. This behavior could indicate a potential range in which the protons are stable in the structure. Note that due to the very broad nature of the peaks and the overlapping resonances, the deconvolution of the peaks is rater complicated and an error of ±25% is expected.

Finally, in order to get more information on the electrochemical kinetics, cyclic voltammetry has been performed on both siloxene and germanane. Figure IV.28-a displays the CV curves for siloxene during the first six cycles. The first cycle presents a different profile, with no defined reduction peak while for oxidation a peak at 0.52 V is observed. The second cycle has a reduction peak at 0.3 V and an oxidation one at 0.40 V with a shoulder at 0.47 V while for the third cycle an additional reduction peak appears at 0.14 V. As the CV proceeds, for the reduction the peak at 0.3 V shifts to 0.38 V and the one at 0.14 V to 0.16 V, whereas the oxidation peak at 0.40 V shifts to 0.37 V and the shoulder appears at a slightly higher potential. These processes show a decrease of the polarization in reduction and oxidation and are in good agreement with the galvanostatic curve of siloxene presented in Chapter III, in which the plateau that appears after the second discharge and the potential of the pseudo-plateau during charge continuously shifts. Figure IV.28-b,c shows the CV performed at different scan rates from 0.1 to 100 mV/s. For scan rates between 0.1 and 1 mV/s, the profile of the CV curve shows the aforementioned peaks, as the scan rate increases the peak current is increased and the peaks become broader and less defined. Indeed, for scan rates superior to 5 mV/s the peaks become very broad and the peak current continuously rises, indicating a high polarization.

The presence of capacitive or diffusion controlled processes was calculated from the relationship between the peak current (ip) and the scan rate (v): [293] i= av b (1)

Where b is determined by the slope between log(ip) vs log(v) as in indicates a diffusion-controlled process while b=1 represents a capacitive process. For the siloxene for scan rates between 0.1-1 mV/s b=0.56 while for scan rates of 10-100 mV/s b= 0.63, suggesting that at low rates the process is diffusive but at high rates there is a slight capacitive contribution. Possibly, this last one will increase at higher scan rates. These observations agree with the rate capability curves presented in Chapter III where at high C/rates there is a decrease in the capacity, related to a diffusion limitation, although the capacity retention is significant. Contrary to siloxene, the profile of the CV curve for germanane slightly evolves after the second cycle, a reduction peak at 0.32 V appears and the peak at 0.04 V is more defined and shifts to 0.08 V as the cycling proceeds. Meanwhile the oxidation peak gradually shifts to 0.43 V while the peak at approximately 1 V decreases in intensity. These processes are in line with the galvanostatic curve presented in Chapter III, in which two pseudo-plateaus are found upon discharge and charge. In general, the profile of the CV curve for germanane suggests a good stability and reversibility for the system. Figure IV.29-b,c show the CV curves at different scan rates from 0.1 to 100 mV/s. Similarly to siloxene, for scan rates between 0.1 and 1 mV/s the CV curve displays the previously mentioned peaks, which become broader and less defined and the peak current increases with the scan rate. For scan rates from 10 to 100 mV/s, the peaks are very broad and the peak current continues to augment, indicating a high polarization. Nevertheless, peak current are much smaller than that of Siloxene, meaning germanane are less kinetically limited that siloxene. Lastly, the b has been calculated from the log (ip) vs. log (v) curve, and the value obtained is 0.59 for both 0.1 to 1 mV/s and 5 to 100 mV/s. This value indicates that most of the capacity is produced by a diffusion-controlled process with a small capacitive contribution regardless the scan rate. The presence of a mainly diffusion-controlled mechanism is probably behind the improved capacity retention properties of germanane even at high C/rates, as observed in Chapter III.

These results indicate that the storage mechanism in siloxene and germanane have a contribution from diffusion controlled mechanism, in which the Faradic reaction occurs in the bulk of the material, and a capacitive mechanism in which the storage only takes place at the surface of the layers and not in the bulk. [293]

Analysis of possible electrochemical mechanism

The electrochemical lithiation mechanism of siloxene and germanane is puzzling and the analysis of their lithiated intermediates and discharged compound showed a slight deviation from the typical features for the lithiation of bulk Si and Ge. For instance, one of the most characteristic lithiated phases for the alloying reaction of bulk Si and Ge, the crystalline Li15Si4 (Li15Ge4/Li17Ge4), identified by in situ XRD, is absent in the siloxene (germanane). The Raman spectroscopy provided more detailed information, and suggests the preservation of the Si and Ge planes after cycling. Specifically for siloxene, two additional features were identified; the loss of the Si-H and Si-OH bonds after cycling and a reversible Raman shift of the Si-Si peak associated with a change in the interlayer distance and the bond nature between the Si-Si planes and the -OH/-H or Li (Na). For germanane, the Ge-Ge planes signal shifts but it is not reversible, probably the preferential bonding of Ge with -H over -OH affects the characteristics of the interlayer space resulting in a different behavior. Lastly, a comparison of the Raman spectra of lithiated siloxene with a series of lithiated phases expected during the lithiation of bulk Si, coincides only with LiSi. Interestingly, this phase presents various layered polymorphs, one of them with the same basic planes of Si6 rings as in siloxene.

The preservation of the layered morphology was later corroborated by SEM and TEM. A decomposition layer composed of LiF, Li2O and Li-carbonates was found for siloxene while the one of germanane mainly consists of Li2O. An additional lithiated component was identified but given the significant SEI layer and the limited penetration depth of EELS, it was not possible to assign it to the presence of Li in the siloxene/germanane layers or to additional species in the SEI layer. Although, the analysis of the Si L-edge revealed a change in the profile for the lithiated siloxene, suggesting a possible interaction with Li.

The analysis of the surface of siloxene continued with the XPS. The results showed an oxidation of the siloxene during the preparation of the electrode, with a decrease in the Si-Si and increase of the Si-O signal. The first one is associated with the formation of a significant SEI layer that hinders the observation of the Si-Si, while the second one to the formation of LixSiOy, Li-OH and decomposition products from the SEI layer (Li-carbonates, Li-alkoxides, Li-esters). A comparison of the XPS spectra with the Li15Si4 reference indicates the presence of a peak associated with LixSi, not necessarily corresponding to Li15Si4. Following, 6 Li MAS NMR was performed and the lithiated/delithiated siloxene for the first and second cycle were compared with the Li15Si4 and LiSi references. The results indicate that the lithiated siloxene resonances (4.8 ppm LixSi and 0.7 ppm SEI) can include the contributions of both Li15Si4 and LiSi. The analysis at different depths of discharge, showed an evolution of the NMR resonances during the lithiation from 0.8 to 2 ppm and to 4.8 ppm at the end of discharge. Upon delithiation, the peaks shift to lower frequencies and only the SEI signal is observed. For the second discharge, the region of the plateau in the galvanostatic curve was explored. Before, during and after the plateau, the peak at 0.8 ppm shifts to slightly higher frequencies, suggesting a surface related processes. As the potential decreases, this peak shifts until reaching 4.8 ppm. These results are slightly different from the reported NMR for the lithiation of bulk Si (Chapter II). Here, the silicon network is broken into small clusters until forming isolated Si 4-, the frequency of the NMR shift decreases with the size of the Si cluster. For the siloxene, the resonances at low frequencies are formed first and later shift to higher values, no peak at high frequency (<8 ppm) was found for siloxene. Then, possibly a surface reaction (2 ppm) occurs before the formation of the LixSi phases (4.8 ppm), this statement is supported by the capacitive contribution identified in the kinetic study (CV). Additional information was obtained by 1 H NMR, the total amount of protons in overall increases with the cycling, probably due to the formation of the SEI layer. Although, the protons of siloxene tend to decrease during the lithiation, in line with the loss of Si-H and Si-OH bonds observed by Raman and IR spectroscopy. The region of the plateau in the galvanostatic curve (0.43 V) showed a decrease in the protons, followed by an increase in the protons with high mobility, suggesting a voltage region in which the protons are more stable. Therefore, this plateau might be related with Li surface reactions.

Even though the gathered information provides hints about the phases that could be formed for the lithiation of siloxene, it is not possible to attribute all the phenomena to a conventional alloying mechanism. Indeed, the features in the galvanostatic and derivative curve for siloxene are slightly different from the one of bulk Si. This last one, along with the preservation of the layered morphology after cycling, suggest that a different electrochemical mechanism or at least different lithiated intermediates are involved during the lithiation/delithiation. These intermediates present an amorphous nature, given the loss of short-range order for the siloxene and germanane upon cycling, and possibly have a similar composition to Li15Si4. The Raman spectroscopy indicates the presence of a compound type LiSi while the 7 Li NMR the formation of a first intermediate related to a surface reaction, possibly Li close at the edges of the Si-Si layers, followed by the formation of another LixSi, with a similar environment to Li15Si4. An additional Li and proton surface phenomena is found at the plateau at 0.43 V during the second discharge. The fact that siloxene could form different lithiated intermediates compared with bulk Si is related to the sensitivity of Si to the cycling conditions. As previously observed, the structure plays an important role in the cyclability and possibly in the nature of the intermediates involved. This behavior is even more pronounced for Ge-based compounds. The nature of the ligands of siloxene (-OH and -H) and germanane (-H) influences the SEI decomposition products, which as reviewed in Chapter I play an important role in the reversibility, capacity retention and CE of the silicon and germanium-based anodes. Moreover, the layered morphology serves a buffer for the volume variations during lithiation/delithiation.

On the other hand, the intercalation of Na/Li into a layered Si-based material has only been theoretically predicted for a single layer of siloxene (silicene), with no experimental record. The calculations foresee a high coverage of the silicene with alkali ions like Na, Li and K due to the nature of their interactions and the most stable site in the hollow of the Si6 ring. The full sodiated/lithiated state of silicene corresponds to X1Si1 (X=Li/Na), the predicted binding energies and diffusion barriers indicate that their intercalation is achievable without the kinetic limitations (higher diffusion coefficient for silicene compared to bulk Si), structure degradation and volume expansion of bulk Si. [130,292,[294][295][296] Recent studies have also predicted the intercalation of Li and Na in germanene with good diffusion values and no structural degradation. [161,162] These calculations are very promising, but they are performed for mono-or by-layer silicene and germanane. The study of the multi-layered material and the impact of the -OH and -H ligands in the structure, over the possible intercalation must be studied. Therefore, in order to better understand the experimental findings for the lithiation/delithiation mechanism of siloxene and germanane, and properly link the observed phenomena with an alloying or intercalation mechanism, theoretical calculations must be performed.

Conclusion

In summary the lithiation/delithiation mechanism of siloxene and germanane was explored by a combination of characterization techniques such as XRD, Raman and Infrared Spectroscopy, TEM, EELS, XPS and NMR. The results show the absence of the characteristic c-Li15Si4 (Li15Ge4/Li17Ge4) formed at the end of discharge for a conventional alloying of bulk Si (Ge) with Li and the preservation of the layered morphology in siloxene (germanane). As observed in the galvanostatic curve the siloxene and germanane produced a significant irreversible capacity due to the formation of a SEI layer. The components of such decomposition layer were identified as Li2O, LiF, Li-carbonates and Li-alkoxides for siloxene while only Li2O for germanane. Thus, the nature of the SEI layer might be related with the ligands in the siloxene (-H and -OH) and germanane (-H) and is probably one of the reasons behind the improved performance compared to bulk Si or Ge. The study of the intermediated compounds indicates the initial formation of a surface lithiated compound followed by the LixSi phases with similar composition to LiSi/Li15Si4 but not necessarily these two last ones; although the exact composition could not be determined. The different the lithiated intermediates for siloxene and germanane reflects the sensitivity of the Si-and Ge-based materials towards the cycling conditions and certainly the layered morphology plays an important role in buffering the associated volume changes. Based on the obtained experimental data it was not possible to assign the electrochemical mechanism of siloxene and germanane to an alloying or intercalation. Studies in the literature predict the intercalation of alkali ions in silicene and germanane over an alloying reaction, though the calculations are performed for mon-or by-layers. Thus, the behavior of the alkali insertion in bulk siloxene and germanane and its interaction with the -OH and -H ligands remains unknown. In this sense, the theoretical calculations would assist the interpretation of the lithiation/delithiation of siloxene and germanane. A deep knowledge of their mechanism is necessary to comprehend the advantages of such materials as anodes and address the possible limitations in order to extract their maximum potential.

General Conclusion

Silicon and germanium are attractive candidates for different battery technologies (LIB, NIB and KIB) due to their high theoretical capacities 5 . Chapter I provided an overview of the advantages and challenges of such electrode materials. For instance, Ge has a superior electronic conductivity, Li-ion diffusion and capacity retention than Si, but its practical applications are hindered by its relative high cost. Si presents a high theoretical capacity, it is abundant on the earth crust and relatively cheap, but it has a poor capacity retention. Both Si and Ge suffer from volume variations due to the alloying reaction with the alkali ion. The constant expansion/contraction upon alkali insertion/deinsertion leads to particle cracking, detachment from the current collector and constant exposure of new surface to the electrolyte causing an uncontrolled formation of SEI layer. These combined processes eventually lead to a capacity fading. Various strategies are proposed in the literature to address these challenges, including specific electrode architectures that buffer the volume variations, the effective electrode formulation that stabilizes the SEI layer or the use of alloys and composite materials. This thesis particularly focused on the use of alloys type Si1-xGex and layered structures derived from Zintl phases. These last ones are characterized by a vast structural diversity and their properties can be easily modified. Several lithiated silicides and germanides belong to the Zintl family, hence their study is crucial for the understanding of the phenomena taking place in a battery.

Addressing the challenges of Si-and Ge-based electrodes does not only includes the aforementioned, but also the comprehension of the processes taking place during the alloying reaction. Chapter II presented the investigation of the Si, Ge and Si1-xGex electrochemical lithiation mechanism. The sodiation and potassiation are not considered, given the few information and limited studies available. The mechanism of Si is highlighted based on the literature review, following the Ge and Si0.5Ge0.5 mechanism were studied by a combination of Operando spectroscopy and diffraction techniques. The Si0.5Ge0.5 was selected based on previous studies. [73] The alloying reaction of Si, Ge and Si0.5Ge0.5 with Li induces amorphization of the crystalline network. Various amorphous lithiated intermediates are formed, their basic Si (Ge) building units are related with the Li content. Initially, big Si (Ge) clusters are built (M5 rings, M4 stars like in Li12M7, M=Si, Ge), subsequently small clusters (dumbbells Li7M3, Li9M4), dumbbells and isolated Si 4-(Li13M4). At the end of discharge, for voltages <100 mV, the most lithiated phase crystallizes, Li15M4 composed of isolated Si 4-ions. Additionally, a process of overlithiation of Li15M4 into Li15+δM4 was identified and assigned to the extra charges storage in the electron deficient Li15M4 phase. Upon delithiation an amorphous product was obtained. For Ge some of the lithiated phases were identified (LiGe, Li7Ge3) and found to be dependent on the cycling conditions. Indeed, at rates between C/5-C/8 Li15Ge4 was formed while at C/14 Li17Ge4. For the Si0.5Ge0.5 there is no segregation into Si and Ge during the cycling, the synergy between both elements improves the electrochemical performance in terms of capacity retention and reduction of volume variation.

Following the impact of the morphology over the electrochemical performance was evaluated in Chapter III. The layered Si and Ge derived from the CaSi2 and CaGe2 Zintl phases were tested as anodes for LIB, NIB and KIB. These phases present the characteristic Si 4-(Ge 4-) Zintl ion that forms planes of interconnected Si6 (Ge6) rings. Hence, in the CaSi2 and CaGe2 the Ca 2+ can be topotactically deintercalated and replaced by -OH and -H to form siloxene, and by -H to form germanane (preferential bonding of Ge with -H). The obtained reversible capacity for siloxene were 2300, 311 and 296 mAh/g for Li, Na and K, respectively, while for Germanane, 2150, 495 and 205 mAh/g for Li, Na and K, although all the systems presented a considerable irreversible capacity. The lithiation capacities and capacity retention are higher compared to those ones of bulk Si or Ge, even at high rates (>1200 mAh/g at C and 2C). The sodiation capacities are significant, particularly for germanane. Note that the reaction between Si and Na is partially limited, especially for the crystalline matrix. The potassiation reaction is limited for siloxene and germanane; in this system the electrolyte plays a key role in the cell reversibility and stability towards the K 0 counter electrode. Then, a full cell configuration and an optimal electrolyte formulation are necessary to properly evaluate the potassiation properties of siloxene and germanane. On the other hand, following the propositions in the literature for the modification of the Zintl phases, the interlayer space of siloxene was expanded by the intercalation of alcohols with various carbon chain length. The relationship between the interlayer distance and the electrochemical behavior was evaluated, foreseeing an enhanced alkali insertion owed to the increase in the interlayer spacing. Unfortunately, this was not the case and the presence of the organic molecule introduces a steric effect that impedes the access of the alkali ion to the bulk. Consequently, an optimal method for the removal of the organic molecule without degrading the siloxene structure must be found. Additionally, given the miscibility between Si and Ge, the layered Si1-xGex counterpart, the siligane, was prepared from the Ca(Si1-xGex)2 Zintl phase. The products exhibited a complex structure and their exact composition could not be determined. The electrochemical performance of the siliganes did not present an improvement compared to siloxene and germanane, nonetheless the theoretical predictions are very promising. (532, 1064 and 532 mAh/g for Li, Na and K, respectively [274] ).

The reports in the literature concerning these phases are very scarce, thus a deep structural characterization is necessary before addressing their usage as battery materials. This information is valuable in order to determine their specific composition and understand their electrochemical behavior.

Finally, the electrochemical mechanism for the siloxene and germanane was explored. Most of the experiments were performed on the lithiation of siloxene, for convenience of comparison with other reports in the literature. The results suggest that the electrochemical cycling of siloxene is somewhat different from the one of bulk-Si. The c-Li15Si4 characteristic of the alloying mechanism is absent in siloxene. Likewise, for germanane no c-Li15Ge4 or c-Li17Ge4 were found. The microscopy analysis indicates that both siloxene and germanane preserve their layered morphology upon cycling and form a significant decomposition layer. This last one is related with the irreversible capacity and is composed of LiF, Li2O and different Li-carbonates for siloxene and only Li2O for germanane, meaning that the nature of the ligands plays an important role in the decomposition products. Since the formation of an optimal SEI layer is mandatory to ensure a good cyclability, the study of the electrolyte formulation and its relationship with the siloxene and germanane ligands on the electrolyte decomposition products must be performed. The investigation of the lithiated intermediates was performed by a combination of ex situ spectroscopy and in situ diffraction techniques. After cycling a possible change in the interlayer distance associated with an interaction between the siloxene/germanane layers with the alkali ion was identified by Raman spectroscopy. The NMR results indicate the formation of a lithiated intermediate at the layers surface, followed by a LixSi phase, with a similar composition to Li15Si4. The exact composition of the LixSi was not determined, since various Li environments were possible. The presence of different lithiated intermediates puts in evidence the sensitivity of the Siand Ge-based materials to the cycling conditions, and probably the layered morphology has an impact on this behavior. Indeed, the theoretical studies suggest that due to the presence of Si-Si planes an intercalation of Li is favored over an alloying reaction. Nevertheless, this study did not provide sufficient proofs to prove this theory neither all the results can be totally assigned to an alloying mechanism. Thus, this study sets the basis of future researches, hopping to understand how differently silicon and germanium behave as a bulk and as a layered material, and unveil the impact of the morphology over the electrochemical mechanism.

In conclusion, this work presented two strategies to address the challenges associated with the Siand Ge-based electrodes, the use of alloys that combine the advantages of two materials, like in Si0.5Ge0.5 and the use of a layered morphology that buffers the volume variations and shortens the diffusion pathways, like in siloxene and germanane. Both strategies presented an improvement in the electrochemical performance and this approach can be translated into other elements, like Sn, Sb, etc. This work highlights as well the importance of the study of the processes taking place during cycling.

The understanding of these phenomena is crucial to determine the reasons behind an improved performance or failure mechanisms. This information is essential to extract the best properties of a given material. Finally, this research does not only serve for battery applications but also provides significant information for the implementation of these materials in other areas.

As perspectives, the Zintl family present other layered phases based Sn, Sb, Bi, As, P (Stanene, Antimonene, Bismuthene, Arsenene, Phosphorene), with the same X6 rings interconnected to form planes as in siloxene and germanane. [START_REF] Matsuda | Basics and Families of Monatomic Layers[END_REF]297] Recent theoretical studies indicate their potential applications as battery electrodes. [292,[298][299][300] Other layered Si and Ge phases can be prepared from different Zintl phases like LiSi, LiGe, Li12Si7, Li7Ge12, LiAlSi, LiBaSi, [START_REF] Nesper | [END_REF] while the layered allo-Si and allo-Ge can be obtained from the Li3NaSi6 [301] and Li7Ge3. [START_REF] Zeilinger | Synthesis , Characterization and Reactivity of Lithium-Containing Silicides , Germanides and Borosilicides[END_REF] All these phases are candidates for battery applications and their study is of great interest due to their layered nature, which as indicated in this thesis, happened to be beneficial for the electrochemical performance. The versatility and structural variety of the Zintl phases is enormous and opens the possibility for the investigation of new electrode materials.

Annex I Experimental details

Raman Spectroscopy.

In situ Measurements. For the in situ analysis a Si0.5Ge0.5 self-supported film was placed in a Raman operando cell with quartz window (Figure S1). [170] The half-cell was assembled in a glove box filled with argon, with a lithium foil as a negative electrode and a glass fiber separator (Whatman TM filter paper) wet with ~60 µL of LP30 electrolyte (1 M LiPF6 in EC: DMC (1:1)). A drilled copper grid was used as the current collector to enable the laser light to reach the Si0.5Ge0.5 electrode. The focal plane was adjusted at the surface of the Si0.5Ge0.5 electrode, slightly separated from the quartz window of the cell by the Cu foil. The pinhole and the optical slits were adjusted to probe a large volume of the electrode, increasing the laser cross section. This, unfortunately, led to the presence of a small signal (peak at 490 cm -1 ) of the quartz window in the spectra.

In order to prevent the electrode from heating, the laser power was appropriately limited (power measured on the sample < 10mW). Additionally, some Raman spectra present a high fluorescence effects, thus, to compare them at different stages of lithiation the baseline has been subtracted.

X-ray Absorption Spectroscopy

For the XAS measurements, the intensity of the monochromatic X-ray beam was measured by three consecutive ionisation detectors. The in situ electrochemical cell [172] was placed between the first and the second ionisation chambers, and the homogeneity of the sample was checked before running the electrochemical test. The latter was carried on continuously for the first lithiation. Successive spectra were collected at a rate of 2 Hz and averaged out over periods of 20 minutes.

PCA and MCR-ALS analysis details.

PCA is a factor analysis, i.e., a multivariate technique for reducing matrices of data to their lowest dimensionality by the use of an orthogonal factor space: in such methods, the analysed matrix corresponds to the experimental spectra, and each spectrum is considered as an n-dimensions vector with n corresponding to the number of points within each XAS spectrum. The first principal component has the largest initial variance, i.e., alone it takes into account as much as possible the whole set of data. The following principal components are built orthogonal to the first one and allow the reproduction of the experimental spectra with the highest possible variance via adapted linear combinations. This supposes that the group of analysed spectra is intrinsically bilinear, which means that all spectra can be expressed as linear combinations of an orthogonal basis set of uncorrelated spectra. This basis set has a lower dimensionality than the number of original spectra and is obtained from the calculation. In the different linear combinations that are built to reconstruct each experimental spectrum, the multiplication factors of the principal components are usually called scores.

Even though the so-obtained principal components are only orthogonal mathematical functions and not real XAS spectra, their determination reveals the number of independent spectral components that can be used to describe the whole set of experimental data. Several methodologies are then available to reconstruct the corresponding real XAS components as well as their evolution such as the MCR-ALS (Multivariate Curve Resolution -Alternating Least Squares) analysis tool. [302,303] A detailed description of this method from a theoretical point of view is given by Tauler et al. [175,304] who were the first to propose it for the analysis of in situ spectroscopic data. The intrinsic limits of this method and of its application are discussed in ref [START_REF] Ruckebusch | [END_REF] .

In the present study, the MCR-ALS analysis was performed with the following constraints: (i) nonnegativity of the intensity of the components, (ii) unimodality for all components and (iii) closure (sum of the components always equal to 100 of the intensity).

Annex II Supplementary information Chapter II. ont été étudiés par une combinaison de techniques de spectroscopie Operando et de diffraction. Le Si0.5Ge0.5 a été sélectionné sur la base d'études antérieures. [73] La réaction d'alliage du Si, du Ge et de Si0.5Ge0.5 avec le Li induit une amorphisation du réseau cristallin. Différents intermédiaires lithiés amorphes sont formés, leurs unités de construction à base de Si (Ge) sont liées à la teneur en Li.

Alloying Electrochemical Mechanism

Résumé de la thèse

Initialement, de grands groupes de Si (Ge) sont construits (anneaux M5, étoiles M4 comme dans Li12M7, M = Si, Ge), puis de petits groupes (haltères Li7M3, Li9M4), des haltères et du Si 4-(Li13M4) isolé. En fin de décharge, pour des tensions <100 mV, la phase la plus lithiée cristallise, Li15M4, elle est composée d'ions Si 4-isolés. De plus, un processus de recouvrement de Li15M4 en Li15+δM4 a été identifié et attribué au stockage de charges supplémentaires dans la phase de Li15M4 déficiente en électrons (Figure 1). Lors À titre de perspectives, la famille Zintl présente d'autres phases lamellaires à base de Sn, Sb, Bi, As, P (stanène, antimonène, bismuthène, arsénène, phosphorène), avec les mêmes anneaux X6 interconnectés pour former des plans comme dans le siloxène et le germanane. [START_REF] Matsuda | Basics and Families of Monatomic Layers[END_REF]297] Des études théoriques récentes indiquent leurs applications potentielles comme électrodes de batterie. [292,[298][299][300] D'autres phases de Si et de Ge lamellaire peuvent être préparées à partir de différentes phases de Zintl telles que LiSi, LiGe, Li12Si7, Li7Ge12, LiAlSi, LiBaSi, [START_REF] Nesper | [END_REF] les allo-Si et les allo-Ge lamellaires peuvent être obtenus à partir de Li3NaSi6 [301] et Li7Ge3. [START_REF] Zeilinger | Synthesis , Characterization and Reactivity of Lithium-Containing Silicides , Germanides and Borosilicides[END_REF] Toutes ces phases sont des candidates aux applications pour les batteries et leur étude présente un grand intérêt en raison de leur nature lamellaire qui, comme indiqué dans cette thèse, s'est avérée bénéfique pour les performances électrochimiques. La polyvalence et la variété structurelle des phases de Zintl sont énormes et offrent la possibilité d'étudier de nouveaux matériaux pour les électrodes.

Résumé

De nos jours, les batteries jouent un rôle clé dans presque toutes les technologies qui entourent le genre humain. Afin de répondre à la demande croissante, la conception d'appareils plus efficaces avec une densité d'énergie et une durée de vie plus élevées est cruciale. Dans ce contexte, le silicium et le germanium apparaissent comme des candidats prometteurs pour les matériaux d'électrodes en raison de leurs capacités théoriques élevées. Bien avant une mise en oeuvre de ces matériaux au niveau industriel, plusieurs défis doivent être relevés. Les capacités élevées délivrées se font au détriment d'une expansion volumique lors de l'insertion des ions lithium par exemple. Ces changements de volume dans les particules de Si et de Ge entraînent la pulvérisation des particules, le détachement du collecteur de courant, la formation excessive et incontrôlée de la couche de SEI et une chute de la capacité. Différentes stratégies ont été rapportées dans la littérature pour surmonter les défis susmentionnés. Dans ce travail, deux approches ont été considérées, d'une part l'étude des alliages Si1-xGex et d'autre part l'étude de composés lamellaires. Dans le premier cas, la formation de la solution solide Si1-xGex améliore la rétention de capacité et la conductivité électronique. Dans le second, les matériaux lamellaires Siloxene et germanane, dérivés des phases de Zintl CaSi2 et CaGe2, amortissent les changements de volume et améliorent la cinétique du système. Une étude fondamentale des mécanismes électrochimiques a été réalisée pour comprendre les processus mis en jeu dans ces deux approches.

Mots-clés : batterie Li-ion, Na-ion, K-ion, anode, silicium, germanium, siloxène, germanane, silicium et germanium lamellaire.

Abstract

Nowadays, the batteries play a key role in almost all of the technologies that surround humankind. In order to satisfy the increasing demand, the design of more efficient devices with higher energy density and cycle life is crucial. In this context, silicon and germanium appear as promising candidates for electrode materials due to their high theoretical capacities. Although, prior to the implementation of these materials at an industrial level, several challenges must be addressed. The high delivered capacities come at the expense of a volume expansion and contraction upon alkali insertion and deinsertion. These volume changes in the Si and Ge particles, lead to particle pulverization, detachment from the current collector, excessive and uncontrolled formation of SEI layer and eventual capacity fade. Different strategies have been reported in the literature to overcome the aforementioned challenges. In this work, two approaches are considered, the study of the Si1-xGex alloys and the use of a layered morphology. In the first one, the formation of the Si1-xGex solid solution improves the capacity retention and the electronic conductivity. In the second one, the layered Siloxene and germanane, derived from the CaSi2 and CaGe2 Zintl phases buffers the volume changes and improves the kinetics of the system. On the other hand, the fundamental study of their electrochemical mechanism is crucial to understand the reasons behind an improvement and a failure. Thus, in this work we have studied the electrochemical lithiation mechanism of the Ge and the Si0.5Ge0.5 alloy, certain lithiated phases were identified and the impact of the cycling conditions (morphology, C/rate, electrode formulation) was assessed. Additionally, the electrochemical mechanism of the layered compounds was explored to identify the intermediates and establish the differences with their bulk counterparts. Keywords: Lithium-, sodium-and potassium-ion battery, anode, silicon, germanium, siloxene, germanane, layered silicon and germanium.

  GENERAL INTRODUCTION ................................................................................................... 1 CHAPTER I. STATE OF THE ART ......................................................................................... 5

Figure I. 2 .

 2 Figure I.2. Solutions for fundamental problem of Si(Ge) anodes. Volume expansion and contraction of a) Core/shell, b) Yolk shell and the pomegranate assembly and c) Nanowires during lithiation and delithiation.The continued cycling leads to an uncontrolled growth of SEI layer in the core/shell and nanowires while the yolk/shell-type structures can stabilize a uniform SEI layer. Adapted from ref.[3,68]

Figure I. 7 .

 7 Figure I.7. Structural representations of a) Si, b) LiSi, c) Li12Si7, d) Li7Si3, e) Li13Si4, f) Li15Si4 and g) Li21Si5.

Figure I. 8 .

 8 Figure I.8. Structural representations of NaSi (left) and NaGe (right).

Figure I. 10 .

 10 Figure I.10. Structural representations of a) CaSi2, b) Siloxene (Si6H3(OH)3) and c) Silicene.

  originates three modifications for the siloxene structure (Figure I.11); i) the Wöhler siloxene, a quasitwo-dimensional structure similar to the bulk c-Si (111) planes, except that each Si atom is bonded to other three Si and one OH or H. This structure can lower its energy by oxygen insertion in the planes to form ii) the Kautsky siloxene, in which the Si6 rings are interconnected by oxygen bridges and iii)

Figure I. 13 .

 13 Figure I.13. Structural model of a) CaGe2, b) Germanane (GeH)n and c) Germanene.

Figure II. 3 )

 3 , the reaction of c-Si with Li place below 110 mV, three resonances at 22-20 (very broad at 105 mV), 18-16 and 6 ppm are observed, corresponding to Li ions in a Si environment of small clusters, dumbbells and isolated Si ions respectively. The signature of the isolated Si atoms dominates over the clusters one, due to the preferential breaking of the clusters, as

  The results presented in this section have been adapted from the publication: Loaiza, L. C.; Louvain, N.; Fraisse, B.; Boulaoued, A.; Iadecola, A.; Johansson, P.; Stievano, L.; Seznec, V.; Monconduit, L. Electrochemical Lithiation of Ge: New Insights by Operando Spectroscopy and Diffraction. J. Phys. Chem. C, 2018, 122, 3709-3718. DOI 10.1021/acs.jpcc.7b11249

Figure II. 5 .

 5 Figure II.5. SEM Pictures for a) Pristine Ball milled Ge and b) Pristine self-supported Ge film used for the operando measurement In order to observe in depth the cycling processes, a series of operando XRD experiments were performed at the rates of C/5 and C/14. The obtained electrochemical signatures were fully consistent with the coin cell experiments above.

II. 6 .

 6 Figure II.6. Operando XRD and associated galvanostatic discharge-charge curves at a) C/5 and c) C/14 cycling rate of the Ge electrodes. b), d) magnifications of the [18-28°] and [38-44°] 2θ ranges.

Figure II. 7 .

 7 Figure II.7. XRD pattern corresponding to the product at the end of discharge of the Operando XRD performed at C/14 on a Ge self-supported electrode.

  at low voltages. c-Li15Ge4 and Li15+δGe4 are in fact structurally related and has until now only been differentiated by solid state NMR, with the defects in the lattice playing an important role in accommodating the extra charges. The existence of a solubility domain affects also the hypothetic Li7Ge3 (or Li7Ge2) (C/5) and Li17Ge4 (C/14). The different hypothesis brought by XRD were further studied by operando confocal Raman Spectroscopy, performed at C/5 rate to gain complementary information about the phases formed at the end of the discharge and charge processes. The evolution of the Raman spectra during the first lithiation/delithiation cycle is shown in Figure II.8, together with the galvanostatic curve. The mode of c-Ge at 298 cm -1 is clearly visible for the pristine electrode measured under open circuit voltage (OCV)

Figure II. 11 .

 11 Clearly, components 1 and 4 represent the pristine c-Ge and the final product of the first discharge. Components 2 and 3 show maxima after 30 and 65 measured spectra from the beginning of the discharge, respectively, corresponding to the reaction of about 1 and 2 mole of Li per mole of Ge.

Figure II. 10 .

 10 Figure II.10. Electrochemical curve for the operando XAS experiment of a Ge electrode performed at a C/8 rate. b) PCA analysis of the components. c) Evolution of the component scores in the PCA analysis.

Figure II. 11 .

 11 Figure II.11. Evolution of the concentration of the MCR-ALS components during the lithiation of Ge.Comp1= c-Ge, comp2= ∼1Li/Ge, comp3= ∼2Li/Ge, comp4= Final lithiated phase.

Figure II. 12 .

 12 Figure II.12. Fit of the EXAFS portion of the MCR-ALS components in (a) k and (b) R space. The black line corresponds to the data and the red one to the fit.

  Figure II.13 shows the XRD pattern for Si0.5Ge0.5. Broad peaks are observed, due to the crystallite size reduction and loss of crystallinity because of ball milling. The time required to obtain the Si0.5Ge0.5 phase was monitored (Figure II.13-a). After 8, 20 or 37 hours of ball milling, the (111) peak from the Si0.5Ge0.5 phase has appeared, but contributions from unreacted c-Ge and c-Si are still present; only after 40 hours the desired phase was pure. Long milling times induce agglomeration of the products on the walls of the jar, hindering a complete reaction, thus a two-step process (two treatments of 20 h) was adopted. The SEM micrographs of Si0.5Ge0.5 (Figure II.13-c) show particles with sub-micrometer size and formation of agglomerates (<6 μm) and the EDX mapping indicates that Si and Ge are well dispersed in the powder. The refined cell parameters were calculated for pure Si (a=5.43(3) Å), Si0.5Ge0.5 (a=5.54(8) Å) and pure Ge (a=5.64(7) Å) and are in agreement with those reported in ICSD database (a=5.4309 Å and a=5.6576 Å, ICSD 000271402 and 000040545, for Si and Ge, respectively). The values follow a linear increase with increasing Ge content which is in agreement with Vegard's Law (Figure II.13-b), thus a solid solution with composition 50-50 was successfully formed. In parallel, lithiated phases were mechanically synthesized to serve as references (Figure II.13-d). For Li15Si4 (a= 10.6546 Å, ICSD 237945) and Li15Ge4 (a= 10.72(5) Å, ICSD 43235) the refined cell parameters were 10.65(8) Å and 10.768(1) Å, respectively. The Li15(Si0.5Ge0.5)4 phase was obtained either by mixing stoichiometric amounts of Li and previously synthesized Si0.5Ge0.5 (sample A) or stoichiometric amounts of Li, Si and Ge (sample B). The obtained main phase presents a refined cell parameter a= 10.753(1) Å that is in between those of Li15Si4 and Li15Ge4. Then, these last ones follow the Vegard's law and lead to the formation of Li15(Si0.5Ge0.5)4. For sample A an additional non-identified

Figure II. 13

 13 Figure II.13. a) Si0.5Ge0.5 XRD patterns as a function of the milling time, zoom over the (111) peak. b) The refined cell Si1-xGex parameters reported as a function of the Ge content. c) EDX mapping of the Si0.5Ge0.5 phase. d) XRD patterns for Li15Si4, Li15Ge4 and Li15(Si0.5Ge0.5)4.

Figure II. 14 .

 14 Figure II.14. Raman Spectra of Si, Si0.5Ge0.5 and Ge (top) and of Li15Si4, Li15(Si0.5Ge0.5)4 and Li15Ge4 (bottom).

  For the Ge K-edge XAS spectrum of pristine Si0.5Ge0.5 shown in Figure II.15 the results of the fit (Table Annex II-1) indicate slightly different Ge-Ge and Ge-Si bond distances in the nearest neighbor shell and an average enrichment in Ge neighbors (2.5 rather than the 2 expected from the nominal stoichiometry). Bond distances and numbers of neighbors are more similar for Ge-Ge and Ge-Si second neighbors, and in line with the expected stoichiometry. The enrichment of Ge neighbors at the shortest atomic scale suggests the presence of small sub-nanometer sized clusters of Ge in pristine Si0.5Ge0.5, while at the middle range level a globally homogenous composition is observed, in agreement with the XRD results.

Figure II. 15

 15 Figure II.15. a) EXAFS spectrum of pristine Si0.5Ge0.5 (top) and b) relative Fourier transform. The experimental spectrum is drawn in blue, and the fitting result in red. The spectrum was fitted by imposing a total number of nearest and second neighbors equal to those expected from the crystal structure determined by XRD (4 and 6, respectively), and by leaving variable their nature (Si or Ge) and their bond distance.

Figure II.

  Figure II.16 shows the 7 Li MAS NMR experiments were performed on mechanically synthesized

Figure

  Figure II.16.7 Li MAS NMR spectra for the mechanically synthesized references Li15Si4 (top) and Li15Ge4 (bottom) Figure II.17 shows the7 Li MAS NMR spectra of the Li15(Si0.5Ge0.5)4 references samples A (Li+ Si0.5Ge0.5) and B (Li+Si+Ge). Note that the samples were dispersed in dry KBr to manage spinning. The NMR data indicate that sample A is not pure, as hinted by the XRD findings (Figure II.13-d). Two main contributions at 35 ppm (54%) and 0 ppm (30%) are observed. Minor contributions are at 9 ppm, 3 ppm and -9 ppm. Several fits are possible for these MAS spectra, therefore the areas should be considered with care. The shifts however are always within 1 ppm. Sample B, has peaks at -2 ppm (66%) and -10 ppm (27%), with minor contributions at 9 and 3 ppm. Since Li15(Ge0.5Si0.5)4 was detected by XRD in both samples, we assign the only common peak (at -9/-10 ppm) to c-Li15(Si0.5Ge0.5)4. The contribution at -2 ppm in sample B must be amorphous, as it was not detected by XRD. It could be associated with amorphous lithiated phases issued from the decomposition of the crystalline ternary phase (aging of the sample).

Figure II. 17

 17 Figure II.17.7 Li MAS NMR spectra for samples A (Li+Si0.5Ge0.5), B (Li+Si+Ge), C (Si0.5Ge0.5 discharged and allowed to rest at OCV) and D (Si0.5Ge0.5 discharged and held at 0 V). The experimental spectra (blue) are overlaid with the fit (dotted red). Individual components of the fit are shown below each spectrum.

Figure II. 18 .

 18 Figure II.18. Galvanostatic curves and derivative curves for a) and d) Silicon, b) and e) Si0.5Ge0.5 and d) and f) Ge cycled vs Li/Li + at the rate of C/14.

Figure II. 19

 19 Figure II.19 shows the in situ XRD measurement for Si0.5Ge0.5 cycled at a rate of C/14. At the beginning of discharge, the XRD is composed of peaks from the Si0.5Ge0.5 phase. There is a shift of the Si0.5Ge0.5 (111) peak towards higher angles, in agreement with a decrease in the cell parameter from 5.539(5) to 5.507(4) Å, before the phase amorphization. The material remains amorphous until voltages <100 mV where diffraction peaks coming from a crystalline Li15X4 phase can be observed. Just before the crystallization of this phase (Figure II.20), huge humps appear at ca. 24 and 42°, indicating

Figure II. 19 .

 19 Figure II.19. in situ XRD for a Si0.5Ge0.5 self-supported film cycled vs Li/Li + at the rate of C/14.. Additionally, during the course of the discharge the Li15(Si0.5Ge0.5)4 (112), (022), (013) and (332), (422), (431) peaks shift to lower angles (Figure II.20, Figure II.22). Such shifts correspond to an increase of the crystallographic refined cell parameters from 10.764(5) Å (xLi=3.42) to 10.772(5) Å (xLi=4.07) (Figure II.22), in contrast with the results of Duveau et al.[73] . Just before the end of discharge the

  to 10.695(8) Å. The same evolution is observed during the second discharge (Figure II.20) and at full lithiation, the refined cell parameter for Li15(Si0.5Ge0.5)4 (xLi=4.256) is 10.725(2) Å (Table Annex II-3).

Figure II. 20 .

 20 Figure II.20. Zoomed area of the in situ XRD of a self-supported Si0.5Ge0.5 electrode performed at C/14.

Figure II. 21 .

 21 Figure II.21. EDX analysis for the Si0.5Ge0.5 self-supported electrode at the end of discharge.

Figure

  Figure II.22. a) Evolution of the refined cell parameter as a function of the xLi content for Li15(Si0.5Ge0.5)4. at the end of discharge/beginning of charge

Figure II. 23 .

 23 Figure II.23. In situ XRD experiment for a self-suported Si0.5Ge0.5 electrode performed at a C/5 rate.

Figure II. 25 .

 25 Figure II.25. Extracted slices in the first discharge and first potentiostatic step at 0 V of the in situ 7 Li NMR of a Si0.5Ge0.5 self-supported film cycled vs Li/Li + at the rate of C/5. At the end of the first discharge, the cell is left in a potentiostatic step at 0 V to observe the kinetically-limited transformations (Figure II.25). Exposing the samples to low voltages enhances the formation of the overlithiated Li15+δ(Si0.5Ge0.5)4, as evidenced by the increase in intensity of the peak at -17 ppm. Subsequently, the cell is left in open-circuit (Figure II.26-left). During this time, the peak at

  Figure II.27-rigth), the peak at -17 ppm totally disappears. Only a small contribution in the 0 ppm region is observed, the Li15(Si0.5Ge0.5)4 phase might convert into a phase not detected by NMR (conductive or with slower relaxation). No significant evolution of the spectrum is observed apart from the 250 ppm region, where a shoulder emerges next to the peak for metallic lithium that indicates deposition of dendritic or mossy Li. The second discharge (Figure II.27-left) resembles the first one, with the main difference of an incomplete conversion of the 10 ppm peak into the -17 ppm. Overlithiation happens during the potentiostatic step at 0 V (Figure II.27-rigth) and contrary to the first discharge, it is still ongoing at the end of the potentiostatic step (the intensity of the -17 ppm peak still increases slightly) and the peak

Figure II. 28

 28 Figure II.28. a) Cartographic view of the EXAFS data (k 2 X(k)). b) Variance plot obtained by the PCA analysis of the EXAFS data. c) MCR-ALS reconstructed spectral components. d) Evolution of the intensity of the MCR-ALS reconstructed spectral components with advancing lithiation. Additional information on the lithiation mechanism of Si0.5Ge0.5 can be obtained by analyzing the XAS measurements at the Ge K-edge. The evolution of the EXAFS signal shown in Figure II.28-a. The results of PCA analysis (variance plot, Figure II.28-b) indicate that at least four components are necessary to express 99.9% of the variance within the whole set of EXAFS spectra with the residual corresponding to experimental noise. The application of the MCR-ALS analysis resulted in the reconstruction of the four components (Figure II.28-c), which were then treated as common EXAFS spectra. Figure II.28-d shows the evolution of the individual components during the lithiation of Si0.5Ge0.5 (see the fitting results, Figure Annex II-1-3). The first component is virtually identical to pristine Si0.5Ge0.5, with a Ge enrichment in the first coordination shell. The second component gradually replaces the first one during the beginning of the discharge, whereas both components are relatively similar, a decrease of the total amplitude is detected and attributed to the formation of Ge-Li bonds.

  5 Å -1 to remove the oscillations at higher values of k, which are attributed to the residuals of the glitch peaks in the initial EXAFS spectra and were filtered out before the PCA analysis. The lack of second neighbors indicates a relatively disordered character of the final lithiated species. No specific formation of the overlithiated phase Li15+δM4 is observed. The relatively large uncertainty in the fitting parameters of the Ge-Li shell, indeed, makes it very difficult to detect small variations in the number of Li neighbors.

Figure II. 29 .

 29 Figure II.29. Evolution of the normalized amount of Ge-Ge, Ge-Si and Ge-Li bonds with advancing lithiation. The value 1 corresponds to the number of Ge-Ge and Ge-Si bonds in pristine Si0.5Ge0.5, or to the theoretical number of Ge-Li bonds in Li15(Si0.5Ge0.5).

  Figure II.29 shows the process of lithiation of Si0.5Ge0.5 with normalized Ge-Ge, Ge-Si and Ge-Li bonds calculated from the evolution of the intensity of the components and the number of different nearest neighbors. The Ge-Ge bonds disappear more rapidly than the Ge-Si bonds during lithiation and simultaneously Ge-Li bonds form in an almost linear fashion.

Finally, the

  depletion of Ge-Ge in Si0.5Ge0.5 is in agreement with the decrease in the refined lattice parameter detected by in situ XRD during the first part of the lithiation (Table Annex II-3) and with the Si-Ge and Ge-Ge Raman bands' evolution. Additional details of the four-component fitting can be found in the Annex II.

Figure II. 30 .

 30 Figure II.30. In situ Raman spectroscopy measurement on a Si0.5Ge0.5 self-supported film cycled vs Li/Li + at the rate of C/5. Raman spectroscopy is useful to complete the study of the local atomic modifications occurring during lithiation/delithiation. Figure II.30 shows the in situ Raman spectra for Si0.5Ge0.5. During the first discharge a decrease in the intensity of the different bands: Ge-Ge, Si-Ge and Si-Si, is observed at the initial lithiation stage and assigned to the formation of a-Lix(Si/Ge). Subsequently, at the end of discharge, the Ge-Ge, Si-Si and Si-Ge signals disappear completely and no other peaks associated to the binary or ternary lithiated phases are observed. The loss of the Raman signal of Si0.5Ge0.5 when fully lithiated could be related with the change in the optical penetration depth occurring upon lithiation or

  5Ge0.5 electrode (Figure II.31) in which Si and Ge remain homogeneously distributed.

Figure II. 31 .

 31 Figure II.31. EDX analysis of a self-supported Si0.5Ge0.5 anode after one charge.

  Ca(Si1-xGex)2. The CaSi2, CaGe2 and Ca(Si1-xGex)2 Zintl phases were mechanically synthesized by placing stoichiometric amounts of Ca (99% ACROS Organics) and Si (325-Mesh, 99%, ACROS Organics) or Ge (325-Mesh, 99%, ACROS Organics) in an airtight 50 mL WC jar, with a ball:mass ratio of 10:1 under an inert atmosphere. The jar was rotated at 600 rpm for 18h, in a planetary ball mill (RETSCH PM 100). The recovered powder was loaded into stainless steel tubes sealed by Arc welding under Ar atmosphere and annealed for 50 h at 800°C. Xene: Siloxene, Germanene and Siligane. The siloxene, germanene and siligane were synthesized in air atmosphere, by dissolving ~300 mg of CaSi2/CaGe2/Ca(Si1-xGex)2 in ≈30 mL of HCl (35% VWR Chemicals) at -20°C for 8 hours. The product was washed with distilled water and acetone to removed the residual soluble products, like CaCl2, separated by centrifugation and vacuum dried for 30 min. All synthesis products were kept in an argon atmosphere to avoid further oxidation.

  diffractograms were collected with a zero-background single Si crystal XRD sample holder. The diffraction patterns were refined (Le Bail profile matching) using the Fullprof software. Confocal ex situ Raman Spectroscopy was performed using a DXR 2 Raman Microscope (Thermo-Fisher Scientific, 532nm excitation). The samples were placed in a glass slide and sealed with Kapton tape in order to avoid exposure to air. The Fourier transformed Infrared Spectra (Nicolet IS10, Themo-Fisher Scientific) were performed using KBr methods, an airtight cell with NaCl window was used for air sensitive samples.The thermogravimetric analysis were done with a STA 449 F3 Jupiter NETZSCH thermal analyzer placed in an Ar-filled glove box. Particle morphology, size and microstructure were studied by SEM (SEMenvironmental FEI Quanta 200 FEG microscope) while the composition was determined by Energy Dispersive X-ray EDX (X INCA Oxford). TEM was performed in a TEM-FEI Tecnai F20 S-TWIN, operating at 200 kV fitted with a Gatan Image Filter Tridiem in post column. The electron diffraction patterns were obtained with Selected Area Electron Diffraction (SAED). The air sensitive samples were prepared in an Ar filled glove box and transferred to the TEM under an Ar flow.In situ temperature controlled XRD. The XRD with temperature were performed in a Bruker D8 Advanced diffractometer with Cu radiation (λ1 = 1.54056 Å, λ2 = 1.54439 Å) equipped with a HTK 1200°C Anton Parr Chamber.Electrode preparation. The electrodes were prepared with a 1:1:1 mass ratio of active material, carbon super C45 and Carboxyl Methyl Cellulose (CMC-DS = 0.9, Mw = 700 000 Aldrich). The carbon additive and the active material were pre-mixed by SPEX for 10 min, in an inert atmosphere, using 1 ball per 100 mg of total mass. The resulting powder was recovered and placed in a mortar, where the binder was added. The mixture was dissolved in distilled water with 0.1% of Triton X and magnetically stired for one night. The slurry was casted into copper foil using a doctor blade (50 µm thickness) and dried at ambient conditions for one day. 11.1 mm disks were cut and dried under vacuum at 70°C overnight.Electrochemistry. The electrodes were tested in Li, Na and K half-cells using CR2032 coin cells, 1 M LiPF6 (NaPF6) in Ethylene Carbonate (EC)/Dimethyl Carbonate (DMC) [1:1] electrolyte with 1% Fluoroethylene carbonate (FEC) as additive, and a Whatman glass fiber separator (GF/ D, 675 μm). Alternative solvent mixtures were tested, EC/DEC, EC/PC (polycarbonate) and tetraglyme while the salt XPF6 (X=Na, Li) concentration and ratio of FEC additive was kept constant. For K several electrolyte formulations and the impact of FEC addition were tested, 0.8 M KFSI in 1:1 EC/DEC, 0.8 M KFSI 1:1 EC/DMC, 0.8 M KPF6 1:1 EC/DEC and 0.8 M KPF6 1:1 EC/DMC. Galvanostatic cycling was performed at 25°C between 0 and 2 V versus Li/Li + , Na/Na + or K/K + at C/20 or C/10 (C/n = 1Li, Na or K/n h), using a VMP3 or MPG2 (Biologic) device.3 Silicon-based layered phasesThe silicon-based layered phases have been known for a long time and have found applications mainly in the optoelectronics and semiconductors domain. In this study, the layered siloxene obtained from topotactic deintercalation of Ca from CaSi2 was used as an anode material for Li, Na and K-ion batteries. Reversible capacities of 2300, 311 and 296 mAh/g were obtained for Li, Na and K, respectively, with good capacity retention and coulombic efficiency for Li and Na. Given the high versatility of the synthesis method, it was possible to chemically intercalate organic molecules with different carbon chain length in siloxene, the implications of the change of the interlayer space in the electrochemistry have been evaluated. Certain results from this chapter have been adapted from the publication: Loaiza L. C. , Monconduit L. , Seznec V. Siloxene : A potential layered silicon intercalation anode for Na , Li and K ion batteries, J. Power Sources. 417 (2019) 99-107. doi:10.1016/j.jpowsour.2019.02.030.

  planes in bulk Si in which the Ca 2+ ions are replaced by other Si atoms. Several polymorphs of various stacking sequences are reported (Figure III.1); i) 1-layer by trigonal (1T) or hexagonal (1H) unit cell (space group P3m1), ii) the 2-layer by hexagonal unit cell 2H (P63mc) and iii) the 3-(3R) and 6-(6R) layers by rhombohedral unit cell (R-3m

Figure III. 1 .

 1 Figure III.1. Crystal structures for the CaSi2 polymorphs 6R, 3R and 1T/1H.

Figure III. 2 -

 2 Figure III.2-a shows the diffraction pattern for the synthesized CaSi2, the Le Bail refinement confirms the presence of the 6R and 3R polymorphs with a=3.8578(3) Å, c= 30.687(5) Å and V=395.52(1) Å 3 as cell parameters for 6R and a=3.826(3) Å, c= 15.917(5) Å and V=201.83(2) Å 3 for 3R, in agreement with the literature (ICSD 248515 and 193539).[229,232] Note that the peak intensity calculation presents some

  Figure III.2-c) confirms the compact layered morphology of CaSi2 without any gaps.

Figure III. 2 a

 2 Figure III.2 a) Le Bail refinement of the CaSi2 XRD pattern, b) Raman Spectrum and c) SEM picture of CaSi2.

1 )Figure III. 3 .

 13 Figure III.3. Topotactic transformation of CaSi2 into Siloxene (left), picture of the synthesized siloxene (right).

Figure III. 4

 4 Figure III.4. a) XRD pattern for siloxene, b) SEM and c) TEM images displaying the layered character of siloxene.

Figure III. 5

 5 Figure III.5. a) XRD and b) electron diffraction pattern for siloxene.

  corrugated silicon layers with three Si-Si bonds and alternating -H or -OH substituents as in Wöhler siloxene, two-dimensional Si sheets made of Si6 rings linked by oxygen bridges, as in Kautsky Siloxene and one-dimensional Si-Si chains interconnected by Si-O bonds. The Wöhler siloxene gradually decomposes into Kautsky siloxene through a hydrolysis process of Si-H to

Figure III. 6 -

 6 Figure III.6-a shows the Raman spectra of siloxene composed of a main band at 497 cm -1 ,

Figure III. 6 -

 6 b shows several vibration bands at 465, 518, 643, 809, 875, 895, 1600, 2110 and 2250 cm -1 , with the most intense bands at 1060 and 3400 cm -1 (Table

Figure III. 6

 6 Figure III.6. a) Raman and b) Infrared spectra of siloxene.

Figure III. 7 , 2 .

 72 two different peak deconvolutions are possible, their corresponding binding energies and atomic quantification are displayed in Table III.The first possibility (Figure III.7-a), includes three components: Si 0 from the residual c-Si from the synthesis as observed in the XRD and Raman experiments and from Si-H, Si 1+ possibly from Si-OH and Si 3+ from Si(OSi)2(OH)Si and/or other sub-oxides. The second possibility (Figure III.7-b) comprises the presence of four components, the three first ones identical to the first possibility and an extra one with

III. 7 .

 7 Figure III.7. XPS spectra of siloxene, two different peak deconvolutions are proposed (left and right).

Figure III. 8 .

 8 Figure III.8. DSC and TGA curves for siloxene at 10 K/min performed under air conditions.

Figure III. 9

 9 Figure III.9 shows the in situ temperature-controlled XRD of siloxene, the experiment was performed in an oven with Kapton foil, which originated a series of additional diffraction peaks, thus only the(100) peak from siloxene and the (111) from c-Si impurity, were observed. No change is detected until 50°C, after which the intensity of the (100) siloxene reflection starts to decrease until complete disappearance at 300°C. Meanwhile, the c-Si(111) diffraction peak prevails all along the experiment and presents a slight broadening and shifting towards higher angles from 100-450°C, possibly due to a thermal expansion phenomenon. At the end of the experiment, only the Bragg peak from the c-Si and those ones from the Kapton foil are observed. Consequently, the siloxene is thermally stable up to 250°C, yet the decomposition starts at around 100°C. The high temperatures could intensify the oxygen cross-linking of the planes until their final destruction and formation of different sub-oxides or SiO2.Since the experiment was performed under vacuum, the oxygen source are the -OH ending groups.These results are in line with the TGA (performed under air, thus the oxygen source is different) in which the siloxene gains mass after 250°C, possibly due to the formation of amorphous sub-oxides (not observed in the XRD pattern Figure III.9), while the DSC shows an exothermic peak at 320°C in line with an oxidation process.

Figure III. 12 .

 12 Figure III.12. Capacity retention of siloxene vs a) Li, c) Na and e) K at C/20. Rate performance at different C/rates for siloxene vs b)Li and d) Na. The capacity values correspond to discharge capacity.

  delivers significant capacity values (Figure III.12-b). On the other hand, for sodium at high rates the electrode delivered poor capacities, (Figure III.12-d) probably this kinetic limitation is associated with the higher Na radius size and/or its sluggish diffusion. In both cases, once the C/rate is reestablished to C/20 similar capacity values to the initial ones are recovered, meaning that high currents are not detrimental.

Figure III. 13 .

 13 Figure III.13. Galvanostatic curves for siloxene cycled vs Li at the rate of a) C/10, b) C/5, c) 1C and d) 2C.

Figure III. 14 .

 14 Figure III.14. For C/10, C/5, 1C and 2C 1340, 1142, 1114 and 796 mAh/g are delivered after 20, 47, 66

Figure III. 14 .

 14 Figure III.14. Capacity Retention for siloxene cycled vs Li at a) C/10, b) C/5, c) 1C and d) 2C. The capacity values correspond to discharge capacity.

Figure III. 15

 15 Figure III.15 the presence of this additive is crucial to ensure a good performance, since poorer capacities are obtained with no FEC.

Figure III. 15 .

 15 Figure III.15. Galvanostatic curve of siloxene cycled vs Li using standard LP30 (1 M LiPF6 in EC/DMC) electrolyte with no FEC additive.

Figure III. 16 .

 16 Figure III.16. Galvanostatic curves for siloxene cycled vs a) Li and 1 M LiPF6 in EC/DMC with 1 and 5% FEC and b) 1 M NaPF6 in EC/DMC with 1 and 5% FEC. The optimal percentage of FEC additive was determined by cycling cells at the same conditions with 1 M LiPF6 or NaPF6 in EC/DMC with 1 and 5% of FEC additive, the results are shown in Figure II.16.

  galvanostatic curves are shown in Figure III.17. For Li, (Figure III.10-a and Figure III.17

  Na, (Figure III.10-c and Figure III.17-b,d,f) the first cycle delivered 750, 815, 880 and 920 mAh/g for EC/DMC, EC/DEC, EC/PC and tetraglyme, respectively. Correspondingly, the second cycle displayed capacities of 311, 350, 320 and 400 mAh/g. No significant changes were found in the galvanostatic curve during the discharge/charge cycles, except for tetraglyme, where the first discharge presents lower sodiation voltages and the desodiation has a different sloping profile. The particular behavior of tetraglyme for the lithiation and sodiation could be an indicator of a different reactivity of ether-based electrolytes with siloxene, possibly the presence of -OH and -H groups impacts the nature of the SEI components.

Figure III. 18

 18 Figure III.18 shows the capacity retention for these electrolytes formulations. ForNa 191, 277 and 253 mAh/g are delivered after 80, 48 and 62 cycles for the EC/DEC, EC/PC and tetraglyme mixture, corresponding to a capacity retention (respect to the second cycle) of 66%, 87% and 62%, respectively.The initial Coulombic efficiency (CE) increases from 35 to 88% in EC/DEC, 35 to 87% in EC/PC and 52 to 100% in tetraglyme; following they attain an average value of 90-98%. For Li, only the EC/PC mixture was tested for longer term (in comparison with the cycling previously presented) and after 40 cycles a capacity of 1795 mAh/g is obtained, corresponding to a capacity retention of 90%, the Coulombic efficiency increases from 56 % to 91% from the first to the second cycle keeping an average of 94%.The rate capability properties of EC/DEC and EC/PC solvent mixtures have been examined for Na(Figure III.19) with the aim to detect any improvement of the cyclability as a function of the electrolyte formulation; however, the behavior is not very distant from the cells cycled with EC/DMC mixture.

Figure III. 18 .

 18 Figure III.18. Capacity retention for siloxene cycled vs. Na at C/20 using 1 M NaPF6 in a) 1:1 EC/DEC with 1% FEC, b) 1:1 EC/PC with 1% FEC and c) tetraglyme with 1% FEC. d) capacity retention for siloxene cycled vs Li using 1 M LiPF6 in 1:1 EC/PC with 1% FEC. The capacity values correspond to discharge capacity.

Figure III. 19 .

 19 Figure III.19. Capacity retention for siloxene cycled vs. Na using 1 M NaPF6 in a) 1:1 EC/DEC with 1% FEC and b) 1:1 EC/PC with 1% FEC.

  from the study due to its very poor performance. Their respecting galvanostatic curves at C/20 are shown in Figure III.21, at a first sight there is an improvement in the cyclability, and the first and second discharge delivered 1186, 471, 629 and 790 mAh/g and 373, 193, 133 and 295 mAh/g for KPF6 EC/DMC, KFSI EC/DMC, KPF6 EC/DEC and KFSI EC/DEC, respectively. The profile of the galvanostatic curve is different from the previous experiments, and for the first discharge the plateau at 1.5 V is absent and instead there is a plateau at ≈1.0 V for the EC/DEC mixtures and a gentle pseudo-plateau at approximately the same potential for EC/DMC. Following the potential decreases until the end of discharge with a smooth sloping profile. The charge profiles are characterized by a rapid increase in the potential and a region with different slope between 0.4-0.5 V that gradually disappears. The second discharge presents a constant decrease in the voltage with a very slight change in the slope between 0.5-0.7 V for the EC/DEC.

Figure III. 21 .

 21 Figure III.21. Galvanostatic curves for siloxene cycled vs K at C/20 using a) 0.8 M KPF6 in 1:1 EC/DMC, b) 0.8 M KFSI in 1:1 EC/DMC, c) 0.8 M KPF6 in 1:1 EC/DEC and d) 0.8 M KFSI in 1:1 EC/DEC electrolyte.

Figure III. 22 .

 22 Figure III.22. Capacity retention for siloxene cycled vs K using a) 0.8 M KPF6 in 1:1 EC/DMC, b) 0.8 M KFSI in 1:1 EC/DMC, c) 0.8 M KPF6 in 1:1 EC/DEC and d) 0.8 M KFSI in 1:1 EC/DEC electrolyte. The capacity retention (Figure III.22) has been improved compared to the previous experiments and 94, 130, 88 and 220 mAh/g are retained after 32, 30, 30 and 47 cycles for the KPF6 in EC/DMC, KFSI in EC/DMC, KPF6 in EC/DEC and KFSI in EC/DEC, respectively; these values correspond to a capacity retention of 25, 67, 47 and 74% with respect to the second cycle. The initial Coulombic efficiency is in the range of 17-38% and after a few cycles for EC/DEC it attains an average value of 88%, while for EC/DMC more than 10 cycles are necessary to reach 80%. Clearly, the electrolyte formulation highly impacts the electrochemical performance and in this case the presence of FEC in the potassium-based electrolytes (Figure III.20) is unfavorable for the cyclability, contrary to the observations for Li and Na.

Figure III. 24 .Figure III. 25 .

 2425 Figure III.24. SEM images for the different modifications of siloxene.

2257 cm - 1 .

 1 Infrared spectra for the modified siloxene, shown in Figure III.26, resemble the one of pristine siloxene. The IR bands indicate the presence of Si-H bonds at 643 and 2110 cm -1 and Si-O-Si at 809 and 1060 cm -1 , Table III.6 depicts a complete assignment of the bands. The rectangles in Figure III.26 highlight bands that are related with the presence of C-H bonds. The first group of bands centered at

  characteristic vibrations; instead, only the -OH, Si-H and Si-O-Si associated bands are observed. For the Si-O-Si band (≈1000 cm -1 ) there is band splitting, consequence of the oxygen incorporation in the planes to form Kautsky siloxene and the chain polymorph. The absence of the C-H vibrations could indicate the intercalation of the dodecyl sulfate with no chemical bonding, some of these molecules might be removed during the washing process while some others are irreversibly trapped in the layers.

Figure III. 26 .

 26 Figure III.26. Infrared spectra for the pristine and modified siloxene, from bottom to top, pristine, eth-siloxene-, prop-siloxene, but-siloxene, pent-siloxene, dec-siloxene and exfoliated siloxene.

Section 2 )

 2 and the electrolyte was 1 M LiPF6 (NaPF6) in EC/DMC 1%FEC. The results are shown in Figure III.27, for all the modifications the overall capacities are lower compared to the pristine siloxene for both Li and Na, still, Li delivers higher capacities compared to Na. For the first discharge, the obtained capacities for the eth-siloxene, pent-siloxene and dec-siloxene are 2199, 1656 and 2589 mAh/g for Li and 414, 336 and 291 mAh/g for Na, respectively. Likewise, the second discharge delivered 840, 750 and 1019 mAh/g for Li and 145, 123 and 123 mAh/g for Na. The profile of the galvanostatic curves resembles the one of pristine siloxene (Figure III.10, section 3.2.1)

Figure III. 27 .

 27 Figure III.27. Galvanostatic curves for the eth-siloxene cycled vs a) Li and b) Na, pent-siloxene cycled vs c) Li and d) Na and dec-siloxene cycled vs e) Li and f) Na. The cells were cycled at a rate of C/20.

Figure III. 28 .

 28 Figure III.28. Capacity retention for the eth-siloxene cycled vs a) Li and b) Na, c) pent-siloxene cycled vs Li and d) dec-siloxene cycled vs Li.

Figure III. 29 .

 29 Figure III.29. XRD diffraction patterns for the eth-siloxene, pent-siloxene and dec-siloxene annealed from 250-300°C. The galvanostatic curves for the annealed eth-siloxene and pent-siloxene are presented in Figure III.30. The lithiation capacities have been improved, while the sodiation present no significant change. For the first discharge, for Li the capacities were 2470, 1952 and 2609 mAh/g for eth-siloxene 250°C, eth-siloxene 300°C and pent-siloxene 250°C, respectively, while for Na the capacities were 410, 544 and 465 mAh/g. The second discharge produced for eth-siloxene 250°C, eth-siloxene 300°C and pent-siloxene 250°C 1002, 848 and 986 mAh/g for Li and 139, 146 and143 mAh/g for Na, respectively. The discharge/charge profile in the galvanostatic curve is not distinct from the non-annealed modified siloxene. Interestingly, for Li, the plateau at 0.4 V during the second discharge, is visible for the eth-siloxene 250°C and pent-siloxene 250°C; and the profile of the second and subsequent discharges resembles the one of siloxene. For the eth-siloxene 300°C no significant change was observed compared to the non-annealed sample.

Figure III. 30 .

 30 Figure III.30. Galvanostatic cycling for the eth-siloxene annealed at 250°C and 300°C and cycled vs a), c) Li and b), d) Na, respectively. Pent-siloxene annealed at 250°C and cycled vs e) Li and f) Na.

Figure III. 31 .

 31 Figure III.31. Capacity retention for the eth-siloxene annealed at 250°C and cycled vs a) Li and b) Na, eth-siloxene annealed at 300°C cycled vs c) Li and d) Na and pent-siloxene annealed at 250°C and cycled vs e) Li and f) Na. The capacity retention properties (Figure III.31), were evaluated for only a few cycles for Li and after 16 cycles 975, 705, 961 mAh/g are delivered for eth-siloxene 250°C and 300°C and pentsiloxene 250°C, respectively and corresponding to a capacity retention with respect to the second cycle of 97, 83 and 97%. For Na after about 20 cycles the cell delivered 103 and 98 mAh/g for eth-siloxene 250°C and pent-siloxene 250°C and 125 mAh/g after 50 cycles and for eth-siloxene 300°C corresponding to a capacity retention of 74, 68 and 85%, respectively.

Figure III. 33

 33 Figure III.33. a) Le Bail profile matching for the CaGe2 phase, b) Raman spectrum for CaGe2, c) SEM pictures of the CaGe2. The CaGe2 was synthesized by ball milling, mixing stoichiometric amounts of Ca and Ge followed by an annealing process at 800°C in a sealed tube under Ar. The diffraction pattern and the corresponding Le Bail profile matching are presented in Figure III.33-a. The analysis indicates the presence of at least two polymorphs, the 3R and 6R, but other polymorphs could be present. The annealing temperature (800°C) allows the stabilization of both 3R and 6R, although if increased to 950°C only the 6R should be formed. The profile matching provided refined cell parameters of a/b=4.008(2) Å, c=10.320(7) Å and V= 143.64(1) Å 3 for the 3R-CaGe2 and a/b=3.999(2) Å, c=30.714(3) Å and V= 425.52(5) Å 3 for the 6R-CaGe2, in line with the literature.[148,229,270,271] Likewise CaSi2 the CaGe2 presents a turbostratic

Figure III. 33 -

 33 Figure III.33-c indicate a layered morphology, with particles of different lateral size, agglomerates of several stacked layers and some particles with different morphology.

Figure III. 34

 34 Figure III.34. a) SEM picture of germanane, b) TEM picture of germanane

  Figure III.36. a) Raman and b) Infrared spectra for germanane. The characterization of germanane continues with the Raman and the infrared spectroscopy, display in Figure III.36. The Raman spectrum (Figure III.36-a) shows a change with respect to the CaGe2, and

227 cm - 1 ,

 1 probably the polarization of the incident light deactivates this Raman mode in our germanane. The Infrared spectrum in Figure III.36-b, shows the presence of Ge-H vibrations, centered

830 cm - 1

 1 corresponds to the Ge-O-Ge stretching mode, indicating a partial oxidation of the planes or the presence of glassy/amorphous oxides. No -OH vibrations are found for germanane, thus the bonding is preferentially done with -H.

Figure III. 37 .

 37 Figure III.37. In situ XRD with temperature from 30-450°C for germananeIn order to evaluate the stability of germanane, an in situ controlled temperature XRD was performed from 30 to 450°C (Figure III.37). Note that a sample with high c-Ge content was deliberately chosen in order to survey the germanane diffraction peaks evolution. The (002) and (100) Bragg reflections of germanane are present until 150°C, temperature after which they shift towards higher angles and intensity decreases until complete disappearance at ≈300°C. Meanwhile, the peak at 27° assigned to c-Ge produced as a synthesis by-product and the additional small peaks from the Kapton window remain unchanged. According to the literature, the amorphization of germanane starts at 75°C

Figure III. 39 )

 39 , the germanane still delivers 1500, 354 and 150 mAh/g for Li, Na and K after 30, 35 and 30 cycles corresponding to a capacity retention of 67, 71 and 73%, respectively. Note that the capacity retention percentages are referred to the second cycle. The germanane showed great cyclability at high C/rates for Li as observed in Figure III.39-b,

  at high C/rates was tested, the results are presented in Figure III.40.

Figure III. 39 .

 39 Figure III.39. Capacity retention of germanane vs a) Li at C/10, c) Na and e) K at C/20. Rate performance at different C/rates for siloxene vs b) Li and d) Na. The capacity values correspond to discharge capacity.

Figure III. 40 .

 40 Figure III.40. Galvanostatic and their corresponding capacity retention curves for germanane cycled vs Li at the rate of a), c) 1C and b), d) 2C. As observed in Figure III.40, the profile of the galvanostatic curve for germanane cycled at C and 2C

Figure III. 41 .

 41 Figure III.41. Comparison of the a) galvanostatic and b) derivative curves for germanane and bulk Ge.

Figure III. 42 .

 42 Figure III.42. Capacity retention of germanane and bulk Ge cycled vs a) Li and b) Na under the same conditions at a rate of C/10. The electrochemical behavior of bulk Ge and germanane were compared (Figure III.41). In bulk Ge

Figure III. 43 .

 43 Figure III.43. Galvanostatic and their respective capacity retention curves for the germanane cycled at C/20 vs K using a), c) 0.8 M KPF6 in 1:1 EC/DEC and b), d) 0.8 M KFSI in 1:1 EC/DEC. Finally, the potassiation of germanane was evaluated (Figure III.43), only two electrolyte formulations were tested; the 0.8 M KPF6 in 1:1 EC/DEC and 0.8 M KFSI in 1:1 EC/DEC; since they exhibited the best performances for siloxene. The obtained capacities for the first discharge were 922

  mechanically synthesized following the protocol described in Chapter II, by mixing stoichiometric amounts of Si and Ge. The XRD diffraction patterns and Raman spectra shown in Figure III.44 confirm the formation of the Si1-xGex solid solution.

Figure III. 44

 44 Figure III.44. a) XRD Patterns and b) Raman Spectra for the Si0.1Ge0.9, Si0.5Ge0.5 and Si0.9Ge0.1 solid solution.

Figure III. 45 .

 45 Figure III.45. XRD patterns for CaGe2, Ca(Si0.1Ge0.9)2, Ca(Si0.5Ge0.5)2, Ca(Si0.9Ge0.1)2 and CaSi2 from bottom to top.

For

  the Raman spectra (Figure III.46), the phases Ca(Si1-xGex)2 phases share many features with the Si1-xGex solid solution, for instance, the Ge-Ge bond presents a band approximately at 285-295 cm -1 , the Si-Ge at 395 cm -1 and the Si-Si between 497-505 cm -1 . An additional band for Ca(Si0.1Ge0.9)2 and Ca(Si0.5Ge0.5)2 appears at 230 cm -1 , a similar band is observed at the same frequencies for CaGe2, then probably both bands have the same origin. Likewise Si1-xGex, the presence of the Si-Ge band is an indicator of the formation of a solid solution.

Figure III. 46 .

 46 Figure III.46. Raman Spectra for Ca(Si0.1Ge0.9)2, Ca(Si0.5Ge0.5)2 and Ca(Si0.9Ge0.1)2

Figure III. 48 .

 48 Figure III.48. XRD patterns for the germanane, siligane Si0.1Ge0.9, siligane Si0.5Ge0.5, siligane Si0.9Ge0.1 and siloxene from bottom to top.

Figure III. 49 .

 49 Figure III.49. Raman spectra for siligane Si0.1Ge0.9, Si0.5Ge0.5 and Si0.9Ge0.1. The Infrared spectra of siliganes (Figure III.50) shows the presence of Si-O-Si band at ≈1070 cm -1 , the H-Si-O at 885 cm -1 and the H-SiO3 at 2259 cm -1 , most of these vibrations are observed also in the siloxene while no evidence of the Ge-H band of germanane at 2000 cm -1 was found. Additionally, the siligane Si0.9Ge0.1 exhibited the Si-OH at 1600 cm -1 , the Si-H at 2116 and 636 cm -1 and the -OH bands

Figure III. 51 . 9 Figure III. 52 .

 51952 Figure III.51. SEM picture of the siligane Si0.1Ge0.9

Figure III. 53 .

 53 Figure III.53. SEM picture siligane Si0.9Ge0.1 (top left), SEM and the corresponding EDX for siligane Si0.9Ge0.1 (top right and bottom)

Figure III. 54 .

 54 Figure III.54. Galvanostatic curves for the siligane Si0.9Ge0.1 cycled vs a) Li and b) Na and the siligane Si0.5Ge0.5 cycled vs c) Li and d) Na. The previously described processes are in line with the features in the corresponding derivative curves. For the lithiation of siligane Si0.9Ge0.1 (Figure III.55-a) during the first cycle there is a reduction peak at 1.2 V accompanied by a double peak at 0.01 and 0.19 V and one double oxidation peak at 0.3 and 0.45 V. The second cycle presents two reduction peaks at 0.32 and 0.01 V, while the oxidation mirrors the first charge. For the siligane Si0.5Ge0.5 (Figure III.55-c), the first cycle presented three reduction peaks at 1.2, 0.08 and 0.01 V and one very sharp oxidation peak at 0.42 V. For the second cycle only the discharge presented a different profile and three reduction peaks can be observed at

Figure III. 55 .

 55 Figure III.55. Derivative curves for the first and second cycle of siligane Si0.9Ge0.1 cycled vs a) Li and b) Na and the siligane Si0.5Ge0.5 cycled vs c) Li and d) Na.

Figure III. 56 .

 56 Figure III.56. Capacity Retention for a) siligane Si0.9Ge0.1 and b) siligane Si0.5Ge0.5 cycled vs Na. Only the sodiation was evaluated for long-term cycling (Figure III.56). The siligane Si0.9Ge0.1 retains 188 mAh/g after 40 cycles while the siligane Si0.5Ge0.5 144 mAh/g, corresponding to a capacity retention (with respect to the second cycle) of 79% and 84%, respectively. The capacity values for the sodiation and lithiation of siliganes are much lower than those ones obtained for germanane and siloxene.

Figure IV. 1 and

 1 Figure IV.1 and Figure IV.2, show the in situ XRD for a self-supported siloxene and germanane anode cycled vs Li. For the siloxene (Figure IV.1) the configuration of the in situ XRD cell induces a high background and the diffraction peaks proper from the phase are no longer visible. The resulting pattern is amorphous and only a very huge hump is observed at around 15°, as the lithiation proceeds no phase crystalizes, as expected from the formation of amorphous lithiated intermediates for bulk Si.

Figure IV. 1 .

 1 Figure IV.1. In situ XRD for a self-supported siloxene electrode cycled vs Li at C/40.

Figure IV. 2

 2 Figure IV.2 shows the in situ XRD experiment for germanane, in this case the increased number of electrons in Ge allowed a stronger diffraction. Thus, despite the high background from the in situ XRD cell, it is possible to observe the (002) Bragg reflection from germanane. This reflection is visible until approximately 1.5 V, following no other diffraction peaks are visible, until ≈0.4 V, where humps starts to appear at around 23° and 40°. These humps increase in intensity as the potential decreases and by the end of discharge they have transformed into very broad peaks that become broader and lose intensity upon delithiation. Note that there is a shift of these peaks, before the end of discharge, towards low and thereafter to high angles.

Figure IV. 2 .

 2 Figure IV.2. In situ XRD for a self-supported germanane electrode cycled vs Li at C/40.

Figure IV. 3 .

 3 Figure IV.3. Raman spectra of a) CaSi2, siloxene and lithiated/delithiated siloxene and b) CaSi2, siloxene and sodiated/desodiated siloxene. c) IR spectra of siloxene and lithiated/ delithiated siloxene and d) siloxene and sodiated/desodiated siloxene

(

  de)sodiation/(de)lithiation the -H and -OH bonds in siloxene are probably replaced by the alkali ion, resulting in a reversible Raman shift. The size of the intercalated ion might be involved as well, for CaSi2 only one Ca 2+ is intercalated, and later replace by two -H and -OH above and below the planes in siloxene. This change in ligand derives in an increase in the interlayer distance and a blue shift of the Si-Si Raman signal. This peak continues to blue shift with the introduction of Li or Na and to redshift with their extraction, indicating a change in the interlayer distance. Note that several experiments were performed in order to rule out possible interferences due to sample inhomogeneity, in all cases the trend was the same. The IR spectra for the cycled electrodes are shown in Figure IV.3-c,d; for both Li and Na the Si-O-Si, H-Si-O bands are preserved while the -OH ones have decreased in intensity for the sodiated/lithiated sample and almost disappeared for the desodiated/delithiated one. Additional bands appear at 1445 and 1310 cm -1 , probably originated from carbonate-based products from the electrolyte decomposition deposited on the electrode surface, reports in literature locate these vibrations at 1650 (C=O), 1398 (C-H), 1300 (C=O)cm -1 . [277] Likewise the Raman spectroscopy observations, the Si-H bands have almost disappeared for the cycled electrodes, suggesting a change in the bond nature, possibly due to new interactions with Na/Li ions, whereas the preservation of the H-Si-O and Si-O-Si bands could indicate the stability of the Si planes with no degradation into silicates upon reaction with the alkali.

Figure IV. 4 .

 4 Figure IV.4. Raman spectra of a) germanane (GeH) and lithiated/delithiated germanane and b) germanane (GeH) and sodiated/desodiated germanane. Given the reversible Raman shift observed for the siloxene experiment, an in situ experiment was performed to rule out possible differences in composition in the electrode. Figure IV.5 shows the Raman spectra for the lithiation, this cell configuration presented an increased noise/signal ratio and the intensity of the Raman bands had decreased. Still, it is possible to observe a band at 495 cm -1 , the G and D bands from the carbon conductive additive and the signal from the cell quartz window at ≈464 cm -1 . This last one impedes the proper assignment of the band at 495 cm -1 to the Si-Si in siloxene.

Figure IV. 5 .

 5 Figure IV.5. In situ Raman spectroscopy for a siloxene self-supported anode cycled vs Li at C/8.

Figure IV. 6

 6 Figure IV.6 shows the same experiment performed for the sodiation, an intense peak at 518 cm -1 ,

Figure IV. 6 .

 6 Figure IV.6. In situ Raman Spectroscopy for the sodiation of a self-supported siloxene anode, a) corresponding galvanostatic curve, b) Raman spectra (inset zoom of the main siloxene Raman peak)

Figure IV. 7 .

 7 Figure IV.7. Raman spectra of different lithiated references and the lithiated siloxene.The Raman spectroscopy results suggest the preservation of the Si-Si planes, for this reason SEM imaging was performed on the lithiated/sodiated and delithiated/desodiated siloxene and germanane.

Figure IV. 8

 8 Figure IV.8 shows the SEM pictures for siloxene; in both lithiation/sodiation and delithiation/desodiation the layered morphology is preserved. As previously observed, the particle size distribution is inhomogeneous and several lateral sizes can be found. The different stacks of layers are surrounded by the conductive carbon additive. These results are in line with the reports of Pazhamalai et al.[243] who observed a preservation of the layered morphology of siloxene after cycling.

Figure IV. 8 .Figure IV. 9 .

 89 Figure IV.8. SEM pictures of a pristine self-supported siloxene anode and after (de)lithiation/(de)sodiation.

Figure IV. 9 ,

 9 Figure IV.9, shows the SEM pictures for a germanane self-supported anode after lithiation/sodiation and delithiation/desodiation. In analogy with siloxene the layered morphology is preserved and the particle size distribution is inhomogeneous. Some of the layers have been exfoliated during the cycling process and the edges of the layers have been deteriorated, probably due to the formation of passivation products on the surface or to the degradation induced by the electron beam. These processes are observed in siloxene also but are more pronounced in germanane since it is easily exfoliated and more sensible to the electron beam.

Figure IV. 10 .

 10 Figure IV.10. TEM picture for a sodiated siloxene self-supported electrode (left), electron diffraction of the sodiated siloxene (right).

Figure IV. 11 .

 11 Figure IV.11. TEM picture of a lithiated self-supported siloxene electrode. A similar analysis has been performed on the lithiated siloxene (Figure IV.11), likewise for the sodiated sample, the layered morphology is preserved but the edges of the layers are less sharp and present a decomposition layer. The electron diffraction and EELS in this decomposition layer indicates the presence of polycrystalline LiF (Figure IV.12). Interestingly, the LiF is found as an assembling of

Figure IV. 12 .

 12 Figure IV.12. TEM picture of a lithiated self-supported siloxene electrode (left). Electron diffraction (center) and Energy Loss Spectra (right) of the surface layer around the particles. Further EELS analysis was performed on different areas of the passivation layer on the particle surface. The results in Figure IV.13, indicate the presence of various signals, in certain cases, the Li K-edge is composed of a double contribution at 60.1 and 64.5 eV, in agreement with Li2O. In other

Figure IV. 13 .Figure IV. 14 .

 1314 Figure IV.13. TEM images of the lithiated siloxene with regions of different composition of the SEI layer, a) type-LiF, b) type-Li2O and c) type-Li2CO3, d) their corresponding EELS spectra and e) EELS spectra of the reference compounds.

Figure IV. 15 )

 15 , likewise the previous observation of SEM, the layered morphology is preserved, with the presence of staggered layers. The loss in the sharpness of the layers and the formation of the passivation layer is more pronounced for germanane, compared to siloxene. Contrary to the pristine germanane, the lithiated compound is amorphous as indicated by the electron diffraction. An EELS analysis of the surface of the layers revealed mainly the presence of Li2O (Figure IV.16). Possibly, the preferential bonding of germanane with -H over -OH has an impact on the decomposition products that are formed upon cycling. Likewise, it is not possible to differentiate the Li inside the layers from the Li in the SEI, due to the thickness of the electrolyte decomposition layer, which is exacerbated in germanane and the limited penetration depth of the EELS.

Figure IV. 15 .

 15 Figure IV.15.TEM images for a lithiated self-supported germanane electrode. Orange edge, electron diffraction of the corresponding image.

Figure IV. 16 .

 16 Figure IV.16. TEM image of a lithiated self-supported germanane electrode and the correspondingEELs spectra for the degradation layer at the particle surface.

Figure IV. 17 .

 17 Figure IV.17. XPS spectra of pristine siloxene with different peak deconvolution a) three components and b) four components, c) XPS spectra of the fresh siloxene composite electrode (active material, conductive additive, binder). Two possible peak deconvolution and assignments are possible for the pristine siloxene (Figure IV.17-a,b) (As discussed in Chapter III). The Si2p states of siloxene can be resolved in two components; Si-Si bonding and Si-O. For the fresh electrode the Si-Si component has clearly decreased in intensity while the Si-O has increased. If the first deconvolution possibility is taken into account (Figure IV.17-a), then the Si 3+ peak splits into two peaks for the fresh electrode, one corresponding to Si 3+ but at lower energy and the other one to Si 4+ . If the second possibility (Figure IV.17-b) is

Following, the lithiated

  and delithiated electrodes of siloxene were analyzed (Figure IV.18). After the lithiation (Figure IV.18-b), the peaks of Si 0 and Si + are not visible anymore and only the Si 3+ peak is preserved. At this stage probably the significantly thick SEI layer (as observed in TEM) hinders the access of the X-ray beam to the siloxene and only the signal from the SEI layer is obtained. After delithiation, the Si 3+ and Si + peaks are observed. For the second discharge (Figure IV.19-left) similarly

Figure IV. 18 .

 18 Figure IV.18. XPS spectra of a) pristine siloxene electrode, b) lithiated siloxene and c) delithiated siloxene.

Figure IV. 19 .

 19 Figure IV.19. XPS spectrum of the (left) second lithiation of a siloxene electrode (green), the red line indicates the pristine siloxene, and (right) sputtered Li15Si4 reference synthesized by ball milling.

  7 ppm corresponds to a LixSi phase formed during the reduction. The lithiated siloxene presents both resonances, while for the delithiated sample only the one close to 0.7 ppm, meaning that most of the signal is coming from the SEI or from Li stored or trapped in the surface of the layers. A comparison of the LixSi from siloxene with the references suggests that both Li15Si4 and LiSi resonances can be comprised in the lithiated siloxene, but other phases with similar Li environments can coexist at this stage. Interestingly, a resonance at about 2 ppm is observed in both lithiated references and might be included as well in the lithiated/delithiated siloxene.

Figure IV. 20 .

 20 Figure IV.20.6 Li NMR of the first discharge/charge cycle of a self-supported siloxene electrode and the LiSi and Li15Si4 references synthesized by ball milling.

Figure IV. 21 .

 21 Figure IV.21.6 Li NMR spectra for the Li15Si4 and LiSi references synthesized by ball milling.

Figure IV. 21

 21 Figure IV.21 shows a more detailed view of the lithiated references. The Li15Si4 NMR signal is composed of a sharp peak at 2 ppm accompanied by a very broad signal centered between 5-6 ppm while LiSi presents three main bands at 19, 7 and 2 ppm. For the Li15Si4, the band at 5-6 ppm (78%) indicates the presence of isolated Si 4-ions surrounded by Li.[110,181,191,192,291] A deconvolution of this peak suggests an

Figure IV. 24 .

 24 Figure IV.24.6 Li NMR of the Li15Si4 and LiSi references synthesized by ball milling and the siloxene self-supported electrode stopped at 0.18 V during the first discharge and 0.10 V during the second discharge.

Figure IV. 24

 24 Figure IV.24 shows a comparison of the NMR spectra at 0.10 V with the Li15Si4 and LiSi references.

Figure IV. 25 . 1 H

 251 Figure IV.25. 1 H NMR for the CMC binder, pristine siloxene, fresh self-supported composite electrode (CMC, C45, Siloxene), cycled self-supported electrode for the 1 st and 2 nd lithiation/delithiation.

Figure IV. 26 . 1 H

 261 Figure IV.26. 1 H NMR for the CMC binder, pristine siloxene, fresh self-supported composite electrode (CMC, C45, Siloxene), self-supported siloxene electrode at different depths of discharge/charge.

The 1 H

 1 NMR spectra has been obtained at different depths of lithiation/delithiation (Figure IV.26).

  Figure IV.28-d. A b value of 0.5

Figure IV. 28

 28 Figure IV.28. a) Cyclic voltammetry of siloxene from first to sixth cycle at a scan rate of 0.1 mV/s. b) Cycling voltammetry at different scan rates from 0.1 mV/s to 1 mV/s. c) Cycling voltammetry at various scan rates from 5 mV/s to 100 mV/s. d) Determination of b-value using the relationship between the peak current and the scan rate.

Figure IV. 29

 29 Figure IV.29. a) Cyclic voltammetry of germanane from first to sixth cycle at a scan rate of 0.1 mV/s. b) Cycling voltammetry at different scan rates from 0.1 mV/s to 1 mV/s. c) Cycling voltammetry at various scan rates from 5 mV/s to 100 mV/s. d) Determination of b-value using the relationship between the peak current and the scan rate.

Figure IV. 29

 29 Figure IV.29-a, shows the CV curves for the first cycles of germanane, during the first cycle one reduction peak at 0.94 V and two oxidation ones at 0.47 and 1.06 V are observed. For the second cycle, there are three reduction peak at 0.55, 0.3 and 0.04 V while for oxidation two peaks at 0.45 and 1.1 V.

Figure Annex I- 1 .

 1 Figure Annex I-1. Scheme of the cell configuration used in operando confocal Raman spectroscopy.

Annex II- 1 .Figure Annex II- 2 . 0207 Figure Annex II- 3 .

 1202073 Figure Annex II-1. Fitting results of the second reconstructed MCR-ALS component in the k (left) and in the R (right) space. The experimental spectrum is drawn in blue, and the fitting result in red. R-factor = 0.0085

  De nos jours, nous sommes confrontés à une crise climatique caractérisée par les effets du réchauffement climatique. La grande majorité des gaz à effet de serre à l'origine du réchauffement climatique proviennent de l'activité humaine, comme l'agriculture, les transports et l'énergie. Ces deux derniers sont principalement dus à la consommation de combustibles fossiles, dont les ressources sont limitées, responsable du rejet d'énormes quantités de CO2 et autres gaz à effet de serre dans l'atmosphère. Afin de réduire ces émissions, la manière dont l'énergie est produite et stockée doit être modifiée. La prise de conscience accrue liée au changement climatique a conduit à la mise en place de nouvelles politiques pour la transition énergétique, des énergies fossiles vers les énergies alternatives, comprenant l'utilisation entre autres : de cellules solaires, d'éoliennes ou de géothermie. Ces sources d'énergie sont renouvelables et peuvent être adoptées à l'échelle mondiale. Néanmoins, leur production d'énergie est intermittente. Par conséquent, un dispositif de stockage d'énergie est nécessaire et les batteries jouent un rôle clé dans ce contexte. Elles sont utilisées dans des applications fixes pour répondre à la demande en énergie lorsque la production est faible. Pour cette application spécifique, les dispositifs ne sont pas limités par la taille et plusieurs technologies (NaS, flux redox et autres Li-ion) avec différentes densités d'énergie et durées de vie ont été mises en oeuvre avec un coût relativement bas. D'autre part, le secteur des transports devrait passer des énergies à base de combustibles fossiles à l'électricité dans les prochaines années. Les batteries font partie des candidats les plus prometteurs pour participer à cette substitution et ont commencé à être mises en oeuvre dans différents moyens de transport, tels que les camions, les bus ou les voitures personnelles. Le marché des voitures électriques est dominé par les batteries lithium-ion (LIB), avec des compositions chimiques souvent à base d'oxydes de nickel-manganèse-cobalt (NMC), ou d'oxydes de nickel-cobalt-aluminium (NCA), ou de phosphate de fer et de lithium (LFP), le tout couplé à une anode en graphite. [1] D'autres chimies lithium-ion ou même les batteries sodium-ion (NIB) sont égalements à l'étude. Le nombre d'appareils nécessitant un dispositif de stockage d'énergie augmente de manière exponentielle. Pour satisfaire cette demande, une seule technologie ne suffira pas. Par exemple, si toutes les voitures en circulation sont transformées en voitures électriques utilisant des LIB, les sources de production des batteries seraient limitées. Plusieurs technologies différentes doivent être mises en oeuvre en tenant compte de chaque application. Ces technologies pourraient inclure les piles à combustible, le stockage de l'hydrogène et les batteries. Parmi ces dernières, il faut explorer d'autres systèmes que LIB, tels que les batteries à base de sodium, de potassium, de calcium ou d'aluminium, ainsi que les systèmes à flux redox. En outre, des études sur les technologies actuelles sont en cours et germanides lithiés appartiennent à la famille Zintl, leur étude est donc cruciale pour la compréhension des phénomènes se déroulant dans une batterie. En outre, le mécanisme électrochimique de Ge, Si1-xGex, du siloxène et du germanane a été étudié, l'objectif étant de découvrir les différents processus à l'origine de l'amélioration des performances de ces matériaux.

Figure 1 .

 1 Figure 1. RMN7 Li in situ d'un film autosupporté de Si0.5Ge0.5 a cyclé en fonction de Li/Li + à la vitesse de C/5. a) Tracé du contour des spectres RMN7 Li en fonction du temps. b) tracé de l'empilement des spectres de7 Li RMN avec le temps croissant. c) tension et courant en fonction du temps. Le pic à -17 ppm correspond au processus de surexposition de Li15 (Si0.5Ge0.5)4 en Li15+δ(Si0.5Ge0.5)4 à bas potentiel. S'attaquer aux problèmes posés par les électrodes à base de Si et de Ge implique également la compréhension des processus en cours pendant le cyclage. Le Chapitre II présente l'étude du mécanisme de lithiation électrochimique du Si, du Ge et de Si1-xGex. La sodiation et la potassiation ne sont pas prises en compte, étant donné le peu d'informations et les études limitées disponibles. Le mécanisme du Si est mis en évidence à partir de la littérature. Les mécanismes du Ge et de Si0.5Ge0.5

  de la délithiation, un produit amorphe a été obtenu. Pour le Ge, certaines des phases lithiées ont été identifiées (LiGe, Li7Ge3) et se sont révélées dépendantes des conditions de cyclage. En effet, à des vitesses comprises entre C/5 et C/8, Li15Ge4 était formé alors qu'à C/14, on obtient Li17Ge4 (Figure 2). Pour le Si0.5Ge0.5, il n'y a pas de séparation entre le Si et le Ge pendant le cycle, la synergie entre les deux éléments améliore les performances électrochimiques en termes de rétention de capacité et de réduction de la variation de volume.

Figure 2 .

 2 Figure 2. DRX et courbes de charge de décharge galvanostatique associées à a) C/5 et c) C/14 (vitesse de cyclage des électrodes de Ge). b), d) grossissements des gammes [18-28 °] et [38-44 °] 2θ ranges. Dans le Chapitre III, nous avons évalué l'impact de la morphologie sur les performances électrochimiques. Les couches de Si et de Ge dérivées des phases CaSi2 et CaGe2 Zintl ont été testées en tant qu'anodes pour LIB, NIB et KIB (Figure 3). Ces phases présentent l'ion Si 4-(Ge 4-) Zintl caractéristique qui forme des plans de cycles Si6 (Ge6) interconnectés. Ainsi, dans CaSi2 et CaGe2, le Ca 2+ peut être désintercalé topotactiquement et remplacé par -OH et -H pour former le siloxène et par -H pour former le germanane (liaison préférentielle de Ge avec -H).

Figure 4 .

 4 Figure 4. Courbes galvanostatiques pour le siloxène cyclé vs a) Li, b) Na et c) K à une vitesse de C / 20.

Figure 5 .

 5 Figure 5. Courbes galvanostatiques pour le germanane cyclé vs a) Li, b) Na et c) K à C / 10. Enfin, le mécanisme électrochimique du siloxène et du germanane a été exploré. La plupart des expériences ont été effectuées sur la lithiation du siloxène, pour faciliter la comparaison avec d'autres rapports de la littérature. Les résultats suggèrent que le cyclage électrochimique du siloxène est quelque peu différent de celui du Si. Le composé caractéristique, c-Li15Si4, du mécanisme d'alliage est absent dans le siloxène. De même, pour le germanane, aucun c-Li15Ge4 ou c-Li17Ge4 n'a été trouvé. L'analyse par microscopie indique que le siloxène et le germanane conservent leur morphologie en lamelles lors du cyclage et forment une couche de décomposition significative. Cette dernière est liée à la capacité irréversible et est composée de LiF, Li2O et différents carbonates de Li pour le siloxène et uniquement de Li2O pour le germanane, ce qui signifie que la nature des ligands joue un rôle important pour les produits de décomposition. La formation d'une couche de SEI optimale étant indispensable pour assurer une bonne cyclabilité, il est nécessaire d'étudier la formulation de l'électrolyte et sa relation avec les ligands siloxène et germanane sur les produits de décomposition électrolytique. L'étude des intermédiaires lithiés a été réalisée par une combinaison de techniques de spectroscopie ex-situ et de diffraction in-situ. Après cyclage, un changement possible de la distance inter-lamellaire associé à une interaction entre les couches siloxène / germanane et l'ion alcalin a été identifié par spectroscopie Raman. Les résultats de RMN (Figure 6) indiquent la formation d'un intermédiaire lithié

Ge)

				Cycling stability
			Specific capacity (mAh g -1 )	Cycle number	Current/rate
			LIB
		Si Nanoparticles	1180	500	3 Ag -1
		Ge Nanoparticles	1100	100	C/10
		Hollow Si	1420	700	C/2
		Porous Si	1200	600	C/2
	Nanosized Si (Ge)	Si Nanowire	3500	20	C/5
		Ge Nanowire	1000	20	C/20
		Si Nanofiber	1821	1000	C/5
		Si Nanotube	2000	50	C/5
		Ge Nanotube	1002	50	C/5
		Si nanofilm	3000	1000	12C
	Micrometer Si (Ge)	3D porous Si Pomegranate-like Si	2800 1160	100 1000	1C C/2
		Simple Ge/C	700	50	0.15 Ag -1
		Simple Si/C	1200	30	C/10
	Si (Ge)/C composites	Core-shell Si/C Yolk-shell Si/C Si/graphene	1490 1500 840	1000 1000 300	C/4 1C 1.4 A g -1
		Ge/graphene NW	1059	200	4C
	Other composites	Si/conductive polymer	550	5000	6 A g -1
		Si-Ge alloy	1600	50	2C
			NIB
	Nanosized Si (Ge)	a-Si Nanoparticle a-Si thin film	270 (after 1 st cycle) 600 (after 1 st cycle)	100 (248 mAh g -1 ) 100 (240 mAhg -1 )	20 mA g -1 C/10
	Si(Ge)/C composite	Si/C fiber	283 (after 1 st cycle)	200 (298 mAh g -1 )	50 mA g -1

Table I . 4 .

 I4 Comparison of the properties of carbon and silicon

	Carbon	Silicon
	Group 14	Group 14
	Four valence electrons in the outermost shell	Four valence electrons in the outermost shell
	Four valence orbitals	Four valence orbitals
	Formation of four two-electrons covalent bonds	Formation of four two-electrons covalent bonds
	Pauling electronegativity	Pauling electronegativity
	X =2.55 for C and 2.2 for H	X =1.90
	Polarization of the bond C-H C δ--H δ+	Polarization of the bond Si-H Si δ+ -H δ-
	Strong C-H bonds	Reactive Si-H bonds
	Large energy difference between s and p orbitals (10.60 eV) originating sp, sp 2 and sp 3 hybridization.	Smaller energy differences between the s and p orbitals (5.66 eV), lower hybridization energies. Silicon tends to use all its three p orbitals resulting in sp 3 hybridization.
	No energetically accessible d orbitals to expand its coordination or valence shell.	d orbitals (or other virtual orbitals) are energetically low alloying Si to expand its coordination sphere from four to five or six.
	Small atomic size, efficient π-π overlap, formation of double and triple bonds. C-C, C=C and C≡C bond distances are 1.54, 1.34 and 1.20 Å respectively.	Larger atomic size, poor π-π overlap Si-Si and Si=Si bond length are 2.35 and 2.16Å respectively.
	C=C results in planar structures	Si=Si nonstable in planar structure, often folded or twisted.

7 Li solid state NMR: Ex situ 7 Li solid

	state NMR measurements were done using a 4.7 T Bruker Avance
	III HD spectrometer operating at 77.7 MHz ( 7 Li Larmor frequency) with a Bruker 1.3 mm double
	resonance CPMAS probe. The mechanically synthetized Li15(Si0.5Ge0.5)4 (sample A and B) were
	measured. Additionally, two samples were obtained by cycling the self-supported Si0.5Ge0.5 film vs Li/Li +
	in a Swagelok cell, until complete discharge (C/5) followed by an OCV period (sample C) and a
	potentiostatic step at 0 V (sample D). The electrolyte was 1 M LiPF6 in EC:PC:DMC with 1% VC and
	5% FEC. The cells were disassembled in an Ar filled glove box, the electrodes were washed with DMC,

Table II . 2 .

 II2 Fitting parameters of the EXAFS portion of the MCR-ALS components in (a) k and (b) R space

	Comp.	Phase	Shell	S0 2	N	Rtheory (Å)	Rmeasured (Å)	σ 2
			Ge-Ge		4	2.449	2.461(3)	0.0052(2)
	1	Ge	Ge-Ge	0.93	12	4.000	4.03(1)	0.019(1)
			Ge-Ge		12	4.690	4.73(2)	0.022(3)
			Ge-Ge		4	2.449	2.462(5)	0.0045(6)
		Ge	Ge-Ge	0.32(5)	12	4.000	4.00(1)	0.0091(8)
	2		Ge-Ge		12	4.690	4.69(2)	0.012(2)
		LiGe	Ge-Ge Ge-Li	0.61(5)	3 5	2.554 2.64-2.99	2.57(3) 2.70(6)	0.015(8) 0.015(8)
	3	Li7Ge3 (Li9Ge4)	Ge-Ge Ge-Li	0.93	1 7	2.444 2.55-2.98	2.57(2) 2.8(1)	0.008(1) 0.020(9)
			Ge-Li		3	2.576	2.54(2)	0.007(2)
	4	Li15Ge4	Ge-Li	0.93	9	2.78-2.90	2.81(2)	0.015(2)
			Ge-Ge		5	4.56-4.64	4.54(5)	0.030(3)

  Loaiza, L. C.; Salager, E.;

	Louvain, N.; Boulaoued, A.; Iadecola, A.; Johansson, P.; Stievano, L.; Seznec, V.; Monconduit, L.
	Understanding the Lithiation/delithiation Mechanism of Si 1-x Ge x Alloys. J. Mater. Chem. A 2017, 5
	(24), 12462-12473. DOI 10.1039/C7TA02100C

Table III .

 III

	First possibility		Second possibility
	Name	Position	%Atomic concentration	Name	Position	%Atomic concentration
	C 1s	285	13.19	C 1s	285	13.21
	C 1s CO	286.46	3.25	C 1s CO	286.46	3.25
	C 1s CO3	289.57	0.37	C 1s CO3	289.57	0.37
	O 1s (CO2/3 + SiOx)	532.65	34.70	O 1s (CO2/3 + SiOx)	532.65	34.76
	O 1s (CO)	533.82	3.10	O 1s (CO)	533.82	3.10
	Si 2p Si 0	99.78	9.59	Si 2p Si 0	99.72	8.90
	Si 2p	100.41	4.79	Si 2p	100.35	4.45
	Si 2p Si +	100.63	10.48	Si 2p Si +	100.50	9.86
	Si 2p	101.26	5.24	Si 2p	101.13	4.93
	Si 2p Si 3+	102.81	10.20	Si 2p Si 3+	102.21	6.22
	Si 2p	103.44	5.10	Si 2p	102.84	3.11
				Si 2p Si 4+	103.29	5.23
				Si 2p	103.92	2.61

2. Peak deconvolution for the siloxene XPS spectra with two different options for the peak assignment and quantification

Table III . 3 .

 III3 Atomic percentages for the elements constituting siloxene determined by EDX and XPSThe EDX analysis (TableIII.3) revealed a composition of 43.99% Si, 36.77% O and traces of Cl and Ca. These values agree with the quantification from XPS analysis in which Si represented 44%, O 35.9%

	Element	Atomic %
		EDX	XPS
	C	17.34	18.7
	Si	43.99	44
	O	36.77	35.9
	Cl	1.82	1
	Ca	0.08	0.4

Table III . 4 .

 III4 Discharge capacities for siloxene cycled vs Li at different C/ratesGiven the exceptional rate capability properties of lithiation of siloxene, higher cycling rates like C/5, C, and 2C were tested for long cycling, their respective galvanostatic curves are shown in Figure III.13.

	C/Rate	Current density		Discharge capacity (mAh/g)
		(A/g)	1 st	2 nd	5 th	10 th
	C/10	0.012	3725	1969	1747	1625
	C/5	0.024	3095	1622	1502	1411
	1C	0.12	3485	1836	1693	1585
	2C	0.25	3551	1888	1684	1556

Table III . 5 .

 III5 Peak position and d value calculated for the (001) reflection from Bragg equation for the pristine and modified siloxene.

	Sample	2 θ (°)	d (Å)
	Pristine siloxene	14.02	6.31
	Eth-Siloxene	10.03	8.81
	Prop-Siloxene	9.80	9.01
	But-Siloxene	9.082	9.73
	Pent-Siloxene	8.87	9.96

Table III . 7 .

 III7 Diffraction peaks in 2θ angle for siloxene, siliganes and germanane (blue shows the hkl reflections from siloxene and purple those ones for germanane)

	Siloxene	Siligane Si0.9Ge0.1 Siligane Si0.5Ge0.5 Siligane Si0.1Ge0.9	Germanane
	14° (001)	14° (001)	18.67°	-	14.5° (002)
	22.5°	22.5°	22.5°	22.5°	-
	27.25° (100)	27.25° (100)	-	-	26.8° (100)
	28.4° (c-Si)	28.4° (c-Si)	28.4° (c-Si)	-	-
	30° (011)	30° (011)	30° (011)	30° (011)	31°
	-	-	-	31.4°	-
	35.5°	-	35.7°	35.7°	35.7°
	39.5°	39.7°	39.7°	39.7°	-
	44.6°	44.6°	44.6°	44.6°	-
	-	46.25°	46.25°	46.25° (110)	46° (110)
	47.4° (110)	47.4° (110)	47.4° (110)	48.3° (112)	48.3° (112)
	49.6°	49.6°	-	-	-
	56° (c-Si)	56° (c-Si)	56° (c-Si)	-	53.8° (200)
	57.2°	57.4°	57.4°	57.4°	-

Table IV . 1 .

 IV1 EDX analysis of the sodiated siloxene

	Element	Weigth (%)	Atomic (%)	Uncert . (%)
	Si (K)	44.03	33.45	1.0
	O (K)	38.55	51.40	1.974
	Na (K)	14.34	13.31	1.171
	Cl (K)	2.76	1.66	1.063
	Ca (K)	0.29	0.15	1.083

Table IV . 2 .

 IV2 Peak assignment and quantification for the XPS spectra peak deconvolution of a deconvolution a fresh siloxene electrode.

	Name	Position	Atomic Concentration
			(%)
	C=C	284.06	5.43
	C-C, C-H	285.05	16.65
	CO2	288.16	7.13
	CO	286.60	35.05
	CO3	289.53	0.43
	Na 1s	1071.62	1.88
	O1s CO2 (CMC)	531.39	6.19
	O1s (CO (CMC)+CO+SiOx)	533.05	25.80
	Si 2p Si 0	99.61	0.04
	Si 2p	100.24	0.02
	Si 2p Si 1+	100.50	0.02
	Si 2p	101.13	0.01
	Si 2p Si 3+	101.83	0.44
	Si 2p	102.46	0.22
	Si 2p Si 4+	103.16	0.45
	Si 2p	103.79	0.23

Table IV . 3 .

 IV3 Peak assignment and atomic concentration of the sputtered Li15Si4 reference

	Name	Position	Atomic Concentration (%)
	C 1s	285.22	4.26
	C 1s CO	286.82	2.19
	C 1s	287.40	19.96
	C 1s	283.04	1.98
	C 1s	288.84	1.20
	O 1s	530.74	3.94
	O1s	532.68	9.19
	O 1s CO	534.45	1.43
	Si 2p	98.95	3.73
	Si 2p	99.58	1.87
	Si 2p	99.62	1.50
	Si 2p	100.25	0.75
	Si 2p	100.70	1.56
	Si 2p	101.33	0.78
	Si 2p	102.10	1.21
	Si 2p	102.73	0.60
	Si 2p	103.30	0.49
	Si 2p	103.93	0.24
	Li 1s	55.70	17.81
	Li 1s	56.92	25.32

Table IV . 4 .

 IV4 Peak assignment for the6 Li NMR of the self-supported siloxene electrode cycled at different depth of lithiation/delithiation.The identification of the intermediate lithiated phases for siloxene continues with the investigation of the cells stopped at different depth of discharge/charge. Their spectra are displayed in Figure

	Name	%	Intensity I/m/NS	Name	%	Intensity I/m/NS
	1 st Discharge 0.28 V			2 nd Discharge 0.43 V
	SEI 0.8 ppm	100	25146.36	4.45	SEI 0.8 ppm 100	105065.47	8.1
	1 st Discharge 0.18 V			2 nd Discharge 0.25 V
	SEI 0.8 ppm	45.84 93104.41	3.89	SEI 0.8 ppm 100	58352.64	7.74
	LixSi 1.7 ppm Total Li	54.16 110024.16 100 203128.57	4.61 8.5		2 nd Discharge 0.10 V
	1 st Discharge 0.0 V		SEI 0.8 ppm 43.85 51331.2 LixSi 4.8 ppm 12.4 14509.89	5.9 1.66
	LixSi	64.67 151934.01	7.8	LixSi 2.2 ppm 41.84 48981.58	5.63
	SEI 0.8 ppm	31.29 73504.72	3.76	LixSi 11 ppm 1.91	2237.47
	11 ppm	4.05	9506.51	0.5	Total Li	100	117060.14 13.46
	Total Li	100	234945.21	12
		1 st Charge 0.4 V			2 nd Discharge 0.0 V
	SEI 0.8 ppm	51.06 19721.29	3.55	LixSi	46	46275.35	6.2
	LixSi 2.2 ppm	48.94 18901.24	3.42	SEI 0.8 ppm 49.29 49586.68	6.6
	Total Li	100	38622.53	6.97	Total Li	100	100601.9	12.8
		1 st Charge 2.0 V			2 nd charge 0.6 V
	SEI 0.8 ppm	100	47506.21	5.71	SEI 0.6 ppm 100	58716.83	5.39
	2 nd Discharge 0.45 V			2 nd Charge 2.0 V
	SEI 0.7 ppm	87.25 86410.08	8.41	SEI 0.8 ppm 100	21325.04	6.59
	LixSi 1.7 ppm	12.75 12627.14	1.24
		100	99037.22	9.65

IV.23, at the

  beginning of the lithiation (0.28 V) a main peak at 0.8 ppm corresponding to the formation of the SEI species is observed. As the potential decreases, this peak starts to shift towards higher frequencies and at 0.18 V it is centered at 1.7 ppm. By the end of the discharge a new peak has emerged at 4.8 ppm accompanied by a shoulder at 0.8 ppm. Both peaks at 1.7 and 4.8 ppm can be attributed to LixSi phases

Table Annex II-1

 Annex Fitting parameter for pristine Si0.5Ge0.5

	Shell	S0 2	N	Effective N (S0 2 x N)	R (theory) R (measured)	σ 2
	Ge-Ge	0.52(3)	4	2.5	2.398	2.452(5)
							0.0043(3)
	Ge-Si	0.30((3)	4	1.5	2.398	2.390(9)
	Ge-Ge	0.40(5)	6	2.9	3.916	3.92(2)
							0.016(1)
	Ge-Si	0.42(5)	6	3.1	3.916	3.91(3)

Table Annex II

 Annex -2. Fitting parameter for four reconstructed MCR components

	Component Shell	S0 2	N	Effective N (S0 2 x N)	R (theory)	R (measured)	σ 2
		Ge-Ge	0.53(3)	4	2.6	2.398	2.451(6)
								0.0043(3)
		Ge-Si	0.29(3)	4	1.4	2.398	2.39(1)
	Comp. 1
		Ge-Ge	0.40(5)	6	2.9	3.916	3.92(2)
								0.016(1)
		Ge-Si	0.42(5)	6	3.1	3.916	3.91(3)
		Ge-Ge	0.36(4)	4	2.3	2.398	2.457(7)
								0.0041(4)
		Ge-Si	0.26(3)	4	1.7	2.398	2.37(1)
	Comp. 2	Ge-Ge	0.28(6)	6	2.7	3.916	3.90(2)
								0.011(2)
		Ge-Si	0.34(6)	6	3.3	3.916	3.89(3)
		Ge-Li	0.19(5)	12	-	2.576	2.6(3)	0.02(1)
		Ge-Ge	0.08(6)	4	1.2	2.398	2.40(1)
								0.004(1)
		Ge-Si	0.21(4)	4	2.8	2.398	2.37(3)
	Comp. 3	Ge-Ge	0.11(6)	6	2.3	3.916	3.85(4)
								0.01(1)
		Ge-Si	0.18(5)	6	3.7	3.916	3.85(4)
		Ge-Li	0.53(4)	12	-	2.576	2.6(1)	0.015(8)
	Comp. 4	Ge-Li		6	-	2.576	2.452(5)	0.016(4)
			0.82(3)
		Ge-Li		6	-	2.856	2.75(3)	0.018(6)

Table Annex II

 Annex -3. Refined cell parameters for the different crystalline phases identified in an in situ XRD experiment performed on a self-supported Si0.5Ge0.5 film at C/14.

		First Discharge		First Charge
	Potential (V)	x Li	Cell Parameter (Å)	Potential (V)	x Li	Cell Parameter (Å)
		Si0.5Ge0.5			Li15(Si0.5Ge0.5)4
	2.25	0.286	5.539(5)	0.186	4.145	10.732(8)
	2.32	0.354	5.540(5)	0.293	4.073	10.730(5)
	0.934	0.426	5.539(5)	0.356	4.00	10.727(2)
	0.499	0.499	5.539(5)	0.393	3.937	10.724(2)
	0.415	0.57	5.540(5)	0.416	3.862	10.721(4)
	0.383	0.638	5.539(5)	0.427	3.790	10.719(4)
	0.356	0.713	5.530(4)	0.434	3.730	10.717(2)
	0.334	0.782	5.530(4)	0.442	3.649	10.714(5)
	0.318	0.847	5.530(4)	0.444	3.611	10.711(4)
	0.306	0.848	5.529(4)	0.451	3.501	10.709(4)
	0.297	0.995	5.529(4)	0.454	3.461	10.707(2)
	0.289	1.062	5.526(5)	0.456	3.437	10.704(3)
	0.282	1.141	5.529(4)	0.458	3.348	10.702(5)
	0.276	1.205	5.524(4)	0.462	3.243	10.670(5)
	0.27	1.286	5.524(3)	0.468	3.156	10.699(4)
	0.264	1.354	5.520(4)	0.471	3.100	10.697(5)
	0.259	1.429	5.518(3)	0.478	3.001	10.696(6)
	0.253	1.489	5.516(4)	0.483	2.92	10.693(5)
	0.248	1.555	5.515(6)	0.484	2.864	10.694(8)
	0.242	1.64	5.514(3)	0.493	2.796	10.695(8)
	0.237	1.711	5.510(3)
	0.232	1.783	5.509(6)		Second Discharge
	0.228	1.856	5.508(6)		Li15(Si0.5Ge0.5)4
	0.224	1.912	5.507(4)	0.127	3.530	10.757(4)
	0.219	1.995	5.540(6)	0.121	3.609	10.756(5)
	0.215	2.056	5.536(6)	0.114	3.685	10.754(6)
	0.211	2.131	5.535(6)	0.108	3.752	10.754(2)
	0.207	2.2	5.530(6)	0.102	3.820	10.754(5)
	0.204	2.275	5.53(1)	0.094	3.897	10.754(5)
	0.2004	2.276	5.532(9)	0.085	3.967	10.754(2)
	0.196	2.416	5.525(9)	0.072	4.042	10.740(2)
	0.192	0.286	5.529(6)	0.054	4.112	10.733(2)
	0.188	0.354	5.523(9)	0.029	4.183	10.727(2)
	0.184	0.426	5.52(2)	0.0025	4.256	10.724(2)
	0.18	0.499	5.519(9)
	0.176	0.57	5.51(1)		Second Charge
		Li15(Si0.5Ge0.5)4		Li15(Si0.5Ge0.5)4
	0.119	3.419	10.764(5)	0.196	4.201	10.722(2)
	0.115	3.489	10.760(4)	0.301	4.130	10.720(2)
	0.11	3.490	10.756(4)	0.363	4.058	10.717(2)
	0.105	3.638	10.758(3)	0.402	3.988	10.715(2)
	0.099	3.704	10.763(6)	0.424	3.915	10.712(2)
	0.091	3.782	10.764(8)	0.439	3.845	10.709(2)
	0.081	3.854	10.767(3)	0.453	3.772	10.707(2)
	0.069	3.923	10.768(2)	0.48	3.702	10.704(2)
	0.053	4.000	10.769(5)	0.616	3.629	10.705(8)
	0.035	4.071	10.772(5)	0.864	3.564	10.704(4)
	0.017	4.139	10.757(2)	0.989	3.515	10.705(4)
	3.48E-04	4.214	10.745(2)	1.5	3.319	10.704(4)

Table Annex II-4.

 Annex Values for the refined cell parameter for Li15(Si0.5Ge0.5)4, Li15Ge4 and Li15Si4. Fit parameters for the7 Li ex situ NMR spectra.

	Phase	Cell Parameter (Å)	Phase	Cell Parameter (Å)	Phase	Cell Parameter (Å)
	Li15(Si0.5Ge0.5)4		Li15Ge4		Li15Si4
	Mechanically	10.753(1)	mechanically	10.768(1)	Mechanically	10.65(8)
	synthesized		synthesized		synthesized
	Li15(Si0.5Ge0.5)4 From in situ XRD(C/14)	10.745(2)	Li14Ge4 ICSD 43235	10.72(5)	Li15Si4 ICSD 237945	10.6546
	Table Annex II-5. Component	Position (ppm)	Width (ppm)	Gaussio-	Contribution (%)
				lorentzian ratio
			Sample A
	1	35	50	0	54
	2	0	23	0	30
	3	9	15	1	7
	4	-9	9	1	5
	5	3	5	0	4
			Sample B
	1	-2	45	0	66
	2	-10	14	0	27
	3	9	8	0	4
	4	3	4	0	2
			Sample C
	1	-11	17	0	72
	2	-12	6	1	14
	3	10	13	0	9
	4	1	5	0	3
	5	-1	1	1	2
			Sample D
	1	-4	22	0	54
	2	-10	9	0	45
	3	0	2	1	1
			Li15Si4
	1	10	7	0	42
	2	16	34	0	36
	3	3	6	0	22
			Li15Ge4
	1	3	12	0	40
	2	5	33	0	33
	3	-7	14	0	22
	4	7	5	0	4

Silicon: Li15Si4 3579,[37] NaSi 954,[299] KSi 995[START_REF] Feng | [END_REF] mAh/g Germanium: Li15Ge4 1384,[37] NaGe 369,[41] KGe 369[305] mAh/g

Note that the terms silicene, siloxene, siloxane and polysilane are often misled in the literature; silicene refers to the monoatomic single layer of Si6 honeycomb rings, siloxene to the bulk OH, H terminated silicene, siloxane to compounds made essentially from Si-O-Si bonds and polysilane to the H terminated silicene.

This experiment was performed using the experimental protocols mentioned in section 2 of the present chapter.

This experiment was performed using the experimental protocols mentioned in section 2 of the present chapter.

Silicon: Li15Si4 3579,[37] NaSi 954,[299] KSi 995[START_REF] Feng | [END_REF] mAh/g Germanium: Li15Ge4 1384,[37] NaGe 369,[41] KGe 369[305] mAh/g

afin d'améliorer leur efficacité, leur durée de vie et leur résolution. Il est donc nécessaire de rechercher de nouveaux matériaux d'électrodes pour différents systèmes, Li, Na, K, etc., qui fournissent des densités d'énergie plus élevées avec un cycle de vie plus long.Dans ce contexte, le silicium et le germanium apparaissent comme de bons candidats en raison de leurs capacités théoriques élevées.6 Chaque élément présente ses avantages : le silicium est abondant dans la croûte terrestre et relativement peu coûteux, le germanium est plus cher mais présente une diffusion des ions Li et une conductivité électrique améliorées par rapport au Si. Avant la mise en oeuvre de ces matériaux au niveau industriel, plusieurs défis doivent être relevés. Les capacités élevées délivrées se font au détriment d'une expansion et d'une contraction du volume lors de l'insertion et de la désinsertion des ions alcalins. Ces changements de volume dans les particules de Si et de Ge provoquent; i) la pulvérisation des particules, l'isolation du réseau conducteur et le détachement du collecteur de courant, ii) les fissures exposent une nouvelle surface à l'électrolyte, provoquant une formation excessive et incontrôlée de couche de SEI (Solid Elecrtolyte Interface), iii) la couche importante de SEI augmente la résistance de l'électrode et piège de manière irréversible les ions Li disponibles.[START_REF] Feng | [END_REF][3][4] Donc, la cellule perd de sa capacité. L'expansion volumique est un processus inhérent à la réaction d'alliage et ne peut être évitée. Par conséquent, des stratégies permettant d'éviter la défaillance des cellules ont été décrites dans la littérature. Ces travaux comprennent l'amélioration du contact entre les particules par l'introduction de liant capable de réduire l'expansion volumique et de maintenir la stabilité structurelle. On peut rapporter, la réduction de la taille des particules à l'échelle nanométrique et la conception d'architectures spécifiques réduisant les contraintes mécaniques suivant une orientation spécifique et évitant la fissuration.[5,6] La mise en oeuvre de formulations d'électrolytes spéciales qui stabilisent la couche de SEI[8] et la conception de réseaux efficaces en composites de type Si (Ge) -graphite / graphène. L'utilisation d'alliages bénéficiant des propriétés souhaitées de deux ou plusieurs éléments. L'utilisation de structures 2D ou stratifiées offrant un espace suffisant entre les couches pour s'adapter aux changements de volume tout en offrant une cinétique de transport rapide.[9] Ce travail a proposé deux approches différentes pour relever les défis associés aux électrodes de Si et de Ge par la mise en oeuvre d'une solution solide de type Si1-xGex, améliorant la rétention de capacité et la conductivité électronique, et par l'utilisation d'une morphologie lamellaire amortissant les variations de volume. Ces phases lamellaires à base de Si et de Ge sont appelées siloxène et germanène. Elles sont dérivées des phases CaSi2 et CaGe2 Zintl. Ces dernières se caractérisent par une grande diversité structurelle et leurs propriétés peuvent être facilement modifiées. Plusieurs siliciures6 Silicium: Li15Si4 3579[37] , NaSi 954[299] , KSi 995[START_REF] Feng | [END_REF] mAh/g. Germanium: Li15Ge4 1384[37] , NaGe 369[41] , KGe 369 mAh/g[305]

à la surface des lamelles, suivie d'une phase LixSi, de composition similaire à celle de Li15Si4. La composition exacte de LixSi n'a pas été déterminée, car différents environnements du Li étaient possibles. La présence de différents intermédiaires lithiés met en évidence la sensibilité des matériaux à base de Si et de Ge aux conditions de cyclage, et la morphologie en lamelles a probablement un impact sur ce comportement. En effet, les études théoriques suggèrent qu'en raison de la présence de plans Si-Si, une intercalationde Li est préférée à une réaction d'alliage. Néanmoins, cette étude n'a pas fourni suffisamment de preuves pour prouver cette théorie et tous les résultats ne peuvent pas être totalement attribués à un mécanisme d'alliage. Ainsi, cette étude jette les bases des recherches futures, en espérant comprendre comment le silicium et le germanium se comportent différemment lorqu'ils sont sous forme lamellaire ou non, et dévoiler l'impact de la morphologie sur le mécanisme électrochimique.Figure 6. Spectres 6 Li RMN d'électrodes de siloxène auto-supportées cyclées à différentes profondeurs de décharge/charge. En conclusion, ce travail a présenté deux stratégies pour relever les défis associés aux électrodes à base de Si et de Ge, l'utilisation d'alliages qui combinent les avantages des deux matériaux, comme dans Si0.5Ge0.5, et l'utilisation d'une morphologie en lamelles qui atténue les variations de volume et

Chapter III. Layered Zintl Phases

Chapter IV. Layered Zintl Phases Electrochemical Mechanism Exploration and higher coordination number tend to resonate at lower frequencies. [110,181,191] Considering this information, the resonance at 2 ppm is likely associated with a Li with high coordination number. 6 Li NMR spectra of the 1 st and 2 nd lithiation/delithiation of a siloxene self-supported electrode.

Figure IV.22.

Figure IV

. 22 shows the 6 Li NMR spectra for the 1 st and 2 nd lithiation/delithiation of a self-supported siloxene anode. All of the spectra are mainly composed of two resonances, at 4.8 ppm and 0.7 ppm.

The second lithiation differs from the first one; while the peak positions are the same, the intensity of the 0.7 ppm resonance for the second discharge increases by a factor of 1.5. No significant differences are found between the first and second delithiation. The variation between the first and second lithiation could be originated from a difference in the electrochemistry of the cells, despite their general reproducibility, rather than an additional formation of SEI layer. Indeed, in all the cells, for the first discharge, the majority of the capacity is obtained below 0.2 V, except for the cell disassembled at the first discharge, which showed more capacity at high potential. This cell reached higher reversible capacity values with lower formation of SEI layer. This assumption is consistent with the peak deconvolution and quantification of the different components presented in Table IV.4. For the first discharge the LixSi accounts for 64% (7.8 a.u.) of the total intensity while the SEI for 32% (3.76 a.u.), a very weak signal at 11 ppm represents 4%. This last one could be attributed to the FEC reduction taking place at approximately 1.5 V and that contributes with 3-4% of the total capacity. For the second discharge, the LixSi accounts for 46% (6.2 a.u.) while the SEI for 49% (6.6 a.u.), in line with a lower discharge capacity. Despite the low formation of SEI in the first discharge (3.76 a.u.) for the first charge the amount of Li SEI species remaining was 5.71 a.u., whereas for the second cycle this change is less pronounced, from 6.6 to 6.59 a.u.

Annexes Electron Energy Loss Spectroscopy

The EELS were performed using a transmission electron microscope FEI TECNAI F20 S-TWIN fitted with a Gatan Image Filter Tridiem in post column. The measurements were done in low dose mode to minimize the effect of the electron beam on the sample. The EELS measurements were done with a 1-1.2 eV energy resolution determined by measuring the full width at half maximum of the zero-loss peak. The following conditions were used to acquire the EELS spectra; a dispersion of 0.2 eV/ch, a convergence angle of 5.8 mrad and a collection angle of 2.23 mrad. The energy loss near edge structure (ELNES) were performed inSTEM mode in defocus to limit the electron beam interaction. For the Li Kedge, the background was subtracted considering a first order log-polynomial law, the multiple scatterings were removed using the Fourier Log method and the energy correction with respect to the zero-loss peak was done. For the Si L-edge, the background was subtracted using a power law.

X-ray Photoelectron Spectroscopy.

XPS analysis was performed using an Escalab 250 Xi spectrometer with a monochromatized Al Kα radiation (hν =1486.6 eV). Prior the analysis, the samples were kept at 10-8 mbar for one night.

The analysis was performed using the standard charge compensation mode and an elliptic 325 x 650 µm X-ray beam spot. Core spectra were recorded using a 20 eV constant pass energy with a 0.15 eV step size and short time iterative scans. The quantification was done using CasaXPS software. The core peaks were calibrated from the 285 eV peak (C-C/C-H) while the fitting was performed using non-linear Shirley-type background with 70% Gaussian -30% Lorentzian Voigt peak shapes, full width at half-maximum constraint range while the quantification was done using the relative sensitivity factor provided with the Escalab machine. respectivement [274] ). Les rapports dans la littérature sur ces phases étant très rares, une caractérisation structurelle approfondie est nécessaire avant de traiter de leur utilisation en tant que matériaux de batterie. Ces informations sont précieuses pour déterminer leur composition spécifique et comprendre leur comportement électrochimique.

Annex III Supporting information Chapter III Layered Zintl phases