Tailoring SnO2 Defect States and Structure: Reviewing Bottom-Up Approaches to Control Size, Morphology, Electronic and Electrochemical Properties for Application in Batteries

Tin oxide (SnO2) is a versatile n-type semiconductor with a wide bandgap of 3.6 eV that varies as a function of its polymorph, i.e., rutile, cubic or orthorhombic. In this review, we survey the crystal and electronic structures, bandgap and defect states of SnO2. Subsequently, the significance of the defect states on the optical properties of SnO2 is overviewed. Furthermore, we examine the influence of growth methods on the morphology and phase stabilization of SnO2 for both thin-film deposition and nanoparticle synthesis. In general, thin-film growth techniques allow the stabilization of high-pressure SnO2 phases via substrate-induced strain or doping. On the other hand, sol–gel synthesis allows precipitating rutile-SnO2 nanostructures with high specific surfaces. These nanostructures display interesting electrochemical properties that are systematically examined in terms of their applicability to Li-ion battery anodes. Finally, the outlook provides the perspectives of SnO2 as a candidate material for Li-ion batteries, while addressing its sustainability.


Introduction
The transition to carbon-neutral energy production, storage and use is at the forefront of the EU's decarbonized energy transition. To that end, lithium-ion batteries (LiBs) are currently at the forefront of energy storage devices because of their high energy density, high specific energy, high output voltage, low self-discharge, wide operational temperature range and rechargeability [1][2][3][4][5]. Even though LiBs serve as an energy source for smaller electronic devices, their performance is still insufficient for certain applications, such as electric vehicles (EV), which require good cyclic performances and low self-discharge. In fact, the choice of the anode material in LIBs is essential in determining the storage capacity or energy density of the battery. Currently, transition metal oxides, e.g., Fe 2 O 3 , Co 3 O 4 , NiO and TiO 2 , and composites of graphite with Sn, Sb and Al have raised interest as new-generation electrode materials [6][7][8]. Among them, Sn-based materials (e.g., SnO 2 , SnS 2 , SnSO 4 , Cu 6 Sn 5 and Ni 3 Sn 4 ) possess a higher specific capacity and a lower potential hysteresis than other transition metal oxides [9,10]. Due to their high conductivity and transparency, SnO 2 materials are used in solar cells, catalytic supports, solid-state sensors and electrode materials for battery applications. SnO 2 is also an important n-type widebandgap (3.6 eV) semiconductor, and its most stable polymorph is the natural and abundant cassiterite ore that crystalizes in the tetragonal rutile structure (P42/mnm). Additionally, compared to the commercialized graphite anode, SnO 2 anode materials have demonstrated a much higher theoretical specific capacity of 1494 mAhg −1 against the 372 mAhg −1 of graphite [2,4,[6][7][8][9], proving their applicability in commercial anode materials. In addition, the theoretical specific capacity of SnO 2 is higher for reversible electrochemical half-cell reactions. Even though the alloying and dealloying of Li with metallic Sn are completely reversible, reduction of SnO 2 into Sn during lithiation is considered as mainly irreversible. Nevertheless, the theoretical specific capacity still remains high at~780 mAhg −1 owing to the formation of various Li x Sn compounds after the first cycle [11][12][13][14]. Both the theoretical calculations [15] and cyclic voltammetry studies [16,17] confirm the presence of these intermediate phases resulting from SnO 2 lithiation. In addition, the ternary phase diagram of Sn-O-Li [18] also indicates that both the Li 2 SnO 3 and Li 8 SnO 6 intermediate phases are likely. Lithiation of metallic Sn presents certain disadvantages, such as a large volume expansion (~200-300%) that creates internal stress at the anode and consequently leads to its degradation [9,13,19].
One way to overcome these irreversible reactions that reduce the specific capacitance of the electrode is by using nanomaterials. In fact, nanomaterials are applied to a wide range of fields, e.g., medicine [20], food [21], the environment [22], textiles [23], cosmetics [24], electronics [25] and energy [26]. As the particle size decreases, the surface-to-volume ratio increases, leading to a high specific surface. Recent studies [5,12,27] have shown that shrinking the SnO 2 nanoparticle size to less than 11 nm [28] enables the reversibility of the lithiation-delithiation processes. In fact, the large surface area in very small nanoparticles generates a large number of reactive sites for Li 2 O nucleation on the surface of SnO 2 that improves inter-diffusion kinetics. On the other hand, Sn-metal nanoparticles tend to enlarge in order to reduce their surface-to-interface energy, i.e., Gibbs free energy, leading to the decrease in the number of active sites for Li 2 O nucleation and impeding the full particle conversion of Sn to SnO 2 . In fact, Sn coarsening is likely to be the major reason for the decay in energy capacity [5,28,29]. In addition, Li 2 O coarsens simultaneously and acts as an electron insulator because of its poor electronic conductivity and, thus, obstructs electron shuttling between electrodes. Therefore, several strategies have been adopted in order to hinder Sn-nanoparticle coarsening [5]. These methods include combining SnO 2 with carbonaceous materials such as graphite [5,30], graphene [5,31,32] or carbon nanotubes [5,33]. In turn, these strategies promote reversible reactions owing to a better dispersion of nanoparticles in the matrix. While the carbon matrix usually enhances electrical conductivity, incorporating grain boundaries with a hybrid interface improves Li + insertion that then deters the coarsening of Sn nanoparticles. Another strategy consists of alloying with Co metal that generates intermetallic phases [34][35][36] with high theoretical capacities, i.e., up to 851 mAhg −1 for CoSn 3 , 796 mAhg −1 for CoSn 2 , 663 mAhg −1 for CoSn and 569 mAhg −1 for Co 3 Sn 2 [37]. SnO 2 can also be grown as hierarchical structures, such as nanorods, nanowires and nanoflowers, with a high surface-to-volume ratio. These morphologies tend to reduce the strain and coarsening of SnO 2 caused by repeated lithiation-delithiation processes owing to their high-aspect ratio and specific surface that provide ample sites for Li 2 O lithiation. Even though there are reports on their syntheses [1,2,[38][39][40][41][42], controlling the final morphology via synthesis parameters still remains a challenge. Furthermore, SnO 2 bandgap varies as a function of the crystal structure and defects, and both depend upon the synthesis conditions [2,3]. Several high-pressure polymorphs of SnO 2 possess interesting electronic properties; however, stabilizing them as single-phase free-standing nanostructures has not yet been achieved. Nevertheless, thin-film deposition methods have introduced novel ways to achieve the stabilization of these high-pressure polymorphs owing to substrateinduced strain and dopants [4,5,7,8,43]. For instance, the pure orthorhombic phase is mainly stabilized in SnO 2 thin films via chemical deposition methods resulting in an epitaxial growth induced by substrate strain. On the other hand, the pure cubic phase was only achieved by direct-current sputtering using nitrogen and antimony dopants [44].
Therefore, this review compiles SnO 2 synthesis methods and correlates them to the structural, morphological, electronic and optical properties. In fact, the bandgap, crystal structure, morphology and defects can be controlled by the synthesis conditions, such as synthesis temperature, duration, precursor and solvent. These properties are then further correlated to the electrochemical properties, such as energy capacity, redox mechanisms and cyclability of the anode materials, that are compared to commercial batteries. This paper aims at providing a reliable survey of the state of the art on the control of SnO 2 properties.

Stabilization of SnO 2 Polymorphs
The tetragonal rutile structure (P42/mnm) is the most common polymorph of SnO 2 . Other polymorphs of SnO 2 have been stabilized by various growth techniques and through the use of dopants. Furthermore, the insertion of dopants not only modifies its crystal structure but also its physical and chemical properties. However, to date, the electrochemical performance for each SnO 2 phase has not been systematically quantified for battery applications. Polymorphs of SnO 2 vary in their polyhedral stacking and are a result of pressure-induced phase transitions of SnO 2 , shown in Figure 1, starting from the rutile structure, studied by the density functional theory (DFT) [45,46]. In fact, SnO 2 undergoes a phase transition from the rutile-type tetragonal P42/mnm to the CaCl 2 -type orthorhombic Pnnm phase ( Figure 1). The Pnnm polymorph of SnO 2 usually stabilizes at a pressure of 12 GPa but can also exist in its metastable form of α-PbO 2 -type orthorhombic or the scrutinyite structure belonging to the space group of Pbcn [47]. The scrutinyite phase is under-stoichiometric in oxygen and is produced by an oxygen-vacancy-mediated transformation that not only increases the unit cell volume but also provides them with interesting physical properties, such as enhanced gas sensing. Oxygen vacancies can also be created by introducing trivalent dopants that substitute Sn in the structure, whereupon generating oxygen vacancies in the structure. For example, both Co and Mn are capable of stabilizing the scrutinyite phase, as they possess several oxidation states [43,48]. However, single-phase scrutinyite SnO 2 is difficult to stabilize and often exists as a mixture of the rutile and orthorhombic Pnnm phases. Subsequently, the transformation of scrutinyite into the pyrite-type cubic Pa3 occurs at 17 GPa. Several works claim that high-pressure cubic phases were in fact stabilized from a pressure of 21 GPa onward. Similarly, pressures as high as 48 GPa were used to stabilize the Pa3 phase with a lattice constant of a = 4.87 Å. However, theoretically, a phase transition from Pa3 to fluorite-type cubic Fm3m was obtained at a lower pressure of 24 GPa. The differences in the pyrite-and fluorite-type cubic phases lie mainly in the oxygen co-ordination. In the pyrite cubic structure, the Sn atoms are coordinated to six oxygen atoms and two more situated further away. On the other hand, in the fluorite structure, eight oxygen atoms are coordinated equidistantly from Sn. The release of pressure reverses the phase transformation to orthorhombic and then to tetragonal [49]. However, at a lower pressure of~28 GPa, a ZrO 2 -type I orthorhombic phase transformation at 18 GPa from the pyrite structure belonging to the Pbca space group is stabilized, along with a 2% increase in the volume of the unit cell. The volume expansion is due to the presence of a higher number of Sn +4 -coordinated oxygen, which increases to 7. Similarly, the fluorite-type cubic Fm3m has a Sn +4 cation coordinated to eight oxygen anions. Finally, the cotunnite-type orthorhombic phase II with a Pnam space group appears at 33 GPa with the Sn cation being coordinated to nine oxygen anions [45,49,50]. Although the formation of the orthorhombic phase usually requires high pressure, i.e., extremely energetic conditions, orthorhombic SnO 2 can nevertheless be stabilized as thin films by controlling the deposition conditions, such as temperature and pressure, or through doping. Several physical deposition techniques are reported in the literature, such as sputtering [52] and pulsed laser deposition (PLD) [53,54], in addition to chemical techniques, such as plasma-enhanced atomic layer deposition (PE-ALD) [55,56] or mist chemical vapor deposition (CVD) [57]. The addition of transition-metal ions or rare-earthion dopants has also shown promising results in the stabilization of the orthorhombic phase. In general, comparable-sized or smaller-radius metal ions usually substitute the Sn 4+ cation in the lattice. For instance, manganese (Mn 3+ radius: 0.65 Å and Mn 4+ radius: 0.54 Å) [43], zinc (Zn 2+ radius: 0.74 Å) [58] or cobalt (Co 2+ radius: 0.58 Å) [59] ions, depending on their oxidation number, possess very similar atomic radii to Sn 4+ ions (0.69 Å). Nevertheless, larger-radius ions can also stabilize the orthorhombic phase. In fact, ionic radii of Ce 4+ (0.87 Å) and Ce 3+ (1.01 Å) are much larger than the Sn 4+ ionic radius; however, via the generation of lattice disorders and structural defects, the orthorhombic phase with a molar content of 41% can be stabilized in the solid solution of Sn 0.7 Ce 0.3 O 2 [60]. Other than lattice disorders, orthorhombic phase stabilization can be obtained by doping SnO 2 with Zn 2+ , Mn 3+ , Co 2+ or Ce 3+ , they having lower oxidation states that trigger the formation of oxygen vacancies in the structure and, in turn, trigger the stabilization of other high temperature phases. For smaller cations such as Mn 3+ , the lattice distortions induced by changes in bond length are responsible for stabilization of the lower symmetry orthorhombic phase of SnO 2 [43]. However, in the case of Sb doping, since Sn and Sb have similar atomic radii, Sb substitutes Sn without modifying the crystal structure even though they have different valences [33]. The presence of Sb 3+ creates oxygen vacancies in the structure; nevertheless, a critical number of these oxygen vacancies is needed to stabilize the cubic phase.
The SnO 2 cubic structure, i.e., pyrite or fluorite types, is a higher-pressure phase than the orthorhombic ones. In order to stabilize these phases, the substitution of oxygen by nitrogen atoms in epitaxially grown SnO 2 films has been carried out [44,61]. The mechanism, once again, is based on oxygen-vacancy creation, where a N 3− anion substitutes an O 2− anion, creating an oxygen-deficient structure. Meanwhile, the co-doping with Sb 3+ cations further exacerbates the formation of oxygen vacancies [44].
Even though many studies describe the stabilization of other SnO 2 metastable phases via thin-film deposition techniques, the synthesis of orthorhombic or cubic SnO 2 phases remains a challenge. In addition, other metal oxides also show a similar behavior. For example, the stable polymorph of HfO 2 at room temperature is the monoclinic phase, but the tetragonal and cubic phases can be stabilized by varying certain synthesis parameters, e.g., doping and defects [62]. On the other hand, a reductive atmosphere at a synthesis temperature of 300 • C tends to induce oxygen vacancies, enabling the synthesis of the HfO 2 cubic phase without dopants [63]. To the best of our knowledge, there are no reports available describing the stabilization of high-pressure single-phase SnO 2 , i.e., cubic or orthorhombic at atmospheric pressure, as these usually exist as mixed phases.

Electronic Structure of SnO 2
Bulk SnO 2 has a bandgap of ∼3.6 eV; however, experimental bandgaps range from 1.7 eV to 4 eV, thereupon widening its range of applications to photovoltaics and photocatalysis [42,64,65]. Bandgap engineering is widely studied in SnO 2 , as it belongs to the family of transparent conducting oxides (TCO). Additionally, bandgaps can be controlled via parameters, such as synthesis routes and the application of a substrate-induced strain [66] for thin-film growth that simultaneously produce intrinsic defects and structural changes. First-principles calculations indicate that the bandgap can be narrowed by increasing the distortion in the SnO 6 octahedra, provoking changes in the bond length and bond angles in the unit cell [67]. The molecular-orbital bonding diagram of Figure 2a   Experimental and theoretical results have demonstrated that SnO 2 possesses both a direct and indirect band gap [42,75,76]. The allowed direct transition corresponds to a 3.68 eV direct bandgap, while the two forbidden direct transitions correspond to 3.03 eV and 3.50 eV [42]. In addition to direct bandgap transitions, there exist also indirect bandgaps corresponding to indirect transition of 2.62 eV and 2.90 eV, respectively [42]. In fact, the fundamental band gap of SnO 2 is estimated to be much lower at~3 eV [77], but certain bandto-band optical transitions are dipole forbidden, which lead to a higher optical bandgap of 3.8 eV [78]. In summary, the bandgaps, i.e., direct and indirect, are affected by the presence of defects in the volume of the SnO 2 structures and the distortion of the oxygen polyhedron owing to polymorphism, as well as surface defects generated as a result of size reduction. This implies that certain transitions that are forbidden in bulk SnO 2 may be allowed in defective or nano SnO 2 because of the breaking of the long-range ordering of the crystal lattice at the surface of the nanoparticle [79], which in turn favors the generation of bandgap states.

Bandgap Engineering in SnO 2
Direct bandgaps are systematically located at the high-symmetry Γ point. Besides, the crystalline quality of thin films related to defects and impurities has been shown to influence the bandgap [80]. The phase transition of SnO 2 to higher-pressure-induced phases also encourages a decrease in the direct bandgap, which is the result of a more compact lattice, ensuing higher orbital overlapping [81]. In agreement with these observations, Table 1 compiles the structural and electronic properties of all the SnO 2 polymorphs. However, experimentally, the synthesis and stabilization of higher-symmetry SnO 2 polymorphs is complicated. Few studies report the presence of other SnO 2 phases for Fe-doped SnO 2 , owing to the substitution of Sn by Fe ions [82], where the rutile phase co-exists with a small amount of α-PbO 2 (Pbcn) secondary phase. After annealing at 800 • C, they observed that only traces of the SnO 2 orthorhombic phase remained, which confirms the low stability of the orthorhombic phase at high temperatures compared to the SnO 2 rutile structure. In addition, the possible presence of iron oxide lowers the SnO 2 band gap to 2-3 eV [83]. In fact, according to the crystal field theory, Fe 3+ ions are placed in an octahedral configuration in the presence of a weak field ligand (oxide) in a high-spin configuration. In addition, Fe 3+ ions possess an ionic radius (0.645 Å) that is slightly smaller than Sn 4+ ions (0.69 Å), which decreases the lattice parameters and, consequently, increases orbital overlapping between Sn 4+ and O 2− ions, leading to a lower bandgap. Radaf et al. succeeded in stabilizing the orthorhombic SnO 2 structure by adding Cr 3+ dopant [84]. The crystallite size decreased with the increase in Cr concentration, and the bandgap consequently decreased from 3.6 eV for the undoped SnO 2 thin film to 3.28 eV with 5% of Cr. While the addition of those metals leads to the stabilization of the orthorhombic phase, Keskenler et al. [85] have demonstrated that W incorporation also stabilizes the Pbca cubic phase until a doping threshold of 2.0 at. %. Here, W 6+ is likely to substitute Sn 4+ , which shrinks the lattice owing to the lower ionic radius. When the W concentration exceeds 2.0 at. %, lower oxidation states of tungsten could also substitute Sn 4+ sites, which counteract the unit-cell shrinkage [85]. This can be explained by the Moss-Burstein effect, where materials with high carrier concentration, such as W, fill unoccupied states deep within the conduction band. Consequently, the Fermi level of the n-type SnO 2 shifts into the conduction band. The increase in the optical bandgap is due to the excited electrons transitioning from the valence band to empty states in the conduction band localized at higher energy levels [85,86]. Table 1. Crystal structure of SnO 2 polymorphs, volume of the unit cell (Å 3 ), its density (in g.cm −3 ) and direct bandgaps (eV).

Point-Defect Engineering in SnO 2
The n-type conductivity of undoped rutile SnO 2 materials can be explained by the defects present in the structure. Among the four different intrinsic defects, i.e., oxygen vacancy V O , tin interstitial Sn i , tin antisite Sn O and oxygen interstitial O i , the predominant and combined occurrence of V O and Sn i leads to electron donor properties [50,89]. SnO 2 nanostructures exist in diverse morphologies (e.g., nanorods, nanocubes, nanosheets, nanowire and nanospheres) as a result of the synthesis route. Interstitial and vacancymediated defects are specific to each morphology, as the shape and size of the nanoparticle influence the surface and volume defects generated [90,91]. Hence, engineering SnO 2 nanoparticles via controlled synthesis conditions allows the tailoring of their size, shape, morphology, intrinsic and surface defects. These properties play an important role in their electrochemical properties and redox mechanisms, especially for LiB applications. Defect engineering in semiconductors, more particularly in nanomaterials, is important for several applications. In fact, surface defects in nanomaterials are capital for surface-related phenomena in catalysis. Surface-defect engineering of SnO 2 has already been studied for photocatalytic [92] and gas sensing [93,94] applications. Furthermore, oxygen-related defects generated in an oxygen-poor environment create exposed Sn 4+ cations, as well as oxygen vacancies at the surface leading to abundant reactive sites [95]. In general, point defects such as surface-oxygen vacancies are common in nanomaterials owing to the high surface-to-volume ratio [96,97]. These defects are tailored via synthesis conditions, i.e., oxygen-rich or oxygen-poor conditions, synthesis temperature and annealing atmospheres. In addition, synthesizing faceted nanoparticles and exposing certain crystal facets to enhance catalytic activity are important topics in catalysis [98]. Furthermore, doping with foreign atoms to create V O and V Sn , as explained before, stabilizes higher-symmetry polymorphs through the production of oxygen vacancies in the structure. On the other hand, in optoelectronic and electronic applications, passivating surface defects is necessary for enhancing the conductivity of SnO 2 , as in the case of F-doped SnO 2 [99]. The principle is to eliminate defect states within the bandgap of the material by reducing these surface traps and, consequently, increasing the charge mobility and conductivity of SnO 2 . Since an oxygen anion is doubly ionized, the depletion of oxygen leads to a general enhancement of the charge-carrier concentration. Stoichiometric defect mechanisms do not interfere with electronic properties of SnO 2 nanoparticles, unlike nonstoichiometric defects. Surface defects in bulk materials have an insignificant influence on their physical and chemical properties because of their low proportion. However, surface defects in nanomaterials can drastically change catalytic and electronic properties, as a result of their high surface-to-volume ratio. •• + Sn i ⁄⁄ , giving rise to the n-type conductivity of SnO 2 with electron mobility from the Sn Sn ⁄⁄ to Sn Sn ⁄⁄⁄⁄ sites. The mechanism for n-type conductivity involves the hybridization of the Sn-5s and O-2p states near such vacancies, facilitating electron transfer from the valence to the conduction band [101]. The only possibility to obtain p-type conductivity is via the introduction of V Sn ⁄⁄⁄⁄ + 4 h. In the case of Sn-deficient SnO 2 , V Sn are the predominant defects that are created. An opposite reaction occurs where four electrons are taken from the valence band to form holes in order to generate interstitial site Sn Sn ⁄⁄⁄⁄ . The production of Sn vacancies can be mediated in Sn-poor conditions or by doping SnO 2 with tri-valent elements substituting the Sn Sn × that create V Sn [102,103]. Simultaneous doping with elements, such as N, creates acceptor states that then facilitate p-type conductivity. The computational and experimental studies on co-doping suggest the replacement of approximately four Sn atoms by four Al atoms and one O atom by one N atom [104]. In the case of metal excess, the defect equation governed by this mechanism involves the formation of Sn interstitial atoms, Sn i x , which further act as electron donors and can be successively doubly ionized to Sn i •• or quadruply ionized to Sn i •••• . These doubly ionized Sn 2+ states can act as traps that restrict the possibility of transition from the conduction band minimum to holes just above the valence band maximum. Therefore, passivation of the Sn i •••• on the surface of the SnO 2 nanoparticles tends to enhance the oxygen-vacancy-related transitions. Lastly, in oxygen-rich conditions, O i ⁄⁄ have the lowest formation energy and are therefore abundant.
Among all these point defects described above, oxygen vacancies caused by oxygenpoor conditions are the most abundant intrinsic defects occurring in SnO 2 nanomaterials because of the lowest formation enthalpy [100]. Moreover, many studies [105][106][107][108][109][110] have probed these new energy levels via photoluminescence (PL) spectroscopy. As previously mentioned, unstable V O × vacancy acts as a donor level and is located at 0.03 eV, just under the conduction band. In addition, ionized V O • is also considered a shallow donor, as it is located 0.15 eV below the conduction band, while V O •• is an acceptor level located at 1.4 eV above the valence band [109,110]. In SnO 2 nanomaterials, the surface-oxygen vacancy is doubly ionized (or V O •• ) and is the most dominant emission [111]. Wang et al. have investigated the photoluminescence mechanisms under a 255 nm excitation wavelength, resulting in band-to-band and defect excitations. Each peak was successfully identified and energy levels in the band diagram of SnO 2 also corroborate them. For example, the transition between the V O x donor level to the V O •• acceptor level is attributed to the 467 nm (2.65 eV) emission peak, whereas the electron transition from V O • to the valence band can be assigned to the 439 nm (2.83 eV) emission peak. In addition, Sn i x is a shallow donor, as Sn interstitials tend to occur exclusively in the +4 state located very near the conduction band, contributing to the n-type semiconductor properties of SnO 2 , even though Sn interstitials are not abundant. The band-to-band transition is identified by the 328 nm emission, corresponding to an energy of 3.78 eV. Habte et al. [112] have demonstrated that the addition of Zn 2+ cations to the SnO 2 lattice leads to PL emission peak shifts, shown in Figure 3d,e, toward lower energies. There could be two reasons for the optical bandgap reduction. Since Zn 2+ cations are smaller, their insertion should promote orbital overlapping because of a reduction in the lattice parameter. The other reason could be the shift toward longer wavelengths corresponding to the ZnO bandgap (3.37 eV). However, they highlighted that the lattice structure remains unchanged with Zn 2+ ; therefore, the decrease in the optical bandgap can be attributed to the presence of Zn-O complexes. Salem et al. [113] have observed similar changes in the bandgap with Ni-doped ZnO. Nevertheless, the addition of Zn 2+ should also enhance emission peak intensities, since the substitution of a smaller and lower valency cation encourages the formation of oxygen vacancies.
Since SnO 2 is an n-type intrinsic semiconductor, the most prominent defects are, therefore, V O and Sn i because of the lowest formation enthalpy. They are present in the volume of the material and contribute to the electronic conductivity. In nanomaterials, these defects are present on the surface and are instrumental in several catalytic reactions, including oxygen evolution reaction, hydrogen evolution reaction, gas sensing or electrocatalytic CO 2 reduction [114]. On the other hand, for applications in electronic devices, these surface states are detrimental to the device's functional properties. Photoluminescence spectroscopy is commonly used to identify these defects by providing information on optical transition between defect levels and band edges [115]. Depending on the application, these surface defects need to be either passivated or exacerbated. The importance of doping SnO 2 with acceptors lies in the possibility of obtaining a p-type semiconductor that would eventually lead to a SnO 2 homojunction diode. In general, surface defects act as trap states that enhance defect-level emission from the bandgap states. Furthermore, these defect states also extend the photo absorption of the materials to the visible region. Several studies report the stabilization of different phases of SnO 2 as thin films through the optimization of growth conditions. In fact, epitaxial growth can promote the stabilization of high-pressure phases of SnO 2 through substrate-induced strain [66]. Table 2 resumes SnO 2 thin-film growth parameters by physical and chemical vapor deposition techniques along with the phase stabilized. Physical vapor deposition (PVD) techniques demonstrate some advantages over chemical deposition techniques as different polymorphs of SnO 2 as thin films can be grown through PVD more easily. For PLD, the most important parameter is the deposition temperature [53,54,116]. At low temperatures (~150-300 • C), thin films are mainly amorphous because of low atomic mobility [54,116]. At a higher temperature (700 • C), orthorhombic and tetragonal phases coexist; however, at very high temperatures (1150 • C), the atomic diffusion is extremely high and SnO 2 therefore tends to rearrange itself in its most stable structure, i.e., rutile [54]. Similarly, for DC and RF sputtering, authors demonstrated that it is possible to stabilize the rutile, orthorhombic or cubic phases depending on the synthesis parameters. Ham et al. [117] succeeded in growing a polycrystalline thin film with a co-existence of the orthorhombic and tetragonal phases. A small amount of rutile SnO 2 at the interface diminishes the mismatch strain to 0.42%, and in turn, the remaining substrate-induced strain enables a heteroepitaxial growth of orthorhombic SnO 2 on the c-plane of sapphire. The high-resolution TEM image in Figure 3a reveals a smooth interface between the substrate and the film oriented [11][12][13][14][15][16][17][18][19][20] and [001], respectively, implying that the growth direction of tetragonal-phase SnO 2 is <100>. Whereas the orthorhombic phase can be obtained by sputtering at a high temperature [117,118], the cubic phase is mainly obtained by substituting O atoms by N atoms with N 2 gas [44,61,119]. Similar to other metal oxides, such as CeO 2 [120], ZrO 2 [121] and HfO 2 [122], SnO 2 can be doped with nitrogen, where the substitution of O by N atoms involves an increase in ordered oxygen vacancies within the crystal structure. By virtue of the epitaxial strain at the interface of SnO 2 and the substrate, crystallization of the orthorhombic or cubic phase can occur. However, at exceedingly high temperatures, the stabilization of the orthorhombic or cubic phase is hindered and the tetragonal rutile phase, where the Gibbs free energy is the lowest, is promoted [54]. Only a few reports of chemical deposition techniques describing the stabilization of the orthorhombic SnO 2 phase are available. Bae et al. [57] have successfully grown the tetragonal and orthorhombic structures via mist-CVD, using two different solvents, i.e., methanol and acetone. The difference in the boiling point of these two solvents favors the stabilization of one phase over the other. As acetone has a lower boiling point, it would supply oxygen atoms more readily than methanol, leading to oxygen-rich conditions. Their results from DFT calculations were also consistent with the experimental results [57], and they concluded that the SnO 2 orthorhombic phase is thermodynamically more favorable under Sn-and O-rich conditions. Another deposition using plasma-enhanced atomiclayer deposition for epitaxial growth of the orthorhombic phase by Kim et al. [55] on an yttrium-stabilized zirconia (YSZ) substrate ( Figure 3b) with a lattice mismatch of 2% or less has been realized. Furthermore, the deposition of SnO 2 thin films using metal-organic chemical vapor deposition techniques has also been investigated on different substrates. Deposition on both sapphire and MgF 2 (Figure 3c) substrates leads to the tetragonal rutile SnO 2 structure [55,123,127]. On the other hand, the orthorhombic phase is the result of the deposition at similar temperatures (around 500 • C) but on different substrates, i.e., 6H-SiC (Figure 3d) or yttrium-stabilized zirconia (YSZ) (Figure 3e) [124,125]. Furthermore, in that work, Sb doping was responsible for the stabilization of the orthorhombic structure. However, phase stabilization is unlikely because of Sb doping alone, which creates oxygen vacancies, but is also the result of substrate orientation, enabling the epitaxial growth of  (Figure 3f) can be controlled by varying the lattice mismatch between the substrate and the film through substrate orientation. Furthermore, Kong et al. report that the [100] interplanar spacing of orthorhombic SnO 2 film and YSZ substrate are comparable, with a lattice mismatch between c and a lattice parameter of SnO 2 and YSZ being equal to 1.3% [125], which allows epitaxial thin-film growth. Chemical deposition methods

Hydrothermal Synthesis of SnO 2 Nanomaterials
The structure-property relationship in SnO 2 nanomaterials is controlled by the synthesis conditions. This signifies that the bandgap, size of the nanoparticle, morphology and chemical properties can be tuned through the synthesis conditions. Gavaskar et al. [133] have synthesized nanoparticles of SnO 2 with the rutile structure using a SnCl 4 ·5H 2 O precursor in ethanol at 200 • C for 20 h. They obtained quasi-spherical nanoparticles with diameters ranging from 50 to 90 nm and with a direct bandgap (3.35 eV) narrower than most SnO 2 bandgaps. NaOH is also often used as an oxygen or hydroxyl source, since SnO 2 crystals grow from stannate Sn(OH) 6 2− that then condense to SnO 2 and water [134][135][136]. Vuong et al. [134] added an aqueous solution of NaOH to SnCl 4 ·5H 2 O dissolved in ethanol, which led to the growth of nanoflowerlike structures in Figure 4a under equivalent synthesis conditions. In addition, SnO 2 anisotropic crystal growth depends on the diffusion of Sn(OH) 6 2− at the surface of preferred orientations of the rutile crystal structure, which occurs after a sufficient amount of time (24 h) and temperature (190 • C). Guan et al. [135] realized complementary studies, in which they investigated the influence of the ratio of the SnCl 4 ·5H 2 O precursor to NaOH on the morphology and structure. At a high ratio of NaOH to SnCl 4 , the nanoparticles are larger compared to low NaOH concentrations. They also explained that a high NaOH concentration promotes the growth of SnO 2 along preferential crystallographic directions, leading to the flowerlike SnO 2 nanorod bundles in Figure 4b [135]. Despite having a completely different shape and crystallite size, direct bandgap energies still range from 3.68 eV to 3.72 eV at both low and high NaOH:SnCl 4 ratios, respectively. The addition of a surfactant, such as PEG [137], has been shown to further improve the previous synthesis protocol by precipitating the uniform flowerlike SnO 2 nanorod bundles in Figure 4c. Cao et al. confirmed the production of nanoflowers under a high NaOH:SnCl 4 ratio and reduced the synthesis time from 24 h to 12 h [139]. By using sodium dodecyl sulfate catalyst and Na 2 SnO 3 precursor, they also succeeded in synthesizing SnO 2 nanocubes by shortening the reaction time to only 4 h, whereupon limiting the growth to preferential crystallographic directions. Runa et al. [140] also managed to synthesize SnO 2 nanocubes using SnCl 4 precursor with urea as a surfactant and absolute ethanol as a solvent in an acidic solution. Similar temperatures and time were employed, but an acidic medium and short synthesis time led to a cube morphology. These conditions limit the growth along the thermodynamically preferential [001] direction of SnO 2 nanorods and lead to a pseudo-cubic shape. Similarly, Xi et al. [138] also synthesized SnO 2 nanocubes, as shown in Figure 4d, from a SnCl 4 precursor with urea and HCl. They suggest that hydroxyl ions from the ammonia-water reaction are in equilibrium with ammonium and hydroxide, thus favoring uniform growth morphologies.

Prospects of SnO 2 Nanomaterials as Anode Materials in LiB: Correlating Their Morphology Obtained from Synthesis Routes to Their Electrochemical Performance
Nanocomposites of SnO 2 have gained attention as anode materials for LiB applications because of their high theoretical specific capacity. It represents one of the most promising anode materials because of its high energy capacity, low cost and high energy density. Among common metal oxide compounds used as anodes in LiB [141], SnO 2 exhibits the highest theoretical energy capacity (1494 mAhg −1 ), against 890 mAhg −1 for Co 3 O 4 [142], 1007 mAhg −1 for Fe 2 O 3 [143] and 1230 mAhg −1 for MnO 2 [144]. In addition to higher energy capacity and density, SnO 2 possesses a low overall potential, i.e., charge and discharge voltages of 0.3 V and 0.5 V vs. Li [147]. SnO 2 nanomaterials could solve issues faced by other metal oxides related to lithium alloying, which leads to irreversible capacity reduction and volume changes [17]. A commercial battery is assessed by different criteria, such as energy density (amount of energy that can be stored per unit mass of the battery), battery power (rate at which the electrical current can be moved through the battery), cycle life (number of charge-discharge cycles until its performance drops), watt hours (amount of power deliverable in an hour), charge speed and resistance or impedance. To solve challenges concerning energy storage systems and their sustainability, measurable quantities are crucial. While energy capacity is a direct indicator of the energy storage performance of the anode, the coulombic efficiency quantifies the reversibility and the efficiency of electron transition from the anode to the cathode over a cycle. Capacity retention evaluates the cycle-life performances of a battery, as it provides the ratio between the discharge energy capacity of successive cycles to the initial one [148].
Hu et al. demonstrated that a spherical nanoparticle with a size under 11 nm is required for a completely reversible lithiation-delithiation reaction [28]. Table 3 summarizes different synthesis routes of SnO 2 nanostructures and nanocomposites with their electrochemical performances. Hydrothermal methods are a common way to synthesize SnO 2 nanoparticles because of their low cost and simple synthesis protocols that can be easily scaled up. Yin et al. [149] developed a hydrothermal method based on the utilization of HCl, SnCl 4 ·5H 2 O and ammonia in an aqueous solution heated at 160 • C for 30 min that produced SnO 2 nanoparticles with sizes ranging from 9 to 21 nm. These nanoparticles display an irreversible discharge capacity of 22.8% after 50 cycles (from 1196.6 mAhg −1 to 217.0 mAhg −1 ) with a current density of 100 mAg −1 (Figure 5a,d). Although the short synthesis time enables the growth of small nanoparticles, aqueous synthesis does not allow proper control of growth kinetics. On the other hand, nonaqueous synthesis methods are usually preferred for the production of small-sized metal-oxide nanoparticles, as they offer better control of the synthesis and tailor particle shape and size. Etacheri et al. [150] also prepared ultrathin SnO 2 nanoparticles by reflux in an aqueous solution. High-capacity retention is obtained after calcination of the SnO 2 nanopowder post-synthesis, referred to as "ordered interconnected SnO 2 nanoparticles" in Figure 6a vs. pristine materials called "disordered SnO 2 nanoparticles". Whereas heat-treated SnO 2 nanoparticles demonstrate a high-capacity retention (81.9%) and specific capacity (∼500 mAhg −1 ) after 100 cycles, pristine SnO 2 nanoparticles exhibit poor electrochemical properties, i.e., 35.2% of capacity retention and a specific capacity of about 92 mAhg −1 . Both SnO 2 nanostructures are agglomerated, but the reason behind the enhancement of their electrochemical properties dwells in the nature of the agglomerates. In fact, aqueous synthesis tends to produce uncon-trolled agglomerates with numerous and large electro-inactive clusters, while calcination rearranges the structure into a porous network because of the release of hydroxyl groups. Porous networks are commonly tested for battery applications, such as sub-microtube [151], hollow microspheres [152] and porous nanotube [153] structures (Figure 6b-d), because of their potentially high electrochemical properties. Therefore, reaction reversibility is possible by reducing the lithium-ion diffusion pathway by creating a highly porous structure made of nanorods, nanoflakes, nanobelts and nanosheets [154]. Narsimulu et al. [155] prepared~10 nm thick SnO 2 nanosheets by a microwave-assisted synthesis method using a stannic chloride precursor and citric acid in an aqueous medium. In theory, the thin SnO 2 nanosheets, along with their porous nature, should enable high electrochemical reversibility; however, at a current density of 100 mAg −1 , the discharge energy density decreases from 1350 mAhg −1 to 257.8 mAhg −1 after 50 cycles (Figure 6b,e), representing a capacity retention of only 19.1%. In fact, the ordered interconnected SnO 2 network mentioned before by Etacheri et al. [150] possesses a four-times larger active surface area (204 m 2 g −1 ) than the nanosheets (59.28 m 2 g −1 ). Nano-Köhler theory [156] refers to the activation of inorganic cluster growth by spontaneous condensation [157], but in the case of a prolonged hydrothermal synthesis, it tends to promote a longer coagulation time and, thus, leads to larger nanostructures, such as nanoflower bundles [158] or nanorod arrays [159]. Wen et al. [158] have synthesized flowerlike structures similar to Guan et al. [135] and Cao et al. [139] using NaOH and SnCl 4 ·5H 2 O precursors in aqueous media. However, since cycling performances are carried out at a 10-times lower current density, i.e., 78 mAg −1 instead of 782 mAg −1 , the initial and final energy capacities are not directly comparable. Liu et al. [159] have grown organized nanorods on an Fe plate with similar conditions, using stannic chloride and NaOH precursors heated at 200 • C (vs. 180 • C) for 24 h in an autoclave. In this case, long nanorods with dimensions of 60 nm × 670 nm were produced that exhibited greater initial energy charge and discharge capacities of 1128 mAhg −1 and 1918 mAhg −1 , respectively. On the other hand, for the same current density, nanoflower bundles [158] display charge and discharge energy capacity of 815 mAhg −1 and 1673 mAhg −1 , respectively. Higher electrochemical performances are again due to the nanoarrays presenting a larger organized network facilitating Li + ion diffusion. Therefore, reaction reversibility is possible on reducing the lithium-ion diffusion pathway, such as in highly porous structures made of nanorods, nanoflakes, nanobelts, nanosheets and hallow nanospheres (Figure 5d-f) [144]. While the synthesis of 0D nanostructures generally requires a SnCl 4 precursor, hierarchical nanostructures are synthesized by using lower oxidation states, Sn(II), of the SnCl 2 precursor. Sharma et al. [160] and Ding et al. [161] have, respectively, synthesized 1D SnO 2 nanowires (Figure 5d) and 3D hollow nanospheres (Figure 5f) using the template-assisted synthesis route. These nanostructures tend to retain their energy capacity for a higher number of cycles compared to free-standing SnO 2 nanoparticles. Wu et al. [162] and Narsimulu et al. [155] have both, respectively, prepared~35 nm (Figure 6e) and~10 nm thick SnO 2 nanosheets by a hydrothermal method. Both types of nanosheets exhibited good cycling capacity and charge retention. Wiley [163], (e) 2D SnO 2 nanosheet, 2012 Elsevier [164] and (f) 3D SnO 2 hollow nanospheres, 2010 American Chemical Society [161]. Insets correspond to their respective SEM or TEM images. The coupling of carbonaceous materials with SnO 2 nanoparticles has been widely studied because of the possible enhancement of cycling capacity and conductivity through their combination. Liu et al. [165] recently illustrated the efficiency of carbon nanotubes and carbon nanotube hairballs coupled with SnO 2 nanoparticles. These carbonaceous materials and a tin precursor were mixed and stirred in an absolute ethanol (nonaqueous) medium before heating at 150 • C for 10 h. Once the product was washed and dried, the resulting powder was calcinated at 360 • C for 10 min. The initial specific discharge capacity tripled (from 768.1 mAhg −1 to 2255.2 mAhg −1 ) when coupled with carbon nanotube hairballs. In addition to higher initial charge and discharge capacities, capacity retention also increases after 100 cycles at 200 mAhg −1 from 29.9% (244.8 mAhg −1 ) to 74.2% (809.2 mAhg −1 ), highlighting the high contribution of carbon nanotubes to anode performance. Deng et al. [166] have compiled a large number of syntheses involving SnO 2 nanostructures coupled with graphene and have evaluated their respective electrochemical performances. The study concludes that SnO 2 -graphene nanocomposites are promising as they could significantly improve electrochemical and electrical properties, even though some issues persist, such as the high cost of graphene fabrication and the large irreversible capacity loss of the initial cycle.
Another solution to improve SnO 2 properties is through doping with transition metals to enhance the chemical, defect and structural properties of SnO 2 . Mueller et al. [167] measured the electrochemical properties of Fe-doped SnO 2 nanoparticles in a carbon matrix. Fe-doped SnO 2 nanoparticles were synthesized via hydrothermal synthesis from tin acetate, sucrose, acetic acid and iron (II) gluconate at 150 • C for 10 h. The doping of SnO 2 nanoparticles with smaller Fe cations decreases the cassiterite lattice structure through the incorporation of oxygen vacancies, which results in the production of smaller nanoparticles (7 to 8 nm vs. 15 nm). Mueller et al. [167] also illustrate that the Fe-ion doping creates smaller nanoparticles with surface defects. The utilization of a carbon matrix coupled with Fe doping doubles the energy capacity, i.e., from 764 mAhg −1 to 1519 mAhg −1 after 10 cycles at a current density of 50 mAhg −1 . This is likely due to the enhanced conductivity of the carbon matrix allowing a better ionic diffusion. In addition, Ma et al. [168] have distinctly quantified the contribution of Co doping on SnO 2 nanoparticles and the carbon matrix to the system, separately. This synthesis differs from the usual hydrothermal synthesis carried out inside an autoclave. They applied a one-pot synthesis route at 180 • C until complete evaporation occurred, followed by heat treatment at 450 • C for 3 h, which lead to the production of ultra-small nanoparticles doped with Co less than 10 nm in diameter. After 50 cycles at 100 mAg −1 , Co-doped SnO 2 nanoparticles show a reversible capacity of 493 mAhg −1 (Figure 5c,f), which is higher than undoped SnO 2 nanoparticles having a reversible capacity of 242 mAhg −1 . The reversible capacity was then significantly enhanced with the addition of the carbon matrix, where the specific energy capacity remains above 1000 mAhg −1 . Ou et al. [169] investigated the properties of Ni-doped SnO 2 nanoparticles synthesized using urea as a surfactant. The protocol requires a SiO 2 nanosphere template, on the surface of which a thin layer of SnO 2 nanoparticles is deposited using urea as a surfactant, hindering the excess growth of SnO 2 despite the long synthesis time of 36 h at 170 • C. The SiO 2 template is then removed via HCl etching before being calcinated at 400 • C for 4 h under an Ar atmosphere. Although the initial capacity loss is still prevalent after 300 cycles, they showed that Ni doping enhances the specific discharge capacity by more than 600%. SnO 2 metal-oxide nanocomposites demonstrate significant improvement in the overall electrochemical properties. W. Zhou et al. [170] synthesized SnO 2 -Fe 2 O 3 nanowire composites that possess a two-times higher energy density than SnO 2 nanowires. The high-aspect ratio and the compatibility in the electronic structure are responsible for the enhancement. Wang et al. successfully synthesized SnO 2graphene-oxide-Co 3 O 4 nanocomposites that exhibit long cycling stability (641 mAhg −1 at 1000 mAg −1 after 900 cycles) with complete reversibility (CR = 100%, 1038 mAhg −1 ) after 100 cycles at a lower energy density (100 mAg −1 ) [171]. Both SnO 2 -graphene and SnO 2graphene-metal-oxide nanocomposites tend to significantly improve the electrochemical and electrical properties of SnO 2 .
The synthesis of SnO 2 -based nanoparticles is mainly carried out by hydrothermal synthesis routes because of the ability of operating at low synthesis temperatures, including room temperature. Nevertheless, in order to eliminate reaction by-products and improve the properties of the as-synthesized SnO 2 , post-synthesis annealing is usually required in the range of 350-500 • C. Annealing can also induce the formation of porous structures and networks with higher specific surfaces. Other methods such as template-assisted syntheses are also applied to create SnO 2 porous network nanostructures. In fact, porous connected networks such as SnO 2 nanotubes provide pathways for Li diffusion, which increases the overall charge retention. In addition, the coupling of SnO 2 nanostructures with carbonaceous materials, especially graphene, which is a well-known strategy to enhance the electrochemical properties of SnO 2 nanomaterials, is still limited by the fabrication cost of graphene. For other hierarchical structures, such as nanorods, nanoflowers and nanosheets, an overall enhancement in the electrochemical properties is also observed. Additionally, using dopants in these syntheses, such as Fe, Co and Ni, clearly enhance charge retention and coulombic efficiency. Furthermore, when doped SnO 2 is combined with carbon-based nanomaterials, its charge capacity is at least doubled. All these strategies starting from simple hydrothermal SnO 2 synthesis followed by doping and combining with nanocarbons are important technological steps toward increasing the electrochemical properties of SnO 2 . Morphologies exhibiting high surface areas, such as nanoflowers, 3D porous nanostructures or any three-dimensional nanomaterials can promote lithium diffusion within the network owing to the significantly increased surface area. However, 1D and 2D nanomaterials are generally preferred because of their isotropic configuration that enables the stacking of nanomaterials while delivering free active sites.

Conclusions and Outlook
SnO 2 is a versatile material whose bandgap can be modified by several strategies, including higher symmetry polymorph stabilization, doping and formation of defects in the structure through the synthesis process. Since the SnO 2 compound is a potential candidate for applications in energy storage, studying SnO 2 phase stabilization is, therefore, of interest. Amongst the available strategies, it has been shown that metal dopants can affect the morphology, decrease lattice parameters and simultaneously create oxygen-related defects. The use of metal dopants also tends to enhance the energy capacity, and when combined with a carbon matrix, e.g., graphene, graphite or carbon nanotubes, the reversibility of the reaction can be improved. In order to significantly improve the electrochemical properties of SnO 2 nanostructures, the stabilization of higher symmetry polymorphs, i.e., high-pressureinduced phases, such as orthorhombic or cubic, is being intensively investigated. However, due to the difficulties encountered in stabilizing these phases, the electrochemical properties cannot be systematically probed. Nevertheless, the use of dopants such as nitrogen or metals to produce oxygen vacancies appears to be the most promising solution to stabilize the orthorhombic or cubic structures. However, the doping of SnO 2 nanoparticles is still unable to stabilize single-phase SnO 2 , and several polymorphs tend to co-precipitate. However, in epitaxial thin films, the substrate-induced strain is capable of surpassing the activation free-energy barrier and leading to the stabilization of single-phase SnO 2 .
Today, there are several challenges related to introducing new nanomaterials for LiB electrode application. In particular, for the SnO 2 compound, the main drawbacks, such as the huge initial capacity loss and extensive capacity fading after prolonged cycling, make their commercial breakthrough challenging compared to the well-established and omnipresent graphite. These drawbacks are mainly due to the volume expansion during charge-discharge processes that are still unaccounted for, despite consistent progress in the field. These stresses induced by the volume expansion have been alleviated to some extent by the addition of carbonaceous materials or by creating hollow structures that allow slightly higher flexibility for volume expansion. Therefore, hollow and hierarchical nanostructures are at the forefront of SnO 2 research. Tailoring SnO 2 nanostructures into desired morphologies is also a strategy to increase the active surface area, which in turn enhances the Li+ activity. In that regard, it has been shown that optimal electrochemical performances are obtained using ultra-small nanoparticles with a highly active specific surface. However, the scaling up of SnO 2 synthesis processes in order to meet LiB demands needs to be cost effective and is one of the technical challenges to solve. Even if these aforementioned issues are surmountable, the technology transfer of nanomaterials to battery technologies involves many logistic issues. In fact, the weight of the active anode materials can attain 16% of the total battery weight in an electric vehicle, representing over 70 kg only for the anode material. Although SnO 2 nanomaterials appear promising for increasing the general lifespan of a battery, their environmental impact is only succinctly addressed in the literature. The end-of-life recycling of materials has been shown to release free-standing ultrafine nanomaterials into the environment during the shredding process. Therefore, combining SnO 2 nanoparticles with carbonaceous materials, including the presently employed graphite, could curb their release into the environment, in addition to improving the electrochemical properties of batteries. However, for a successful technology transfer, the sustainability of SnO 2 is an important issue. In 2021, the production of Sn exceeded 370,000 tons. Being the 49th most-abundant element on Earth implies that Sn is quite scarce, and this would eventually lead to massive environmental and sustainability issues. To that end, growing nanostructures of SnO 2 with large surface-to-volume ratios will be an advantage through the reduction in the volume of the material. This increases the sustainability of the SnO 2 compound in LiB owing to the incorporation of lower quantities of Sn in the nanostructures.
Battery energy storage systems (BESS) are an important part of the net-zero energy transition. They have a wide range of power and storage capacities for both small-scale devices, including mobile phones, and large-scale devices in industrial utilities. LiBs have powered up to 90% of BESS globally. Present-day anodes employ cost-effective and light-weight carbon-based materials that tend to disintegrate after a finite number of chargedischarge cycles. In addition, nanomaterials have been proposed as anode materials to first overcome issues of volume expansion, which leads to the terminal degradation of the anode. Second, the high surface-to-volume ratio provided by nanostructured morphologies offers a higher number of active sites for Li storage that also facilitate reversible reactions. Therefore, integrating the SnO 2 compound into the existing graphite-based anodes for LiB would offer a quick adaptation of this nanomaterial in battery applications. In fact, graphite-based anodes currently being employed already exhibit a strong network for SnO 2 -nanomaterial integration with a potential for a better battery cycle life and higher capacity, while making related processes more sustainable and cost effective.
Funding: This research has been supported by the European Regional Development Fund project grant number TK134 "EQUiTANT", Eesti Maaülikool (EMÜ Bridge Funding (P200030TIBT)).

Informed Consent Statement: Not applicable.
Data Availability Statement: All figure were reproduced with permission or via creative commons attribution.

Conflicts of Interest:
The authors declare no conflict of interest.