This chapter analyses the main advances made in the field of sodium–vanadium fluorophosphates as cathodes for Na-ion batteries and tries to clarify some discrepan‐ cies and common errors published about these compounds. The sodium–vanadium fluorophosphate family can be divided in two main members: Na3V2(PO4)2F3 (V+3 extreme phase) and Na3V2O2(PO4)2F (V4+ extreme phase). Na3V2O2x(PO4)2F3-2x, where 0 < x < 1 would correspond to intermediate V3+/4+ mixed valence phases. Among them, the V3+ extreme has demonstrated to be difficult to isolate, whereas the V4+ and mixed valence phases can be more easily prepared by different synthesis methods and from different vanadium sources. In terms of electrochemical performance, mixed valent compound provides good performance, with high specific capacity at moderate/high cycling rates, and long cycle life. The future perspectives for this family of compounds are discussed in terms of raw materials availability, price, and performance relative to other cathode systems for Naion batteries.

**Keywords:** Na-ion batteries, cathodes, fluorophosphates, vanadium, High voltage

#### **1. Introduction**

Energy production is a key issue for the support of our society and way-of-life. Driven by the need to reduce the emissions of CO2 and increase energy security, policy makers have imple‐ mented different measures in order to shift to low-carbon energy resources. One of these policies is to promote the use of renewable energy sources at the expense of the carbon-based fuels [1,2]. However, the increasing use of these clean energy sources entails other problems, such as modulating time-variable energy production from renewable resources due to their depend‐

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enceontheweathertointegratethemintothegrid.Thus,inordertosolvetheproblemconcerning the intermittency of these resources, additional energy storage devices are needed. This way, electrochemical energy storage is a strategic research area that will help to increase the weight of renewable energy sources in the energy mix.

In this field, batteries are one of the most promising systems to act as buffer to regulate the variable energy income that can be obtained from the wind, sun and from the ocean waves. For this purpose, one of the main technologies under development is Na-ion batteries. In the last 20 years, Li-ion batteries have attracted all the attention in energy storage, but recently, Na-based compounds have made a comeback because of controversial debates regarding the size of reserves and higher cost to obtain Li [3]. The use of Na instead of Li in rocking chair batteries could mitigate the feasible shortage of lithium in an economic way for many reasons. The first one, the unlimited sodium sources, being the sodium content in the earth's crust and water of 28,400 mg·kg−1 and 11,000 mg·L−1 compared to 20 mg·kg−1 and 0.18 mg·L−1 for lithium [4]. Moreover, sodium presents lower price than lithium and it is easy to recover [5].

Although sodium (Na)-ion batteries are considered a promising alternative to lithium-ion systems, it is usually said that they possess limited electrochemical activity compared to Li due to two intrinsic reasons. First, Na has a lower ionization potential than Li (5.139eV vs. 5.392 eV) [6], leading to lower operating voltages and thus lower energy densities. Second, Na+ ions are heavier and larger than Li+ ions, leading to slow diffusion within a solid electrode during cycling and larger volume expansion of the electrode [7]. These two factors would induce lower gravimetric energy densities, so renewed research on these systems has been focused on their utility as buffering systems in the electrical grid, where battery volume is not so crucial due to the stationary nature of the practical use of these batteries.

However, and in spite of the great efforts devoted to the development of higher performance Li-ion batteries, if practical energy densities achieved up till now for both Li- and Na-based systems are compared, it can be seen that Na-ion batteries are catching up with the practical energy densities of Li-ion in a quite short time (**Figure 1**).

This great advance in gravimetric energy densities for Na-ion batteries is due to the extensive research that has been done in the last years to identify and optimise the best materials for these systems, so they can be soon commercially viable. As it has been said, the incorporation of this kind of systems to the electrical grid would lead to a significant increase of the propor‐ tion of renewable energies on the energy mix, with the positive environmental consequences that this fact would result in. Thus, it is a priority to continue searching for and optimising the best-performing materials to be used in the different components of a battery: cathode, electrolyte and anode.

In the case of the cathode, a good performing material should provide high energy density to the battery, that is, it should be able to charge and discharge the highest possible specific capacities with also high operating voltage and for many cycles (long cycle life). In this sense, sodium–vanadium fluorophosphates have demonstrated to fulfil all of these characteristics.

**Figure 1.** Theoretical and (estimated) practical energy densities of different rechargeable batteries: Pb–acid (lead acid), NiMH (nickel metal hydride), Na-ion (estimate derived from data for Li-ion assuming a slightly lower cell voltage), Liion (average over different types), HT-Na/S8 (high temperature sodium–sulphur battery), Li/S8 and Na/S8 (lithium–sul‐ phur and sodium–sulphur battery assuming Li2S and Na2S as discharge products), Li/O2 and Na/O2 (lithium–oxygen battery, where theoretical values include the weight of oxygen and depend on the stoichiometry of the assumed dis‐ charge product, i.e., oxide, peroxide or superoxide). Reproduced from reference [8].

#### **2. High-voltage cathode materials for Na-ion batteries**

In the last decade, three dimensional frameworks built on transition metals (M) and polyanions (XO4) n- have become a hot topic in the research field of electroactive materials for lithium and Na-ion batteries. The smaller theoretical gravimetric capacity achieved due to the presence of polyanion groups is compensated by the positive features presented by these compounds, such as very stable frameworks and high inductive effect [9]. The "inductive effect" is a concept proposed by Goodenough *et al.* consisting on the tune of the redox potential of the 3d metal centre (M) depending on the electronegativity of the X-O bonding [10].

On one hand, sulphate-based cathodes offer a good combination of sustainable syntheses and high energy density owing to their high-voltage operation due to electronegative SO4 2− units [11]. Na2Fe2(SO4)3 has demonstrated to be the most promising compound being operable at 3.8 V *versus* sodium and offering a reversible capacity of over 100 mAhg−1 with a 50% capacity retention at high rates (20C) [12]. However, the main issues concerning sulphur-containing cathodes are the low utilisation of active material, volume expansion during extraction/ insertion of Na ions and the dissolution of polysulphide into the electrolyte [13].

On the other hand, framework materials based on the phosphate polyanion have also been identified as promising electro-active materials for sodium metal and Na-ion battery applica‐ tions. It is the strong inductive effect of the PO4 3− polyanion that moderates the energetic of the transition metal redox couple to generate relatively high operating potentials for these compounds [14]. NaMnPO4 is one of the most studied compounds in this field due to the high operating voltage provided by its lithium analogue: LiMnPO4 which works at 4.1 V *versus* lithium [15]. However, contrary to LiMnPO4, NaMnPO4 possesses two structure modifications: marycite-type and olivine-type structures being this latter one the electroactive one. Boyadz‐ hieva *et al.* have recently reported, for the first time, the reversible lithium and Na ion inter‐ calation into the olivine NaMnPO4 phase when testing it *versus* lithium. However, the electrochemical performance of this material is not yet competitive compared to the wellknown LiMnPO4 analogue so that further optimisation is needed [16].

Na4Co3(PO4)2P2O7 pyrophosphate is another promising candidate due to its high working potential region (between 4.1 and 4.7 V *versus* sodium). In the work presented by Nose *et al.* [17], this compound offers 95 mAhg−1 reversible capacity with negligible fading after 100 cycles and 84% of capacity retention at rates of 25C. Moreover, the battery presented may be capable of running at even higher potentials (~5.1 V) in order to utilise the Co3+/ 4+ redox couple leading to the extraction of 4 Na ions from the Na4Co3(PO4)2P2O7 phase. Something similar is proposed by Xu *et al.* in the case of Na3V2O2(PO4)2F sodium–vanadium fluorophosphate. Due to the difficulty to access the third sodium in the formula unit some theoretical calculations suggested the substitution of oxygen with chlorine to form Na3V2Cl2(PO4)2F which would increase the energy density from 520 to 758 Whkg−1[18]. However, a suitable electrolyte would be needed in both cases in order to work in such high-voltage conditions.

Among the framework materials fluorophosphates possess even higher operating voltages than phosphates, because the inductive effect of fluorine is added to the effect of phosphate. This latter feature makes them a key to solve the energy density issue of sodium-based batteries. Fe, Mn, and V have been the most investigated transition metals but the three of them present different structures: Whereas the sodium–iron fluorophosphate possesses a twodimensional layered structure [19], the sodium–manganese fluorophosphate presents a threedimensional tunnel structure [20]. However, manganese compound has demonstrated to be poorly electrochemically active so that more studies have been performed for the iron-based compound. For the moment, one of the best results achieved has been for carbon coated porous hollow spheres of Na2FePO4F phase. This nanostructured material contained about 6–8 *wt.*% of carbon and presented a specific surface area of 8.7 m2 g−1 which led to a specific capacity of 90 mAhg-1 at C/10 and a good cycling stability for 100 cycles [21].

Regarding sodium–vanadium fluorophosphates, three phases have been described in the literature: NaVPO4F, Na3V2O2(PO4)2F and Na3V2(PO4)2F3 with theoretical specific capacity values of 143, 130 and 128 mAh·g−1, respectively. Barker *et al.* first proposed NaVPO4F as a tetragonal structure with *I4/mmm* space group [22]. On the other hand, several groups reported also the existence of a NaVPO4F polymorph which they indexed with *C2/c* space group [23– 26]. The Na3V2O2(PO4)2F phase has been described as a tetragonal symmetry compound with *I4/mmm* space group by Sauvage *et al.* [27] and Massa *et al.* [28] but, more recently, Tsirlin *et al.* [29] proposed the *P42/mnm* space group for the room temperature phase among the diverse polymorphs of this material at different temperatures. The third of the sodium–vanadium fluorophosphates mentioned in the literature, Na3V2(PO4)2F3, is also described as a tetragonal symmetry with *P42/mnm* space group [30–32] but more recently some authors have proposed a different space group for this compound: *Amam* [33,34].

A deep study of the bibliographic data related to these three sodium fluorophosphates leads to doubt about the real existence of these three different compounds. It is worth noting that the existence of the NaVPO4F phase has already been questioned by authors such as Sauv‐ age *et al.* [27]. Furthermore, there are no structural data about the suggested monoclinic pol‐ ymorph of NaVPO4F [23–26]. Besides, the diffractogram shown in their works as well as the electrochemical curves obtained match with the one of Na3V2(PO4)3 NASICON compound. In this case, the formation of a NASICON type phase instead of NaVPO4F could be possible because the samples were prepared by the ceramic method and sublimation of VF3 could oc‐ cur, as it has been described for other lithium–vanadium fluorophosphates [35]. Thus, struc‐ tural data of only two compounds are found in the literature: Na3V2O2(PO4)2F [29] and Na3V2(PO4)2F3 [30]. The diffractograms of both phases are quite similar that it is possible to establish a relationship between them.

Electrochemical studies on the three mentioned tetragonal structure compounds have been performed under different conditions *versus* lithium or sodium [22–24,27,36–39] and in all cases the charge–discharge curves display two voltage *plateaux* of similar length at the same voltages with little differences depending on the anode material used (*ca.* 3.6 and 4.1 V *ver‐ sus* Na/Na+ ).

#### **3. Sodium–vanadium fluorophosphates as cathode materials for Na-ion batteries**

The study of sodium–vanadium fluorophosphates is especially relevant because of the high operating voltages offered by these compounds that could lead the way to high energy Naion batteries.

The similarity of the X-ray diffraction patterns as well as the almost identical electrochemi‐ cal data for Na3V2O2(PO4)2F and Na3V2(PO4)2F3 suggested that both materials could belong to the same family of compounds, where the fluorine content is modulated by the presence of V3+ and VO2+ (V4+) leading to the following general formula Na3V2O2x(PO4)2F3-2x which was proposed by our group for the first time [40]. The extreme members (x = 0 and x = 1) would correspond to the mentioned phases, Na3V2(PO4)2F3 for x = 0 (**Figure 2a**) and Na3V2O2(PO4)2F for x = 1 (**Figure 2b**), whereas intermediate compounds (0 < x < 1) would be V3+/V4+ mixed valence phases (**Figure 3**). The existence of a mixed valence family of compounds between Na3V2(PO4)2F3 and Na3V2O2(PO4)2F phases was also reported by Park *et al.* [41] and more re‐ cently, by Qi *et al.* [42].

**Figure 2.** Crystallographic structure of a) Na3V2(PO4)2F3 and b) Na3V2O2(PO4)2F. Oxygen in red and fluorine in green. Vanadium octahedron is represented in blue and phosphate tetrahedron in grey.

**Figure 3.** Vanadium octahedron for mixed valent Na3V2O2x(PO4)2F3-2x (0 < x < 1) compounds. Vanadium is represented in blue, oxygen in red and fluorine in green.

In order to understand the structure and properties of this family of compounds, a series of sodium–vanadium fluorophosphate samples were hydrothermally prepared, varying the type and amount of carbon used as reductive agent during the synthesis [40]. This way, the possible influence of that carbon on both the final properties of the sample and the attainment of different compositions in the family of general formula Na3V2O2x(PO4)2F3-2x (0 < x < 1) was analysed. Intermediate amounts of carbon present in the Na3V2O2x(PO4)2F3-2x final product (1 – 50 *wt.*% C) led to V3+/V4+ mixed valence phases (0 < x < 1) while a sample containing almost only V4+ was obtained in a composite with less than 1 *wt.*% of carbon. The electrochemical study of these samples showed the presence of two reaction voltages at 3.6 and 4.1 V *versus* Na/Na+ and good specific capacity values up to 100 mAhg−1 at 1C when an electrochemical grade carbon was employed. The material containing a moderate amount of this type of carbon (~6 *wt.*%) showed the best electrochemical results. A deeper analysis of this sample with an average oxidation state of +3.8 for the vanadium, showed excellent specific capacity values from 100 to 75 mAhg−1 for charging rates of C/20 and 5C, respectively. It is worth mentioning the high capacity retention up to 95% after 200 cycles at 1C [43]. A combined experimental and theoretical study on this family of compounds was also carried out by Park *et al*. [44,45]. Different samples with compositions x = 0, 0.2, 0.5, 0.8 and 1 were synthesised by solid state reactions. The obtained compounds were characterised by the use of different techniques such as X-ray Diffraction (XRD), Transmission Electron Microscopy (TEM), Nuclear Magnetic Resonance (NMR), etc. The electrochemical study showed specific capacity values of approx‐ imately 100 mAhg−1 only in the case of the V4+ composition. The specific capacity value was maintained for all the compounds during 150 cycles at C/2. Recently and in order to achieve different compositions of the family of sodium–vanadium fluorophosphates a new synthesis based on the solvothermal method at low-temperature (60-120°C) has been employed [42]. The size of the particles obtained with this method was in the range of 20–50 nanometers. The electrochemical analysis of these compounds showed a rate capability of 73 mAhg−1 at 10C rate and long cycle stability over 1200 cycles at 2C.

Regarding the V4+ phase, Na3V2O2(PO4)2F, the first electrochemical results were presented by Sauvage *et al.* getting a capacity of 87 mAhg−1 at C/100 through a two voltage-step behaviour at 3.6 and 4.0 V vs. Na/Na+ . They found that the obtained electrochemical curve was almost identical to the ones reported for Na3V2(PO4)2F3 and NaVPO4F phases giving rise to doubt about the real existence of this latter one in light of lack reliable information about its structure [27]. Some years after, a novel single-step hydrothermal treatment was discovered to obtain the Na3V2O2(PO4)2F phase [46]. The absence of *in-situ* carbon in the synthesised product made it necessary to employ an *ex-situ* carbon coating in order to improve the conductivity of the material and hence its electrochemical performance. However, control of the thermal treatment process resulted determinant to avoid the degradation of the material and the appearance of secondary phases such as NASICON Na3V2(PO4)3. The carbon coated material prepared by impregnation and a flash thermal treatment, exhibited good rate capabilities between 100 and 65 mAhg−1 for rates of C/20 and 1C respectively and good cycling stability (**Figure 4**).

**Figure 4.** Comparison of discharge specific capacity values between raw Na3V2O2(PO4)2F and carbon coated flash ther‐ mally treated Na3V2O2(PO4)2F/C samples.

Additional works showed that the reversible capacity of Na3V2O2(PO4)2F/graphene electrodes exceed 100 mAhg−1 after 200 cycles at a C/20 current rate [37]. Apart from that, a theoretical and experimental study on a solvothermally obtained Na3V2O2(PO4)2F phase was presented were structural and electrochemical characterisation of this material was carried out [18]. Moreover, computational studies revealed the possibility of extracting the third sodium of the mentioned phase at 5.3 V *versus* sodium. However, this would be experimentally complicated due to the overcoming of the stability window of the electrolyte (4.8 V). However, a hypo‐ thetical new vanadium fluorophosphate phase, Na3V2Cl2(PO4)2F, where the oxygen is substi‐ tuted by chlorine , was proposed from the theoretical study. This latter compound would provide 188 mAh·g−1 theoretical specific capacity and 758 Wh·kg−1 energy density. Finally, electrochemical results of RuO2-coated Na3V2O2(PO4)2F nanowires were presented were a reversible capacity of 120 mAh·g−1 was achieved at 1C rate during more than 1000 cycles [47].

Concerning the V+3 phase, Na3V2(PO4)2F3 , the electrochemical mechanism of the reversible Na extraction was described by Chihara *et al.* [48] and Shakoor *et al.* [38]. Both groups reported that the electrochemical reaction occurs through a single-phase reaction with negligible volume variation (~2%). More recently, the ion migration mechanism and diffusion capability of the Na3V2(PO4)2F3 phase have been explored [49,50]. However, the presented results could correspond to any of the sodium–vanadium fluorophosphates included in the general formula Na3V2O2x(PO4)2F3-2x. Moreover, one of the three discharge *plateaux* that they achieve in the phase (the one at 3.4 V) could correspond to the non fluorinated NASICON Na3V2(PO4)3 phase as recently reported[46]. On the other hand, an *ex situ* high resolution 23Na and 31P solid state NMR study on this Na3V2(PO4)2F3 phase was reported [51]. It is mentioned that 23Na NMR spectrum of pristine Na3V2(PO4)2F3 phase exhibits two major resonances at 92 and 146 ppm, which are assigned to partially filled Na(2) site and fully occupied Na(1) site, respectively. However, in recent works concerning ssNMR of the mixed valent sodium–vanadium fluoro‐ phosphate phase the peaks found in the spectra were assigned to the presence of different vanadium environments instead of the occupancy of the Na cations [44,52]. Similar results were found in the case of the Na3V2O2(PO4)2F phase containing only V4+ . The results of the 23Na ssNMR spectrum obtained in this phase show the presence of a unique peak with an occupancy factor of the two Na cations of 1 for Na(1) site and ½ for Na(2) site (see **Figure 5**) [53].

**Figure 5.** 23Na ssNMR spectrum of Na3V2O2(PO4)2F sample (black line). Fittings for Na(1) and Na(2) are shown in green and blue, respectively.

The crystal structure of the Na3V2(PO4)2F3 phase at high temperature (400 K) obtained from high resolution synchrotron radiation measurements in *operando* revealed the existence of a combined solid solution region and four intermediate phases during the charge process [33,34]. In addition, a new space group (*Amam)* was achieved for the Na3V2(PO4)2F3 phase . On the other hand, the electrochemical results obtained in a full cell using Na3V2(PO4)2F3 as cathode, *hard carbon* as anode and an optimised electrolyte showed a capacity of 97 mAh·g−1 after more than 120 cycles at C/5 and 70 mAh·g−1 at 5C rate. [39].

Concerning *ex-situ* structural study of the aforementioned phases, a detailed study was recently presented [44,45] to determine the sodium extraction/insertion mechanism of the mixed valent fluorophosphate materials. Examination of the *ex-situ* XRD patterns of V3.8+ Na3V2O1.6(PO4)2F1.4 compound revealed the coexistence of a two-phase region in the first half of the charge indicating the presence of a stable intermediate phase between the two *plateaux* at 3.6 and 4.1 V *versus* Na/Na<sup>+</sup> . It was also reported that the V3+ present in the V3.8+ Na3V2O1.6(PO4)2F1.4 phase creates a local environment similar to the one presented by Na3V2(PO4)2F3 phase [38] leading to a high-voltage V3+/V4+ redox reaction. It is assumed that 2.4 electrons are supplied from the V3.8+ / V5+ redox couple leading to a theoretical specific capacity of 155.6 mAhg−1. On the other hand, X-Ray Absorption Near Edge Spectroscopy (XANES) studies revealed that an oxidation process happens while charging and that the reverse process occurs when discharging but no evolution of the value of the vanadium oxidation state was determined.

Concerning Na3V2O2(PO4)2F phase, different space group symmetries have been described in the literature to define it [28,29]. In terms of sodium insertion/extraction, the *I4/mmm* is shown to reversibly cycle 1.12 Na ions per formula unit (Na3V2O2(PO4)2F) [27] while the *P42/nmn* polymorph with a carbon-coated synthetic procedure, reversibly cycles 1.56 Na ions per formula unit [46].

Finally, different time-resolved *in-situ* synchrotron X-ray powder diffraction studies have been carried out in order to analyse reaction mechanism evolution, lattice parameters ,and sodium extraction of V3.8+( Na3V2O1.6(PO4)2F1.4), V4+ Na3V2O2(PO4)2F and V+3 Na3V2(PO4)2F3 materials [34,54,55].

The electrochemical behaviour of the two electrode materials studied *versus* Na/Na<sup>+</sup> showed similar electrochemical reaction mechanism evolution where a simultaneous solid solution and two-phase region were observed. However, the maximum sodium amount that can be extracted when a potentiostatic step is applied (2 hours at 4.3 V) is higher for V3.8+ (Na3V2O1.6(PO4)2F1.4) sample than for the V4+ Na3V2O2(PO4)2F one, 2.23 *versus* 1.75, respectively. This fact can be due to several reasons, such as better intrinsic conductivity, the presence of *in-situ* carbon in the mixed valent sample or to the more rigid structure of the only V4+ containing sample. Nevertheless, both materials present high reversibility and flexibility during charge/discharge process. The effects of overcharging up to 4.8 V and cycling-related degradation of the same two compounds, V3.8+( Na3V2O1.6(PO4)2F1.4) and V4+ Na3V2O2(PO4)2F [56] have been recently reported. The V3.8+(Na3V2O1.6(PO4)2F1.4) phase presents less stress than V4+ phase due to the more disordered structure and the presence of *in situ* carbon. Thus, this phase is able to reach the maximum theoretical Na+ extraction limit when charging up to 4.8 V while in the case of V4+ Na3V2O2(PO4)2F phase this limit is established at the usual 4.3 V cutoff voltage. Moreover, mid-term cycling shows the formation of secondary products due to the degradation of the materials in both cases.

To finish with, the structural evolution of V3+ Na3V2(PO4)2F3 phase has recently been analysed by using high angular and intensity resolution synchrotron radiation [34]. The two voltage domains (at 3.7 and 4.2 V *versus* Na/Na+ ) have been studied separately. In the case of the low voltage *plateau* (3.7 V) three biphasic reactions have been observed while in the higher voltage region (4.2 V) a small solid solution zone together with a biphasic zone is shown (**Figure 6**). Unexpectedly, the structure of the NaV2(PO4)2F3 end member is different from the previous ones revealing a different space group (*Cmc21*) and two different vanadium environments (V3+ and V5+).

#### **4. Conclusions**

Polyanions (XO4) n- based on transition metals (M) are one of the most promising candidates for high-voltage Na-ion batteries. Among them, sulphate and phosphate based framework materials stand out. Regarding phosphate-based materials, and as it has been previously shown in this chapter, there has been some controversy about the structural and electrochem‐ ical features of sodium–vanadium fluorophosphates. This family of materials show very high

**Figure 6.** Potential−composition electrochemical curves obtained upon Na+ extraction from Na3VPF. The first cell was cycled (GITT) between the compositions Na3VPF and Na2VPF while recording 58 5 min diffraction scans; the second one (galvanostatic cycling) between the compositions Na2VPF and Na1VPF while recording 95 5 min scans. Significant single-phase compositions are highlighted by coloured circles. For the first cell they are obtained by the electrochemis‐ try. Cycling of the second cell resulted in Δx > 1 because of electrolyte oxidation, hence in this case reported composi‐ tions are those obtained from Rietveld refinement. Top line describes the nature of the observed deintercalation mechanism. Reproduced from reference [34].

energy density due to the high voltage provided by the adding inductive effect of fluorine. Moreover, these materials have demonstrated to be very stable and present outstanding electrochemical properties. V+3 Na3V2(PO4)2F3, V+3.8 Na3V2O1.6(PO4)2F1.4 and V+4 Na3V2O2(PO4)2F phases included in the family of compounds Na3V2O2x(PO4)2F3-2x have been structurally and electrochemically deeply analysed. Despite the similarity of all these materials, up to date studies show that the electrochemical behaviour of these electrode materials is clearly de‐ pendant not only on the sodium extraction/insertion mechanism, the occupancy and distri‐ bution of sodium and the electrochemical cycling history but also on the synthesis process employed for the obtaining of the starting material that determines its properties.

#### **5. Future perspectives for sodium–vanadium fluorophosphates as cathodes for Na-ion batteries**

As it has been presented, sodium–vanadium fluorophosphates are more than promising cathodic materials for near future commercial Na-ion batteries. However, the appearance of real commercial Na-ion batteries involves the development of the whole battery, that in‐ cludes advancing in finding the appropriate components for a Na-ion cell based on a so‐ dium–vanadium fluorophosphate cathode. In the first place, an anode matching this cathode must be chosen. This anode material should present relatively low operating volt‐ age in order to provide the battery with the maximum available voltage. In the second place, the specific capacity of this anode and its rate capability should be in balance with the cath‐ ode performance in order to make an easier mass balance to prepare the cells. These two conditions will be necessary to assure that the battery, and thus its components, present the best possible performance in terms of battery voltage and power.

On the other hand, a special effort must be done to search for a long-lasting and high-voltage working electrolyte for these systems. Since the recovery of Na-ion battery research, great efforts have been directed to the search of new electrode materials whereas studies dealing with the electrolyte are much scarcer. One of the reasons for this trend can be that it has been shown that the SEI formed on carbonaceous electrodes is markedly different for sodium- and lithium-based electrolytes even using the same solvent [57,58]. Non-aqueous Na electrolytes presently used are mainly based on NaClO4 or NaPF6 salts in propylene carbonate or other solvents and mixtures such as ethylene carbonate: dimethyl carbonate. In general, a good electrolyte should exhibit: i) good ionic conductivity, ii) a large electrochemical window (i.e., high and low onset potential for electrolyte decomposition through oxidation and reduction at high and low voltages, respectively), iii) no reactivity towards the cell components, iv) thermal stability (i.e. melting point and boiling point lower and higher than the standard temperatures for the cell utilization, respectively). Finally, it should be intrinsically safe, have as low toxicity as possible and meet cost requirements for the targeted applications. All these features are intrinsically dependent on the nature of the salt and the solvent(s) and the possible use of additives [59].

Another argument pointed out as a drawback for Na-ion batteries based on sodium–vanadium fluorophosphates is the use of a vanadium compound in them. It is said that vanadium is not a green element that needs careful handling after battery life cycle, but the same can be said from the electrolytes used in both sodium- and lithium-based systems. In any case, a battery will always need to be collected to be treated as a residue after its cycle life, as it is nowadays done. Apart from that, it is also commented that vanadium cost is another drawback for the development of these batteries. But it must be recalled that there exist commercial vanadiumrich energy storage systems, such as redox flow batteries, that comprise the use of high amounts of vanadium for their two electrodes [60]. Thus, a careful cost analysis should be done to confirm or discard this statement once the appropriate anode and electrolyte materials are chosen and a viable cell prototype is proposed.

Finally, and now focusing our attention on the commercial development of only sodium– vanadium fluorophosphates compounds, a last challenge that must be overcome is to develop a synthetic method that allows an easy and economical mass production of these materials. For this purpose, different parameters will have to be taken into account: the different possible synthetic methods (hydrothermal, microwave or ceramic processing); the different possible vanadium starting reactants; and the yields obtained by crossing these latter two parameters (**Figure 7**).

**Figure 7.** Main parameters that must be adjusted to get a commercially viable production method for sodium–vanadi‐ um fluorophosphates.

To sum up, sodium–vanadium fluorophosphates family is one of the main possible future cathodic materials for high-voltage Na-ion batteries. Their excellent electrochemical proper‐ ties and performance have been perfectly well demonstrated by different research groups, as it has been told in this chapter. Now, the great defy for these materials lays on the indus‐ trial pace, that is, in the assembly and building of a real Na-ion full cell based on these mate‐ rials and the development of a commercially viable process to produce these compounds in great amounts.

#### **Acknowledgements**

This work was financially supported by the Ministerio de Economía y Competitividad (MAT2013-41128-R) and the Gobierno Vasco/Eusko Jaurlaritza (IT570-13). University of the Basque Country (UPV/EHU) is acknowledged for funding under project UFI11/15.

#### **Author details**

Paula Serras1,2\*, Verónica Palomares1 and Teófilo Rojo1,3

\*Address all correspondence to<EMAIL_ADDRESS>
1 Inorganic Chemistry Department, University of the Basque Country UPV/EHU, Leioa, Bizkaia, Spain

2 Nuclear Engineering and Fluids Mechanics Department, Faculty of Engineering, Universi‐ ty of the Basque Country UPV/EHU, Bilbao, Bizkaia, Spain

3 CIC Energigune, Alava Technology Park, Miñano, Álava, Spain

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http://dx.doi.org/10.5772/62759

#### **Abstract**

In this book chapter, we have discussed the recent results on vanadium oxide-based materials for energy storage applications. Primarily, we present the new results from our own research group on V2O5-layered nanostructures that are made from a facile wet chemistry synthesis. By fine control of the synthetic condition, the morphology, crystallinity, and layer-to-layer distance of V2O5 nanostructures can be tuned. Particu‐ larly, highly disordered V2O5 nanolayers which have an interplanar distance up to 1.1 nm, offering a fast transport of K-ion between layers in an aqueous electrolyte and hence the high-energy storage capacity and power density. Uniqueness of our results in‐ cludes materials characterization and measurements using multiple spectroscopic tools, including synchrotron X-ray pair distribution function (PDF) analyses and in situ X-ray diffraction (XRD). Combined with half-cell and button-cell electrochemical measure‐ ments, the complementary results provide insight on the ionic transport of ions between the layers of V2O5 nanostructure.

**Keywords:** Vanadium pentoxide, Electrochemical energy storage, X-ray diffraction, Pair distribution function, *in situ* XRD

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

#### **1. Introduction**

#### **1.1. Electrochemical energy storage**

Rising demands for energy coupled with concerns over environmental pollution and depend‐ ency on petroleum fuel have contributed to a great need for new energy conversion and storage technologies [1, 2]. According to the United States Vehicle Technologies Office (VTO) at the Office of Energy Efficiency and Renewable Energy (EERE) within the Department of Energy (DOE), transitioning to a light-duty fleet of hybrid electric vehicles and plug-in electric could reduce US foreign oil dependence by 30–60% and greenhouse gas emissions by 30–45%. Thus, two approaches, both involving electrochemical processes, are generally suggested for solving this energy dilemma for the United States: (i) Development of highly efficient electrochemical energyconversiondevices, suchas fuel cells, as analternative tothe internal combustionengine. Such devices can use gas or liquid fuels rather than gasoline to reduce CO2 emissions; (ii) Development of highly efficient electrochemical energy storage (EES) devices, such as Li-, Na-, orK-ionbatteries, andelectrochemical capacitors,to accelerate the adoptionof electricalvehicle and renewable non-carbon-emitting energy from solar and wind energy sources [3, 4].

**Table 1.** Properties of different electrical energy storage devices.

Many EES devices, including most batteries and the majority of electrochemical capacitors,

$$\rm{MO}\_X + \rm{A}^+ + \rm{e}^- \leftrightarrow \rm{MO}\_X \rm{A} \tag{1}$$

where MOX is a transition metal oxide, and A+ denotes the alkali ions (e.g., Li+ , Na+ or K+ ). Upon the insertion or de-insertion of A+ , metal (M) cations will be reduced or oxidized to balance the charge. When an EES device is charged, a voltage (V) will build up between the anode and the cathode. The energy density (E) depends on the capacitance (C) and voltage of the device, while the power density (P) depends on the voltage and equivalent internal resistance (RS) of the device, as following equations show:

$$\mathbf{E} = \mathbb{V}\_2 \mathbf{C} \mathbf{V}^2 \tag{2}$$

Vanadium Pentoxide (V2O5) Electrode for Aqueous Energy Storage: Understand Ionic Transport using Electrochemical, X-Ray, and Computational Tools http://dx.doi.org/10.5772/62759 175

$$\mathbf{P} = \mathbf{V}^2 / \left(\mathbf{4R\_s}\right) \tag{3}$$

To provide clean and efficient energy solutions, a new genre of EES devices that are able to fill the gap between high-energy-density battery devices and high-power-density electrostatic capacitor devices are in great demand. From Eqs. (2) and (3), it is clear that to increase energy density and power density of EES devices, one needs to:

#### **1.2. Current state of knowledge of Na-ion-based aqueous EES**

Current EES can be classified into aqueous and non-aqueous storage according to the chem‐ istry of the electrolytes. Non-aqueous electrolytes can achieve wider voltage window (usually >3.0 V), and thus, non-aqueous EES usually exhibits much higher-energy density than aqueous counterparts. However, non-aqueous systems are expensive and often require the use of highly toxic and flammable solvents, which can become hazardous if used improperly, such as overcharging or short-circuiting. Non-aqueous EES devices, including lithium ion batteries (LIBs), require strict preparation standards at high cost, to allow for water- and oxygen-free cell fabrication. On the other hand, aqueous EES has narrow voltage window, typically around 1.2 V, limited by oxygen and hydrogen evolution reactions of water. However, aqueous electro‐ lytes are generally inexpensive. Water- and oxygen-fee manufacturing environment (lower than 2 ppm [part per million]) required by non-aqueous EES is unnecessary for aqueous EES, which lowers the cost of cell packing dramatically. Also, aqueous EES can use electrolytes with high salt concentration and hence have nearly two orders of magnitude higher ion conduc‐ tivities compared to non-aqueous electrolytes. Aqueous lead acid batteries (LABs) are one of the most commonly used aqueous EES devices and have great advantages such as low cost and high surge current compared to LIBs. However, LABs have a short cycle life, and their components such as lead and sulphuric acid are not environmentally friendly. Other aqueous EES devices, especially rechargeable ones using earth-abundant and non-toxic electrode materials and Na- and/or K-ions as charge carriers, appear more appropriate for stationary energy storage to be used with renewable energy sources due to lower cost of raw material, longer cycling life, and better environmental friendliness compared to Li-ion and lead acidbased counterparts [5–7].

Na-ion- and K-ion-based aqueous energy storage devices have showed great potential as a new class of energy storage devices as alternatives to Li-ion storage devices, since the electro‐ chemical potential of Na and K is almost identical to that of Li-ions. Although radii of the bare ions is in the order of Li+ < Na+ < K+ (**Table 2**), due to the effect of hydration number (number of water molecules that migrate with each cation) of alkali metals, total radius of Na- and Kion and their bound water molecules (Stokes radius) actually are smaller in the order of K+ < Na+ < Li+ , resulting in a higher chemical diffusion coefficient of Na- and K-ions in the aqueous electrolyte compared with Li-ions. Moreover, current mineable Li resources can be sustained for approximately only 65 years at most. In contrast, Na and K are about 1,000 times more abundant in earth crust and hence more economical. Thus, aqueous energy storage devices based on Na- and K-ions holds great promise for economical and sustainable energy storage devices with high stability, better ionic conductivity, and low cost. In this context, to find a suitable electrode materials and architecture, which guarantees a reversible intercalation and deintercalation of Na- and K-ions, while maintaining a stable electrode and electrolyte interface, becomes imperative to the design of new generation of energy storage devices.

**Table 2.** Radii and ionic conductivities of various alkali metals in water at 298 K.

#### **1.3. Layered vanadium oxides nanomaterials for EES**

Charge storage mechanisms involving intercalation and deintercalation of alkaline ions has been commonly found in ESS devices with a high-energy density, such as batteries. Although intercalation process is highly reversible and suitable for the high capacity energy storage, intercalation materials usually are not suitable for high rate performance. This is largely due to the fact that the diffusion of intercalated ions (e.g., alkaline ions) in the electrode materials is sluggish, having diffusion coefficients in the range from 10−13 to 10−17 m2 /s. To shorten the diffusion time (τ), that is the time required for alkaline ions move through the host electrode material, one has to either reduce the ionic path of alkaline ions within the electrode material (L) or increase the diffusion coefficient of alkaline ions in the solid phase (D) [8], described by the following equation:

$$
\boldsymbol{\pi} = \mathbf{L}^2 / \mathbf{D} \tag{4}
$$

Since the diffusion coefficient (D) is one of the intrinsic properties of alkaline ion-host electrode material couple, the only practical way to decrease the diffusion time (τ), and hence increase the rate performance of the electrode material, is to decrease the diffusion length. In this context, diffusion time can be considered as the minimum time required by the alkaline ions to insert and become homogenously distributed within the host electrode material. Hence, a diffusion time of alkaline ions within host solid electrode is closely related to the C-rate of an EES device, which is a measure of the rate at which an EES device is discharged relative to its maximum capacity (1/n C-rate means that the discharge current will discharge the entire device in n hours).

**Table 3.** Relationship between the diffusion time (τ), C-rate, and diffusion length (L) of transport of alkaline ions within the electrode materials.

**Table 3** describes relationship between the diffusion time (τ), C-rate, and diffusion length (L), assuming the diffusion coefficient of alkaline ions within metal oxide electrode materials is in the order of 10−17 m2 /s. Considering 36,000 s (10 h) a reasonable time to completely charge and discharge a battery device, a maximum diffusion length of 600 nm is needed. And this value appears to be easily achieved through tailoring the conventional preparation routine for the electrode materials. However, for a high-rate EES device that affords charge and discharge time around 36 s, a much shorter diffusion length of 19 nm is required, which might be too small to be prepared via conventional preparation routine, which typically involves a solid state reaction through decomposition of carbonate salts at high temperature (usually >500°C). A more likely prospective method to prepare nanoscaled electrode materials for shorter diffusion time is the synthesis of layered electrode structure. Many dioxides of transition metals consist of [MO6] (M: Mn, Ti, V, Ir, Ru, Cr) octahedra with shared vertices and edges. Stacking of [MO6] octahedra enables the building of two-dimensional layered structures [9– 14]. Water molecules and alkaline ions such as Li<sup>+</sup> , Na+ , or K+ can easily move into or out of the layers without significant structural rearrangements. Particularly, large interplanar distance (usually close to 1 nm) favors the storage of alkaline cations and a higher capacitance [15].

Among various transition metal oxides, vanadium pentoxide (V2O5) has attracted great interest in the fields of electrochromic devices and as electrode materials for batteries and electro‐ chemical capacitors, thanks to its low cost and ability to exist in different oxidation states from V2+ to V5+. Layered vanadium oxides structures have been synthesized using electrodeposition, deposition under ultra-high-vacuum conditions or a top-down solution-phase synthesis [16, 17]. Particularly, solution-phase exfoliation methods that can synthesize layered nanostruc‐ tures in large quantities have been well studied. High-temperature pre-treatment was usually required, and intercalators (e.g., tetrabutylammonium ions) were used to assist the separation of layers and hamper the reassembly of the bulk laminar material. Furthermore, by appropriate control of the chemical nature of the solvent, intercalators, exfoliation time, and temperature, as well as the post-treatment temperature, it is possible to tune the dimension of the interplanar distance in order to improve the electrochemical performance, that is, to make the interlayer distance large enough to be capable of accommodating the alkaline ions and water.

#### **2. Highly crystalline V2O5-layered electrode material for aqueous K-ion storage**

**Figure 1.** (a–c) TEM images of K0.33V2O5-layered nanostructure after being annealed at 450°C and (d) selected-area elec‐ tron diffraction of such materials.

V2O5-layered nanofibers can be synthesized by a facile two-step method involving solutionphase procedure followed by mild thermal treatment in the air [18]. In the first step, nearly amorphous vanadium oxides were prepared by mixing vanadyl (IV) sulfate hydrate with potassium hydroxide under vigorous stirring in water at room temperature. The resulting precipitations were dried in vacuum and then heated to 450°C for 2 hr in air. **Figure 1** shows TEM images of the final product after thermal treatment, showing a layered V2O5 nanofiber with an average diameter of ~120 nm and a length in a scale of micrometres. The "rolling-up" formation of the layered nanofibers is subsequently proposed based on TEM observations, suggesting that [VO6] building units were exfoliated into layered K0.33V2O5 sheets at mild temperature with intercalated K-ions within between the layers with a large interplanar distance of 0.96 nm. The results suggested nanolayers rolled up along <001> direction, which appeared to be the most energetically favorable mechanism due to the large periodicity of the nanolayers.

**Figure 2.** XRD patterns and crystalline schematics of K0.33V2O5-layered nanofibers heat treated at 450°C. Wavelength of X-ray was 0.3196 nm. Diffractions from layered KxV2O5 (JCPDS 81045) and non-layered KxV2O5 (JCPDS 27718) were labeled in black and blue.

**Figure 1c** shows the high-resolution TEM (HRTEM) image of the final layered nanofiber with a uniform lamellar structure. An average repeating distance of 0.71 nm between the layers was observed over almost the entire layered nanofiber in the direction parallel to the longitudinal axis (growth direction). In addition to this periodic stacking of the layers, a clear fringe spacing of 0.33 nm was also observed from HRTEM imaging. **Figure 1d** shows the selected-area electron diffraction (SAED) pattern from the same layered nanofibers shown in **Figure 1c**, showing a single-crystalline nature of the layered nanofibers demonstrated by the discrete diffraction. The fringe distances were calculated from SAED to be 0.67 and 0.32 nm, matching well with the distances calculated from HRTEM. These values also matched well with dspacing values from K0.49V2O5 crystal (JCPDS 81045). The interplanar distance (d001 = 0.71 nm) obtained from HRTEM was smaller than that calculated from X-ray diffraction (XRD; 0.96 nm), which can be explained by the removal of structured water between the nanofibers' layers while under high vacuum and the consequent contraction of the interplanar distance of layered nanofibers during HRTEM analysis.

XRD patterns of the resulting product in which two series of reflections were shown in **Figure 2a**, which contained the conventional (hkl) set of reflection, and the (00l) set of diffrac‐ tion with a high-degree ordering of the lamellar structures. The XRD reflection was in good agreement with potassium-intercalated layered V2O5 (K0.49V2O5, JCPDS 81045), except that the peak at 2θ = 2.46° was indexed as the (002) plane of non-layered KXV2O5 (JCPDS 27718). The strongest diffraction peak at 2θ = 1.98° was indexed to the (001) basal reflection with a corresponding d-spacing (d001) of 0.96 nm. Energy dispersive X-ray spectroscopy revealed the chemical composition of the final product that contained 83.7% V and 13.6% K (atomic percentage) with a chemical formula of K0.33V2O5.

**Figure 3.** (a) *In situ* XRD of K0.33V2O5-layered nanofibers at potentials from 0 to 0.9 V (vs. Ag/AgCl). The diffraction peaks from graphite oxide were marked by black arrows. (b) The evolution of the (001) planes upon the CV cycling.

**Figure 3a** shows a series of time-resolved XRD patterns of K0.33V2O5-layered nanofibers using an electrochemical cell that the authors have reported in detail elsewhere. XRD patterns were obtained during anodic (black curves) and cathodic scans (blue curves) within the potential window between 0 and 0.9 V using a cyclic voltammetry (CV) program at a scan rate of 1 mV/ s. Compared to the powder diffraction shown in **Figure 2**, steady diffraction background from the aqueous electrolyte solution and the commercial carbon paper from the in situ electro‐ chemical cell were observed (marked by black arrows). The K0.33V2O5-layered nanofibers showed varied peak intensities and positions at different potentials, suggesting dynamic changes of crystalline structure during the Faradaic processes, while the diffraction of layered nanofibers remained unchanged after one complete CV cycle. The results suggest that K0.33V2O5-layered nanofibers possessed a reversible surface-controlled capacitive process, as well as a stable layered crystalline structure during the charge/discharge processes.

The changes of (001) diffraction peaks during the CV cycling are shown in **Figure 3b**. During the anodic scan, d-spacing of (001) plan (d001) expanded from 0.93 nm (2θ = 1.98°) to 1.04 nm (2θ = 1.76°), corresponding to a 12% expansion of interplanar distance. On the other hand, d001 contracted from 1.04 to 0.95 nm during the cathodic scan. The expansion and contraction of d001 match very well with the charge storage mechanism of electrochemical extraction and insertion of K-ions into K0.33V2O5-layered nanofibers. During the anodic scan, oxidation of K0.33V2O5-layered nanofibers took place, accompanied with the removal of K-ions from interplanar layers of K0.33V2O5. These processes can be expressed by the forward reaction of following equation:

$$\mathbf{K}\_{0.33}\mathbf{V}\_2\mathbf{O}\_3 \leftrightarrow \mathbf{K}\_{0.33-\mathbf{x}}\mathbf{V}\_2\mathbf{O}\_3 + \mathbf{x}\mathbf{K}^+ + \mathbf{X}\mathbf{e}^- \tag{5}$$

where x is the molar fraction of extracted K-ions. Depletion of K+ weakened the electrostatic interaction between K-ions and negatively charged [VO6] octahedral units within the layers, and increased the interplanar distance (d001). Conversely, during the cathodic scans, K0.33V2O5 layered nanofibers were reduced, accompanying with the insertion of K-ions into interplanar layers [backward reaction in Eq. (5)]. The strengthened electrostatic interaction between [VO6] and K-ions decreased the interplanar distance.

**Figure 4.** Three-electrode half-cell analyses: (a) CVs of K0.33V2O5-layered nanofibers/PEDOT as loading of 5 and 1.25 μg, respectively; (b) gravimetric capacitance of K0.33V2O5, PEDOT, and active carbon as functions of scan rates.

**Figure 4a** shows the cyclic voltammograms (CVs) of the mixture of K0.33V2O5-layered nanofibers and PEDOT, respectively. The PEDOT and active carbon only electrodes showed broad redox peaks (data are not shown here), Indicating sluggish redox kinetics of K-ions. In contrast, the K0.33V2O5/PEDOT electrode showed well-defined anodic and cathodic peaks (**Figure 4a**), representing the K+ extraction and insertion behavior, occurring at ~0.6 (anodic scans) and ~0.5 V (cathodic scans), respectively. The calculated gravimetric capacitance (CMS) was strongly dependent on the scan rate as shown in **Figure 4b**. As the scan rate decreased, CMS increased. At a scan rate of 5 mV/s, the CMS reached the maximum value of 136 F/g (based on K0.33V2O5 mass). In general, the electrical conductivity of bulk V2O5 is low (10−5 ~ 10−3 Ω cm), so the addition of highly conductive PEDOT into K0.33V2O5 improves the rate capability, although the CMS value of the K0.33V2O5/PEDOT electrode was associated strongly with the K0.33V2O5 as shown in **Figure 4b**.

**Figure 5.** Calculated b-values of K0.33V2O5-layered nanofibers from CV analyses at various potentials.

Charge-storage mechanism of K-ion within the K0.33V2O5 electrode can be understood by analyzing current–voltage relationship provided by CV scans at different scan rates. Assuming that the peak current (i) obeyed the power law relationship with scan rate (v) at a given potential and was express as a combination of surface-controlled capacitive effects (i1 = k1v) and diffusion-controlled K-ion intercalation (i2 = k2v½):

$$\mathbf{i} = \mathbf{k}\_1 \mathbf{v} + \mathbf{k}\_2 \mathbf{v}^\vee \tag{6}$$

Accordingly, for the peak current response at a given potential can be described by a simple power law:

$$\mathbf{i} = \mathbf{a}\mathbf{v}^{\mathsf{b}}\tag{7}$$

where a is adjustable parameter and b is equal to 0.5 or 1 when currents are strictly dominated by diffusion-controlled K-ion intercalation or surface-controlled capacitive effects. Thus, the contributions from capacitive charge and diffusion-limited redox charge during the CV cycling were extracted precisely. **Figure 5** shows typical power law regression analyses of i vs ν plot for the both anodic and cathodic CV scans with scan rates ranging from 5 to 200 mv/s. The bvalue was calculated to be 0.738 during the anodic scan and to be 0.708 during cathodic scan, indicating that combination of both surface-controlled capacitive process and diffusionlimited redox process contributed significantly to the redox behavior of K0.33V2O5-layered nanofibers.

### **3. Highly disordered V2O5-layered electrode material for aqueous K-ion storage**

**Figure 6.** (a) TEM image of the K0.2V2O5 nanosheets. (b) SEM image of network formed by the nanosheets. (c) X-ray diffraction (λ = 0.72768 Å) of the disordered nanosheets. (d) X-ray diffraction of bulk V2O5.

Disordered K0.2V2O5 nanosheets were synthesized in an analogous way to the previously mentioned crystalline K0.33V2O5 nanofibers, utilizing an aqueous reaction at room temperature between potassium hydroxide with vanadyl (IV) sulphate hydrate to form an amorphous VOx precursor. The amorphous VOx precursors then underwent a washing and drying process and were then thermally treated at low temperature and short time in air to form the disordered K0.2V2O5 nanosheets. In **Figure 6a, b** TEM and SEM images of the disordered K0.2V2O5 nano‐ sheets are shown the sheet dimensions vary from sheet to sheet but are on the order of hundreds of nanometres. When loaded in more concentrated manner, they form a network of nanosheets. EDS measurements found that the atomistic ratio of the potassium to vanadium ratio was 9.731% K to 9.0269% V, resulting in the chemical formula of K0.2V2O5.

The XRD spectra of the disordered nanosheets in **Figure 6c** contained only a few very broad peaks due to the large degree of disorder within the structure. The strong low angle peak at 3.82° 2θ was indexed as the (001) plane. The d-spacing between the layers of the disordered nanosheets is 10.916, but other structural data are hard to extract due to the severe broadening caused by the disorder. In contrast, **Figure 6d** is the XRD spectra of highly crystalline bulk V2O5. It contains very sharp well-defined peaks and can be easily identified as orthorhombic structure with Pmmn symmetry and lattice parameters of a = 3.563, b = 11.50870, c = 4.375 (**Figure 7**).

**Figure 7.** X-ray atomic pair distribution function analysis: (top left) disordered nanosheets and (top right) the solved crystal structure of bi-layered [VO6] octahedral layers with a potassium intercalate. (Bottom left) bulk V2O5 and (bot‐ tom right) the bulk crystal structure.

Due to lack of crystallinity, there is much broadening of the peaks in the XRD spectra, and the atomistic crystal structure was not able to be accurately identified. Therefore, the local structure of the disordered nanosheets was characterized by total X-ray total scattering experiments and refinement of the atomic pair distribution function (PDF). The PDF is calculated by taking the sine Fourier transform of the total structure factor and is a real-space representation of inter atomic distances through a radial distribution function. PDF takes into account both the Bragg diffraction and diffuse scattering caused by local nonperiodic structural defects allowing for the investigation of disordered materials. Total scattering experiments can be done with both neutrons and X-rays allowing for good resolution between elements.

The major phase identified in the PDF of the disordered nanosheets was a monoclinic structure with a C2/m symmetry. It was made of V2O5 bilayers consisting of distorted [VO6] octahedra and a potassium intercalate within the bilayers. The bulk orthorhombic V2O5 phase consisting of [VO5] polyhedra was also detectable in the PDF of the disordered nanosheets. Based in the refinement, the sample was comprised of 72% monoclinic phase and 28% of the orthorhombic phase by weight. The short range order in the disordered nanosheets is dominated by the monoclinic phase, which is only ordered out to 14 Å, accounting for the initial sharp decrease in the intensity of the PDF. The long range order is dominated by orthorhombic phase. The Xray PDF of the bulk V2O5 is sharper and more intense as compared to the nanosheets due to the high crystallinity. The results of the real-space refinement of the PDF and the Rietveld refinement in reciprocal space of the XRD spectra are nearly identical.

**Figure 8.** Three-electrode half-cell analyses: (a) CVs of K0.2V2O5 disordered nanosheets/PEDOT as loading of 5 and 1.25 μg, respectively. (b) Gravimetric capacitance of K0.2V2O5, bulk V2O5 and as a function of scan rate.

The electrochemical behavior of the layered K0.2V2O5 disordered nanosheets material was investigated utilizing a three-electrode half-cell with a rotating disk electrode. Aqueous inks were made with the 4:1 ratio of disordered nanosheets material to poly (3,4-ethylenedioxy‐ thiophene) (PEDOT), a conductive polymer. CV measurements were conducted in a neutral aqueous 1 M KCl electrolyte solution with a 1 V potential window from −0.1 to 0.9 V versus a saturated KCl Ag/AgCl reference electrode. CV measurements at various scan rates between 5 and 200 mV/s are shown in **Figure 8a**. Three distinct oxidation peaks are observable in the anodic scan at 0.20 V, 0.53 V, and 0.78 V. The corresponding peaks of all three redox couples can be observed in the cathodic scan at 0.16 V, 0.38 V, and 0.79 V. All of the redox peaks remain visible throughout all of the scan rates tested.

The K0.2V2O5 nanosheets exhibit excellent capacitance at low scan rates, 661 F/g (183 mAh/g) at a scan rate of 5 mV/s in 1 V potential window. A significant amount of capacitance was retained even at higher scan rates, and 334 F/g (93 mAh/g) at 200 mV/s. 52% of the specific capacitance was retained with a 40-times increase in scan rate, from 5 to 200 mV/s. Assuming one charge transfer per vanadium atom, the theoretical capacitance of V2O5 is 294 mAh/g, 0.6 electron transfer per vanadium atom. As shown in Eq. (4), the amount of potassium atoms inserted is directly proportional to the charge stored, when fully reduced at 5 mV/s the amount of potassium inserted is equivalent to K1.2V2O5 and at 200 mV/s the amount of potassium inserted is equivalent to K0.6V2O5.

**Figure 9.** (a) Anodic and cathodic redox peaks from the CV measurement of the disordered K0.2V2O5 nanosheets at 0.53 and 0.38 V, respectively, and (b) the corresponding b-value calculation. (c) Infinite sweep rate extrapolation.

To further investigate the charge storage mechanism, the data from CV measurements was further exploited with electrokinetic analysis. The scan rate (*v*) vs. peak current for the redox couple shown in **Figure 9a** was analyzed with the b-value evaluation from Eq. (6). The results of the fitting are shown in **Figure 9c**. The anodic scan had a b-value of 0.814, and the cathodic scan had a b-value of 0.865. Both processes seem to be similar in mechanism having similar bvalues, with the oxidation process associated with the anodic scan being slightly more diffusion limited than the reduction process associated with the cathodic scan. As compared to the crystalline K0.33V2O5 nanofibers, the b-value of the disordered nanosheets is closer to 1; evidence that the disordered nanosheets are less diffusion limited then crystalline nanofibers.

Further information can be obtained by extrapolation to an infinite scan rate. The total charge (qT) stored can be expressed as the sum of the charge stored at the surface by capacitive charge storage process and the charge stored in the bulk by diffusion limited diffusion process. The charge stored on the surface is from capacitive processes, that is, double-layer/under potential deposition, charge separation at the surface, and pseudocapacitive reversible near-surface redox reactions and is assumed to be constant (q∞) with varying scan rate. Semi-infinite linear diffusion is assumed for the intercalation process into the bulk and is approximated with a constant (c) multiplied by the inverse square root of scan rate (*v*−1/2)

$$\mathbf{q}\_{\rm tr} \left( \nu \right) = \mathbf{q}\_{\rm s} + \mathbf{c} \mathbf{v}^{-1/2} \tag{8}$$

The extrapolated specific capacitance for the disordered K0.2V2O5 nanosheets at an infinite scan rate is 308 F/g. Thus, meaning that capacitive processes account for 308 F/g of the charge capacitance. At a scan rate of 5 mV/s, the capacitive processes account for 46% of the capacitance and at 200 mV/s, the capacitive processes account for 92% of the capacitance.

Vanadium Pentoxide (V2O5) Electrode for Aqueous Energy Storage: Understand Ionic Transport using Electrochemical, X-Ray, and Computational Tools http://dx.doi.org/10.5772/62759 187

**Figure 10.** (a) *In situ* XRD (λ = 0.72768 Å) of the K0.2V2O5 disordered nanosheets at potentials from −0.1 to 0.9 V (vs. Ag/ AgCl). (b) The evolution of the (001) plane upon the CV cycling.

The intercalation of potassium ions was studied through *in situ* XRD measurement while electrochemical cycling. In **Figure 10a**, the full XRD spectra for all of the scans collected during the second cycle of the CV measurement are shown. During the anodic scan the (001), peak shifts to lower angle continuously stating an applied potential of 0.4 V all the way to 0.9 V. The d-spacing increased from 0.965 to 1.106 nm. The removal of potassium causes the oxidation of the vanadium from V4+ to the V5+ increasing the van der Waals repulsion between bilayers, which are made of the negatively charged [VO6] − octahedra units, in turn causing the expansion in the interlayer distance.

Conversely, during the cathodic scan, the (001) peak shifts back to higher 2θ angle, decreasing the d001 spacing from 1.106 to 0.965 nm. The shift (001) peak during the cathodic scan is also continuously occurring between 0.4 and −0.1 V. The insertion of potassium causes the reduc‐ tion of vanadium from V5+ to V4+ and causes the van der Waals repulsion between bilayers to decrease, and therefore, the d001 spacing decreases. The (001) peak returns to the starting 2θ angle, suggesting that the intercalation/de-intercalation of the K+ ions are reversible in nature.

The (020) peak also shifts during cycling but at an order of magnitude smaller than that of the (001) peak. It also shifted in the opposite direction with the (001) layered peak, increasing the d020 upon reduction, and decreasing d020 upon oxidation. This shifting is caused by the increase and decrease of the V–O bond distance. The higher oxidation state of vanadium resulted in a shorter V–O bond. As with the (001) peak, the shifting in (020) appears to be completely reversible.

#### **4. Summary**

The emergence of in situ time-resolved techniques, including X-ray, neutron, and electron miscopy tools, provides us with an unprecedented opportunity to "observe," in nanoscale at millisecond time steps, the electrochemical processes in a close-to-real operation conditions. At the same time, the advent of high-performance computational tools and advance compu‐ tational science methods are allowing us to analyze and predict electrochemical reactions at the nanoscale and use big data tool sets to "see" how an electrochemical reaction drives functionality under real operation conditions. In this study, we have shown the power of combining advanced experimental, analytical, and computational tools, which help move beyond current understanding of electrochemistry and reaction engineering, in case of highly crystalline and highly disordered V2O5-layered nanostructure, and help leap forward into complex reaction kinetics to find new ways to convert and store energy.

#### **Acknowledgements**

This work was supported by the US Department of Energy (DOE), Office of Science, Basic Energy Sciences under Award # DE-SC0010286. Use of the National Synchrotron Light Source, Brookhaven National Laboratory, was supported by the US DOE, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-98CH10886. This research used resources of the Advanced Photon Source, a US DOE Office of Science User Facility operated for the DOE Office of Science by Argonne National Laboratory under Contract No. DE-AC02-06CH11357. Research conducted at ORNL's Spallation Neutron Source was sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences, US DOE.

#### **Author details**

Daniel S. Charles and Xiaowei Teng\*

\*Address all correspondence to<EMAIL_ADDRESS>
Department of Chemical Engineering, University of New Hampshire, Durham, USA

#### **References**

### *Edited by Dongfang Yang*

This book covers selected topics in different aspects of science and technology of alkaliion batteries written by experts from international scientific community. Through the 9 chapters, the reader will have access to the most recent research and development findings on alkali-ion batteries through original research studies and literature reviews. This book covers inter-disciplinary aspects of alkali-ion batteries including new progress on material chemistry, micro/nano structural designs, computational and theoretical models and understanding of structural changes during electrochemical processes of alkali-ion batteries.

Alkali-ion Batteries

Alkali-ion Batteries

*Edited by Dongfang Yang*

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