Revitalized interest in vanadium pentoxide as cathode material for lithium-ion batteries and beyond

Author’s Accepted Manuscript Revitalized interest in vanadium pentoxide as cathode material for lithium-ion batteries and beyond Jinhuan Yao, Yanwei Li, Robert C. Massé, Evan Uchaker, Guozhong Cao www.elsevier.com/locate/ensm

PII: DOI: Reference:

S2405-8297(17)30229-5 https://doi.org/10.1016/j.ensm.2017.10.014 ENSM236

To appear in: Energy Storage Materials Received date: 9 June 2017 Revised date: 23 October 2017 Accepted date: 24 October 2017 Cite this article as: Jinhuan Yao, Yanwei Li, Robert C. Massé, Evan Uchaker and Guozhong Cao, Revitalized interest in vanadium pentoxide as cathode material for lithium-ion batteries and beyond, Energy Storage Materials, https://doi.org/10.1016/j.ensm.2017.10.014 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Revitalized interest in vanadium pentoxide as cathode material for lithium-ion batteries and beyond JinhuanYao a, Yanwei Li a,*, Robert C. Massé b, Evan Uchaker b, Guozhong Cao b,* a

Guangxi Key Laboratory of Electrochemical and Magneto-chemical Functional Materials,

College of Chemistry and Bioengineering, Guilin University of Technology, Guilin 541004, China b

Department of Materials Science and Engineering, University of Washington, Seattle, WA

98195, USA

[email protected]

[email protected]

Abstract: Revitalized interest in vanadium pentoxide (V2O5) arises from two very important developments in rechargeable batteries. One is the push on lithium-ion batteries for higher energy density batteries: using lithium metal as anode and searching for higher capacity and high voltage cathode. Using lithium metal anode eliminates the big obstacle for V2O5 cathode that does not come with lithium ions. V2O5 possesses the highest reversible capacity among known cathode materials. Another is the recent intensive research for cathode materials beyond Li-ion batteries (LIBs). In the past several years, interest in complementary alkali-ion battery technologies has seen a tremendous resurgence. Out of the set of alternative chemistries, V2O5 has seen the most 1

considerable and promising gains as a cathode for Na-ion battery (NIB), Mg-ion battery (MIB), and other metal batteries. Unlike LIBs, these systems face a set of new challenges as dictated by the properties of the transported ionic species and the consequent effects on the electrode materials. The purpose of this review is to summarize the most interesting or surprising phenomena, the important questions raised and next experimental and theoretical steps to advance V2O5 cathode materials in the field of metal batteries. This review focused on selected topics covering the influences of surface chemistry, crystallinity, doping, defects, and nanostructures on the lithium-ion intercalation properties and recent developments on other metal batteries including NIBs and MIBs. The perspectives and remaining challenges for V2O5 based cathode materials have been discussed.

TOC graphic

2

Vanadium pentoxide as cathode material for lithium-ion batteries and beyond (a)

O1

(b)

O1

O1

O3 O3 O3 1.88 Å

b

c

(c)

1.54 Å 2.02 Å V 1.77 Å

O2

O2 O3

O3

O2

2.81 Å

a

a

   

Li-ion batteries Na-ion batteries Mg-ion batteries Other metal (Zn2+, Ca2+, and Al3+) batteries

Keywords: Vanadium pentoxide; Cathode materials, Li-ion batteries; Na-ion batteries; Mg-ion batteries

1. INTRODUCTION Research for high performance energy storage devices has steadily been attracting more allure due 3

to the rapidly growing demand for high power and high energy applications such as electric vehicles (EVs) and hybrid electric vehicles (HEVs) [1, 2]. Lithium-ion batteries (LIBs), as today’s most advanced and established energy storage devices, have gained great commercial success for their wide applications ranging from portable electronic and implantable medical devices to electric vehicles and power grid storage [3, 4]. To meet the growing demands of new technologies, the next generation of LIBs must have high energy density and good rate performance, improved safety, and low cost. From a materials standpoint, the storage capacity and affordability of lithium-ion batteries are typically limited by the cathode. Therefore, developing cost-effective and high performance cathode materials is the imperative task for advancing LIB technology [5, 6]. It is for this reason that there is such a tremendous interest in developing enhanced cathode materials. Layered lithium cobalt oxide (LiCoO2) was the first cathode material used in commercialized LIBs [7-9]. Although the complete removal of lithium gives a large theoretical capacity of 274 mA h g-1, the Li1-xCoO2 structure tends to be unstable at low lithium content (x> 0.5) [10]. Therefore, the usable specific capacity of LiCoO2 is below 150 mA h g-1. In addition, LiCoO2 is very expensive and highly toxic, which raise concerns about the environmental impact and manufacturing expense of LIBs [11]. LiNiO2 has high energy density and low cost, but the synthesis of LiNiO2 is much more difficult than LiCoO2 due to the formation of nickel over-stoichiometric phase [12]. In addition, the thermal and chemical stability of LiNiO2 remain inferior to LiCoO2 [13]. Olivine LiMPO4 (M = Fe, Mn, Co) are intrinsically more stable, but exhibit extremely low electronic conductivities ( ∼10−9 s cm−1), which make for slow Li+ extraction/insertion kinetics and modest lithium storage [14]. LiMn2O4 suffers from severe 4

capacity fading due to the dissolution of Mn2+ into the electrolyte, the generation of new phases during cycling, and the related micro-strains [15]. Li-rich oxides xLi2MnO3·(1-x)LiMO2 (M = Ni, Co, Mn, etc.) can achieve high capacities of over 250 mAh g-1 with high operating voltage (>3.5 V vs. Li/Li+ on average) [16]. However, the irreversible first cycle capacity loss, poor rate capability, and severe voltage decay and capacity fade during cycling still keep this type of high-capacity cathode materials from use in practical applications [17].

With the development of nanostructures, many of the insertion materials discarded in the past have found rejuvenated interest [18-21]. Among the promising cathode materials for LIBs, vanadium pentoxide has gained great interest because of its high energy density, low cost, easy preparation, abundant sources, and adequate safety characteristics [22-25]. In particular, the theoretical capacity of V2O5 with de/intercalation of two Li+ is approximately 294 mA h g-1, much higher than those of commonly used cathode materials, making it a very promising cathode material for next-generation LIBs. For the first generation lithium ion batteries, LiCoO2 has been used as the cathode with great commercial success. Although both LiCoO2 and V2O5 are layered materials and both are toxic, V2O5 is far more complex in crystal structure, chemistry, and phase transitions than LiCoO2 [26-29]. V2O5 does not initially contain lithium ion, so it is difficult to pair with graphite anodes for a complete lithium ion battery. However, this daunting obstacle would disappear when lithium metal anode is used, once the growth of dendritic crystals is properly addressed. Current research for high energy density and low cost batteries has directed the focus on lithium metal anodes for LIBs, lithium-oxygen batteries, and lithium-sulfur batteries, and is making impressive progress [30-37]. 5

The complexity and rich chemistry, crystallinity, and morphology have made V2O5 a very interesting material for research [38-42] and for many applications, such as sensors, [43, 44] artificial muscles, [45, 46] electrochromic windows, [47-50] and actuators [51, 52]. It should be mentioned that Prof. Livage, the pioneer of the so-called “chimie douce”, played a key role to promote the development of vanadium chemistry and made outstanding contributions in the above fields [25, 26, 42, 44, 45, 53-55]. However, the most recent and intensive research in V2O5 has been focused on the applications for and beyond lithium ion batteries, as interest in complementary alkali-ion battery technologies has seen a tremendous resurgence over the past several years. While there are a few of alkali based chemistries that have been explored, the most considerable and promising gains have been seen in sodium-ion batteries (NIBs) [56] and magnesium-ion batteries (MIBs) technology [57, 58]. However, unlike their well understood LIB counterparts, these systems face a set of new challenges as dictated by the properties of the transporting ionic species and the resulting effects such has on the electrode materials. Most notably, the larger radius and different bonding features of Na or Mg ions can lead to unexpected electrochemical performance or reaction mechanism [59].

It should be noted that the complex chemistry, crystal structure, thermodynamics and transport properties make the literature scattered and incomplete, with seemingly contradictory results. Such deviations are often due to subtle changes in chemistry, crystallinity, and processing conditions. A lot of work is needed to achieve a better comprehensive understanding of this system. It is not feasible to present a comprehensive coverage of all the related topics; instead the authors focused on selected topics covering the influences of surface chemistry, crystallinity, doping, defects, and 6

nanostructures on the lithium-ion intercalation properties and recent developments on both sodium and magnesium ion batteries. The goal is that this review paper will provide readers with a thorough introduction to the field of V2O5 as a cathode material for alkali-ion and alkaline-earth ion batteries. The authors have also contributed this review paper to help the readers navigate through the incomplete and even contradictory research reported in literature, so that more researchers can spend their time to make more meaningful efforts to drive the field forward. This review is not positioned to draw a final conclusion as an authoritative textbook, but to help and inspire more creative work in this complex compound for and beyond lithium-ion batteries.

Some review papers on vanadium pentoxide have been published in literature. For example, Whittingham [23] published an authoritative review with detailed descriptions of the fundamentals and research done on vanadium pentoxide for lithium-ion batteries about 12 years ago. Wang and Cao [60] published a review about 10 years ago focused on the developments in the fabrications and characterizations of various vanadium oxide nanostructures, such as nanowires, nanobelts, nanorolls, and ordered arrays of nanorods, nanotubes, and nanocables for significantly enhanced lithium intercalation performance. The review had a very limited coverage on the fundamentals and battery performance. Chernova et al. [61] published a review discussing the relationship between the structures and properties of layered vanadium and molybdenum oxides for LIBs and electrochromics about 8 years ago. McNulty et al. [18] recently published a review that detailed the synthesis and structural properties of vanadium oxides, and related polymorphs, bronzes and phases focused mainly on vanadium oxide aerogels and vanadium oxide nanotubes (VONTs), without covering a lot of important experimental strategies, novel nano and hybrid structures, and 7

the impacts of doping and defects. Huang et al. [62] published a review paper in 2015 summarizing the recent progress in the synthesis of V2O5 cathode materials in terms of the fabrication of nanostructured materials, carbon hybridization, and cation doping.

The purpose of this review is to summarize the most interesting or surprising phenomena, the important questions raised and next experimental and theoretical steps to advance the V2O5 cathode materials in the field of alkali-ion and alkaline-earth ion batteries. This review is structured as follows: First, the crystal structure and basic lithium intercalation behaviors of V2O5 are summarized. Secondly, the strategies frequently utilized to improve the electrochemical performance of V2O5 are discussed in detail; these strategies include: (1) nanostructured V2O5 with controlled morphology capable of enhancing the electrode transport kinetics and suppressing the structural breakdown typically induced by volume changes; (2) heterogeneous nanostructures, such as surface coatings, composites, and other complex nanostructures that address the problems related to the ionic conductivity, electronic conductivity, and structural stability of V2O5 by integrating multiple nanocomponents tailored to meet a specific demand; (3) cation doping to improve the structural stability and enhance electrical conductivity; (4) the intentional introduction of defects in the V2O5 crystal structure to improve the electronic conductivity and alkali-ion storage capacity; (5) structural modification through the use of a polymorphic phase, varying crystallinity, or expanded interlayer distance to mitigate degradation while at the same time promoting the overall capacity and reversibility of heavier alkali species. Through many typical examples for each approach, we show how these strategies address the problems encountered when developing V2O5-based cathodes for alkali-ion and alkaline-earth ion batteries. Furthermore, 8

the future perspectives and remaining challenges for V2O5 based cathode materials are also discussed.

2. HOST STRUCTURE OF V2O5 V2O5 is a typical intercalation compound with a layer structure in which the layers are held together by weak van der Waals forces. It most commonly crystallizes with an orthorhombic structure belonging to Pmmn space group. A complete unit cell is comprised of two formula units, the atomic positions and the crystal structure parameters of which are listed in Table 1 [63].

Table 1. Experimental crystal structure parameters of orthorhombic V2O5 [63]. Lattice constants (Å)

a

b

C

Atom O1 O2 O3 V

11.512 x 0.1403(4) −0.0689(3) 1/4 0.10118(8)

3.564 y 1/4 1/4 1/4 1/4

4.368 Z 0.531(1) 0.003(1) 0.001(2) 0.8917(2)

Wyckoff position (4f) (4f) (2a) (4f)

Figure 1 shows the crystal structure of orthorhombic V2O5, which consists of edge- and corner-sharing layered square pyramids, leading to V2O5 sheets linked together via weak vanadium-oxygen interactions orthogonal to the c-direction [64]. In a single-layer slab, there are three crystallographically distinct oxygen centers denoted as O1, O2, and O3. The single coordinated terminal/apical vanadyl oxygen atom, O(1), has a relatively short V-O bond length of about 1.54 Å; the bridging O2 oxygen connects two adjacent vanadium atoms via corner-sharing 9

VO5 square pyramids and the corresponding V-O bond length is approximately 1.77 Å. The triply coordinated chaining oxygen atom, O(3), links three vanadium atoms via edge-sharing VO5 square pyramids and the three corresponding V-O bond lengths are 1.88, 1.88, and 2.02 Å. Alternatively, the orthorhombic V2O5 structure can be described as highly distorted VO6 octahedral building blocks linked together to form the layered anisotropic structure [65]. Adding one more oxygen atom at the base converts the square-pyramids into a distorted octahedra as shown in Figure 1c. The peculiar feature of this polyhedron is that the short vanadyl V-O bond remains nearly unchanged and oriented along the [001] direction, as previously mentioned. The additional oxygen moves into the opposite position and forms a weak V…O van der Waals bond with a length of approximately 2.81Å.

(a)

O1

(b)

O1 O2

O2 O3

b

c

O3

(c)

O1

1.54 Å 2.02 Å V 1.77 Å

O3 O3 O3 1.88 Å

O2

2.81 Å

a

a

Figure 1. Crystal structure of orthorhombic V2O5. Views from (a) ac plane and (b) ab plane. (c) The coordination environment around a single V atom. The boxed region indicates the unit cell and the dotted line shows the sixth, electrostatically much weaker, V…O bond. V atoms are in yellow and O atoms are in blue.

Another typical form of vanadium pentoxide is hydrated vanadium pentoxide (V2O5•nH2O xerogels) [26], which could be converted into orthorhombic V2O5 by heat treatment at temperatures above 320 oC [66]. V2O5•nH2O xerogels consist of V2O5 bilayers (two layers, each 10

having the V2O5 stoichiometry) separated by water molecules as shown in Figure 2 [67]. The distance between the closest approach bilayers stacks is about 11.5 Å. When the V2O5•nH2O xerogel intercalates or extracts guest species, this distance will expand or contract correspondingly. The distance between the two single sheets in the V2O5 bilayer slab is about 2.90 Å. The coordination environment of V atoms in each bilayered slab is taken as octahedral. The VO6 octahedra shares edges to from double chains, which then arrange in parallel and side by side by interchain V−O bonds through sharing the corners of octahedra to form the slab. Water plays a fundamental role in stabilizing the thin bilayers, as proposed by the interpretation of the XRD measurements [68]. Thermal analysis [69] showed that there are three types of water in the V2O5•nH2O xerogels. Approximately 1 H2O per V2O5 is weakly bound and n can be reversibly varied between about 1.6, at ambient conditions, and 0.6-0.5 at low water activity, such as vacuum or heating to l00−120 °C. Below n = 0.5, the dehydration is less reversible. Most of the remaining water can be gradually removed by further heating, but a small amount is chemically bound to vanadium (about 0.1 H2O per V2O5) and is only liberated by heat treatment at temperatures above ~320 °C where the xerogel is converted to orthorhombic V2O5.

~2.90 Å water molecules

~11.5 Å

c a 11

Figure 2. Schematic of the crystal structure of V2O5•nH2O xerogel.

3. LITHIUM ION DE/INTERCALATION IN V2O5 Due to the weak van der Waals forces between the adjacent layers, the V2O5 easily undergoes various phase transitions (α, ε, δ, , and  phases) depending on the degree of Li-ion intercalation (x) [28]. These phase transitions are reflected in the discharge profile of V2O5 as three distinct potential plateaus at about 3.4, 3.2, and 2.3 V that correspond to the α/ε, ε/δ, and δ/ two-phase regions, respectively (Figure 3) [28].

Figure 3. Evolution of phases with the degree of lithium intercalation into V2O5 and the cycling of

-phase; collected at 180 μA cm-2. Reprinted with permission from ref [28]. Copyright Elsevier B.V.

Initially, about 1% of lithium is intercalated in crystalline V2O5 leading to the formation of α-phase LixV2O5 (x< 0.1), which is succeeded by the ε-phase LixV2O5 at 0.35
Both the α-to-ε and ε-to-δ phase transitions can be reversible under reshuffling of V-O bonds [28, 70]. When the lithium intercalation increases up to the range of 1
-LixV2O5 (2
13

(a)

(d)

(b)

(c)

(e)

Figure 4. Comparison of the arrangement of VOn polyhedral in (a) V2O5, (b) a-LixV2O5, (c) ε-LixV2O5, (d) δ-LixV2O5, and (e) -LixV2O5. Reprinted with permission from ref [29]. Copyright 1991 Elsevier B.V.

Table 2 summarizes the space group and lattice parameters for the electrochemically formed a-V2O5, ε-LixV2O5, δ-LixV2O5, and -LixV2O5 [76]. The structure of the V2O5 layers is rather similar in the pristine V2O5 and in the a-V2O5, ε-LixV2O5, and δ-LixV2O5; the decrease in the a parameters can be attributed to the increased puckering of the layers as shown in Figure 3. Moreover, the increase of the number of inserted lithium between the layers is responsible for the increase in the c parameter. Upon incorporation of lithium in the interlayer sites, the apical oxygen of the VO5 square pyramid tends to form coordinative interactions with the lithium, thereby pinching together the apices of the square pyramids and initiating the puckering of the layers [77, 78]. In this structure, the apical V-O bond lengths are extended and the corner shared oxygens are nudged slightly into the interlayer space (Figure 4). For the a-phase, these alterations are slight

14

and yield only modest changes in the unit cell parameters. However, much more pronounced puckering is evidenced with continued lithium incorporation, resulting in stabilization of the ε-phase (Figure 4c). In ε-LixV2O5, there is a considerable increase in the c parameter from 4.47 Å for x = 0.3 (ε phase) to 4.53Å for x = 0.6 (a metastable highly lithiated ε′ phase) [79, 80]. The Li ions occupy cubooctahedral cavities enclosed by eight oxygen atoms of one V2O5 layer and four oxygen atoms of the adjacent layer [81]. There is a concomitant decrease of the a parameter as a result of the puckering of the sheets, which brings the V2O5 polyhedra in closer contact [82]. With further incorporation of lithium, the δ-phase is stabilized as depicted in Figure 4d. The inserted lithium ions have a tetrahedral local coordination geometry involving interactions between downward pointing vanadyl oxygens and distorted chained oxygen that pushes upwards from the layer below [82]. Most notably, alternate layers glide by half a unit cell parameter (b/2) along the crystallographic b axis, thereby resulting in a doubling of the unit cell parameter to 9.91 Å. None of these changes from the parent orthorhombic phase up to the δ-phase depicted in Figure 4 invoke cleavage of the covalent V-O bonds and thus proceed quite readily through atomic displacements and are entirely reversible. However, beyond x = 1, the -phase is stabilized (Figure 4e) with much more severe puckering involving bond-breaking and rotation of the VO5 square pyramids to yield an alternating up and down coordination instead of the alternating pairs of VO5 pyramids observed at the lower extent of lithium incorporation (Figure 4a-d).

Table 2. Lattice parameters for a-V2O5, ε-Li0.3V2O5, ε-Li0.6V2O5, δ-LiV2O5, and -LiV2O5 electrochemically produced phases [76].

15

a (Å) b (Å) c (Å)

a-V2O5 (Pmmn)

ε-Li0.3V2O5 (Pmmn)

ε-Li0.6V2O5 (Pmmn)

δ-LiV2O5 (Amma)

-LiV2O5

11.51 3.56 4.37

11.40 3.56 4.47

11.40 3.56 4.53

11.20 3.56 9.91 (c/2 = 4.95)

9.64 3.60 10.60 (c/2 = 5.3)

(Pnma)

3.1. Initial lithium de/intercalation The room temperature lithium intercalation into V2O5 allows for the successive acquisition of the a-LixV2O5 (x< 0.1), ε-LixV2O5 (0.35
16

Figure 5. (left) Tentative phase diagram of LixV2O5 system and (right) galvanostatic intermittent discharge curve of V2O5 under the current density of 180 μA cm-2 and relaxation criterion of 0.1 mV h-1. Reprinted with permission from ref [28]. Copyright 1994 Elsevier B.V.

3.2. Formation of -LixV2O5 from δ-LixV2O5 The irreversible structural transformation which leads to the formation of a -LixV2O5 from δ-LixV2O5 occurs when discharge is carried out beyond x=1. Figure 6 shows the structural relationship between δ-LixV2O5 and -LixV2O5 [71]. The arrangement of the double chains of VO5 pyramids in δ-LixV2O5 is similar to that of the V2O5 framework, but with a slight puckering of the layers. In the -LixV2O5, layer puckering is more pronounced than that in δ-LixV2O5 and the VO5 pyramids alternate up and down individually instead of forming pairs as in δ-LixV2O5. Puckering of the δ-LixV2O5 layers corresponds to tilting of the VO5 pyramids which tends to bring the apical oxygen (O1) atoms of two neighboring pyramids closer to each other. Modifying one chain out of three in this way leads to the atomic arrangement found within the layers of the -LixV2O5. Obtaining the -LixV2O5 would require a preliminary translation of b/2 of a whole layer with 17

respect to its two neighbors.

Figure 6. Structural relationship between δ-LixV2O5 and -LixV2O5, showing the possible shift of apical oxygen atoms and the consecutive rearrangement of the VO5 pyramids. Reprinted with permission from ref [71]. Copyright 1995 Elsevier B.V.

The formation process of -LixV2O5 from δ-LixV2O5 depends on several parameters such as the discharge depth, the cycle number and the crystallite morphology [71]. The δ-LixV2O5 to

-LixV2O5 transformation proceeds gradually beyond x=1 where the amount of -LixV2O5 increases with increasing depth of discharge and cycle number. Figure 7a presents the first, second, and twentieth dis/charges curves for V2O5 to -LixV2O5 [76]. Considering the first cycle, the structural changes induced by Li+ intercalation into V2O5 are well known in the composition range 0 1, the electrochemical lithium intercalation is no longer reversible, as illustrated by the appearance of a new redox process located at higher 18

potential about 3.6 V in the first charge curve. This higher plateau at about 3.6 V can be ascribed to the formation of γ′-LixV2O5, whose structure is isomorphic to the high-temperature form

-LixV2O5 [29, 86]. While the first discharge curve is characteristic of V2O5, the first part of the second discharge curve exhibits a higher voltage, showing unambiguously that a new phase (γ′) is present within the electrode. Lithium is then intercalated or deintercalated in a mixture of α and γ phases during further cycling of the electrochemical cell. The complex phase transition mechanism of V2O5 is also reflected in the cyclic voltammogram as shown in Figure 7b. The three main processes in the 4–2.2 V potential range for lithium intercalation in the oxide are clearly pointed out as reduction peaks located at 3.4, 3.2, and 2.3 V, corresponding to the two-phase regions α/ε, ε/δ and δ/γ, respectively. Three related peaks appear during the anodic scanning, located at about 3.5, 3.4 and 2.5 V, respectively. However, the electrochemical behavior is not completely reversible, as indicated by the emergence of a new redox system observed after the first intercalation process. Indeed, an additional cathodic peak is observed in the high potential region (at 3.6 V), in accordance with the irreversible formation of the γ/γ′ system as evidenced in Figure 7a.

(a)

(b)

Figure 7. (a) Charge-discharge cycles collected at 0.2 mA cm-2 and (b) cyclic voltammogram 19

collected at a sweep rate of 10 μV s-1 for LixV2O5 cells. Reprinted with permission from ref [76]. Copyright 2012 John Wiley & Sons, Inc.

3.3 De/intercalation to -LixV2O5 The deep discharge of V2O5 leads to the intercalation of more than two lithium per unit formula. Delmas et al. [87] investigated the electrochemical performance of -LixV2O5 in 1991. They found that the electrochemical intercalation of lithium in V2O5 down to a voltage smaller than 1.9 V (vs. Li/Li+) leads irreversibly to the formation of -LixV2O5 (x~3), but otherwise exhibits very good electrochemical behavior. The new structural framework characteristic of -LixV2O5 is maintained during lithium deintercalation, as shown by the occurrence of a solid solution and by the very good reversibility of the intercalation process. Leger et al. [88] also investigated the structure and electrochemical performance of -LixV2O5 in 2005. The first and second discharge-charge cycle curves obtained in a potential window between 3.8 and 1.5 V are shown in Figure 8a. The curves are almost identical to those reported by Delmas et al. In addition to the 0.4 moles of lithium trapped in the -LixV2O5 after the first charge, the second discharge does not allow recovery of the maximum lithium uptake achieved at C/20 rate in the first discharge, so that the final composition -Li2.65V2O5 is obtained. This evidence that some of the lithium intercalated during the first discharge is removed during the first charge indicates that the -LixV2O5 phase is indeed irreversible. A typical cyclic voltammogram of V2O5 is presented in Figure 8b. The four main processes for lithium intercalation in the oxide are clearly pointed out as current peaks located at 3.4, 3.2, 2.35, and 1.95 V, corresponding to the two-phase regions α/ε, ε/δ, δ/γ, and γ/, 20

respectively, while one broad peak at 2.85 V appears in the anodic scan, consistent with the irreversible formation of the  phase. They also performed XRD characterization on the two composite electrodes with the -Li3V2O5 and -Li0.4V2O5 compositions. Constrast to what happens in the α, ε, δ, and γ phases of LixV2O5 system (x < 2), the -LixV2O5 system exhibits disordered structure (and/or small particle sizes). For -Li3V2O5 compound, all the diffraction peaks can be indexed on the basis of a tetragonal symmetry with the following unit cell parameters: a=b= 9.20 Å, c= 4.09 Å, which is in good accord with that obtained by Delmas et al. [28, 87] Clearly, the -LixV2O5 has a tetragonal symmetry and not a cubic one, as suggested in the previous report [89]. The new tetragonal structure (-LixV2O5) electrochemically formed during the first discharge is retained after the oxidation process. The -Li0.4V2O5 phase exhibits a tetragonal structure with slightly smaller parameters a=b= 9.17 Å and c= 4.09 Å. This fact most likely means that lithium ions are responsible for the ordering of the structure.

Figure 8. (left) First and second discharge-charge curves of LixV2O5, (right) cyclic voltammogram of LixV2O5 (sweep rate 10 mV s-1). Reprinted with permission from ref [88]. Copyright 2005, The Electrochemical Society.

21

4. Approaches to improve the electrochemical properties The theoretical capacity of V2O5 with two Li-ion de/intercalations is approximately 294 mA h g-1. Moreover, V2O5 has the advantages of low cost, easy preparation, abundant sources, and good safety properties, which make it a very promising cathode material for next-generation LIBs [62, 90, 91]. However, the practical application of V2O5 has been hampered due to its poor cycling stability, low lithium ion diffusion coefficient, and moderate electrical conductivity [24, 92-94]. These disadvantages slow down the kinetics of both ion and electron transports in electrodes, and the kinetics problems would be more severe in the case requiring high reversible capacity at large current densities. In the past decades, inspiring progress have been accomplished relying on the fabrication of nanostructured materials, doping, heterogeneous structures, and the introduction of point defects in crystal structure. Generally, the construction of nanostructured materials can effectively shorten the diffusion path length for both ions and electrons and increase surface area for de/intercalation reactions, while heterogeneous structures and cation doping significantly improve the electrical conductivity and cycling stability through the reduction of polarization. The key of these strategies is enhancing the ionic and electronic conductivity and structural stability of V2O5. Consequently, the electrochemical performances of V2O5 as cathodes are significantly improved in terms of reversible capacity, high rate capability, and long-term cycling stability.

4.1. Nanostructured V2O5 Nanostructuring has demonstrated itself to be one of the most effective strategies towards 22

producing high-performance electrode materials for LIBs [95-98]. The high specific surface area of nanomaterials offers a high contact area with the electrolyte, which may decrease electrochemical polarization, provide large active area, and improve the rate capability [99]; nanostructured morphologies significantly accelerate the transport kinetics in electrodes because of the shortened lithium ion and electron transport distances [95, 100]. The nanosized electrode active materials can sustain much higher stresses before pulverization compared to the bulk counterparts, which benefits the cycling stability of electrode materials [101]. Various nanostructured V2O5 cathode materials have been prepared and these materials showed much better electrochemical performance than micron-sized, or larger, V2O5 crystals. These nanostructured V2O5 materials typically exhibit nano-confinement along different dimensions, and can be classified into zero-dimensional (0D), one-dimensional (1D), two-dimensional (2D), and three-dimensional (3D) nanomaterials.

4.1.1. Zero-dimensional nanostructured V2O5

Zero-dimensional nanomaterials are the most commonly synthesized structure for LIB electrode materials [102]. Due to the shorter lithium diffusion path length in the nanoparticles when compared to the microsized particles, 0D nanostructure delays significant influence of the concentration polarization in the solid state in the dis/charge process and therefore results in the higher specific charge capacity in the same voltage range[103]. In addition, the nanostructured V2O5 could increase the contact area between electrode material and electrolyte, which is very helpful for obtaining a higher activity. V2O5 nanoparticles also showed improved cycling stability

23

and better rate capabilities compared with the commercial microstructured V2O5[104]. Generally speaking, the smaller diameter of nanoparticles means the larger specific surface area and shorter lithium diffusion pathways [105]. From this viewpoint, extremely small V2O5 nanoparticles are highly preferred when only concerning the advantages resulting from the nanometer size effect. However, one noticeable issue for these 0D nanoparticles is their high interfacial resistance arised from the largely increased boundary resistance between adjacent particles [106]. Furthermore, extremely small nanoparticles are likely to get agglomerated and lose their original nanostructured morphology in the process of electrochemical cycling [47, 107, 108]. The agglomeration of nanoparticles would also impose a great challenge for electrode fabrication using the conventional method. To solve these problems, preventive methods such as coating the surface of the nanoparticles with carbon or anchoring dispersed nanoparticles on the surface of a conductive matrix, such as graphene and carbon nanotubes, are commonly used. These approaches will be discussed further in section 4.2. It should be noted that all the preventive methods would obviously add the complexity in materials synthesis and processing, and result in increased uncertainties in materials properties and the cost as well.

4.1.2. One-dimensional nanostructured V2O5

Compared to the 0D nanostructures, 1D nanostructures (nanowires, nanorods, nanotubes, and nanobelts) provide a direct pathway for efficient charge transport along a micro-scale axis, meanwhile the two radically smaller nano-scale dimensions have the advantages fast rate capabilities [109, 110]. Moreover, one-dimensional nanostructures need only make a few points of

24

contact to ensure efficient electron transport, whereas nanoparticles would create many point contacts between adjacent nanoparticles and may easily become disconnected as the particles expand and contract upon charge/discharge [111]. Owing to their unique merits, one-dimensional nanostructures have been intensively studied for applications in electrochemical energy storage devices [60, 112-115]. The commonly used methods for the preparation of one-dimensional V2O5 nanostructures are template growth method [116-119], hydrothermal combined with post calcinations treatment [74, 75, 120-125], electrospinning technique [126-132], chemical vapor transport [133-135], chemical vapor deposition (CVD) [136], and heating of ball-milled V2O5 powders in air [137, 138].

Template growth method involves deposition of the material of interest, or a precursor for that material, into the pores of a microporous template membrane. After deposition, the template will be eliminated by either chemical etching or thermal annealing. By using polycarbonate filtration membrane (PC) as template, Patrissi et al. [116] have fabricated polycrystalline V2O5 nanorod arrays. They found that the resultant V2O5 nanorods delivered three times the capacity of a thin-film V2O5 electrode at a high rate of 200 C, which increased to four times higher at discharge rates between 500 and 1190 C. Single-crystalline V2O5 nanorod arrays have been successfully fabricated by employing electrophoretic deposition with radiation track-etched hydrophilic PC membrane as template [47, 117, 118]. Uniform V2O5 nanorods with a length of about 10 μm and diameters of 100~200 nm were grown over a large area with near unidirectional alignment (Figure 9a). These nanorods possess single-crystalline structure with a growth direction of [010]. The as-prepared single crystal V2O5 nanorod array electrodes have approximately five times higher 25

applicable current density than sol-gel derived film, and in a given current density, a nanorod array electrode can intercalate up to 5 times higher concentration of lithium than sol-gel derived film as shown in Figure 9a. The differences in electrochemical performance observed in single crystal V2O5 nanorod arrays and sol-gel derived film can be attributed to the differences in microstructure and nanostructure of these two different electrodes as schematically illustrated in Figure 9b. V2O5 nanorod arrays grown by electrochemical deposition are single crystal, with V2O5 layers parallel to the nanorod axis. Such structure is highly favorable for lithium intercalation and deintercalation because the surface oxidation and reduction reactions occur along the side surface of nanorods and the solid-state diffusion distance is very short (~100 nm), approximately half of the diameter of the nanorods. Moreover, such structure allows the most freedom for dimension change in the process of intercalation and extraction reactions. Sol-gel V2O5 films are polycrystalline and consist of platelet V2O5 grains with [001] exposure perpendicular to the substrate surface. Thus, the lithium intercalation and deintercalation processes would be comprised of lithium diffusion through grain boundaries, reduction and oxidation reactions at the surface of individual crystal grains, and diffusion inside individual grains. The difference in microstructure would have similar effects on the charge transport provided that the charge carrier properties do not change.

(a)

(b)

Figure 9. (a) Plot of discharge capacity versus current density for both single-crystalline V2O5 26

nanorod arrays (solid dots) and sol-gel derived V2O5 films. (b) Schematics showing the diffusion paths of Li+ during the intercalation process in both single-crystalline V2O5 nanorod arrays (left) and sol-gel derived V2O5 films (right). Reprinted with permission from ref [117]. Copyright 2004 American Chemical Society.

Hydrothermal method is an effective method to fabricate one-dimensional nanostructures [114, 139]. For example, Li et al. [140] have synthesized V2O5 single crystalline nanobelts on a large scale by hydrothermal treatment of aqueous solutions of V2O5 and H2O2. Such nanobelts grow along the [010] direction with width, thickness, and length in the range of 100−300 nm, 30−40 nm, and tens of micrometers, respectively. In the potential range of 1.5-4.0 V, the V2O5 single crystalline nanobelts display a specific discharge capacity of 288 mA h g-1 in the first cycle, which is much higher than that of commercial V2O5 powders (132 mA h g-1). The higher capacity of V2O5 single crystalline nanobelts can be ascribed to their larger specific surface area. The specific capacity of V2O5 single crystalline nanobelts reduces from 288 to 191 mA h g-1 and then remains stable in the subsequent two to six cycles.

In contrast to hydrothermal routine which needs heating step (normally120-200 °C) and high pressure (~10 MPa), Rui et al. [141] demonstrated a very facile method for the large-scale synthesis of single-crystalline V2O5 nanobelts under ambient condition by simply vigorous stirring the commercially available V2O5 powder in an aqueous NaCl solution (Figure 10). Different from the prevailing Ostwald ripening mechanism, Rui’s method involves the dissolution of large crystals followed by the formation of nanobelts with smaller size but larger surface area. When

27

used as cathode material for LIBs between 2.0 V and 4.0 V (vs. Li+/Li), the V2O5 nanobelts electrode exhibits higher reversible specific capacities, better cycling stabilities, and much enhanced high rate capability as compared to the commercial V2O5 counterpart. Additionally, binder-free bulky papers can be prepared by the intertwining V2O5 nanobelts with the acid-treated multi-walled carbon nanotubes (MCNTs). The binder-free and bulky-paper cathodes also show superior lithium storage performances at fast charge/discharge rates. The superior lithium storage properties of V2O5 nanobelts can be attributed to the following aspects: (1) the 1D nanobelts structure can not only provide large interfacial contact areas between active material but also offer a shorter lithium ion diffusion distance, resulting in an improved electrochemical kinetics; (2) the 1D nanobelts structure can alleviate the strain of volume expansion and offer excellent electronic conductivity along the longitudinal direction.

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Figure 10. (a) Schematic diagram of the synthesis and assembly of V2O5 nanobelts as cathode materials for batteries. (b–f) Structural characterizations of the as-synthesized V2O5 nanobelts. (b) Rietveld refined XRD pattern. (c) FESEM (inset: digital photograph of the solution after the reaction). Reprinted with permission from ref [141]. Copyright 2016 Elsevier B.V.

Electrospinning has been recognized as a convenient and versatile method for synthesizing ultralong nanowires with controllable diameters, lengths, compositions, and complex architectures [142, 143]. Electrospinning allows the production of nanofibers from various materials, for example organics and inorganics in different configurations and assemblies [144]. This is highly beneficial for energy devices, where inorganic materials especially metal oxides, can be prepared and electrospun. Electrospinning of vanadium oxide nanowires is attracting more and more 29

attentions, since nanostructured vanadium oxides have the potential to offer high capacities for LIBs [126-131, 145]. Wang et al. [126] fabricated porous V2O5 nanotubes, hierarchical V2O5 nanofibers, and single-crystalline V2O5 nanobelts by using electrospinning method and subsequent annealing. The preparation process was schematically illustrated in Figure 11a. Typically, the precursor nanofibers were electrospun from vanadium (IV)-acetylacetone/PVP/DMF at an appropriately high voltage. The obtained non-woven film was annealed at 400, 500, and 600 °C to afford porous V2O5 nanotubes, hierarchical V2O5 nanofibers, and single-crystalline V2O5 nanobelts, respectively. The uniqueness of this strategy was that the morphology of the nanostructures could be tailored by changing the annealing temperature (Figure 11b). The morphological evolves from nanofibers (below 400 °C) into porous nanotubes (400~450 °C), then hierarchical nanofibers (500 °C), and finally into nanobelts (600 °C). Clearly, the annealing temperature was crucial in controlling the morphology of the V2O5 nanostructures. As potential cathode materials for LIBs, the as-formed V2O5 nanostructures showed highly reversible capacities and superior cycling stablity. This excellent electrochemical performance can be attributed to their unique 1D nanostructure, which improved electrolyte penetration and facilitated lithium diffusion in the electrode.

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Figure 11. (a) Schematic of the preparation of porous V2O5 nanotubes, hierarchical V2O5 nanofibers, and single-crystalline V2O5 nanobelts. SEM images of V2O5 nanostructures annealed at: (b) 350 °C, (c) 400 °C, (d) 450 °C, (e) 500 °C, (f) 550 °C, and (g) 600 °C. Reprinted with permission from ref [126]. Copyright 2012 John Wiley & Sons, Inc.

Chan et al. [133] synthesized V2O5 nanoribbons by a chemical vapor transport method. The nanoribbons were single crystalline, with the length along the b-axis, width along the a-axis, and the c-axis corresponding to the layers making the height of the nanoribbons. Studies on the chemical, structural, and electrical transformations of V2O5 nanoribbons at the single nanostructure level revealed that the transformation of V2O5 into the ω-Li3V2O5 phase depends not only on the width but also the thickness of the nanoribbons. Transformation can take place within 10 s in thin nanoribbons, corresponding lithium diffusion constant 3 orders of magnitude faster than in bulk counterpart. Complete delithiation of ω-Li3V2O5 back to the pristine V2O5 nanoribbon was observed for the first time. The authors attributed this new observation to the ability of facile strain relaxation and phase transformation at the nanoscale. Moreover, efficient electronic transport can be maintained to charge a Li3V2O5 nanoribbon within less than 5 s. These exciting nanosize effects provide clues for fabricating high-performance LIBs for applications in EV and HEVs. Yin et al.[136] prepared porous V2O5 micro/nano-tubes via a CVD process. Such porous V2O5 micro/nano-tubes exhibited great improvement of the lithium storage performance such as a high reversible capacity of 290 mA h g-1, excellent cycling stability and reversibility. Further cyclic voltammetry measurements under different scanning rates confirmed that the intercalation/deintercalation is of high efficiency and the kinetic performance of the battery is 31

significantly improved. The improved electrochemical performance can be ascribed to the porous structure and large surface area of V2O5 micro/nano-structures, which reduces the diffusion distance of the solid-state lithium ion.

Glushenkov et al. [137, 138] prepared V2O5 nanorods from V2O5 powder by the two-stage procedure of ball milling and annealing in air. The V2O5 nanorods have a single crystalline structure elongated along the 〈010〉 direction as shown in Figure 12. The nanorod has a rectangular cross section and its surface is occupied mostly by [001] facets. Another pair of side surfaces is [100] surfaces and their relative width is smaller. Compared with the commercially available V2O5 powder (Sigma Aldrich), the nanorods show a more stable cyclic behavior in a cathode. The suppressed size and beneficial morphology (shape and crystal orientation of nanorods) helps to avoid drastic disintegration when the material is transformed into the ω-LixV2O5 phase and cycled within this phase as a cathode material. Although there are many unknown parameters when comparing lab synthesized samples with commercial products, this work does shed some light on the impacts of crystallite sizes and more importantly the facets on the lithium ion intercalation reactions and the subsequent ion diffusion. For a layer structured crystal like V2O5, the surface energies at different facets vary appreciably and thus the reactivity would be very different from one facet to another. It is also easy to understand that ion diffusion parallel to the layers would be far easier than diffusion perpendicular to the layers. Obviously more research is needed to have a better understanding of the effects of facets and the specific morphology on the intercalation reaction in V2O5.

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(d)

Figure 12. TEM characterization of a typical nanorod. (a) Bright-field image and the corresponding SAED pattern of a typical nanorod. (b) HRTEM image of the nanorod lattice. (c) 3D model revealing the typical shape of V2O5 nanorods. (d) The charge and discharge capacities of electrode based on V2O5 nanorods and on the reference powder as functions of cycle number. The cells were cycled between 1.5 and 3.5 V at current rate of 10 mA g-1. Figure 12a-c were reprinted with permission from ref [137]. Copyright 2008 American Chemical Society; Figure 12d was reprinted with permission from ref [138]. Copyright 2010 Springer Science and Business Media.

Due to the unique structural flexibility, 1D nanomaterials can be fabricated into flexible free-standing electrodes, which could provide more supports for the development of flexile batteries [146-148]. Recently, some novel flexible free-standing cathodes based on 1D nanostructured V2O5 have been reported and these intriguing flexible V2O5 cathodes exhibited quite promising electrochemical performance and potential applications [141, 149-158]. For example, Wang et al. [155] fabricated a novel free-standing self-assembled V2O5 nanobelt membrane by hydrothermal reaction combined with post calcination treatment. The as-made V2O5 nanobelt membrane can be directly integrated into lithium ion batteries serving as cathode, and 33

exhibits enhanced electrochemical performance, including good rate capability, high specific capacity, and good cyclic stability. Combination of 1D nanostructured V2O5 with highly conductive materials, like carbon nanotubes, graphene sheets, and conducting polymers, to fabricate a nanocomposite electrode is a promising strategy to enhance the mechanical property and electronic conductivity. Such composite architecture provides effective charge transport and electrode integrity, and therefore endowing the electrodes with high capacity, high rate-capability and excellent cycling stability. Seng et al. [150] prepared free-standing and flexible V2O5/MCNTs composite electrodes (Figure 13a,b) by a membrane filtration technique. The V2O5/MCNTs film electrodes displayed good rate capability and excellent cycling stability (Figure 13c,d). The superior lithium storage performance can be attributed to the completely reversible phase transitions of α-V2O5 to δ-LiV2O5, good lithium diffusivity in V2O5, and increased electronic conductivity and electrolyte penetration from the incorporated MWCNT web. Wang et al. [157] designed a flexible free-standing V2O5 nanowire/reduced graphene oxide (rGO) composite membrane for direct use as a cathode for lithium−vanadium batteries. Due to the free binder and integration of the excellent electronic conductivity of rGO and the interconnected network of ultralong V2O5 nanowires, the V2O5 nanowire/rGO composite paper electrode shows an excellent electrochemical performance, high energy density, and good mechanical integrity.

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(a)

(c)

(d)

Figure 13. (a) X-ray diffraction pattern of V2O5 free-standing film; inset: photograph of V2O5 film, which is greenish brown and highly flexible. (b) FESEM images of the film electrodes. (c) Galvanostatic charge/discharge cycling performance of the MWCNT/V2O5 free-standing electrodes at 0.17 C; inset: voltage profiles of the MWCNT/V2O5 free-standing electrodes at the second charge/discharge cycle. (d) Rate capability test of the 30 wt% MWCNT/V2O5 free-standing electrode. Reprinted with permission from ref [150]. Copyright 2011 Elsevier B.V.

When the free-standing electrodes are fabricated for flexible or wearable batteries, it is important that the electrodes be evaluated under bending conditions. From the viewpoint of practical applications, Noerochim et al. [152] designed a soft bendable cell which consisted of a free-standing V2O5–polypyrrole (PPy) cathode, a lithium foil anode, and gel polymer electrolyte (Inset in Figure 14g). The pristine V2O5 film (Figure 14a) shows straight nanowires about 80 to 120 nm in diameter and several microns in length. The the V2O5–PPy film (Figure 14b) shows 35

similar morphology to the pristine V2O5 film, but with the PPy uniformly deposited throughout entire lengths of nanowires with irregular or spherical shapes. The V2O5–PPy film is composed of very dense web-like nanowires joined together in a well packed layer, with the thickness of the flexible electrode around 10 mm (Figure 14c). Bothe the pristine V2O5 and V2O5–PPy free-standing films can be rolled up (Figure 14d,e) or bent to any curvature, and then returned to their original shape, while still maintaining their useful properties. The V2O5–PPy film presents a significantly higher modulus than the pristine V2O5 film (Figure 14f), suggesting that the PPy coating does improve the mechanical strength and greatly enhances the stiffness of the V2O5–PPy film. Electrochemical measurements showed that the V2O5–PPy film electrodes exhibit higher specific capacities and better cycling stability than the pristine V2O5 film electrode in Li batteries because of the improved electronic conductivity and the enhanced lithium-ion accessibility in the cathode. The cell was tested under repeated bending conditions for several cycles. The results show that the battery performance of the repeatedly bent cell was similar to that of the conventional cell (Figure 14g-i).

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(g)

(h)

(i)

Figure 14. FESEM images of free-standing V2O5 film (a) and V2O5–PPy film (b); cross-sectional view of free-standing V2O5–PPy film (c). Photographs demonstrating the flexibility of the V2O5 film (d) and the V2O5–PPy film (e). Stress and strain curves of V2O5 and V2O5–PPy film samples (f). (g) The cycling stability V2O5–PPy film electrode under repeated inward bending at a constant current density of 40 mA g-1 (Inset is the schematic diagram of a flexible cell for bending-state electrochemical testing.); (h) FESEM images of sample before and after repeated-bending tests; and (i) photograph of cell bent inwards at 180_ that was used to power a red LED. The LED glowed even when the battery device was bent, and the demonstration could be repeated over several cycles. Reprinted with permission from ref. [152] Copyright 2012 The Royal Society of Chemistry.

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4.1.3. Two-dimensional nanostructured V2O5

2D nanoarchitectures are of great interest in lithium based energy-storage devices, in particular LIBs, owing to the shortened paths for fast lithium ion diffusion and large exposed surface providing more lithium-insertion channels [159]. 2D nanoarchitectures are ideal frameworks for fast lithium storage, which requires stability, large active surface area, and short transport paths for lithium intercalation/deintercalation [160]. Li et al. [161] prepared two-dimensional leaf-like V2O5 nanosheets by a very facile sol-gel method combined with freeze drying technique followed with annealing in air, the preparation process is shown in Figure 15a. The low-cost raw materials (only commercial V2O5 powder and H2O2 without using any templates), high product yield (nearly 100%), and facile experimental procedures favorably enable the method suitable for large-scale production. The thickness of the V2O5 nanosheets is 60-80 nm (Figure15b). The V2O5 nanosheet actually is polycrystalline and consists of small nanorods. As shown in Figure 15c, the discharge capacities are 303, 273, 251, 219, and 160 mA h g−1 at the current densities of 50, 200, 500, 1000, and 2000 mA g-1, respectively. Even at a very high current density of 5000 mA g-1, the leaf-like V2O5 nanosheets electrode can still deliver high capacity of 104 mA h g-1, showing that the leaf-like V2O5 nanosheets structure favorably reduces the diffusion length for lithium ions and enables the high rate performance of LIBs. After 100 cycles at a current density of 500 mA g-1, the specific discharge capacity can be retained 206 mA h g−1 (Figure 15d) and the corresponding capacity fading rate is only 0.22% per cycle, indicating its outstanding capacity retention upon cycling.

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Figure 15. (a) Schematic illustration of the synthesis route of two-dimensional leaf-like V2O5 nanosheets; (b) FESEM images of V2O5 nanosheet; (c) dis/charge curves at various current densities; (d) cycling performance at a current density of 500 mA g−1, inset shows the charge/discharge curves correspond to different cycles. Reprinted with permission from ref [161]. Copyright 2013 John Wiley & Sons, Inc.

Wang et al. [162] synthesized large-area pure V2O5 nanosheets by a novel and facile dissolution-splitting method (Figure 16a-c). In ammonium persulfate solution, NH4+ ions can readily be intercalated into the interspace between the layers of bulk V2O5 to form intercalated compound (NH4)2V6O16. Meanwhile, sections of the crystal interlayer are dissolved because of the acid environment provided by S2O62-, which leads to the peeling of nanosheets from the parent bulk crystal. After the dissolution-splitting process occurs, the volume of nanosheets expands greatly. The obtained (NH4)2V6O16 nanosheets are dried and then calcinated in air to yield V2O5

39

nanosheets. Such V2O5 nanosheets are 10−20 nm in thickness and several micrometers in lateral width. In sharp contrast to the bulk counterpart, the as-prepared V2O5 nanosheets exhibit greatly enhanced lithium storage performances in terms of high reversible capacity (290 mA h g-1), good cycling stability, and high rate capability (144 mA h g-1 at 10 C and 95 mA h g-1 at 20 C) (Figure 16d,e). However, such large sheets are not preferable considering the intercalation reactions are likely to take place at the edges of the sheets, while the mass and charge transport take place along the sheets. Large sheets would reduce the reactivity and transport properties and suffer from degradation during cycles. The experimental data in Figure 16d and e demonstrated less competitive cyclic stability and rate performance, though much better than that of bulk samples. Yan et al. [163] also prepared V2O5 nanosheets were synthesized by a dissolution-splitting method. To address the deficiencies of V2O5 nanosheets (poor electronic conductivity and a strong tendency to aggregate), they proposed a facile and elegant self-organization strategy to decorate V2O5 nanosheets with Ag nanoparticles and TiO2 nanorods simultaneously. This strategy is based on van der Waals interactions and results in novel two-dimensional hybrid architectures. The addition of TiO2 enhances the Li+ intercalation rate and capacity of V2O5, while the introduction of Ag enables efficient electron transport from the current collector to the electrode, though the introduction of such electrochemically inactive materials compromises the storage capacity. The incorporation of secondary phase dopants can also effectively prevent the V2O5 nanosheets from restacking, and thus preserve the electroactive surface areas from the loss of electrochemical activity. The resulting 2D hybrid architectures exhibit better electrochemical properties than their bulk counterpart. 40

Figure 16. (a,b,c) Schematic illustration of the steps for preparation of V2O5 nanosheets. (d) Comparative cycling performance of the two samples at a current density of 300 mAg-1. (e) Rate performance at different current densities from 0.1 A g-1 to 6 A g-1. All measurements were conducted with a voltage windowof 2−4 V. Reprinted with permission from ref [162]. Copyright 2012 John Wiley & Sons, Inc.

Rui et al. [164] prepared few-layer V2O5 nanosheets with a thickness of 2.1-3.8 nm through direct exfoliation of bulk V2O5 crystals in formamide solvent as illustrated in Figure 17a. Firstly, bulk V2O5 was swelled by intercalation of formamide molecules into the interlayer space, resulting in significantly weakened interlayer attraction. Subsequently, ultrasonication was introduced to impose a transverse sliding force on the swollen phase, leading to the exfoliation of the V2O5 layers. Used as a cathode material for LIBs, these ultrathin V2O5 nanosheets display large reversible capacity, high Coulombic efficiency, and stable cycling. More importantly, it shows ultrahigh rate capability with a capacity of 117 mA h g-1 at a charge and discharge rate of 50

41

C as shown in Figure 17b. However, it is not clear if surface adsorption of lithium ions via formation of the electric double layer has made an appreciable contribution to the storage capacity. The specific surface area and lateral size of the nanosheets would be important information for a more sophisticated interpretation of the results.

Figure 17. (a) Exfoliation of layered bulk V2O5 into few-layered nanosheets. (b) Rate capability of few-layer V2O5 nanosheets in the voltage range of 4.0-2.05 V at various charge and dischargerates (1 C = 294 mA g-1). Reprinted with permission from ref 2. Copyright 2013 The Royal Society of Chemistry.

An et al. [165] prepared ultrathin V2O5 nanosheets through a supercritical solvothermal reaction followed by annealing treatment as shown in Figure 18a. As cathode material for LIBs, the ultrathin V2O5 nanosheets show high-rate and long-life capability at 2.4~4.0 V with initial capacity of up to 90 mA h g-1 at 15 C and excellent cyclic stability with little capacity loss after 200 cycles (Figure 18b,c). The excellent rate capability is partly attributable to the ultrathin two-dimensional nanostructure that facilitates fast lithium diffusion in the electrode. Another important factor is the voltage range used in this study, within 2.4−4 V, only one lithium ion would be de/intercalated into V2O5 with the theoretical storage capacity of 147 mAh/g. There will be a very small change of crystal structure and it is relatively common to achieve excellent cycling 42

stability in nanostructured V2O5.

a

b

c

Figure 18. (a) Schematic illustration of the procedure for fabricating the ultrathin vanadium oxide nanosheets. (b) Rate performance of ultrathin V2O5 nanosheets and V2O5 microsphere cathodes at various rates from 1 C to 15 C. (c) Cycling performance of the ultrathin V2O5 nanosheets cathodes at high rates of 10 C and 15 C.The voltage range is 2.4−4 V. Reprinted with permission from ref [165]. Copyright 2013 PCCP Owner Societies.

Liang et al. [166] developed a facile one-pot solvothermal method to synthesize uniform VO2(B) ultrathin nanosheets with a lateral size over 100 μm, which can be readily transformed into V2O5 nanosheets with good structure retention, including the exceptionally small thickness and large lateral size by calcination in air. As cathode materials for LIBs, the resulting V2O5 nanosheets display high reversible capacity, good cyclic stability, and great rate capability. To further improve the lithium storage performance of 2D nanostructured V2O5, Song et al. [167] synthesized mesoporous V2O5 nanosheets by a facile hydrothermal method and subsequent instantaneous heating and calcination. These V2O5 nanosheets exhibit ultrastable capacity retention at different current density, and also show excellent rate capability, maintaining a reversible capacity of 118 mA h g-1 at 6000 mA g-1 after 1000 cycles. The remarkable performance can be attributed to the unique mesoporous nanosheet structure as well as the presence of noticeable amount of V4+ ions 43

and the attendant oxygen vacancies in V2O5, which have substantially improved electronic-ionic transport and mitigated the internal mechanical stress induced by the volume change of the material upon cycling. These results demonstrate the significant potential of mesoporous V2O5 nanosheets for high power and long life batteries.

Nanostructured thin film electrodes without any additives and binder used in powder-based electrodes have been employed as the idea system for fundamental research due to their low resistance, cleanliness, and purity [168]. The preparation and lithium storage performance of nanostructured V2O5 thin film electrodes have been studied [169-176]. The nanostructured V2O5 thin films were fabricated by a very facile sol-gel method combined with electrochemical deposition or drop casting technique followed by annealing in air. As a representative example, Figure 19a,b shows the morphologies of nanostructured V2O5 thin films prepared by cathodic deposition from V2O5 sol directly on fluorine-doped tin oxide coated (FTO) glasses following by annealing at 500 °C in air [175]. The V2O5 thin films is very uniform and comprised of small “wrinkled” flakes of 0.5−1.5μm in diameter parallel to the substrate. Each flake is composed of smaller particles 20-30 nm in diameter. The cross-section view of the annealed V2O5 thin films (Figure 19c,d) shows similar nanostructures with 20-30 nm nanoparticles and there exist 10 nm gaps separating the adjacent nanoparticles. The V2O5 thin film electrodes exhibit excellent electrochemical properties as cathodes for lithium ion intercalation (Figure 19e,f). The specific energy density is calculated as 900 Wh kg-1 for the first cycle and 723 Wh kg-1 for the 180th cycle when the films are tested at 200 mA g-1. Even at an ultra-high current density of 10.5 A g-1, the thin film electrodes still retain a good discharge capacity of 120 mA h g-1, and the specific power 44

density is over 28 kW kg-1. The high performance of such cathodic deposited V2O5 thin film electrodes can be attributed to the unique nanostructure, in which the 20-30 nm nanocrystallites provide a shorter diffusion path for lithium, and the 10 nm wrinkled gaps offers a higher surface area with more accessible intercalation sites favoring electrolyte penetration and interface reactions. Moreover, the volumetric change can be effectively accommodated by the nanostructure, the film’s mechanical integrity and stability during battery testing is good, leading to excellent cyclic stability. This facile cathodic deposition method provides a new avenue for fabricating additive-free V2O5 electrode and even flexible V2O5 electrode with good electrochemical performance. (e)

(f)

Figure 19. SEM image of the 500 °C annealed V2O5 film: (a, b) top view and (c, d) cross-section view. (e) The cyclic performance of 500 °C annealed V2O5 film cathodes at a current density of 200 mA g-1 (1.3 C). (f) The cyclic performance of 500 °C annealed V2O5 film cathodes at various current densities from 200 mA g-1 (1.3 C) to 12.5 A g-1 (70 C) for up to 177 cycles. Reprinted with permission from ref [175]. Copyright 2011 John Wiley & Sons, Inc.

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Li et al. [177] prepared a uniform V2O5 nanostructured ultrathin film by spin-coating followed by heat treatment. Ostreng et al. [178] prepared nano-structured V2O5 thin film cathodes by atomic layer deposition method. All these films show much enhanced lithium storage activity, improved high rate capability and long-term cycling stability. Such nanostructured films may have a significant contribution of capacitive storage through surface adsorption. It should be noted that the electrochemical performances of these nanostructured V2O5 thin films is directly related to their thickness since no binder and conductive agents are added in the films. Generally, thinner films have high electrochemical activities and facilitate higher gravimetric capacity; thicker films have a higher material mass but worse electrochemical activities and lower gravimetric capacity. To address the critical needs for both high power and high energy with good cyclic stability, it is very crucial to achieve a good balance between higher gravimetric capacity with thinner films and higher material mass with thicker films [179].

4.1.4. Three-dimensional nanoarchitectured V2O5

Although the 1D and 2D nanostructured materials show considerably improved electrochemical performances, a few disadvantages exist. For example, when prepared as electrodes, there is no effective interconnection between these nanostructured materials because of their low dimensionality, which often results in high electrical resistance and low mechanical stability. Although the particle size would not have any impact on theoretical packing density, the practically achievable packing density may be seriously affected by the particle size at the nanometer scale. The smaller size would also impose a great challenge for electrode processing

46

and manufacturing. Three-dimensional hierarchical nanostructures can provide a solution to the above problems due to their quasi-continuous conductive networks for both ions and electrons, which can greatly accelerate the trasport kinetics and improve the structure stability [110, 180-182]. In addition, the micro/nanostructures are believed to have better ability to suppress agglomeration [183], thus leading to improved capacity. Solvothermal and hydrothermal are two most commonly used methods to synthesize three-dimensional hierarchical nanostructures, such as porous microspheres, hollow microspheres, yolk-shelled and multi-shelled microspheres, microflowers, and hollow microflowers.

It is known that the volumetric energy density of the typical dispersed nanoparticles is very low due to its low packing density. Microspheres with porous structure are possible candidates to surmount this major limitation for developing lithium ion battery materials with high volumetric energy density. Wang et al. [184] synthesized monodisperse V2O5 microspheres with a porous structure by hydrolysis method and subsequent reduction/oxidation treatment at high temperatures. Figure 20a shows the morphology of the synthesized porous monodispersed V2O5 microspheres. The microspheres contain broadly distributed pores with sizes less than 30 nm and have a high specific surface area of 31.2 m2 g-1. This porous V2O5 used as a cathode material for LIB shows a much improved electrochemical performance with highly reversible capacity and good cycling stability in the voltage ranges of 2.05 V to 4.0 V and 2.5 V to 4.0 V (vs. Li+/Li) (Figure 20b,c). However, the cycling performance in the voltage range of 1.5 V to 4.0 V (vs. Li+/Li) was very poor due to the irreversible formation of -phase LixV2O5 (2 < x < 3). Shao et al. [185] fabricated hierarchical V2O5 microspheres composed of stacked platelets (Figure 20d-f). The preparation 47

procedure involves a room temperature precipitation of V(OH)2NH2 precursor in aqueous solution and subsequent calcination. The obtained hierarchical V2O5 microspheres display a high discharge capacity, excellent rate capability (223 mA h g-1 at a current density of 2400 mA g-1), and good cycling stability (200 mA h g-1 after 100 cycles) as cathode materials for LIBs in the voltage range of 2.0 V to 4.0 V (vs. Li+/Li). Obviously, the hierarchical structure of V2O5 microspheres accounts for its good cycling performance. The interstices formed between the adjacient primary plates can effectively accommodate the volume variation of V2O5 upon cycling. The secondary microspherical structure can suppress aggregation of the primary plates.

(b)

(a)

(d)

(c)

(e)

(f)

Figure 20. (a) SEM image of the V2O5 porous microspheres; the inset highlights the porous structure. (b,c) Cycling performance of V2O5/Li cells (0.2 C) at the voltage range of 2.05−4.0 V and 2.5-4.0 V, respectively. (d) SEM and (e) TEM images of the as-prepared hierarchical V2O5 microspheres. (f) Schematic illustration of the formation ofprecursor and hierarchical V2O5 microspheres from the side view. Figure 18a-c were reprinted with permission from ref [184]. Copyright 2011 The Royal Society of Chemistry. Figure 18d-f were reprinted with permission from ref [185]. Copyright 2013 American Chemical Society.

48

Zhang et al. [186] prepared three-dimensional porous V2O5 hierarchical microspheres on a large scale by an additive-free solvothermal method followed by annealing at 350 °C in air. The as-synthesized V2O5 microspheres are composed of well-defined porous nanofibers that arrange themselves in an oriented manner and form a highly porous hierarchical structure. When evaluated as a cathode material for LIBs, the V2O5 microspheres display relatively stable capacity retention at different current densities and excellent rate capability (with a capacity of 105 mA h g-1 at the 30 C rate) in the voltage range of 2.5 V to 4.0 V (vs. Li+/Li). An et al. [187] fabricated 3D porous V2O5 octahedrons via a solid-state conversion process of freshly prepared ammonium vanadium oxide (AVO) octahedrons. The porous V2O5 octahedrons diplays a capacity of 96 mA h g-1 at a current density of 2 A g-1 in the voltage range of 2.4 V to 4 V (vs. Li+/Li) and excellent cycling stability. The high porosity within the V2O5 nanoparticle aggregates forms a robust reservoir for Li+ ions, and consequently improves the diffusion kinetics in the electrode. Besides, these hierarchical porous channels offer efficient contact between the surface of the electroactive particles and the electrolyte. However, the lithium storage performances of the two 3D porous V2O5 nanomaterials in wider voltage ranges (such as 2.0 V to 4.0 V and 1.5 V to 4.0 V) were not mentioned.

Hollow micro and nanostructured V2O5 materials have drawn special interest because of their advantageous features for facile lithium ion intercalation and good cycling stability [93, 188-196]. Recently, substantial efforts have been dedicated to the facile synthesis of hollow structures with different complex interiors, such as core-shelled, yolk-shelled, and multi-shelled hollow structures [197]. For example, Pan et al. [190] reported a one-pot template-free solvothermal method for the 49

controllable synthesis of uniform VO2 microspheres with different complex interiors, including yolk-shelled and multi-shelled structures (Figure 21a-h). The interior structure and size of the uniform VO2 spheres can be controlled by changing the solvothermal reaction duration and concentration of the precursor. The resulting VO2 hollow complex structures are robust and could be readily transformed into V2O5 hollow spheres without structural deformation by calcination in air (Figure 21j,k). The formation of such multi-shelled structures might be explained by the repeated Ostwald-ripening process taking place on the pre-formed solid cores. Figure 21i shows the time-dependent interior structural evolution of the VO2 microspheres. Vanadium oxide nanoparticles are first generated by the hydrolysis of VOC2O4 and then aggregation to form solid microspheres in stage I. Then the solid spheres undergo the initial inside-out Ostwald-ripening process and transform to the yolk-shelled structure (stage II). With extended solvothermal reaction, secondary Ostwald-ripening takes place on the pre-formed solid cores, resulting in formation of the multi-shelled structure (stage III). Finally, completely hollow microspheres are obtained as a result of the thorough dissolution and re-crystallization of the less stable interior architectures (stage IV). Hence, the interior structure of the VO2 microspheres could be effectively tailored by simply controlling the reaction duration. When evaluated as cathode materials for LIBs, the resultant V2O5 hollow microspheres exhibit a high initial reversible capacity of 256 mA h g-1 at a current density of 300 mAg-1, and much better cycling performance than the solid V2O5 microspheres counterpart in the voltage range of 2V to 4 V (vs. Li+/Li). The authors attributed the excellent electrochemical performance of such V2O5 microspheres to the unique yolk-shell structure of the microspheres. First, the nanoscale building blocks in the core and the shell 50

facilitate the transport of lithium ions and electrons. Furthermore, the void space within the microspheres and the porous shell ensure efficient electrolyte penetration and increase the contact area between the active materials and electrolyte. Finally, the microspheres would effectively inhibit the aggregation of the primary nanoparticles and their dissolution into the electrolytes.

Figure 21. (a) Time-dependent interior structural evolution of the VO2 microspheres.TEM images of the VO2 microspheres prepared with different reaction durations. (a,b) 2 h; (c,d) 2.5 h; (e,f) 4 h; g,h) 12 h. The scalebars in a, c, e, and g are 0.5 μm, and the scale bars in b, d, f, and h are 200 nm. (j) FESEM image and (k) TEM image of the yolk-shelled V2O5 microspheres obtained by annealing VO2 in air. Reprinted with permission from ref [190]. Copyright 2013 John Wiley & Sons, Inc.

Wang et al. [198] fabricated multi-shelled V2O5 hollow microspheres based on the adsorption of anion by negatively charged carbonaceous microsphere (CMS) and following catalytic combustion process to remove the carbonaceous microsphere templetes (Figure 22). They authors evaluated the lithium storage performance of multi-shelled V2O5 hollow microspheres in the potential between 1.5 V and 4.0 V (vs. Li+/Li) and found that the triple-shelled V2O5 hollow microspheres (3S-V2O5-HMSs) exhibited superb high specific capacity (447.9 mA h g-1 in the first 51

cycle) and outstanding cycling stability (capacity retention of 89.84 % ater 100 cycles) at a current density of 1,000 mA g-1. Even at a high current density of 2,000 mA g-1, a reversible capacity of 331.8 mA h g-1 can still be obtained. It is worth noting that this 3S-V2O5-HMSs showed an extremely high initial specific capacity of 565.4 mA h g-1 (much higher than the theoretical value of 441 mA h g-1) at a low current density of 50 mA g-1. The authors ascribed the extra capacity exceeding the theoretical value to a capacitive behavior due to surface/interface lithium storage mechanism. As proved by spherical aberration-corrected STEM ABF (annular bright-field) and HAADF characterizations, there exist many defects on the surface of 3S-V2O5-HMSs after discharge. These defects could offer a rich supply of surfaces/interfaces and abundant additional capacitive-controlled lithium-storage, therefore giving rise to a higher actual specific capacity than the theoretical value. Though an extremely high specific capacity can be obtained at low current density, the 3S-V2O5-HMSs exhibited poor cycling stability after 50 cycles at 50 mA g-1, which might be related to an intense irreversible phase change combined with structural damage due to excessive Li-insertion.

Figure 22. Schematic of two synthesis routes of multi-shelled hollow microspheres. (i) Cation-adsorption process and (ii) anion-adsorption process. At the end of both processes, a catalytic annealing is carried out to produce the multi-shelled hollow microspheres. Reprinted 52

with permission from ref [186]. Copyright 2016 Nature Publishing Group.

Generally, the dense films can not offer sufficient penetration for electrolyte into electrodes, therefore the Li+ intercalation into thick V2O5 films would be severely hindered by slow solid state Li+ diffusion within large V2O5 particles during charge/discharge process [199, 200]. Fabricating highly porous film electrodes is one of the most effective strategies to achieve high power density of batteries. Porous structure could provide higher specific surface area, sufficient contact between active material and electrolyte, short lithium diffusion path, and effective accommodation for volume variation during charge/discharge process. Liu et al. [172] prepared three-dimensional porous nanostructured V2O5 films on FTO substrate by cathodic deposition from V2O5 sol with the assistance of block copolymer Pluoronic P123 and subsequently annealing in air. As shown in Figure 23a,b, the porous V2O5 films compose of homogeneously distributed pores of less than100 nm in diameter, as well as V2O5 size of around 100 nm. This porous V2O5 films can be readily used as cathode for LIBs without adding any polymer binder and conductive additives, and exhibited excellent electrochemical properties (Figure 23c,d). The porous V2O5 films exhibit a high discharge capacity of 160 mA h g-1 at the ultrahigh current density of 9 A g-1; when cycled at 300 mA g-1, the initial discharge capacity is 294 mA h g-1, and it maintains a high capacity of 240 mA h g-1 at the 40th cycle. Wang et al. [201] prepared carbon- and binder-free porous V2O5 thin films directly on steel substrates by an electrostatic spray deposition method followed by annealing in air at 350 °C. The morphology of the films can be tailored by changing the composition of solvent of the precursor solution, the substrate temperature, and the deposition 53

time. This three-dimensional porous V2O5 films can be directly applied as cathodes for LIBs, and exhibits high disharge capacity, low initial capacity loss at 0.5 C, and ultrahigh rate capability in the voltge range from 2.5 V to 4 V. The highly porous 3D structure is responsible for the excellent electrochemical performance.

(c)

(d)

Figure 23. (a,b) SEM images of 500 °C annealed porous V2O5 film on FTO glasses. (c) The rate performance of 150 nm and 500 nm thick porous V2O5 film cathodes at various current densities. (d) The cyclic performance of 150 nm porous V2O5 thin film cathodes at a current density of 300 mA g-1 for 40 cycles. The voltage window for charge-discharge tests ranges from 0.6 V to 1.1 V (vs. Ag/Ag+). Reprinted with permission from ref [172]. Copyright 2011 Elsevier B.V.

The enhanced electrochemical high-rate and cycling performance of these three-dimensional micro/nanostructures can be attributed to the improved kinetic properties as a result of their morphology. First, the nanoscale building blocks can ensure facile Li+ transportation by increasing the ratio of active material surface area to electrolyte. Second, the micro-sized and robust

54

secondary structure can effectively suppress the aggregation of the primary nano-particles and structural degradation upon cycling. In addition, the effective interconnections of nanoscale building blocks in micro-sized secondary structure can ensure rapid electronic conduction throughout the electrode. Finally, the porous and/or hollow space within micro-sized secondary structure can further benefit the electrochemical performance because of the better electrolyte penetration and improved capability to accommodate volume variation during cycling.

Nanostructured V2O5 offers much enhanced storage capacity, rate performance and cyclic stability. However, the influences and impacts of nanostructures are complex and vary significantly from one research group to another. In addition to the characteristic dimension and morphologies, the crystallinity, surface chemistry, and possible impurities of the nanostructured V2O5 would be sensitively dependent on the synthesis method and subsequent annealing. Variation of crystallinity, surface chemistry and impurities all are known to exert large impacts on the electrochemical properties of V2O5. Making the situation even more complex is the variation of electrochemical conditions used by different research groups. Unlike most other cathodic materials, V2O5 may insert one lithium ion (corresponding to a theoretical capacity of 147 mA h g-1), two lithium ions (a theoretical capacity of 294 mA h g-1), and even more than two lithium ions (ω phase when the voltage is extended below 2 V). This makes the comparison of different research more complicated and difficult. The different amounts of lithium ion inserted into V2O5 inevitably affect the rate performance and cycling stability. In general, insertion of one lithium ion per V2O5 with appropriately designed and fabricated nanostructures would achieve a capacity close to its theoretical capacity of 147 mA h g-1 with good rate performance and cycling stability 55

when the voltage is controlled in 2.5-4 V. When two lithium ions are inserted (with a theoretical capacity of 294 mA h g-1), both rate performance and cycling stability suffer, and nanostructures would demonstrate more pronounced impacts on capacity, rate performance and cyclic stability. In most cases, nanostructuring alone is not sufficient to achieve good electrochemical properties and lithium ion intercalation performance, so other modifications in surface chemistry, crystallinity, and microstructures are explored.

4.2. Heterogeneous nanostructured V2O5 The electrochemical performance of single-phased nanomaterials is still insufficient because of their intrinsic weaker material properties in conductivity, energy density, cycling stability, and mechanical

stability.

In

this

context,

heterogeneous

nanostructured

materials

with

multi-nanocomponents have been proposed and are currently studied as promising electrode active materials because of their synergic properties, which arise from integrating multi-nanocomponents, each tailored to meet a different demand (e.g., high conductivity, high energy density, and good mechanical stability) [110].

4.2.1. Surface coated V2O5

It has been verified that the surface structures of electrode materials are very crucial for the electrochemical performance of LIBs. Surface coating of electrode materials (i.e., core-shell heterogeneous structures) has been intensively investigated to improve the cycling stability and rate capability of LIBs [202]. The typical coating materials are carbon and conducting polymers. The key mechanisms for the positive effect of surface coating on V2O5 materials include the 56

following: (1) electron-conducting coatings facilitate the charge transfer at the surface of particles; (2) suppresses the aggregation and buffers the stress of the inner nanoscale active materials; (3) a physical protection barrier that improve chemical stability; (4) modification of V2O5 surface chemistry that improves performance; and (5) prevents V2O5 from collapsing and being crushed under harsh application conditions.

Currently, surface coating is usually combined with nanotechnology to fabricate core-shell heterogeneous nanostructures, which can suppress aggregation, offer good electronic conductivity, shortened lithium ion diffusion length, and thus guaranteeing a high rate capability [106]. Aurbach’s group [203, 204] prepared carbon-coated V2O5 nanoparticles by burning off carbon-coated V2O3 nanoparticles in air (400 °C). These carbon-coated V2O5 were found to perform better as cathode materials than bare V2O5 nanoparticles or V2O5 particles of micrometer size, in terms of high capacity, very good rate capability and stability upon storage and cycling in alkyl carbonate solutions at elevated temperatures. The authors believed that the carbon coating prevents the electrical isolation of the V2O5 nanoparticles due to interactions with solution species and surface film formation at elevated temperatures. Hence, the core-shell structure of the nanoparticles with the surface carbon layer, allows for a very good interparticle electrical contact in composite electrodes, which enables high capacity at relatively high rates. Because of the carbon coating, it is possible to benefit from the high rate of nanoparticles in lithium intercalation reactions (short diffusion length), without having the usual deficiencies of nanoparticles, such as poor active mass integrity and high surface reactivity. Xu et al.[205] synthesized Mn-doped V2O5 sheet network by a one-step polymer-assisted chemical solution method. The V2O5 particles were 57

covered with thin carbon layers, which remained after decomposition of the polymer, forming a network-like sheet structure. This V2O5 network delivers high capacity and excellent stability due to the continuous pathway for electrons and lithium ions as well as rapid electrolyte diffusion. In this case, carbon acts as a bridge to connect the particles, as well as the protective shell layer to alleviate the volume change of V2O5 during the cycling processes. Yan et al. [206] reported a 3D hollow V2O5 nano/microspheres coated with crumpled reduced graphene oxide (rGO). This V2O5@rGO composite contains only 5 wt% rGO, yields excellent high rate capability potential range from from 2 V to 4V (reversible capacities of 289 mA h g-1 at 100 mA g-1 and 163 mA h g-1 at 5000 mA g-1) and outstanding cycling stabilities (a capacity retention of about 94% after 200 cycles at 2000 mA g-1). In addition to providing high continuous electron pathway, this crumped rGO has the unique advantage of self-adaptive behavior, i.e. spontaneous unfolding-folding synchronized with cyclic volumetric expansion-contraction of core active materials [207].

Zhang et al. [208] reported a unique capillary-induction approach for the preparation of carbon-coated V2O5 nanocrystals using porous carbon as a hard template (Figure 22a). The carbon-coated V2O5 nanocrystals with particle sizes ranging from 10−20 nm aggregate together to form larger secondary particles with a diameter of several hundred nanometers (Figure 24b,c). The carbon-coated V2O5 nanocrystals exhibited dramatically enhanced reversible capacities, excellent rate capability, and googd cycling stability (Figure 24d) in the potential range from 2 V to 4 V (vs Li/Li+). The impressive electrochemical performance of carbon-coated V2O5 nanocrystals was ascribed to its unique nanostructure. The small particle size of V2O5 nanocrystals greatly shortens the diffusion and transport distance of lithium and electrons within the nanocrystals, whereas the 58

voids among the nanocrystals provide fast transport pathways for the electrolyte and lithium ions. Furthermore, the carbon-coating not only increases the electrical conductivity of the electrode but also improves the stability of V2O5 material through reducing the surface reactions of the V2O5 nanocrystals with the electrolytes during charge-discharge cycling and therefore results in the excellent rate capability and cyclic stability of V2O5 material. Zou et al. [209] fabricated a novel composite structure of V2O5 nanoparticles encapsuled in 3D networked porous carbon matrix coated on carbon fibers (V2O5/3DC-CFs) as cathode for LIBs. The V2O5/3DC-CF cathode exhibits outstanding specific capacity, excellent cycling stability (capacity retention close to 100% at 10 C after 5000 cycles), and extraordinary rate capability (capacities of 285 and 115 mA h g-1 at the current densities of 1C and 50 C, respectively) in the potential range from 2 V to 4 V (vs Li/Li+). Though the actual carbon content within the V2O5/3DC-CFs composite was not provided in the work, it can be speculated that the specific capacity based on the whole V2O5/3DC-CFs should be much lower than the reported values.

(a)

(d)

Figure 24. (a) Schematic formation of carbon-coated V2O5 nanocrystals using mesoporous carbon as template. (b) TEM and (c) high resolution TEM images of V2O5 nanocrystals, showing the dimension, carbon coating, and high crystallinity of the material. Inset: the corresponding selected area electron diffraction (SAED) pattern. (d) The cycling performance of the carbon-coated V2O5 nanocrystals at current densities of 1, 2, 5, and 10 A g-1. Reprinted with permission from ref [208]. 59

Copyright 2011 American Chemical Society.

Chemical vapor deposition (CVD) is a powerful technique that can enable a homogeneous coating with desired thickness of carbon layer. Cheah et al. [130] prepared a homogeneous conductive carbon layer coating on polycrystalline V2O5 nanofibers by using plasma-enhanced chemical vapor deposition (PECVD). The carbon-coated V2O5 nanofibers exhibit improved cycling stability and discharge capacity in ambient and elevated temperature conditions over bare V2O5 nanofibers. Recently, Kong et al. [153] fabricated an interwoven nanocable architecture constructed by V2O5 nanoparticles encapsulated into multi-graphitic nanotube (V2O5@G nanocables) by means of electrospinning technique combined with CVD (Figure 25a-e). This integrated V2O5@G nanocable film can be directly used a binder-free cathode for the fabrication of flexible LIBs and exhibits ultrafast and stable lithium storage performance (Figure 25f-i). The authors attributed the excellent performance of the V2O5@G nanocables to a harmonious integration of several distinct advantages via the multiscale system. At macro-scale (battery pack scale), the free-standing and conductive film enables the construction and direct utilization without the extra weight from conducting additives, binders, and metal substrates; At micro-scale (materials ensemble scale), the interconnection of carbon coating nanocables not only provides efficient 1D electron transport pathways but also forms a stable skeleton framework that can tolerate strain relaxation and accommodate the volume change of V2O5, which effectively mitigates the stress and protects the active materials from pulverization during cycling process; At nano-scale (individual material unit), the nanoscale diameter and the uniform dispersion of the 60

V2O5 particles as well as the large contact area between nanocables and the electrolyte enhance the lithium ion transport due to the shortened diffusion paths, leading to excellent rate ability and high power density.

Figure 25. (a) Schematic diagram of a single nanocable where V2O5 is uniformly encapsulated into the graphitic nanotube. (b) Schematic of the configuration of self-supported flexible V2O5@G membrane interwoven by nanocables. (c) SEM images of the V2O5@G membrane. (d and e) Photographs showing the self-supported flexible V2O5@G membrane under flat and bent states. (f) The scheme of the fabrication of a flexible prototype Sn-V2O5 battery. (g) The cycle performance of the full cell at a current density of 0.5 C. (h) Stability of full cell to resist bending. The open circuit voltage over 200 bending cycles. (i) Photographs of a prototype flexible Sn-V2O5 lighting a white LED device under flat and bent states. Reprinted with permission from ref [153]. Copyright 2016 The Royal Society of Chemistry.

Other than carbon coatings, some inorganic coatings, such as SiO2 and TiO2, are also reported to improve the lithium storage performance of V2O5. Xu et al. [210] fabricated SiO2-coated V2O5 nanoflake arrays on nickel foam via a hydrothermal process. The SiO2 layer could enable the

61

internal V2O5 to swell inward into the hollow space owing to its mechanical rigidity, and therefore prevent the SiO2-coated V2O5 electrode from collapsing and agglomeration under harsh application conditions. Moreover, the porous silica layer not only offered a high surface area for lithium ions and electron transfer at the electrode/electrolyte interface but also promoted the lithium ions to shuttle back and forth through the electrolyte and functional material. As a result, when evaluated as a binder-free electrode for LIBs, the as-prepared sample exhibits excellent lithium storage performance with a large reversible capacity, excellent cycling stability, and high rate capability. Atomic layer deposition (ALD) is based on sequential self-limiting reactions and provides precise control in film uniformity, thickness, composition and morphology [211]. Very recently, Xie et al. [212] prepared amorphous V2O5 (a-V2O5) thin films onto the surface of hydroxyl functionalized multi-walled carbon nanotube (CNT) paper by ALD. TiO2 as a protective coating is applied onto the surface of V2O5 by ALD. They found that the 15 cycle TiO2/50 cycle V2O5/CNT paper electrode delivers a discharge capacity of 400 mA h g-1 at 100 mA g-1, approaching the theoretical value of V2O5. The dissolution problem of vanadium which is a major hurdle that limits V2O5 for cathode applications has been fully addressed by TiO2 ALD coating, without sacrificing capacity and rate capability.

Conductive polymers such as polyaniline (PANI) [213, 214], polyprrole (PPY) [152, 215, 216], and poly(3,4-ethylenedioxythiophene) (PEDOT) [217, 218] have also been coated on V2O5 to decrease direct contact with the electrolyte, which increases their structural stability and suppresses the strain involved during charge-discharge cycles. It is important to note that for surface coating, optimizing the thickness of coating layer is very crucial since too thick of coating 62

would act as battier for lithium ion diffusion while a small amount of coating can hardly form a full coating layer on active materials [219]. The hybridization effects associated with the carbon coating, particularly the sp2/sp3 and disordered/graphene (D/G) ratios, could influence the electronic conductivity of the active materials [220]. These factors should be considered when designing carbon coated V2O5 heterogeneous nanostructures. Conducting polymer coatings were confirmed to improve the life of LIBs in the short term, while its stability and long-term performance still needs to be investigated. In particular, the roles of different conductive polymers on the electrochemical performances of V2O5 electrode remain elusive and poorly understood. Ab initio calculations and molecular dynamics may be another set of powerful tools to investigate the effect of polymer coating on the kinetics of charge transfer and mass transfer through the surface of V2O5 materials. It is known that the most important step during the charging/discharging process is the transferred electron must be reciprocally compensated by extraction/insertion of lithium ion to keep the charge balance. If the lithium ion diffusivity can not attain the requisite diffusivity of the transferred electron, it will limit the transference of electron, leading to deteriorated electrochemical performance. Therefore, an ideal coating layer for electrode material should play multi-functional roles. For example, it should be highly conductive for both lithium ion and electrons, should facilitate easy solvation/desolvation of lithium ion, and provide reliable protection against electrolyte, and so on [221]. However, such a perfect coating can hardly be found in the existing coating materials. Instead, the combination of different coating materials, e.g., to build a “hybrid coating layer”, might be an alternative in the next stage.

4.2.2. V2O5/MCNTs composites 63

Carbon materials with well-defined nanostructure can also be directly employed for the preparation of V2O5/carbon hybrid composite. Among those carbon nanomaterials, carbon nanotubes (CNT) are one of the most attractive systems [222]. Due to the covalent sp2 bonds between individual carbon atoms, CNT has extremely high tensile strength and can tolerate huge strains before mechanical failure. CNT also possess nanometer size diameter, high specific surface area, structural integrity, high electrical conductivity, and chemical stability [223]. All these fascinating properties make it an ideal matrix for hybridizing with electrode active materials to achieve LIBs with high rate capability and long cycle life [224]. For example, Zhou et al. [225] synthesized multi-walled carbon nanotubes (MWCNTs)/V2O5 integrated composite with nanosized architecture through hydrothermal treatment combined with a post-sintering process. Electrochemical tests revealed that MWCNTs/V2O5 could deliver a superior specific capacity (402 mA h g-1 during initial discharge at a current density of 100 mA g-1 between 1.5 and 4 V versus Li/Li+), good cycling stability (222 mA h g-1 after 50 cycles) and high rate capability (194 mA h g-1 at a current density of 800 mA g-1). The authors attributed the superior lithium storage performance to its porous and integrated nanostructure, which provides a large specific surface area, a good conducting network and effective buffering against the strain upon cycling. Jia et al. [226] prepared a three-dimensional nanoarchitecture consisting of mesoporous V2O5 and penetrating CNTs by an aerosol-spray drying process followed by two-step thermal annealing. Due to their nanostructured features, such nano-composites display effective charge transport, thus offering batteries with high capacities and rate-capabilities. Cao et al. [227] prepared hybrid films composed of V2O5 nanoparticles and single-walled carbon nanotubes (V2O5/SWNT) with 64

mesoporous structure using a floating CVD method followed by controllably hydrolytic deposition of V2O5 nanoparticles. As cathodes for LIBs, this V2O5/SWNT hybrid films exhibit a high-rate capacity. They also found that SWNTs can improve Li+ diffusivity by 2 to 4 orders of magnitude to significantly facilitate lithium-ion transport, achieving an excellent electrochemical performance of the hybrid cathode system. Yu et al. [228] prepared hierarchical networks with interconnected V2O5 nanosheets anchored on the skeletons of CNTs by a facile hydrothermal method and a following calcination. As shown in Figure 26a,b, flake-like V2O5 nanosheets are highly interconnected and cover the CNT skeleton (the inset of Figure 26a). The interconnected V2O5 nanosheets are in intimate contact with the CNTs, which is favorable for the enhancement of conductivity. The individual nanosheets are single crystalline with the thickness of about 10-20 nm and diameter of 50-100 nm (Figure 26c,d). The as-prepared CNTs@V2O5 showed dramatically improved electrochemical performance in terms of long cyclic stability at high rate and very good rate capability when they are used as cathode material for LIBs (Figure 26e-h). The authors speculated that the CNTs and two-dimensional V2O5 nanosheets could work synergistically to provide sufficient tolerance for the volume variation during Li+ intercalation/deintercalation, accelerate the kinetic process of ion diffusion, and exhibit higher electrical conductivity and structural stability. Cheng et al. [229] made a porous sheet-like V2O5-CNTs nanocomposite with an ordered brick-and-mortar arrangement of CNTs and ribbon-like V2O5 nanoparticles by using a simple ice-templating approach. The porous sheet-like structure with an internal CNT conducting network ensures that the electrode materials have high electronic conductivity, efficient electron transportation and large electrode/electrolyte contact area. Benefiting from the unique structure, 65

the sheet-like V2O5-CNTs nanocomposite possesses superior electrochemical performance when employed as a cathode material for LIBs. (e)

(f)

(g)

(h)

Figure 26. (a) SEM,with inset showing higher magnification, (b) TEM, and (c) HRTEMimages of CNTs@V2O5. (d) TEM images of CNTs@V2O5, the inset shows the corresponding selected area electron diffraction pattern. (e) Cycling performance and Coulombic efficiency of CNTs@V2O5 and V2O5 microflowers (V2O5-mf) in the voltage range of 2−4 V at the current rate of 1 C. (f) Cycling performance of CNTs@V2O5 and V2O5-mf in the voltage range of 2.0-4.0 V at the current rates of 20 C and 30 C. (g) Rate capability of CNTs@V2O5 and V2O5-mf at various current rates. (h) Electrochemical impedance spectra (EIS), fitted EIS curves, and the equivalent circuit (inset) used to fit the impedance data. Reprinted with permission from ref [228]. Copyright 2013 American Chemical Society.

4.2.3. V2O5/Graphene composites Graphene (a monolayer of sp2-bonded carbon atoms) is another novel carbon material with well-defined nanostructures and excellent physicochemical properties [230, 231]. It has ultrahigh specific surface area, extraordinary electrical conductivity, outstanding mechanical flexibility, and

66

good chemical stability [232-236]. These unique properties render graphene one of the most attractive carbon-based materials used for fabricating hybrid electrodes with large reversible capacity and long-term cycling stability [237-246].

More recently, various V2O5/graphene hybrid nanocomposites have been prepared and studied as cathode materials for LIBs. For example, Choi et al. [247] prepared reduced graphene ball (RGB)-V2O5 composite powders by a spray-pyrolysis process (Figure 27a). The ultrafine V2O5 nanoparticles were uniformly distributed within the porous graphene ball in the RGB-V2O5 ball (Figure 27b-f). The RGB-V2O5 composite, with graphene sheets and small V2O5 crystallites, proved to be more advantageous for lithium storage than microporous V2O5 powders (P- V2O5) without graphene sheets prepared by the same process (Figure 27g,h). The RGB-V2O5 composite had higher initial charge and discharge capacities and better cycling and rate performances than the microporous V2O5 powders (P-V2O5). The impedance analysis showed that the RGB-V2O5 composite was more amenable to lithium ion diffusion than the microporous V2O5 powders (P-V2O5). The authors believed that the high electrical conductivity and small grain size of the RGB-V2O5 composite enhanced the lithium ion diffusion rate and facilitated a fast charge/discharge process.

67

(a)

(g)

(h)

Figure 27. (a) Schematic diagram of formation mechanism of the RGB-V2O5 composite powder by spray pyrolysis. (b,c) FE-SEM images, (d,e) TEM images and (f) SEAD pattern of the RGB-V2O5 composite powder. (g) Initial charge/discharge curves at a constant current of 1 A g-1 and (h) rate performances at various current densitiesfor the RGB-V2O5 composite powders and the microporous V2O5 powders (P- V2O5). Reprinted with permission from ref [247]. Copyright 2014 John Wiley & Sons, Inc.

Liu et al. [73] fabricated a hybrid material constructed from two-dimensional graphene nanosheets (GNS) and one-dimensional V2O5 nanowires by using hydrothermal approach. The ultralong well-ordered single crystalline V2O5 nanowires were supported on transparent GNS substrate and exhibited a higher reversible capacity and better rate capability as compared with the bare V2O5 nanowires in a potential window between 1.5 V to 4.0 V vs. Li/Li+. The authors ascribed the good electrochemical properties of V2O5/GNS to a cooperative effect between the nanowires and GNS, i.e., the continuous GNS layers supporting or wrapping homogeneously around the single crystalline V2O5 nanowires serve as smooth and fast paths for electron 68

transportation during the charge/discharge process, while the well-ordered single crystalline surface of the nanowires facilitates the migration of lithium ions between the active material and electrolyte.

Nanomaterials, especially nanoparticles, generally tend to aggregate due to the very high surface energy, which results in the loss of nanomorphology and decay of electrochemical performances. In fact, the dispersion of V2O5 nanomaterials on graphene can effectively inhibit the self-aggregation of nanomaterials while the dispersed nanomaterials can also suppress the re-stacking of graphene. Rui et al. [248] introduced a facile solvothermal approach to synthesize reduced graphene oxide supported highly porous V2O5 spheres consisting of nanoparticles. The process is driven by the heterogeneous nucleation of vanadium oxide on graphene oxide sheets and subsequent thermal pyrolysis in air to yield the V2O5/rGO composite. The V2O5 spheres were strongly anchored on the surface of rGO sheets, which facilitates the fast charge transport through the highly conductive rGO sheets. As a cathode material for LIBs, the V2O5/rGO nanocomposites exhibit a significant improvement in electrochemical performance. Han et al. [249] constructed V2O5 quantum dots/graphene hybrid nanocomposite by a two-step solution phase synthesis method. The V2O5 quantum dots have an average diameter of about 2 nm and are uniformly dispersed on the rGO. The V2O5 quantum dots/graphene exhibits an excellent cycling performance with 89% capacity retention after 300 cycles in the voltage ranges of 2.0–4.0 V vs. Li/Li+. The enhanced electrochemical properties could be ascribled to the short lithium ion transfer distance, two-dimensional electron channels, homogeneous dispersion and immobilization of V2O5 quantum dots. Cheng et al. [250] fabricated self-assembled V2O5 nanosheets/rGO hierarchical 69

nanocomposite by a simple solvothermal method (Figure 28a). In this nanocomposite, the V 2O5 nanosheets assembling on the rGO constitute a three-dimensional hierarchical nanostructure with high specific surface area and good electronic/ionic conducting path (Figure 28b,c). When used as the cathode materials, the V2O5 nanosheets/rGO nanocomposites exhibit highly reversible capacity and good rate capability relative to the bulk material. The V2O5 nanosheets/rGO nanocomposite anode was able to charge/discharge at high current densities of 3000 mA g-1(10 C), 6000 mA g-1(20 C), and 15 A g-1(50 C), with discharge capacities of approximately 138, 112, and 76 mA h g-1, respectively (Figure 28d). However, the cycling stability is still unsatisfactory (Figure 28e).

(d) (a)

(e)

Figure 28. (a) Schematic illustration of the formation of three-dimensional V2O5 nanosheets/rGO hierarchical nanocomposite. (b,c) SEM images of V2O5 nanosheets/rGO hierarchical nanocomposite at different magnifications. (d) Rate capability of V2O5 nanosheets/rGO nanocomposite electrode and V2O5 microsphere electrode at different current densities (1C = 300 mA g-1). (e) Cycling performance and Coulombic efficiency of V2O5 nanosheets/rGO nanocomposite electrode and V2O5 microsphere electrode at a current rate of 2 C. Reprinted with permission from ref. [250]. Copyright 2013 The Royal Society of Chemistry. 70

The electrochemical performance of V2O5/graphene composites could be further enhanced by introducing materials with another dimensionality, such as 1D CNT. The CNT in the ternary composites can avoid the restacking of graphene sheets, enhance the electron transport by constructing 3D electrical conductive networks, and dramatically increase the surface area by forming hierarchically porous assembly structure. Palanisamy et al. [251] fabricated a 3D V2O5/RGO/CNT ternary composite with hierarchial porous structure by using 2D RGO and 1D CNT as conductive network by microwave-assisted hydrothermal method. Compared to 2D V2O5/RGO as a control group, the 3D V2O5/RGO/CNT delivers a high capacity of 220 mAh g-1 at 1 C (294 mA h g-1) after 80 cycles and an excellent rate capability of 100 mAh g-1 even at a ultra-high rate of 20 C in the voltage range from 2.0 V to 4.0 V vs. Li/Li+. Moreover, the porous 3D V2O5/RGO/CNT structure is favorable to the Li+ diffusion into bulk and the capacitive Li+ storage on surface, suggesting that the hierarchically porous 3D V2O5/RGO/CNT architecture can be a rational design to enhance the additional surface Li+ storage along with the bulk reaction.

The presence of graphene with nanostructured V2O5 is helpful for providing a large surface area, high electrical conductivity, and fast charge transport, and therefore they significantly improve the performance of the electrode materials. It should be noted that the functionalization of graphene and control of the interface between the nanoparticles and graphene are crucial. The functionalization involves a charge transfer process that affects the performance of the composite materials. Thus, a further understanding of the surface chemistry between graphene and V2O5 nanoparticles is needed. Moreover, the mechanism of improved performance is partially unclear, 71

and the relationship between the number of graphene layers, defects, disordered features, and the lithium storage mechanism is also critical.

4.2.4. Other heterogeneous structures

Hu et al. [252] fabricated a novel nanoarchitectured electrode composed of an efficient mixed-conducting network (Figure 29a,b). In the naocomposite, most of the V2O5 nanoparticles were encapsulated within the carbon tube-in-tube (CTIT). When used as cathode material for LIBs, the V2O5/CTIT nanocomposites exhibited fast kinetics for lithium intercalation/deintercalation, highly reversible capacity, good cycling stability, and excellent high rate capability (Figure 29c). The significantly improved lithium storage performane can be attributed to the mixed-conducting nanostructure, which provides both an electronic pathway and a lithium-ion pathway.

(a)

(b)

(c)

Figure 29. (a) Schematic representation of the desired design based onan efficient mixed-conducting network. (b) Typical TEM image of the V2O5/CTIT nanocomposites showing that most of the V2O5 nanoparticles are encapsulated within CTIT. The V2O5 nanoparticles indicated by red arrows and CTIT indicated by black arrows. (c) Cycling and discharging/charging rate performance of the V2O5/CTIT nanocomposite electrode. Reprinted with permission from ref [252]. Copyright 2009 John Wiley & Sons, Inc.

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Yan et al. [253] manufactured a free-standing heterogeneous nanostructured electrode composed of SnO2/V2O5 core/shell nanowires (SVNs), which combined the advantages of better conductivity of SnO2 nanowires and short diffusion distance of the thin V2O5 layer. Figure 30a shows the synthesis strategy for the SVNs. Firstly, the SnO2 nanowires were synthesized using chemical vapor deposition on a stainless-steel substrate. Then, V2O5 was coated on the SnO2 nanowires by pyrolysis of vanadyl-acetylacetonate (VO(acac)2). The SnO2 nanowires had a smooth surface with a diameter around 40 nm (Figure 30c). The as-synthesized SVNs had a larger diameter (about 100 nm) and showed a rougher surface and particles-decorated configuration (Figure 30b,d). When used as cathode material for LIB, this SVN-based electrode showed an outstanding high rate capability and excellent cycling stability (Figure 30e,f). The improved performance was ascribed to the unique core/shell nanoarchitecture, in which the thin V2O5 layer is beneficial for fast lithium intercalation/deintercalation and the SnO2-core nanowire provides a fast path for electron transportation.

a

e

f

Figure 30. (a) Schematic diagram showing the strategy for coating V2O5 on SnO2 nanowires 73

grown on a stainless-steel substrate. The left part corresponds to the photograph of the SnO2 nanowires and the SVNs sample (the scale bar is 1 cm). (b) Typical TEM image of a single SVN. (c) FE-SEM image of SnO2 nanowires. (d) FE-SEM image of the typical SVNs (the inset indicates the diameter is about 100 nm). (e) The galvanostatic charge/discharge curves of the SVNs electrode at different current densities. (f) The cycling performance tested at a current density of 5 A g-1 (7.5 A g-1), and the inset showing the corresponding Coulombic efficiency profiles. Reprinted with permission from ref [253]. Copyright 2011 John Wiley & Sons, Inc.

Zhou et al.[254] produced a novel TiSi2/V2O5 hetero-nanostructure based on the highly conductive two-dimensional TiSi2 nanonet as a platform through simple chemical synthesis. The key to their design is the capability to control the properties of materials on multiple scales concurrently as shown in Figure 31a. At atomic scale, the Ti-doping can stabilize V2O5 during lithiation and delithiation and therefore dramatically improves the cycle life-time; at nanoscale, each component in the heteronanostructure is designed for a specific function: the TiSi2 nanonets provide pathways for fast charge transport, the Ti-doped V2O5 nanoparticles act as the ionic host, and the SiO2 coating can prevent lithium from reacting with TiSi2 nanonets. As shown in Figure 31b-d, the hetero-nanostructures showed a high yield of the nanonet and the V2O5 nanoparticles (typically 20−30 nm in diameter) were coated on the interconnected TiSi2 nanonets. Used as cathode electrode for LIBs, the resulting TiSi2/V2O5 nanostructures showed a high specific capacity of 350 mA h g-1, a power rate up to 14.5 kW kg-1, and 78.7% capacity retention after 9800 charging/discharging cycles (Figure 31e-g).

74

(a)

(e)

(f)

(g)

Figure 31. (a) Schematic illustrations of the TiSi2/V2O5 hetero-nanostructure design. The TiSi2 nanonet is directly grownon a current collector; V2O5 is deposited onto TiSi2 by ahydrolysis process. Upon annealing at 500 °C in O2, V2O5 nanoparticles form. (b) Top-view SEM, (c) TEM and (d) HRTEM image of TiSi2/V2O5 hetero-nanostructures. (e) Charge capacity and (f) rate-dependent specific capacities of TiSi2/V2O5 heteronanostructures. (g) Long term cycling performance of TiSi2/V2O5 heteronanostructure electrode at the very high rate of 25 C. (1 C = 350 mA g-1). Reprinted with permission from ref [254]. Copyright 2012 American Chemical Society.

Liu et al. [255] designed a single nanotube structure that embeds all the components of an electrochemical storage device within the pores of a nanoscale template (Figure 32d). Ru metal (current collectors) and crystalline α-V2O5 (active storage material) were conformally coated at two ends of anodic aluminum oxide (AAO) nanotube arrays by atomic layer deposition (ALD). V2O5 at one end was prelithiated to serve as the anode, while pristine V2O5 at the other end served as the cathode. Parallel nanobattery arrays with an areal density of 2 × 109 cm−2 were achieved after filling 1 M LiPF6 in ethylene carbonate (EC): diethyl carbonate (DEC) (1:1 by volume) electrolyte into the AAO nanopores. In both half-cell (Figure 32a-c) and full-cell (Figure 32d-g) configurations, such array battery exhibits excellent high rate capability (∼50% at the 150 C rate, 75

relative to 1 C), providing ∼80 mAh g−1 at 150 C (24 s charge–discharge time), and outstanding cycling stability (more than 80% of initial capacity retained after 1,000 cycles at 5 to 25 C). These excellent performances can be attributed to the nanotubular electrode design, in which a thin storage layer (V2O5) and an integrated current-collecting layer (Ru) deposited through ALD facilitate fast ion and electron transport, respectively. The results are important to elucidate ion transport in highly confined electrolyte environments, a key factor that is bound to affect any nanostructure-based energy storage approach. Moreover, the authors found that the specific capacity scales linearly with scan rates, as expected for charged surfaces or electrical double layers. This demonstrates that properly scaled nanostructures can utilize the full theoretical capacity of the charge storage material while their ion insertion processes occur at high power, much like what happens in a double-layer capacitor, indeed blurring the traditional distinction between surface effects and ion insertion.

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(b) (e)

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Figure 32. (a) SEM EDS mappings and schematic diagrams of cathode cross-section with Ru-nanotube current collector. (b) Galvanostatic charge–discharge cycling curves of V2O5/Ru-nanotube half-cell at various C rates. (c) Charge–discharge curves at the second, 1,000th and 1,800th cycles of V2O5/Ru-nanotube half-cell. (d) Schematic of parallel nanopore battery array and cross-section of a single-pore full-cell device. (e) Rate performance (capacity normalized by cathode V2O5 mass) of the symmetric full-cell device. (f) Long cycle life at a 5 C rate with a current density of 0.26 mA cm−2 and (g) charge and discharge curves at the second and 1,000th cycles of the symmetric full-cell device. Reprinted with permission from ref. [255] Copyright 2014 Nature Publishing Group.

Surface coating and admixing with carbon nanotube or graphene have unambiguously improved the lithium ion intercalation properties by 1) providing efficient electronic conduction pathways, 2) serving as a buffer and spacer to alleviate the volume variations during lithium ion de/intercalation, 77

and 3) keeping the apertures open for electrolyte access. The introduction of carbon materials with nanostructured V2O5 may also modify the surface chemistry of V2O5 by introducing oxygen vacancies if the mixtures were annealed at elevated temperatures or if the carbon coating was formed through pyrolysis of organic additives or precursors. In addition, the presence of carbon may exert some catalytic effects to promote the intercalation reactions at the interface between V2O5 and the electrolyte. But in spite of all the benefits, adding carbon materials has reached its limit for improvement of long term cycling stability. The carbon materials only introduce an efficient pathway for charge transfer outside V2O5, but not through the V2O5. The charges can only be delivered to the surface of V2O5, while the low electronic conduction within V2O5 was not tackled.

4.3. Cation doped V2O5 The introduction of heteroatoms could greatly influence the structure and the electronic state of electrode materials [256]. To overcome the capacity fading inherent with undoped nanostructures, doped V2O5 nanostructures have been attracting numerous attentions recently. Various elements, such as Ag [257, 258], Na [259], Cu [260-263], Mg [264], Ni [196], Mn [173, 265-268], Fe [269-271], Tb [272], Ti [254, 273, 274], Y [275], Al [243, 276, 277], Cr [278-280], and Sn [174, 281] have been doped into V2O5. The resultant M-doped V2O5 usually exhibited the following distinct properties: (1) the doping of cations leads to the formation of lower valence state V cations (V4+ and V3+) resulting in an increased electronic conductivity [260, 261, 280]; (2) the inclusion of metal cations into V2O5 layers leads to the formation of [MO6] octahedral units which

78

can stabilize the layer structure of V2O5 phase during electrochemical interaction of Li+, allowing better structural stability during electrochemical cycling [270, 271, 276]; (3) the local structural defects caused by the foreign metal cations may serve as possible nucleation centers for phase transformation during lithium ion intercalation/deintercalation process which enhanced cycling performance as well [278, 282]; Additionally, these structural defects may provide more pathways for Li+ diffusion in V2O5 [205, 261, 273, 282]; (4) the doping of cations could decrease the charge transfer resistance and accelerate the kinetics of Faradic reactions [261]. (5) The doping of cations may decrease the crystallite size, or alter the morphology and/or the electrochemical performances of V2O5 [174, 277].

Iida and Kanno [283] have systematically investigated the effect of various elements (Nb, Ce, Nd, Dy, Sm, Ag, and Na) doping on the growth of V2O5 films. Takeuchi et al. [284] have reviewed the preparation, characterization, and reactivity of silver vanadate bronzes (AgxV2O5) and their application in rechargeable lithium ion batteries. Coustier et al. [285] prepared Ag-doped V2O5 (aerogel-like). Compared with undoped V2O5, the electronic conductivity of Ag doped V2O5 was increased by two to three orders of magnitude. The electrochemical performance of the doped V2O5 was excellent. The enhanced electrochemical properties can be attributed to the formation of conductive metallic silver that was derived from the reduction of doped cations.

In view of practical commercial use, it is more attractive to replace Ag with less expensive elements. Among those metal cations, Cu2+, Mn2+, Cr3+, Fe3+, and Al3+ are most commonly studied as dopants for V2O5. Yu et al. [261] prepared Cu doped hierarchical V2O5 flowers by a simple

79

hydrothermal method followed by an annealing process. Compared with un-doped V2O5 flowers, the Cu doped V2O5 flowers show a higher electrical conductivity, higher lithium diffusion coefficient, and lower charge transfer resistance. In particular, the Cu doped V2O5 flowers exhibit a significant improvement in electrochemical performance in terms of highly reversible capacities, good cycling stabilities, and excellent high rate performance compared to the un-doped material (e.g. 94.0% higher at 20 C). Similar results were also observed in the Cu doped V2O5 samples synthesized by other methods [260, 262, 263]. The authors attributed the improved performance of Cu doped V2O5 to the following reasons. The doping of Cu led to the formation of V4+ cations, giving rise to a higher electronic conductivity to the V2O5; the [CuO6] octahedral chains formed in the framework of CuxV2O5 strength the three-dimensional character of the material were speculated to prevent the deformation of the material structure during charging/discharging cycling.

Yu et al. [173] prepared Mn-doped V2O5 films by a simple H2O2-V2O5 sol-gel process with the direct addition of manganese salt. They found that Mn-doping in V2O5 films results in enhanced discharge capacities with much improved cyclic stability. For example, at a current density of 68 mA g-1, Mn-doped V2O5 films exhibit a discharge capacity of 283 mA h g-1. Excellent cyclic stability with a fading rate of less than 0.06% per cycle is observed even at a very high current density of 680 mA g-1, much better than pure V2O5 films (0.8%). The authors attributed the improved lithium storage performance of Mn-doped V2O5 films to the introduction of oxygen vacancies and the formation of [MnO6] octahedra due to adding appropriate amount manganese into vanadium pentoxide. Zeng et al. [267] proposed a nanostructured Mn-doped V2O5 lithium-ion 80

battery cathode material that facilitates cathodic charge transport. The synthesis strategy uses a layered compound, vanadium (III) jarosite, as the precursor, in which the Mn2+ ions are doped uniformly between the vanadium oxide crystal layers. Through a two-step pyrolysis/oxidation transformation, the vanadium jarosite was converted into Mn2+-doped V2O5. The doping of larger Mn cations in the modified V2O5 structure increases the cell volume, which facilitates diffusion of Li+, and introduces oxygen vacancies that improve the electronic conductivity. Comparison of the electrochemical performance in LIBs of undoped and the Mn2+-doped V2O5 hierarchical structure made from layered vanadium jarosite confirms that the Mn-doping improves ion transport to give a high cathodic columbic capacity (253 mA h g-1 at 1 C, 86% of the theoretical value, 294 mA h g-1) and good cycling stability.

The Pereira-Ramos group conducted a series of studies on electrochemical properties of Cr doped V2O5 (Cr0.11V2O5.16) prepared by sol-gel approach [278, 279, 286, 287]. They found that the Cr0.11V2O5.16 exhibited a higher capacity and better cycling stability than the un-doped V2O5. The orthorhombic Pmmn symmetry of the pristine material is kept upon lithium intercalation in the LixCr0.11V2O5.16 compounds over a wide lithium composition range (0 < x ≤ 2). From the unit cell parameters variation as Li accommodation proceeds, the host lattice can be described on the basis of a ε-related phase characterized by a linear cell volume expansion of 7% in the voltage limits 3.8-2 V. Such a structural evolution strongly contrasts with the un-doped LixV2O5 phase diagram which predicts the appearance of the successive phases α, ε, δ, γ. They ascribe such a single phase behavior of LixCr0.11V2O5.16 to the presence of short chains of chromium oxide in V2O5. These limited structural changes of Cr0.11V2O5.16 during lithium interaction/deintercalation 81

allow the complete recovery of the initial structure after the charge process. Such a moderate structural behavior of Cr0.11V2O5.16 ensures a high electrochemical reversibility. Therefore, strong modification of the Li-V2O5 phase diagram can be induced through appropriate structural modifications of the host lattice. These modifications can be due to the existence of short [CrO6] octahedra chains linking V2O5 layers in the sol-gel Cr0.11V2O5.16 compound, which increases its three-dimensional character and prevents the significant deformations of the framework such as layer puckering, slab gliding, bond breaking, and polyhedra rotation during electrochemical lithium intercalation. Zhan et al. [282, 288] also studied the electrochemical performance of Cr doped V2O5 synthesized by an oxalic acid assisted sol-gel method. They found that the irreversible phase transition of V2O5 was effectively prevented by Cr doping. The Cr doped V2O5 (Cr0.1V2O5.15) showed a better capacity retention and higher lithium ion diffusion coefficient than the un-dopedV2O5. The authors attributed this higher lithium ion diffusion coefficient of Cr0.1V2O5.15 to the formation of a more open structure caused by Cr doping.

Cheah et al. [277] synthesized Al-inserted V2O5 nanofibers by electrospinning technique. The electrochemical lithium intercalation behavior of Al-inserted V2O5 nanofibers between 1.75 and 4.0 V (vs. Li/Li+) were explored using cyclic voltammetry and galvanostatic charge-discharge studies. Compared to native V2O5 nanofibers, the Al-inserted V2O5 showed enhanced high rate and elevated temperature performance. The authors attributed the improved electrochemical performance to the following aspects. First, the addition of Al3+ions into the V2O5 structure leads to the reduction of V5+ to V4+ and even to V3+ during the in situ synthesis of Al-inserted V2O5 nanofibers. The introduction of these low valent vanadium ions (V4+ and V3+) before the 82

electrochemical lithium intercalation/deintercalation decreases the extent of electrochemically induced phase changes, related to the conversion of V2O5 to γ-LixV2O5 phase. Further, Al3+ insertion leads to increased c-spacing in the layered structure and it is believed to occupy between the VO5 layers, thus providing the necessary structural stability for V2O5 phase. Zhan et al. [276] synthesized V2O5 and Al0.2V2O5 nanoparticles by an oxalic acid assisted soft-chemical method and compared their electrochemical performance as cathode material for LIB. Charge-discharge cycling showed that the Al0.2V2O5 nanoparticles exhibited much better capacity retention than the un-doped V2O5, which was attributed to the enhanced structural stability of the material. However, the chemical diffusion coefficients of the Al0.2V2O5 nanoparticles were a bit smaller than those of un-doped V2O5, which may arise because of Al3+ dopants blocking the Li+ diffusion in the material lattice.

Maingot et al. [270] synthesized Fe doped V2O5 (Fe0.12V2O5.16) from V2O5 xerogel through an exchange reaction. In its Rietveld determined structure, [FeO6] apex sharing octahedral linking V2O5 ribbons increase the three-dimensional character of the material. This latter point seems to improve the electrochemical behavior when compared to un-doped V2O5, especially in terms of cycle life. The authors attributed the improved electrochemical performance to the [FeO6] chains which are finite, along the c-axis, strengthen the cohesion between the layers, and maintain the structure through cycling. It should be noted that due to very weak contribution to the overall diffraction diagram, the positions of the two shared apex O atoms in [FeO6] were not introduced. Indeed, they could not be detected on the Fourier map difference. Li et al. [269] prepared three-dimensional porous V2O5 and Fe0.1V2O5.15 thin films by electrostatic spray deposition 83

technique. They found that the Fe0.1V2O5.15 thin film shows much better cycling performance than the non-doped V2O5 thin film. However, the Fe-doping causes a slight decrease in the lithium ion diffusion coefficient. It was reasoned that Fe3+ acts as a stabilizing agent in the layered V2O5 and increase the reversibility of the charge and discharge processes towards deeper depth of lithium intercaltion/deintercalation; the dissolution of the active material in the electrolyte can also be significantly suppressed in the Fe-doped sample. Almeida et al. [272] studied the electrochemical behavior of Tb0.11V2O5 electrode in propylene carbonate and LiClO4. It was found that the Tb doping improves the lithium intercalation kinetic and structural stability of V2O5. They considered that the bonding of the earth-rare ions and the residual H2O structure was critical phenomena for the improved performance of doped V2O5.

Li et al. [174] prepared Sn-doped V2O5 film by sol-gel method followed by annealing in air. This Sn-doped V2O5 film exhibits larger lithium-ion diffusion coefficient, better electrochemical reversibility, and lower electrochemical reaction resistance than pure V2O5 film. Moreover, this Sn-doped V2O5 film displays an excellent cycling performance. The improved electrochemical performance of the Sn-doped V2O5 film was credited to the following aspects: the presence of lower valence vanadium ions (V4+) may improve the conductivity of the Sn-doped V2O5; the occupancy

of

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could

stabilize

the

V2O5

layers

and

benefit

the

facile

intercalation/deintercalation of lithium ion within the layers; the local structural defects may serve as possible nucleation centers for phase transformation during charge/discharge process, and therefore enhancing the kinetics of lithium-ion intercalation/extractions in Sn-doped V2O5 as well as its cycling performance. 84

In addition to the doping of cations, the preparation of V2O5-based mixed oxides was also explored to improve the electrochemical performances. For example, Lee et al. [289] investigated the intercalation properties of V2O5-TiO2 mixture films. The results demonstrated that the addition of TiO2 greatly decreased the particle size and crystallinity of V2O5 in the film, resulting in significant improvement of reversible capacity and high rate performance of the mixed V2O5-TiO2 films. For example, the addition of 20 mol% Ti into V2O5 polycrystalline demonstrated an approximately 100% improvement in lithium intercalation performance as compared to pure V2O5 electrodes. They ascribed such performance improvement to the combined effects of decreased lithium diffusion distance, which suppresses the concentration polarization of lithium in the V2O5-TiO2 electrode and low crystallinity providing more lithium intercalation. Ko et al. [290] prepared amorphous V2O5 powders with zirconia (ZrO2) dopant by a one-pot spray pyrolysis method at temperatures above the melting temperature of V2O5. They found that zirconia played a key role in the formation of fine-sized amorphous V2O5 powders. The presence of a small amount of ZrO2 additive disturbed the crystallization of the melted V2O5 powders during the quenching process. The V2O5 powders with 7 wt% ZrO2 deliver capacities of 309, 269, and 222 mA h g-1 after the first, second, and 50th cycles, respectively, at the current density of 294 mA g-1. The capacity retention is much better than the crystalline V2O5 counterpart. The authors ascribed the good cycle performance of V2O5-ZrO2 composite powders to the faster lithium ion diffusion in the fine-sized amorphous powders. Further, zirconium dopant also improved the electrochemical properties of the amorphous V2O5 powders by acting as a structural pinpoint to preserve the original structure during the cycling process. Mai et al. [291] synthesized bowknot-like RuO2 85

quantum dots@V2O5 (denoted as RQDV) nanomaterials via a hydrothermal synthesis method followed by annealing treatment. The RuO2 quantum dots have a diameter of 0.5-3.5 nm. The RQDV cathode exhibits higher rate performance and better cycling stability in comparison with the pure V2O5 cathode. The enhanced electrochemical performance can be attributed to the unique nanostructure with RuO2 quantum dots dispersed on the V2O5 surface. The unique RQDV nanostructure offers accessible intercalation sites and shortens Li+ and electron transport path-ways, which favors electrolyte penetration and interface reactions; RuO2 quantum dots are dispersed on the V2O5 surface, which enhances the electronic conductivity of this material.

Pre-intercalation has also been recognized as a promising strategy to enhance the diffusion of Li+ (or electrolyte cations) and structural stability of layered metal oxides during the charge/discharge processes [292-296]. For example, Li et al. [297] prepared Na-preintercalated V2O5 (NaxV2O5, x=0.00, 0.005, 0.01, or 0.02) by a simple so-gel approach and freeze-drying method with the direct addition of sodium salt followed by an annealing treatment. XRD results showed that all of the synthesized NaxV2O5 have typical orthorhombic structures without impurity phases. However, the interplanar distance of preintercalated samples increased in comparison with pristine V2O5. In addition, the V4+ concentration in the Na preintercalated V2O5 samples gradually increased as amount of sodium increased. NaxV2O5 samples exhibited improved electrochemical properties compared with those of pristine V2O5. Among all of the samples, Na0.01V2O5 delivered the highest reversible specific capacity, best cycling stability, and excellent rate capability. Further, the Na0.01V2O5 sample also showed the highest lithium-ion diffusion coefficient among all samples, suggesting that a small amount of preintercalated Na can effectively expand the ion 86

channels and improve the reactive kinetics of V2O5 owing to the larger size of Na+ expanding the interplane spacing of the host. Zhao et al. [292] made detailed investigations on the influence of alkali metal ions (A=Li, Na, K, Rb) preintercalaion on electrochemical performance of vanadium oxide cathode material by experimental characterizations and ab inito calculations. The preintercalated materials correspond to Li0.7V6O15 for Li−V−O, Na0.9V6O15 for Na−V−O, and K0.7V6O15 for K−V−O. In Rb−V−O, RbV3O8 is the major phase, and an unknown minor Rb−V−O phase. Structural analysis revealed that the A-V-O (A=Li, Na, K) phases are monoclinic and isostructural and are different from the original orthorombic V2O5 structure, while RbV3O8 exhibits another layered structure (Figure 33d-f). The layer separation increases with increasing size of the preintercalated cations (Li, 7.22 Å; Na, 7.26 Å; K, 7.40 Å; Rb, 7.80 Å) and is significantly larger than that of orthorhombic V2O5 (4.36 Å), suggesting a significant lattice expansion after large alkali metal ion intercalation. Electrochemical characterizations demonstrated that appropriate alkali metal ion intercalation in admissible layered structure can effectively overcome the limitation of cyclability as well as rate capability (Figure 33b,c). This is attributed to the synergistic stabilizing effect between the intercalated ions and the layers with appropriate layer−surface configuration in charge−discharge processes together with the enlarged layer spacing. The V−O layers can selectively capture A ions according to the ionic size, and the large alkali metal ions can act as pillars in between layers to yield a stable interlayer expansion and prevent the destructive distortion or collapse of the V−O layers in charge−discharge processes (Figure 33a). Such strategy provides a facile and effective method to regulate the diffusion channel of some intercalation compounds at the technical level, which are promising and 87

important for the future design and optimization of intercalation compounds as cathodes for rechargeable lithium- and other metal-based or metal ion batteries.

Figure 33. (a) Schematic representation of large alkali metal ion intercalation. (b) Cycling performance of A-V-O nanowires formed by preintercalating large alkali metal ions into vanadium oxides at a charge/discharge current density of 0.1 A/g. (c) Rate performance of A−V−O nanowires. (d−f) Illustration of the crystal structure of V2O5, A−V6O15 (A = Li, Na, K) and RbV3O8, respectively. Li−V6O15, Na−V6O15 and K−V6O15 are isostructural. The red, gray, and purple balls represent the O, V, A, or Rb atoms, respectively. Layer separation D represents the repeat distance in crystal structure. Reprinted with permission from ref [292]. Copyright 2015 American Chemical Society.

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Cation doping and/or pre-intercalation are effective strategies to enhance the electrochemical properties (such as electronic conductivity, structural stability, and Li+ intercalation/deinterclation kinetics) of V2O5 cathode material for LIBs. However, it should be mentioned that an appropriate doping or pre-intercalation level is very crucial for obtaining the optimum electrochemical performance. Excess dopant may block the insertion of Li ions and moreover the inactive dopant would reduce the utilization of active materials and affects electrochemical capacity [254, 298]. In addition, developing doped high-performance V2O5 using earth-abundant elements is more desirable and Fe, Mn, Cu, and Al-doped V2O5 are some of the most attractive choices used in commercial LIBs. Though single-cation doped V2O5 has been studied extensively for its lithium ion storage performance, the double- and multi-cation doped V2O5 has not been reported so for. It is expected that due the synergistic effect of different dopants, the double- and multi-doped V2O5 can exhibit better electrochemical performance than the singly doped V2O5. Although the cation doping has been studied extensively and demonstrated positive impacts on the enhanced electrochemical properties and lithium ion storage performance, there is no or little work dedicated to the possible negative impacts during long term cycling. For example, will the doping cations migrate to and enrich at the surface under repeated cycles? Will they be leached by the electrolyte? Cation doping in V2O5 is typically accomplished by introducing low valence state cations such as Fe (2+/3+), Ni (2+/3+/4+), Mn (2+/3+/4+), Al (3+) and Sn (4+). Such doping would also introduce oxygen vacancies and at the same time compromise the crystallinity. Will the presence of oxygen vacancies and poor crystallinity compromise the long term stability of V2O5?

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4.4. Structural defects Crystalline electrode materials usually assure well-defined Faradic redox reactions in electrochemical active materials during Li+ intercalation/deintercalation. However, those electrode materials often suffer from limited capacity or poor rate capability, due to the restricted Li+ diffusion paths and intercalation sites in the long-range well-ordered structure in perfect crystalline materials. While electrode materials with surface or bulk defects and/or poor crystallinity could offer more open framework, higher energy states, and more reaction sites (which results in higher capacity), better cyclic stability, and improved high rate performance.[299-301] Electrode materials with structural defects or with poor crystallinity all possess higher Gibbs free energy than that of a perfect bulk crystal counterpart; in other words, they are all far from thermodynamic equilibrium. The electrodes away from equilibrium may offer a new direction of research for the advancement of electroactive materials to achieve much improved Li+ intercalation properties and, thus, lead to the development of more efficient LIBs [302, 303].

V2O5 is a low mobility semiconductor with a typical charge carrier mobility of about 1 cm2 V-1 s-1 and having predominantly an n-type conductivity [304]. Electrons are the charge carrier, and an increase in the carrier density is accompanied by reduction in oxygen concentration in the lattice. This oxygen deficiency will determine the electron concentration in V2O5. On the basis of the generally accepted hoping mechanism,[305, 306] an electron in V2O5 migrating from a V4+ center to an adjacent V5+ site can be represented by the following scheme:

90

̶ V4+ ̶ O ̶ V5+ ̶ → ̶ V5+ ̶ O ̶ V4+ ̶

(1)

Thus, the conductivity is proportional to the number of oxygen defects [307].

The oxygen defects in bulk V2O5 can be readily introduced by heating V2O5 in low oxygen partial pressure. Liu et al. [308, 309] fabricated V2O5 xerogel films by casting V2O5 sol onto FTO glass substrates and annealing at 300 °C for 3 h in nitrogen and air. They found that the V2O5 film annealed under nitrogen gas flow atmosphere exhibited less crystallized phase, enhanced optical absorption and electrical conductivity as compared with film annealed in air under otherwise identical conditions (Figure 34a). In particular, the N2 annealed vanadium pentoxide films demonstrated much improved lithium ion intercalation capacity and cyclic stability (Figure 34b,c). Despite that the initial discharge capacity was only 68 mA h g-1 due to the surface defects layer, after 24 cycles, the discharge capacity was as high as 158 mA h g-1. After 50 cycles, the discharge capacity was still 148 mA h g-1. The authors attributed the much improved lithium ion intercalation capacity and cyclic stability to the less crystallized structure, the presence of a surface defects layer (V4+ and/or V3+) which not only enhanced the charge transfer conductivity but also behaved like a protective coating layer to ensure the morphology stability of the V2O5 film and served as possible nucleation centers to facilitate the phase transformation process during the lithium ion intercalation/deintercalation process.

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(a)

(b)

(c)

Figure 34. (a) Absorption spectra of V2O5 films annealed in air and nitrogen at 300 °C for 3 h; the inset presents the photo of films after annealing in air (left) and nitrogen (right). (b) The lithium ion intercalation discharge capacity of V2O5 films annealed in air and nitrogen as a function of cyclic numbers at a current density of 600 mA g-1. (c) The discharge capacities of V2O5 films annealed in air and nitrogen at 300 °C for 3 h as a function of applied discharge current densities. The measurements were carried out in a potential window between 0.6 V and -1.4 V vs. Ag/AgCl as a referenceelectrode. Reprinted with permission from ref [308]. Copyright 2009 The Royal Society of Chemistry.

Peng et al. [310] reported hydrogenated 2D V2O5 nanosheets with super Li storage properties. The V2O5 nanosheets with most of O(II) vacancies were fabricated by hydrogenating V2O5 nanosheets at a relatively lower temperature of 200 oC, enabling easier and faster Li+ diffusion. In the hydrogenation process, hydrogen atoms first adsorbed at the oxygen sites forming OH and then H-V2O5 with most oxygen vacancies at O(II) sites could be produced because the formation of oxygen vacancy in O(II) sites by removing the OH group requires less energy than removing oxygen directly. The H-V2O5 with most oxygen vacancies in O(II) sites has improved conductivity, faster diffusion of Li+, and improved structure stability for Li+ intercalation/deintercalation, resulting in higher capacity, better rate capability, and improved cycling stability. The initial discharge capacity of H-V2O5 is as high as 259 mA h g-1 and remains 55% when the current 92

density is increased 20 times from 0.1 to 2 A g-1. Additionally, the H-V2O5 nanosheets have a stable discharge capacity boasting less than 0.05% decay per cycle after stabilization.

The above reported methods to introduce V4+, including heating V2O5 under vacuum, inert or reducing conditions (e.g., N2 or H2), have the limitations of multiple steps, complicated process and producing a small amount of tetravalent vanadium ions. Song et al. [311] reported facile synthesis of self-doped V4+-V2O5 nanoflakes through combusting the NH4V4O10 and VO2(B) nano flake mixture with the presence of 2-ethylimidazole. As cathode materials for LIBs, the resulting V4+-V2O5 exhibited superior rate capability (a reversible capacity of 293, 268, 234, 214, 196 and 139 mA h g-1 at 100, 300, 600, 1000, 2000 and 4000 mA g-1, respectively) and remarkable cycling performance (95% capacity retention after 100 discharge/charge cycles at 2000 mA g-1). The outstanding performance is ascribed to the uniform nanoflake structure, which offers more surface area and reaction sites for lithium ion intercalation, and the presence of tetravalent vanadium ions accompanied with oxygen vacancies, which might promote and catalyze the electrochemical reactions at the surface, in addition to improve the intrinsic electrical conductivity and Li+ ion diffusion coefficient, as reflected by the smaller polarization and more reversible and easier phase transitions than those of V2O5 without detectable V4+ and oxygen vacancies.

Swider-Lyons et al. [312] demonstrated that point defects may be introduced into metal oxide to increase its lithium ion capacity by using various heat treatments to modify the defect structure of polycrystalline V2O5 and then comparing the lithium capacity of the materials. As shown in Figure 35, the V2O5 treated under O2/H2O at 460 °C showed a 23% higher lithium-ion storage capacity

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than the as-received materials despite no change to its long-range structure. Other heating conditions lower the lithium capacity of the V2O5. The authors inferred that heating under O2/H2O introduces defects, such as cation vacancies associated with lithiated oxygen sites, which can electrochemically exchange lithium ions and serve as additional charge-storage sites.

Figure 35. Charge-discharge profiles of 0.5-mg V2O5 electrodes heated under different atmospheres: (A) first complete charge cycle (10 μA), (B) first discharge cycle (10 μA). Reprinted with permission from ref [312]. Copyright 2002 Elsevier B.V.

It has recently been demonstrated that disordered and amorphous structures in materials may form percolation pathways through the opening of active diffusion channels that could potentially accelerate ionic diffusion [299, 313-316]. In addition, several previous studies have revealed that the diffusion of lithium ion in amorphous materials proceeds more rapidly than in crystalline materials with similar particle size and morphology [317]. Chae et al. [318] prepared a vanadium pentoxide electrode in the amorphous form (amor-V2O5), and its electrode performances are compared to those for its crystalline counterpart (c-V2O5). They concluded that the amor-V2O5 electrode outperforms c-V2O5 in several ways. First, it is free from irreversible phase transitions 94

and Li trapping, which evolve in c-V2O5, probably due to the lack of interactions between the inserted Li+ ions/electrons and V2O5 matrix. Second, the amor-V2O5 electrode outperforms its crystalline counterpart with respect to reversible capacity and cycle performance. They believed that this must be due to lithium ion storage at vacant sites, which are ill-defined but highly populated in amorphous materials. The rate capability is also excellent for amor-V2O5. They ascribed the superior capacity and rate property for this amorphous electrode to the enlarged accessible surface area for lithium ion and rather opened lithium ion diffusion pathways provided by the vacant sites in the amorphous V−O network. Afyon et al. [319] demonstrated vanadate-borate glasses as high capacity cathode materials for rechargeable LIBs for the first time. They found that the composite electrodes of V2O5-LiBO2 glass with reduced graphite oxide (RGO) deliver specific energies about 1000 Wh kg-1 and retain high reversible capacities of about 300 mA h g-1 for the first 100 cycles (between 1.5 V and 4.0 V at 50 mA g-1 rate). The composite electrode also display a notable rate capability with a discharge capacity of about 299 mA hg-1 (35th cycle) at a current density of 400 mA g-1. The remarkable lithium storage performance of vanadate-borate glasses are mainly due to the absence of a long range order that allows for subtle structural adaptations, which cannot occur within crystalline phases. This is very different from the huge capacity loss observed for crystalline V2O5 phases in the first cycle within 1.5–4.0 V due to the irreversible phase transformation to ω-Li3V2O5. Based on the above results, exploring on disordered and/or amorphous electrode materials is expected to be a promising development, and worthy of further research towards developing high performance secondary batteries with high energy density and long cycle-life. 95

V2O5 with oxygen vacancies, particularly on the surface, and low crystallinity possess much improved electrochemical properties, better electronic conductivity and lithium ion intercalation capacity. However, the long term stability of such highly defective surfaces and low crystallinity has not been thoroughly studied. Since the intercalation reaction is a first-order phase transition, the change of surface chemistry (including oxygen vacancies) and the change of crystallinity would change the surface energy and the bulk Gibb’s free energy, which will in turn affect the activation energy required for initial nucleation of the new phase. Fine-tuned control over these parameters would allow us to make the irreversible phase transition reversible. The open question is: will it be possible to reversibly intercalate three lithium ions in V2O5 by appropriately designing and engineering surface energy and bulk Gibb’s free energy?

5. BEYOND LITHIUM-ION BATTERIES Nowadays, LIBs have been largely adopted as the most popular energy storage system for mobile electronic devices, hybrid electric vehicles, electric vehicles, and the smart grids. Despite great commercial success, the high cost of lithium geatly hindered the large-scale application of LIBs, due to its huge requirement and resource scarcity on the earth. Figure 36 shows the abundance (atomic fraction) of the chemical elements in Earth's upper continental crust as a function of atomic number [59]. Clearly, it is more cost-effective to use less expensive and more abundant materials when available. Although there is still opportunity for improvements to Li-ion batteries, research on alternatives to lithium has grown considerably. Over the past several years, sodium-ion batteries (SIBs) and magnesium-ion batteries (MIBs) have been gaining more and

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more attentions as realistic candidates for large-scale energy storage applications [59, 320-324]. Table 3 compares the qualities of various alkali-ion battery chemistries. For SIBs, the sodium intercalation chemistry is very similar to that of lithium. Thus, it is possible to use similar compounds as electrode materials for both SIBs and LIBs systems. However, compared to LIBs, the larger radius (0.98 Å) of the Na+ leads to sluggish electrochemical reaction kinetics and unsatisfactory cycling stability of the electrode due to the lager volume change and collapse of the host structure during Na+ insertion/extraction process. As for the Mg2+, the intercalation process is more complicated. The higher polarizability leads to strong intercalation between Mg2+ and the negatively charge host lattice, which makes the intercalation and diffusion of Mg2+ in electrode more difficult. In contrast to Li+ and Na+ that can lose their solvation shell upon intercalation, Mg2+ intercalation involve co-intercalation of solvent molecules. As a versatile cation host structure, vanadium oxides have been one of the more popular cathode materials for new and novel battery chemistries [325-327]. Here we review its properties and device performances with respect to sodium and magnesium ion intercalation, which are among the most intriguing lithium substitutes on the basis of their natural abundance and comparable physical properties.

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Figure 36. Abundance (atom fraction) of the chemical elements in Earth's upper continental crust as a function of atomic number. Reprinted with permission from ref [59]. Copyright 2015 Springer Science and Business Media.

Table 3. Comparative qualities of lithium, sodium, and magnesium for alkali-ion battery application [328, 329]. Parameter

Lithium

Sodium

Magnesium

Cationic radius (Å)

0.76

1.02

0.72

Atomic weight (g mol-1)

6.9

23.0

24.3

E (V vs. SHE)

-3.04

-2.71

-2.37

Carbonate cost ($/ton)

6000

150

1000

Metallic capacity (mA h g-1)

3862 (Li+)

1166 (Na+)

2205 (Mg2+)

Metallic capacity (mA h cm-3)

2046

1129

3833

Cathode capacity per M2V2O5 (mA h g-1)

294

236

590

Coordination preference

Octahedral and tetrahedral

Octahedral and prismatic

Octahedral

°

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5.1. Na-ion intercalation in V2O5 Compared to the comprehensive body of work on LIBs, research on SIBs is still in its nascent stages. The SIBs works under the same principles as LIBs technology, but with different characteristics as dictated by the properties of the transporting ionic species and the consequent effects such has on the electrode materials. Na+ and Li+ share similar chemical properties and comparable electrochemical performances, which indicates that the development of SIBs can be based on previously applied approaches or methods utilized for LIBs.

SIBs are attractive because sodium resources are seemingly inexhaustible and ubiquitous, and therefore cost considerably less (by a factor of roughly 30−40 times) than lithium; additionaly, the absence of Na-Al alloying reaction at low voltage, as is the case with lithium, permits the use of aluminum, replacing copper as the anodic current collector, which equates to an overall cell cost savings of ~2% [330-332]. The lower operating voltage of SIBs results in enhanced stability of the non-aqueous electrolyte [331], but also leads to lower energy density. Most of the reported electrode materials for SIBs exhibit similar or slightly lower specific capacity and redox potential than the case used in LIBs [333, 334]. Moreover, the higher molar mass and larger radius of Na+ than that of Li+ induce a lower gravimetric energy and sluggish electrochemical reaction kinetics [335]. Therefore, large distortions occur in the host lattice during de/sodiation process, which ultimately lead to pulverization of the electrode and the impending failure of the cell.

Most of the investigations, to date, examining Na-ion battery cathodes have focused on intercalation based materials, especially layered transition metal oxides. Among these layered 99

transition metal oxides, vanadium oxide is an attractive multifunctional material used for widespread application in various fields such as catalysis [336], energy storage [337], and biomedical devices [338]. A wide assortment of V2O5 nanostructures have been synthesized by a variety of methods [188, 339-342]. For Li-ion battery application, the intercalation of lithium into V2O5 has been well documented as reviewed in Section 3. However, reports concerning the use of vanadium oxide for sodium based energy storage systems are relatively limited.

Early investigations involving Na+ intercalation into vanadium pentoxide aerogels found the mechanism of insertion to be very similar to that observed for Li+ [343, 344]. However, there was one apparent difference in that Na+ strongly prefers one insertion site in contrast to the several prominent peaks exhibited by Li+ insertion. This was determined to be the anion-vacancy defect site, at least within the xerogel system [345]. Wu et al. employed XPS to investigate the intercalation of sodium into V2O5 thin films prepared by physical vapor deposition [346]. For a nearly pristine V2O5 starting material, there was an obvious shift in the Fermi level of approximately 0.7 eV due to the strong interaction sodium can have on the host electronic structure. Core level spectra revealed that the vanadium species is effectively reduced to mostly V4+, with even some V3+ states present. A charge transfer of about 0.42 electrons per intercalated Na atom has been calculated, suggesting that the intercalated Na atoms are not completely ionized. Finally, sodiation above Na2V2O5 was shown to lead to the formation of sodium oxides, and eventually metallic sodium, on the electrode surface. Additionally, the work function was noted to decrease with increasing sodium deposition time because of sodium intercalation; cycling behavior and electrochemical performance were not considered. In a separate initial investigation, 100

the diffusion coefficient of sodium in V2O5 was found to range from 5×10−14 to 9×10−12 cm2 s-1 with a maximum insertion capacity of approximately 1.6 moles of Na per V2O5. It was also clear that sodium insertion in V2O5 induces a reversible phase change into a new structure that can be stable over long term cycling [347].

Tepavcevic et al. [348] synthesized nanostructured bilayered vanadium pentoxide by electrochemical deposition from aqueous vanadyl sulfate electrolyte on Ni foil substrate, and then annealed in vacuum at 120 oC to remove intercalated water. The V2O5 structure is composed of 2D bilayered stacks which are separated by large interlayer spacing (~13.5 Å). The V2O5 bilayers made of base-faced square-pyramidal VO5 units that are arranged parallel to one another along regular intervals. However, it was observed that the stacking sequence is imperfect and the spacing between bilayers is more random, i.e., the bilayered structure is atomically ordered only in a short range scale. The bilayered V2O5 material was found to have a higher degree of local symmetry than its orthorhombic V2O5 counterpart due to the decreased apical V-O bond distance. The bilayered V2O5 cathodes demonstrated a specific capacity of 250 mA h g-1 at a current density of 20 mA g-1 and were able to maintain 85% of their initial capacity after 350 cycles with current density varying from 20-630 mA g-1 (Figure 37c) [348]. The sodium intercalation performance of the orthorhombic V2O5 cathode was significantly less at 150 mA h g-1 and rapidly decreased with cycling, despite the long range order that may enable unhindered diffusion of intercalated ions. The mechanism of sodiation in the two different V2O5 structures was also found to be entirely different as revealed through the discharge potential profiles (Figure 37a). The incorporation of Na+ into bilayered structure shows solid state solution intercalation with no apparent phase 101

transition, whereas the incorporation of Na+ into orthorhombic electrode is accompanied by two phase transitions. Ex situ and in situ synchrotron characterizations demonstrated that the intercalation of Na+ into bilayered V2O5 electrode induced organization of overall structure with appearance of both long- and short-range order; The deintercalation of Na+ was accompanied with the loss of long-range order while short-range order is preserved. This behavior was determined to be reversible and could be attributed to the electrostatic interaction between sodium ions and the terminal oxygen of the square-pyramidal VO5 unit. In contrast, the orthorhombic V2O5 electrode experienced rapid deterioration and eventual loss of crystallinity after extended cycling (Figure 37b). The results were reinforced in an imitative study [349, 350]. These findings suggest the advantage of short-range ordered structures deviating from the thermodynamic equilibrium for sodium-ion battery application.

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Figure 37. (a) First four charge–discharge cycles of bilayered V2O5 and orthorhombic V2O5 electrodes at 20 mA/g in 1 M NaClO4/PC electrolyte. (b) Ex situ synchrotron XRD patterns of orthorhombic V2O5 before and after cycling. (c) Cycle life of bilayered V2O5 electrodes. (d) Cycling performance comparison of the same electrochemically grown bilayered V2O5 film: deposited on Ni substrate and pressed on stainless steel mesh current collector. Both cells were cycled at 630 mA/g current densities, within the potential window of 3.8–1.5 V (vs Na/Na+) from 1 M NaClO4 in PC. Reprinted with permission from ref [351]. Copyright 2012 American Chemical Society.

Single-crystalline bilayered V2O5 nanobelts synthesized via hydrothermal method with comparable performance have also been reported, some of the results for which are shown in Figure 38 [351]. The solvent used during synthesis was found to play a critical role in structure

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formation based on the chemisorption of preferred molecules. The interlayer spacing was calculated using Reitveld refinement methods to be approximately 11.5 Å, and TEM revealed that the nanobelts grew along the [010] direction meaning that the primarily exposed facets are (100). At different current densities of 40, 80, 160, 320, and 640 mA g–1 the bilayered V2O5 nanobelt electrode delivered capacities of 235.7, 231.4, 195.7, 151.2, and 134 mA h g–1, respectively. However, two phase changes were noted to take place with the evolution of sodium in the host material; the differences with this work and that of Tepavcevic et al. may arise because of the electrochemical deposition method. Interestingly, ex situ XRD analysis of a sodiated electrode concluded that despite the already enlarged interlayer distance, the interlayer actually further increased to ~15.3 Å. Orthorhombic V2O5 powder, on the other hand, exhibited considerably lower capacity and poor cycling performance. These results clearly establish that this preferential crystal structure provides superior channels for facile Na-ion incorporation during cycling. Zhu et al. [352] prepared a sponge-like V2O5 cathode material for sodium-ion batteries by hydrothermal method combined with freeze-drying. The sponge-like material is self-assembled from nanosheets with exposed (001) facets. The interlayer distance of the material is estimated to be 11.36 Å, very similar to that (11.5 Å) of the above mentioned single-crystalline bilayered V2O5 nanobelts. They also observed a multistep Na+ ion intercalation process in the material based on the analysis of discharge-charge profiles. Contrary to the single-crystalline bilayered V2O5 nanobelts, the interlayer spacing of the sponge-like V2O5 decreases upon the intercalation of a sodium ion. The numerous macrospores in the spongy structure can effectively soak up the electrolyte. Moreover, the exposed (001) facet of the nanosheets is favorable for charge transfer at the 104

electrolyte/electrode interface. The thin nanosheets along the [001] axis provide short diffusion pathways for sodium ions. Thus, when the spongy V2O5 material was used as a cathode material for sodium-ion batteries, a high specific discharge capacity of 216 mA h g-1 was obtained at a current density of 20 mA g-1. In addition, the material retained 73% of the initial capacity after 100 cycles at a current density of 100 mA g-1.

Figure 38. (a) Refined structural model of the bipyramidal layered structure of V2O5. (b) Lattice resolution TEM image, in which the (001) crystal plane can be identified. (c) Atomic resolution HRTEM image of a V2O5 nanobelt, from which the interlayer structure of the bilayered V2O5 can be directly observed. (d) Discharge capacities versus cycle number at current densities of 40, 160, 320, and 640 mA g–1. (e) Rate performance of a bilayered V2O5 nanobelt at varied current densities.

Su et al. [353] went on to examine the implementation of orthorhombic V2O5 in Na-ion battery application by employing a hollow nanosphere aggregate morphology with primarily exposed 105

(110) facets. Initially, a vanadyl ethylene glycolate nanosphere precursor, approximately 800 nm in diameter and 50 nm thick, was prepared through a polyol mediated reflux process. The precursor was made up of smaller, arranged nanocrystals that were easily converted to orthorhombic V2O5 upon sintering. The precursor morphology was maintained as there was no discernible shrinkage or structural deformation in the V2O5 nanosphere. TEM and Reitveld refinement confirmed the predominantly exposed facets of the V2O5 nanocrystals as being (110). This is contrary to bulk V2O5 which typically has (001) preferential facet exposure. The exposure of high energy surface facets, along with the unique architecture, can facilitate Na-ion ion insertion and extraction by lowering the energy barrier for faster and more efficient intercalation mechanism. This phenomena paired with the morphology loaded with intercrystalline voids are credited with the noteworthy electrochemical performance. As revealed through modeling simulations, the nano-sized intercrystalline voids make it possible for the shell to tolerate plastic yielding, because no mechanical constraint is imposed by a core during sodiation; moreover, the hollow morphology means that there are no mechanical constraints imposed by a core and that the shell can expand as needed in the radial direction via elastic flow. When tested electrochemically, the hollow V2O5 nanospheres exhibited capacities of 159.3, 153.4, 150.8, 131.6, and 112.4 mA h g−1 in the second cycle at 40, 80, 160, 320, and 640 mA g−1 current density, respectively, while solid V2O5 nanocrystals demonstrated rather poor performance. Remarkably, a capacity and Coulombic efficiency of 141 mA h g−1 and 95.4%, respectively, could be retained after 100 cycles. Ex situ SEM performed after cycling confirmed that the overall spherical shape and hierarchical architecture were retained. The same material was also used as Na-ion battery anode material 106

where it demonstrated commendable performance [354].

Raju et al. [355] successfully encapsulated orthorhombic V2O5 nanoparticles within nanoporous carbon in order to combat the sluggish sodium-ion diffusion in a compact crystal structure. The conformal coating of 5-7 nm V2O5 nanoparticles was attained by utilizing hydrolysis deposition that featuring water adsorption in nanoporous carbon followed by hydrolysis reaction within the nanopores. The wet chemical deposition process was robust in that it could be manipulated to effectively control the V2O5 loading level with only slight alteration of the V2O5 nanoparticle size. N2 sorption measurements found that with V2O5 loading levels of 55 and 70 wt%, 6.8 and 13.1% of the originally available carbon pore volume was filled, respectively. TGA shows that the V2O5 nanoparticles do not form a network until relatively high loading levels are implemented as noted by observing the catalyzed and intrinsic oxidation of carbon. The discharge potential profiles were seemingly linear with sodium incorporation; for an electrode with 55wt% V2O5 the specific capacity was 276 mA h g-1 while the capacity for the entire nanocomposite was 170 mA h g-1 (Figure 39). However, a noticeable plateau appeared in the potential profile when the electrode was discharged at higher current rates, indicating the adoption of a different de/sodiation process. Correspondingly, CV was conducted at varying rate in order to elucidate the sodiation processes at play. It was concluded that the capacitive process accounts for ∼20% of the total charge storage at 0.5 mV s-1, but increases to ~40% at 5 mV s-1. It is also evident, based on the voltammogram profiles, that a diffusion-controlled redox process dominates from 2.5 to 3 V while a capacitive mechanism reigns outside of this potential window. CV analysis done using large ions, tetraethylammonium (TEA+) tetrafluoroborate (BF4–) in propylene carbonate, revealed that the 107

majority of the capacitive charge storage is due to the pseudocapacitance of V2O5.

Figure 39. Galvanostatic charge-discharge profiles of (a) nanoporous carbon with various V2O5 loadings at 40 mA g-1 and (b) for the 55wt% V2O5 loading at current densities of 40, 80, 160, 320, and 640 mA g-1. (c) CV profiles of bulk V2O5 and 55wt% loaded V2O5 at 0.5 mV s-1. (d) Nyquist plots of nanoporous carbon, carbon with 55wt% V2O5 loading, and bulk V2O5. Reprinted with permission from ref [355]. Copyright 2014 American Chemical Society.

Ali et al. [356] designed an orthorhombic V2O5/C composite by sequential methodology; formation of nanometer-sized V2O5 with the aid of oxalic acid followed by ball-milling with acetylene black. Due to the incorporation of highly conductive carbon-based material, the charge transfer resistance of the compsite was obviously decreased compared to the pure V2O5. As a consequence, the nano sized V2O5/C composite exhibited a superior reversible capacity and high rate capability. They found that the reaction mechanism of V2O5 upon sodium insertion/extraction is as follows: during recharge process, the insertion of sodium ions into V2O5 induces formation of 108

NaV2O5 as a major phase with minor Na2V2O5 phase. The Na2V2O5 exists both in crystalline and amorphous-like phase. During the sodium insertion, the oxidation state of vanadium in V2O5 varies from 5+ to 4+ and the lattice parameter of c is increased by 9.09%, corresponding to the increase in the unit-cell volume of 9.2%. During recharge process, the sodium ions are extracted and V2O5 reverts to its original crystal structure along with NaV2O5 as a minor phase.

The effects of crystallinity on sodium storage performance of V2O5 were also reported (Figure 40) [357]. Amorphous and nanocrystalline V2O5 were prepared through a combination of sol-gel processing paired with electrochemical deposition and investigated as cathode for SIBs. The amorphous V2O5 delievered a capacity of 241 mA h g-1 when examined as cathode material for SIBs, twice the capacity of its crystalline counterpart (120 mA h g-1). Moreover, the amorphous V2O5 demostrated a much higher discharge potential, energy density, and cycle stability. The superior sodium storage performance of amorphous V2O5 arises from the fast Faradaic reactions which occur in amorphous phase deriving from a percolated diffusion network. The discrepancy in performance between the amorphous and nanocrystalline V2O5 is primarily accredited to the low entropic energy associated with the ordering of intercalated atoms and a more open framework. The less structured and more open channels in amorphous phase decrease the diffusion barrier for sodium ions to transition between sites and lead to high rate capability and energy density. A separate study corroborated the positve qualities of amorphous V2O5 use for Na-ion battery, but suggested a new process for electrode prepartion that could effectively preserve their oftentimes unique morphology [358]. Though amorphous structure could significantly improve the electrochemical performance of some electrode materials, the structure stability is an issue need to 109

be considered in particular for long-term charge/discharge cycles since amorphous phase is thermodynamically unstable.

Figure 40. (a) XRD patterns (with V2O5 reference) and (b) CV plots (collected at 1.0 mV s−1) for the amorphous (red) and crystalline (black) V2O5; the rate capability of the (c) amorphous and (d) crystalline V2O5 when discharged at current densities ranging from 23.6 (0.1 C) to 1170 mA g−1 (5 C). Reprinted with permission from ref [357]. Copyright 2014 The Royal Society of Chemistry.

At present, SIBs are an very promising technology brimming with much promise and intrigue. In theory, SIBs present a complementary alternative to LIBs owing to the primary advantages of abundant sodium resources, potentially low cost, and better safety. However, the reignition of interest in this field has revealed that analogous electrode active materials can behave substantially different than expected based on their prior LIB performance. One of the main issues SIB electrodes face is the ability to successfully intercalate/deintercalate sodium ions fastly and

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reversibly and with good stability during long-term cycling because sodium ions are substantially larger than their lithium ions counterparts, and thereby are much more intrusive and disruptive on the host material. Towards combatting this issue, vanadium oxide is well suited for SIB application given its large interlayer spacing, minimal volume expansion upon cycling, and high capacity; however, this is oftentimes not enough and more effort must be devoted in order to make the structure more receptive to the sodiation process. In summary, there is still much work to be done in the development electrode materials better suited for sodium insertion/extraction before SIBs can be implemented on a large scale.

5.2. Mg-ion intercalation in V2O5 and Related Vanadate Materials 5.2.1 Mg battery primer

As with Na-ion batteries, research on Mg-ion batteries is still in the exploratory phase, although interest is rapidly increasing [58]. The electrochemical properties of MgxV2O5 were first reported in 1987 [359-361], although Gregory et al. [362] first studied V2O5 in the context of rechargeable magnesium batteries. In general, these initial studies yielded modest results with respect to either electrochemical or chemical intercalation [362, 363]. Nevertheless, the pioneering effort by Gregory et al. showed some promise for V2O5, which had a fairly high open circuit voltage of 2.66 V vs. Mg/Mg2+ and a capacity of 196 mA h g-1 via chemical intercalation. But limited by a dearth of Mg-compatible electrolytes, Mg-ion battery studies were sparse, especially after the commercialization of Li-ion batteries in the early 1990’s. Aurbach et al. [364-366] have since 111

contributed several important breakthroughs, and in 2000 demonstrated a full Mg-ion battery that was capable of thousands of cycles. A thorough review of this subfield is beyond the scope of this paper, but the interested reader is directed toward a rapidly growing number of cathode [58, 59, 367-369], electrolyte [370, 371], and comprehensive reviews and perspectives [57, 364, 372-379] pertaining to Mg-ion batteries.

In brief, magnesium has several advantages over lithium as a charge carrier. As shown in Table 3, magnesium has a higher reduction voltage, but it is significantly cheaper and more abundant. Since it is divalent, it also delivers twice as much charge compared to Li, and boasts an extremely high volumetric capacity (3832 mA h g-1). Handling, machining and disposal are easier for Mg since it does not react violently with air or water. In addition, its ionic radius is comparable to lithium, meaning no additional mechanical strain is placed on the host lattice during insertion/deinsertion (unlike sodium, for example). Quite significantly, magnesium is not susceptible to dendritic growth during plating/stripping, so it could be used as a high capacity anode without encountering the same safety problems as metallic lithium [380-382]. Considering these factors, one can envision a net gain in energy density at a reduced cost to manufacture.

However, from the perspective of solid-state chemistry, there exist major challenges in the development of commercially viable Mg-ion batteries: (a) unlike Li, surface passivation layers on Mg do not conduct Mg ions, and thus the choice of electrolytes is very limited [362, 370, 383], (b) to get around the insulating surface layers on Mg, anodic materials may need to be developed [384-387], and perhaps most challengingly, (c) suitable cathodic materials possessing high

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insertion voltage and capacity coupled with excellent long-term stability must be developed [58, 59, 368].

With this in mind, it is important to realize that finding a suitable V2O5 cathode material (or any other material, for that matter) will not be the end of the story for rechargeable magnesium batteries. Since vanadate electrodes have almost exclusively been studied in electrolytes that passivate Mg metal, additional compatibility studies will be necessary to combine the cathode with high-voltage stable electrolytes that do not react with Mg. In one of few reported instances of such an experiment, Yoo et al. [364] tested V2O5 in a 2:1 mixture of PhMgCl and AlCl3 in tetrahydrofuran (Figure 41). This electrolyte is knowns as the “all-phenyl complex” or APC and has a wide voltage window in excess of 3V with respect to magnesium [365]. The well-known Chevrel phase cathode materials (Mo6S8) [366] support reversible Mg2+ intercalation in this electrolyte when paired with a Mg metal anode, but no insertion is detected when V2O5 is used as the positive electrode. (However, a lack of any other current signal indicates that the metal oxide is at least thermodynamically stable towards the electrolyte.) Addition of LiCl in the same electrolyte makes Li insertion possible, which implicates the removal of the solvation structure of Mg2+ at the electrode surface as a much more challenging problem to overcome. This also suggests that the ideal Mg electrolyte may depend on the cathode material in question; as another example, TiO2 has performed poorly with APC but gave positive results when tested in a magnesium borohydride/tetraglyme electrolyte [388]. Less desirably, suitable intercalation anodes may be required, although these would invariably fall short of the capacity of metallic Mg. A third option involves developing solid-state electrolytes. V2O5 has been used as a cathode in a number 113

of such studies on polymer gel electrolytes [389-393]. As of now these cells generally suffer from low capacities and/or poor cycling performance, although solid electrolytes capable of efficient Mg stripping and plating have been demonstrated [394].

Figure 41. A V2O5 electrode showed no Mg insertion in the APC electrolyte, but Li insertion peaks are observed upon addition of LiCl. Reprinted with permission from ref [364]. Copyright 2013 The Royal Society of Chemistry.

As alluded to above, insertion of the divalent magnesium ion into positive electrode materials has consistently been more problematic than lithium. This has widely been attributed to larger electrostatic forces and desolvation effects [325, 372, 395]. As a consequence, Mg experiences high kinetic barriers to diffusion within the material, and slow insertion kinetics at the electrode-electrolyte interface. Upon intercalation, it has been observed that most of the Mg-ions are stuck at the surface of single crystal V2O5, and cannot diffuse very far into the bulk [396]. With respect to V2O5, there have been two main approaches used to ameliorate these barriers and improve transport properties. One is tailoring the nanoscale morphology of the cathode materials,

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and the other involves screening the charge of divalent Mg-ion. Vanadium aerogels and xerogels (of the form V2O5·nH2O) explore both approaches simultaneously. At the same time, heterostructures and dopants have been used to improve the electrochemical properties (e.g. electronic conductivity and charge transfer resistance) of vanadium oxides. More recently, computational studies have emerged as valuable tools to help guide materials screening efforts and predict properties. Given its potential as a high energy cathode, V2O5 has often been used as model material; the salient insights from such studies are summarized as well.

5.2.2 Nanostructured V2O5 for Mg-ion Intercalation

Reducing the particle size and increasing the material surface area shortens the diffusion length in the active material. This well-worn path has been used to increase performance, perhaps most notably in the case of LiFePO4 [105]. However, the shortened diffusion length may come at the cost of enhanced parasitic side reactions between the electrode and the electrolyte [368, 379]. Such reactions hurt Coulombic efficiency and must be minimized as much as possible. Additionally, nanosized materials will have considerable challenges scaling up because their low tap density places limits on the volumetric energy density at the cell level [397].

Amatucci et al. [325] demonstrated Mg intercalation into nanocrystalline V2O5. A powder with particle sizes in the 20–50 nm range was synthesized using combustion flame-chemical vapor condensation. BET analysis showed the powder had a high surface area of ~90 m2 g-1, and the electrode prepared from this powder provided a reversible capacity of 180 mA h g-1 for Mg-ion intercalation. However, the large potential hysteresis (Figure 42) indicates slow reaction kinetics. 115

Figure 42. Reversible electrochemical insertion of Mg2+ into nanocrystalline V2O5. The electrolyte was 0.5 M Mg(ClO4)2 in propylene carbonate and the current density was 7.6 mA g-1. The 3-electrode cell contained an activated carbon counter electrode and Ag quasi-reference electrode. Reprinted with permission from ref [325]. Copyright 2011 The Electrochemical Society.

Jiao et al. [398-400] prepared VOx nanotubes that intercalated Mg2+. The nanotubes were obtained after a sol-gel reaction between V2O5 and octadecylamine, followed by hydrothermal treatment in an autoclave at 180°C for one week. The open-ended nanotubes were 1−3 µm long with outer diameters of 60−100 nm, and inner diameters of 15−40 nm, resulting in a large surface area. Electrochemical Mg2+ intercalation was possible, but the initial capacity was only 80 mA h g-1 and quickly faded. Doping the V2O5 gel with copper (Cu/V mole ratio = 0.1) before the hydrothermal treatment enhanced the conductivity of the nanotubes, as confirmed by EIS measurements [399]. The initial discharge capacity for the copper-doped improved to 120 mA h g-1. While the capacity and stability were still low, this result demonstrates that transition metal dopants can positively affect the capacity and kinetic behavior of V2O5.

More recently, Kim et al. [401] used a hydrothermal method to prepare VOx nanotubes using 116

V2O5 and either a low or high concentration of octadecylamine (denoted LT-VOx and HT-VOx, respectively). The energy for the hydrothermal process was provided by a microwave heated at 180°C for 18 hours. The nanotubes obtained from this process measured 1-5 µm in length and 80-100 nm wide (outer diameter). Electron microscopy (both SEM and TEM) revealed no apparent difference in morphology between the two samples. However, XPS showed that while the LT-VOx nanotubes contained roughly an even mixture of V5+ and V4+ metal centers, the HT-VOx nanotubes also contained a detectable fraction of V3+ centers. This difference proved to be an important factor in the cyclic stability (Figure 43b). While both cathodes exhibited high initial capacities (>210 mA h g-1), the HT-VOx nanotubes maintained a capacity of 150 mA h g-1 after 20 cycles, corresponding to ~71% capacity retention. Conversely, the LT-VOx capacity had already decayed to ~70% capacity after only four cycles. The authors attributed the improved cyclic stability and reduced charge transfer resistance (as measured by EIS) to the presence of V3+-O clusters in the nanotubes, which repel Mg ions less than the same clusters at higher oxidation states and enhance the structural stability of the nanotubes. Similar electrochemical performance was reported for a graphene oxide/V2O5 composite [402]. Du et al. obtained a high capacity of 178 mA h g-1 and retained 140 mA h g-1 after 20 cycles at a 0.2C rate using a dichloro complex electrolyte [366].

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Figure 43. (a) Initial charge-discharge voltage profile (0.2 C rate) and (b) cyclic performance and Coulombic efficiency of the HT-VOx and LT-VOx nanotubes. Reprinted with permission from ref [401]. Copyright 2014 The Royal Society of Chemistry.

Gershinsky et al. [403] sought to decouple the effects of shielding and size (which are convolved in the case of xerogels, for example) by investigating thin film V2O5 electrodes. A film ~200 nm thick was vacuum deposited from V2O5 powder onto a Pt foil substrate. By taking the differential capacity of the galvanostatic profile, four distinct events for both insertion and deinsertion were resolved (Figure 44a,b), hinting at the complexity of the mechanisms for (de)magnesiation at the electrode surface. Stable cycling with a capacity of ~150 mA h g-1 was achieved for 35 cycles, with an initial capacity of ~180 mA h g-1 and nearly 100% Coulombic efficiency (Figure 44d). This capacity corresponds to a stoichiometry of Mg0.5V2O5.This study demonstrated the efficacy of thin film electrodes to study Mg insertion mechanisms, which in general are not well understood and significantly more complicated than Li insertion.

Solid-state nuclear magnetic resonance (NMR), and magic angle spinning (MAS) NMR in

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particular, is another compelling but underutilized technique that can help resolve the chemical bonding environment (and by extension, reaction mechanisms) in battery materials. Wang et al. [404] employed 25Mg MAS NMR to study the formation of magnesiated V2O5 and other materials. While the reagents used here (di-n-butylmagnesium and diphenylmagnesium) were not effective for chemical magnesiation, they did produce an amorphous surface layer of MgO. With MAS NMR, this layer is easily discernable from MgxV2O5, which in turn distinguishes between a conversion or intercalation mechanism.

Figure 44. A differential capacity plot obtained for V2O5 thin-film electrodes with 0.1 M Mg(TFSI)2 in acetonitrile as the electrolyte and activated carbon cloth reference and counter electrodes at a current density of (a) 0.5 µA cm-2 and (b) 1 µA cm-2. (c) The galvanostatic dis/charge profile of V2O5 shows a small voltage hysteresis. (d) capacity retention and Coulombic efficiency as a function of cycle number. Reprinted with permission from ref [403]. Copyright 2013 American Chemical Society.

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5.2.3 Mg-ion charge screening

The strong influence of water on the insertion of Mg2+ into V2O5 was first noted by Novák and Desilvestro [405]. Microcrystalline V2O5/carbon black composite electrodes could not cycle in any capacity in dry electrolytes, but the electrochemical performance improved with the addition of water, and reached a maximum when 1M H2O was added to the electrolyte (1M Mg(ClO4)2 in acetonitrile). The initial capacity for V2O5 in this electrolyte was ~170 mA h g-1, but quickly faded below 50 mA h g-1 after only 20 cycles. A similar study demonstrated comparable results with wet and dry propylene carbonate as the solvent; an initial capacity of ~158 mA h g-1 was achieved in the wet electrolyte but the cyclic stability was poor [406]. Co-intercalation of water is the likely explanation for this improvement. Adding the bulk of the partial hydration sphere reduces the charge density and facilitates ion mobility. However, the presence of water in the electrolyte renders it incompatible with Mg metal as an anode.

To attempt to get around this issue, various hydrated vanadium bronzes (MV3O8·nH2O, for which M was an alkali or alkali earth metal) were studied [407, 408]. The active material was prepared by combining appropriate stoichiometric amounts of V2O5 and an aqueous alkali hydroxide solution or an aqueous suspension of the alkali earth oxide, then vacuum drying the resulting material. The bronzes cycled in MgCl/AlCl3/1-ethyl-3-methylimidazolium chloride performed the best after being dried at 100°C to remove a fraction of the water bound to the host lattice. MgV3O8·3H2O gave an initial capacity of ~150 mA h g-1 and maintained 80 mA h g-1 after 60 cycles. Their results indicated the importance of structural water in the bronze, however IR

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spectroscopy indicated that water was expelled from the material during cycling, which adversely affects performance.

Until recently, only two reports by Tang et al. [343] and Le et al. [409] have investigated Mg-ion intercalation into aerogels. In both cases, the aerogel was synthesized by supercritically drying a V2O5 hydrogel using liquid CO2. While up to 4 equivalents of Mg could be intercalated chemically using MgBu2 (i.e. Mg2V2O5was formed) [409], no more than ~0.6 mol of Mg2+ were reversibly inserted electrochemically in either report. Both pseudocapacitive charge storage [343] and regular intercalation [368] have been proposed as the relevant mechanism.

Another study [410] prepared an aerogel composite of V2O5 nanowires and reduced graphene oxide. This configuration allowed for favorable Mg2+ and electron transport. Using galvanostatic intermittent titration, a high capacity of 330 mA h g-1 was obtained, and when cycled at a rate of 1 A g-1, good capacity retention was observed (81% of ~120 mA h g-1 after 200 cycles). The composite also demonstrated a wide operating temperature range: capacities of 40 and 170 mA h g-1 were achieved at -30 and 55°C, respectively. Significantly, they investigated whether the interlayer spacing or crystal water content was the source of the performance boost. The water content can be controlled by annealing at different temperatures (Figure 45a) and the interlayer spacing was measured via XRD (Figure 45b), and gave results that are similar to the earlier work by Wang et al. [170]. The water content was more strongly correlated to performance, whereas the interlayer spacing only varied from 10 to 11 Å. This suggests that shielding the charge of Mg2+ in the host material with water is a better strategy to reduce electrostatic effects than physically

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enlarging the unit cell of the host. Still, the interlayer spacing remained wide in all the samples examined, so it may be that a large spacing is a necessary but not sufficient condition for good performance.

Figure 45. (a) Thermogravimetric analysis of V2O5 aerogel. (b) XRD patterns and (c) corresponding cycling performance of aerogel samples annealed at different temperatures. (d) The effect of water content on the observed capacity and interlayer spacing. Reprinted with permission from ref [410]. Copyright 2015 Elsevier B.V.

Vanadium oxide xerogels, similar to vanadate bronzes, have been more thoroughly studied [411-415]. As mentioned earlier, vanadium oxide gels benefit from combined effects of high active surface areas and the charge screening of the divalent magnesium ion by water in the lattice. 122

These materials are easily obtained using sol-gel processing techniques, which lend themselves to commercial scale up. These gels are composed of bilayers of V2O5 layers separated by water molecules that serve to provide large interlayer spacing in excess of 10 Å [67]. Unlike orthorhombic V2O5, the bilayers are constrained and do not undergo phase transitions, which benefits long-term stability. Gels often show unusually high voltages for Mg hosts as well [409, 412, 414]. This boost is attributed to the co-insertion of the electrolyte; the solvation sphere is not fully shed upon insertion, and the uncompensated solvation energy necessarily pushes the voltage up [367].

Stojković et al. [413] studied a xerogel/carbon composite in an aqueous solution saturated with Mg(NO3)2. A capacity of ~107 mA h g-1 was obtained for the composite, compared to less than 50 mA h g-1 available in the case of just the xerogel or microcrystalline V2O5 by itself. The same authors further examined this composite’s performance in additional metal nitrate electrolytes [415]. Modest capacities were obtained throughout (~100 mA h g-1 or lower), with co-intercalation of water implicated as one of the possible reasons limiting the capacity. However, UV-Vis spectroscopy and visual inspection of the electrolyte upon cycling showed that the degree of dissolution of V5+ from the electrode depended on size and valance of the inserted ion. That is, K+ (1.51 Å) apparently disrupts the host the most and led to poor stability, followed by Na+ (1.09 Å) and Li+ (0.76 Å). But the comparably sized Mg2+ ion (0.72 Å) showed the best capacity retention and rate performance (Figure 46), which the authors attribute to Mg-O interactions “gluing” the bilayers together to enhance stability.

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Figure 46. (a) Intercalation capacities of V2O5/graphite composites obtained from cyclic measurements in aqueous alkali metal nitrate electrolytes. (b) Rate performance at the indicated rates after 10 cycles at 1000 mA g-1. Reprinted with permission from Ref [415]. Copyright 2015 Elsevier B.V.

Imamura et al. [411, 412] prepared a xerogel/carbon composite electrode that could intercalate up to 1.84 mol of Mg2+, corresponding to a capacity of ~540 mA h g-1 of active material. A sol-gel step using metallic vanadium and H2O2 was used to prepare the active material. The cyclic voltammagram resembled that of Li, in that both show two peaks upon charge/discharge. This indicates the presence of two distinct insertion sites in the xerogel, which are identical for Li+ and Mg2+. Intercalation into the “innerlayer” sites (that is, in between the space afforded by four corner-sharing VO5 square pyramids) took place at ~3 V vs. Mg/Mg2+ while the “interlayer” insertion placed Mg ions above apical oxygen atoms and was observed at ~2.5 V vs. Mg/Mg2+. Galvanostatic discharge of the material at 1 A g-1 gave capacities approaching 600 mA h g-1. However, there may be a discrepancy in the reported rate [367] (i.e. 1 mA g-1 compared to 1 A g-1) which would suggest this discharge rate is prohibitively slow for practical purposes. Nevertheless, 124

they also obtained stable cycling above 300 mA h g-1 for 30 cycles.

Lee et al. [414] prepared Mg-doped V2O5 starting from MgV2O6 [416] dissolved in water. This material was treated with an ion exchange resin to obtain a magnesium deficient sol. The final xerogel had a formula of Mg0.1V2O5∙1.8 H2O. The amount of magnesium could be controlled in this low-temperature route, and was tuned such that the material stoichiometry was Mg0.1V2O5. The kinetics of insertion were shown to depend on the electrolyte, with acetonitrile offering more reversible behavior than carbonate-based electrolytes. This result is in agreement with a recent computational analysis on the desolvation of Mg2+ in a wide variety of solvents [395]. Cycling at a rate of C/10 in 0.5 M Mg(ClO4)2 in acetonitrile gave an initial capacity of ~300 mA h g-1, and remained above 250 mA h g-1 for seven cycles (Figure 47).

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Figure 47. Slow scan cyclic voltammograms at a scan rate of 0.1 mV s-1 of Mg-doped V2O5 with Ag/Ag+ as the reference electrode in (A) 0.1M MgClO4 in acetonitrile and (B) 0.1M Mg(TFSI)2 in a 30:70 mixture of ethylene and dimethyl carbonate. (C) Charge-discharge profile for the same material with 0.5 M Mg(ClO4)2 in acetonitrile as the electrolyte for cycles 1, 2, and 5 and (D) cyclic capacity (a = discharge; b = charge). Reprinted with permission from ref [414]. Copyright 2014 The Electrochemical Society.

Inamoto et al. [417] doped orthorhombic V2O5 with sulfur. Sulfur is more polarizable than oxygen, which decreases the bond energy and voltage of sulfides relative to oxides, but at the same time allows for easier insertion/deinsertion kinetics [365]. A sulfur-metal oxide-V2O5 composite was prepared by generating plasma using a microwave and pieces of wetted carbon felt at low pressure. The metal oxide was one of MnO2, MoO3, Fe2O3, NiO, or ZrO2, and was incorporated to prevent dissolution of sulfur. Borrowing from previous results with added water [405, 406], 0.3 M Mg(ClO4)2 with 1.8 M H2O was employed as the electrolyte. A high initial capacity of 420 mA h g-1 was obtained when MnO2 was used as the additive (compared to 300 mA h g-1 with just the sulfur added to V2O5, and 170 mA h g-1 for orthorhombic V2O5), although long-term cycling performance was not reported and the discharge potential was relatively low (~1.5 V vs. Mg/Mg2+). Regardless, doping with sulfur and transition metal cations warrants more attention as a strategy to improve Mg-ion insertion and diffusion kinetics.

Along the same lines of the NaxV2O5 study [348], Tepavcevic et al. [418] also investigated Mg2+ intercalation into bilayered V2O5. The electrodeposited oxide is structurally very similar to xerogels and accommodates good electrical and ionic transport. X-ray fluorescence (XRF) 126

microscopy showed a uniform distribution of Mg ions in the V2O5 host. When paired with HAADF STEM imaging, this provided strong evidence at the atomic-scale resolution for intercalation between the bilayers as the relevant reaction mechanism. The electrode was discharged against Mg metal using Mg(ClO4)2 in acetonitrile as the electrolyte and provided a capacity of ~240 mA h g-1. This corresponds to the reduction from V5+ to V4+ and agreed with the oxidation state change observed from XANES spectroscopy. A full cell was also demonstrated after discharging the cathode to introduce Mg and then replacing the anode with nanocrystalline tin/carbon composite. This cell reached a capacity of ~150 mA h g-1 when cycled at a C/15 rate, and XANES and XRF measurements showed this capacity was limited by the anode. The authors stressed the importance of structural water, as it helps maintain the interlayer spacing for Mg2+ intercalation (or partially solvated Mg2+ co-intercalated with water or other electrolyte components) and reduces electrostatic effects via dipole interactions.

5.2.4 Computational Studies for V2O5

Several recent computational studies have used V2O5 as a model system and are worth reviewing here. Such studies allow us to examine many materials and phenomena in parallel as a rate that far exceeds what can be expected from traditional trial-and-error methods [419].

Wang et al. [420] first used density functional theory (DFT) calculations to investigate Mg-ion batteries using V2O5 as the host structure. They compared the transport properties of Li and Mg in both bulk and single-layered V2O5. They found that binding energy of Mg to V2O5 was much larger for the single layered V2O5 compared to bulk (4.06 vs. 2.20 eV, respectively). While Li 127

diffusion benefitted from the single layer configuration, the calculated diffusion barrier for Mg was ~1.4 eV regardless of the form of V2O5. The number of Mg-O bonds is reduced from four to three in the single-layered case, but this is apparently not enough to significantly affect the energy barrier for diffusion.

Bulk and single-layer V2O5 was also studied using DFT by Vujković et al. to support their experimental results [415]. Although orthorhombic V2O5 was modeled and xerogels of V2O5 were studied experimentally, the trends in one correlated well with the other. They considered the diffusion barrier and structural effects of intercalation for Li+, Na+, K+, and Mg2+. The calculated barrier for Mg diffusion was 0.92 eV for bulk V2O5 and 0.82 eV for single-layer V2O5, and they found that size was a stronger determining factor for bulk diffusion (i.e. the barrier increased in the order: Li < Mg < Na < K) and valance was a larger determinant for single-layer diffusion (Li < Na < K < Mg). Structurally, Mg insertion induced the smallest volume change (<2%) and matched the stability trend in their experimental results. The stability was in part attributed to the strong bond between Mg2+ and apical V=O upon intercalation, suggesting at least some trade-off between high diffusion barriers and high structural stability.

In another computational study, Carrasco [421] explored the role of van der Waals (vdW) forces in thermodynamic and kinetic properties of alkali and alkaline-earth ion insertion into α-V2O5. They found that the insertion energy for Mg into V2O5 is significantly affected by the position of the V4+ ions (Figure 48). Energy differences between global minima and other configurations can be up to 0.85 eV in the case of Mg2+ ion (Fgiure 48b). This large energy difference is associated

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with the atomic relaxation around the V4+ centers, whose size is larger than that of V5+ ions. The configurations with both V4+ ions close to the inserted Mg2+ are preferred (Figure 48). In particular, configurations where each V4+ ion is located in different bilayers are substantially more stable than others with two V4+ ions in the same bilayer. This arrangement minimizes the geometric strain caused by the larger V4+ centers and also equalizes electrostatic repulsion between Mg2+ and V5+ centers from different directions. In addition, configurations with V4+ ions close to each other and far from the inserted ion are particularly unfavorable. Incorporating van der Waals forces increased the calculated Mg-insertion voltage from 2.50 to 2.70 V, which agrees well with the open circuit voltage of 2.66 V observed by Gregory et al.5 This effect apparently arises from the vdW model’s inclusion of nonlocal correlation effects. Additionally, inclusion of vdW forces reduces the spacing between V2O5 layers, which imposes an additional geometrical strain to the transition state structure and, consequently, led to a higher diffusion barrier (0.97 eV) for Mg2+ migration along the b-axis, compared to that (0.77 eV) without vdW forces.

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Figure 48. (a) Supercell containing three V2O5 bilayers with one ion inserted into the most favorable site. (b) Intercalation energy for Li, Na and Mg as a function of the position of the V4+ ions formed upon insertion. (c,d,e) Bilayers 1−3 with vanadium positions identified viewed from above. Yellow: alkali (earth) metal; Blue: V; Red: O. Reprinted with permission from ref [421]. Copyright 2014 American Chemical Society.

Zhou et al. [422] calculated the open circuit voltage, band structure, diffusivity, and α→δ phase transition energy of V2O5 (The two phases differ in how the layers are stacked on top of each other, Figure 49a) by first principles calculations. Their calculated open circuit voltage of a magnesium battery with V2O5 cathode is 3.06 V, with an even higher open circuit voltage obtained when vdW forces were considered (3.60 V). This value is much higher than that (2.7 0V) calculated by Carrasco et al. [421], but still follows the trend that incorporating vdW forces leads to stronger interlayer binding that in turn increases the voltage relative to the 130

Perdew-Burke-Ernzerhof (PBE) functional. This calculated voltage (3.06 V) is even higher than that for a lithium battery by 0.22 V. While such a value has precedent in the literature, this only pertains to vanadate gels rather than orthorhombic V2O5 [409, 412], where in the former, co-intercalation of the solvent has been implicated in raising the voltage [368, 379, 423]. Though V2O5 is known to have conductivity issues, they predict that MgV2O5 is highly conductive, with a very small band gap of 0.02 eV. In α-V2O5, the diffusion barrier for Mg ions is 1.26 eV, whereas in δ-V2O5, the barrier increases to 1.38 eV and shows a double peak configuration. That is, Mg ions diffusing through δ-V2O5 pass through a high-energy activated state, a local energy minimum at an intermediate state, and then another activated state before reaching the final position. They found that the the α→δ transformation is a barrierless process, and it takes place via a rotation of the adjacent [V2O5]n layers as the magnesium content increases (Figure 49b).

Figure 49. (a) Schematic diagram of the α-V2O5 to δ-V2O5 transforamtion. (b) The potential energy curves of the transitions from α-V2O5 to δ-V2O5 and from α-MgV2O5 to ordinary MgV2O5. Reprinted with permission from ref [422]. Copyright 2014 The Royal Society of Chemistry.

Gautam et al. [424] derived the full temperature-composition phase diagram of MgxV2O5 131

(0 < x < 1) by Grand canonical Monte Carlo (GMC) simulations. MgxV2O5 exists as a two-phase equilibrium between α-V2O5 and δ-V2O5. However, they note that a metastable phase (ε-V2O5) undergoes Mg ordering when V2O5 is half magnesiated (Mg0.5V2O5). This phase may dominate since δ-V2O5 requires structural rearrangement of the crystal lattice, a kinetically limited process. Such metastable transitions have been documented in, for example, LiFePO4 and may be present here as well [425]. Contrary to Zhou’s results

408

, they find that the barriers for Mg migration in

δ-V2O5 (~600-760 meV) are significantly lower than in the α-V2O5 (~975-1120 meV), and predict that cycling V2O5 exclusively in the δ phase may improve performance.

The same researchers expanded this work to include additional kinetic and thermodynamic parameters pertinent to multivalent cation insertion in the α and δ polymorphs [426]. Specifically, the structure change, voltage, stability, and ion mobility within the host were examined for the insertion of Li+, Mg2+, Ca2+, Zn2+, and Al3+. In this case they used a functional that considers van der Waals forces (vdW-DF2 [427]) and found that the formation of MgV2O5 expands the interlayer spacing in the δ phase by less than 2%, but causes a 10% increase in α-MgV2O5. They also found that δ-MgV2O5 is only slightly less stable than the ground state (by ~27 meV atom-1) while α-MgV2O5 is roughly four times less stable (~102 meV atom-1). The voltage also depended on the polymorph, as in insertion voltage was estimated as 2.56 and 2.21 V for the δ and α phases, respectively. These values are lower than predicted by the other computational studies, but compares favorably with some experimental results [403].

Notably, they were also able to correlate the energy barrier to diffusion to the coordination

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environment of the cation within the host structure [424, 428]. The nudged elastic band method was applied to identify the diffusion path and its associated energetics for the same cations in the δ-V2O5 polymorph, as well as other compelling cathode materials: spinel Mn2O4, olivine FePO4 and layered NiO2. While they recovered the expected trend with valence state (i.e. higher valance cations experience higher diffusion barriers), their analysis showed a wide dispersion in energy barriers for divalent cation migration.

For example, a low barrier of ~200 meV for Ca2+ is observed in δ, but this becomes as high as ~800 meV for Mg2+ diffusion. However, in α-V2O5 the Ca2+ barrier soars above 1700 meV, and for Mg2+ it increases to ~1000 meV (Figure 50). These differences depend on the preferred coordination number of the cation in question and the topology of the diffusion path. Ca2+ prefers to be coordinated by eight anions and Mg2+ prefers six. In the α phase, the stable cation site is coordinated by eight oxygens, and must pass through a 3-coordinated face corresponding to an activated state (8→3→8). However, in the δ phase, the cations pass through the same path as described by Zhou et al. [422] Along this path, the cations are first coordinated by four oxygens with ~2 Å-long bonds and two additional oxygens by ~2.33 Å bonds (“4+2”) and diffuse by passing through two 3-coordinate activated states separated by a 5-coordinate local minimum before reaching the next stable site “ ( 4+2”→3→5→3→“4+2”). Thus, for the α phase, Ca2+ must start in its most stable coordination environment, then break five bonds before migrating to the next site, resulting in the large energy barrier. On the other hand, Ca2+ in the δ experiences a less severe change in bonding environment and never achieves its preferred coordination number, resulting in the low energy barrier. The reasoning is similar for Mg2+: the “4+2” site is quite stable 133

and requires an appreciable amount of energy to displace it along the rest of the δ-phase diffusion path (~800 meV), but the severe change in α-phase coordination results in a slightly greater barrier (~1000 meV).

Hence, the topology of the structure is shown to play an integral and heretofore underappreciated role in determining the diffusion barrier, which has several implications. First, high-rate multivalent batteries should benefit from matching cathode materials with cations that have dissimilar preferred coordination environments compared to the most stable intercalation sites. Forming such a metastable compound on discharge reduces the voltage but increases safety. At the same time, synthesizing Mg-containing cathode materials directly is not advisable as this likely places Mg in a stable site with high diffusion barriers.

Figure 50. The activation barriers for diffusion for five ions through (a) α-V2O5 and (b) δ-V2O5. Dashed and solidlines correspond to the diffusion when the host material is charged and discharged, respectively. Reprinted with permission from ref [426]. Copyright 2015, The Royal Society of Chemistry. 134

There are some discrepancies to be addressed between these studies that arise from differences in the computational procedures. Current models cannot yet simultaneously capture all the various physical and chemical phenomena and future experimental work should try to confirm these findings. Nevertheless, computational materials science is rapidly becoming a valuable tool that will only become more useful as processing power increases. The conclusions drawn from these studies should inform the design of high energy density cathode materials for Mg-ion batteries.

5.2.5 Perspectives on V2O5 cathodes for Mg-ion Batteries

The many studies on V2O5 have shown potential as a cathode material for Mg-ion batteries, but limited stability and transport kinetics remain unsolved challenges. Tailoring the nanostructure of the active material and a screening the divalent charge of Mg2+ with water have been the most popular approaches to improve electrochemical performance. Cation and anion doping are additional techniques that improve conductivity and charge transfer kinetics. Recent computational studies using V2O5 as a model material have shed light on the mechanisms behind its leisurely transport behavior.

Another issue for V2O5 is that it has been nearly exclusively studied in electrolytes that are not compatible with Mg metal, and it is clear that the electrochemical performance is affected by the choice of electrolyte [395, 414]. The most common electrolyte for these studies has been Mg(ClO4)2 in acetonitrile. However, Mg(ClO4)2 is strongly hygroscopic [429], and the potential 135

effects of trace water (or other impurities for that matter) have generally not been controlled or discussed. Further, co-intercalation of solvents and the inclusion of crystalline water has clearly improved the performance of V2O5 compounds. However, the relative dearth of long-term stability studies leaves open the possibility that solvent co-intercalation may be detrimental to the structural integrity of the host, while appreciable water content may be a factor preventing V2O5 from reaching its maximum capacity.

The electrochemical setup for many Mg-ion batteries cathode studies has also been questionable [59]. Among other issues, it is conceivable that using Mg metal – which is passivated in most electrolytes [383]– as a counter or reference electrode artificially limits cathode performance. New high voltage electrolytes have been proposed [370], and compatibility studies will be needed to show if any in this set of electrolytes can simultaneously support electrochemical reactions at both the cathode and the anode.

5.3. Other cations intercalation in V2O5 and Related Vanadate Materials Inspired by the encouraging results obtained for Na+ and Mg2+ insertion into V2O5, the possibility of reversible intercalation of other cations (Zn2+, Ca2+, and Al3+) are also investigated very recently. Senguttuvan et al. [430] reported reversible Zn2+ intercaltion in bilayered hydrated V2O5 (BL V2O5). To evaluted the feasibility of Zn2+ intercalation into bilayered hydrated V2O5 cathode and its compatibility, ialkt was cycled against Zn metal in 0.5 m AN-Zn(TFSI)2 electrolyte at C/10 in the potential range of 1.5–0.3 V vs Zn2+/Zn (Figure 51a). The BL V2O5 delievers a reversible capacity of ~170 mAh g-1 for 120 cycles at about 0.85 vs Zn2+/Zn (Figure 51b). From 136

the capacity yielded, they find that the amount of intercalated Zn2+ is about 0.59 mol per mole V2O5. During the first discharge/charge cycle, the interlayer spacing of BL V2O5 increased from 11.2 to 13.1 Å during the first discharge and decreased back to 11.5 Å after the end of charge, suggesting the entrapment of Zn2+ in BL V2O5 as evidenced from the low Coulombic efficiency during the first cycle (Figure 51c). The authors speculated that the interlayer expansion (~1.9 Å) after discharge is directly associated with Zn2+ intercalation in the middle of the interlayer space of BL V2O5; the Zn2+ appears to be solely responsible for the increase of of interlayer spacing because solvent (ACN) cointercalation was ruled out based on the Raman experimental results. The BL-V2O5 also shows an outstanding high rate capability for Zn2+ intercalation. Evan at the high rate of 20 C, the BL-V2O5 still delivers a stable capacity of 130 mAh g-1 (Figure 51d). This high rate capability and fast Zn2+ diffusivity in the lattice interlayer provide substantial experimental evidence for the very high mobility of Zn2+ when matched with the correct host material [426, 428]. Kundu et al. [431] reported a vanadium oxide bronze pillared by interlayer Zn2+ and water (Zn0.25V2O5•nH2O) as the cathode materials for an aqueous rechargeable zinc battery. The Zn0.25V2O5•nH2O exhibits a very high reversible capacity of 300 mAh g-1 at a moderate current density of 50 mA g-1 (C/6). From 1C to 8C, the highest delivered capacity decreases only slightly from 282 to 260mAh g-1. Even at 15C, the Zn0.25V2O5•nH2O still offers a discharge capacity of 223mAh g-1 and exhibits capacity retention of more than 80% after 1,000 cycles. The reversible Zn2+ (de)intercalation was proved by ICP-OES, quantitative energy dispersive X-ray spectroscopy (EDX), and XPS analyses. Further insights into the reversible Zn2+ (de)intercalation and dynamic structural evolution of the host material were clarified from an 137

operando XRD investigation. An increase in the interlayer distance from 10.8Å to 12.9 Å was observed when the pristine Zn0.25V2O5•nH2O was immersed in the electrolyte because of water intercalation. The structural evolutions of ZnxV2O5•nH2O during Zn2+ (de)intercalation are complex: there exist two solid-solution regimes (0.25
138

b)

d)

c)

Figure 51. (a) Schematic diagram of a Zn/BL-V2O5 cell (red: oxygen, green: zinc, blue: vanadium). (b) Discharge/charge curves and (c) cycling performance of the Zn/BL-V2O5 cell at C/10. (d) Rate performance of the Zn/BL-V2O5 cell. Reprinted with permission from ref [411]. Copyright 2016 John Wiley & Sons, Inc.

As a trivalent cation, Al3+ can exchange three electrons during redox reaction, which could provide much higher capacity than the monovalent LIBs, SIBs, and divalent MIBs. Moreover, Al is the third most abundant element in the earth's crust [432] and has the advantages of less reactivity, low flammability, and easy handling. However, the reports of cathode materials that can reversibly insert and extract Al3+ are scare because Al3+ has extremely high charge density. In recent years, V2O5 attracted much attention as possible cathode materials for rechargeable Al-ion

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batteries (AIBs). Jayaprakash et al. [433] firstly reported the application of V2O5 cathode in rechargeable Al battery in 2011. They used orthorhombic V2O5 nanowiers as cathode, aluminum metal as anode, and AlCl3 dissolved in 1-ethyl-3-methylimidazolium chloride ([EMIm]Cl) as electrolyte. Under the current density of 125 mA g-1, the battery delivered a discharge capacity of 305 mAh g-1 in the first cycle and 273 mAh g-1 after 20 cycles, showing the good capacity retention. However, the coulombic efficiency of the cell is low and the discharge plateau is only about 0.55 V, which decreases its potential practical application. Reed et al. [434] further investigated the underlying chemistry of this electrochemical system and convincingly concluded that the V2O5 cathode was electrochemically inactive toward aluminum in the acidic AlCl3-[EMIm]Cl electrolyte and the battery-like performance was attributed to reactions between the AlCl4- and iron and chromium in the stainless steel current collector. However, the following work by Chiku et al. [435] again refuted the conclusion of Reed and Menke. To avoid the interference of the reactions between AlCl4- and stainless steel, they replaced the stainless steel current collector with inert molybdenum (Mo) and the ionic liquid electrolyte with a mixture of AlCl3, dipropylsulfone and toluene (1:10:5 in mole ratio). The possibility of reversible Al3+ insertion/extraction into/from amorphous V2O5/carbon composite was convincely verified by cyclic voltammetry, charge-discharge tests, and XPS analysis. The CV of the amorphous V2O5 exhibited reversible reduction/oxidation reaction peaks at 0.8/1.6 V vs. Al/Al3+. Charge-discharge tests demonstrated that the rechargeable Al cell with the V2O5/carbon positive electrode exhibited maximum discharge capacity over 200 mA g-1 with the discharge rate at C/40 (11.05 mA g-1). XPS analysis after discharging and the following charging verified that the redox of vanadium ion in 140

the V2O5 active material occurred, and the average valence of V changed between 4.38 and 4.91 during discharging and the following charging. A parallel study reported by Wang et al. [436] also proved that V2O5 was reactive towards Al metal in acidic AlCl3/[1-butyl-3-methylimidazolium] (BMIM)Cl electrolyte. Very recently, Gu et al. [437] provided clear evidence of reversible Al3+ storage in V2O5 nanowire. Based on the analyses of thermodynamic and kinetic mechanisms and changes of chemical valence states and structures of V2O5 after cycling, they proposed the following Al3+ storage mechanism (Figure 52): In the first discharge process, Al3+ intercalated into crystalized V2O5 nanowires. Simultaneously, this electrochemical intercalation of Al3+ led to the reduction of V5+ and the formation of an amorphous layer on the edge of nanowires. In the subsequent cycling, a new phase forms along the nanowires’ edges and a two-phase transition reaction occurs. In summary, the research on rechargeable Al batteries is still at its infant stage and great efforts need to be made for its further development. Though V2O5 exhibits a high discharge capacity as cathode material for rechargeable aluminum batteries, its cycling stability, rate capability, and average voltage plateaus appear far from practical applications.

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Figure 52. Schematic diagram illustration Al3+ electrochemical insertion/extraction in nanocrystallized V2O5 nanowires. Reprinted with permission from ref [418]. Copyright 2017 Elsevier B.V.

Ca2+ has a bigger ionic radius than Mg2+ and Al3+ and, therefore, exhibits lower polarizing character [438]. Moreover, Ca is the fifth most abundant element in the Earth's crust [432] and its redox potential (2.86 V vs. SHE) is obviously higher than Mg (2.37 V vs. SHE), Al (1.66 V vs. SHE), and Zn (0.76 V vs. SHE), endowing a significantly higher full cell output voltage than that achievable with Mg, Al, or Zn. There have been several reports on V2O5 intercalation of Ca2+. For example, Amatucci et al. [325] investigated the electrochemical reactivity of Ca2+ into orthorhombic V2O5 and observed a reversible capacity of 200 mA h g-1 in the potential from 1 to 142

-0.5 V vs. a Ag quasi-reference electrode. Hayashi et al. [439] reported a discharge capacity up to 450 mAh g-1 for Ca2+ intercalation into orthorhombic V2O5 at a low current density of 50 A cm-2. However, the diffusion of Ca2+ in V2O5 is very slow as reflected from the large overpotential during discharge and the poor performance at high current densities of 150 and 200 A cm-2. This is consistent with theoretical predictions of a high barrier for Ca migration in the thermodynamically 2+

stable α-V2O5 [426, 440]. To confirm the insertion of Ca

2+

into the oxide lattice, they investigated

the structural changes in orthorhombic V2O5 after discharging (Ca2+ insertion) and charging (Ca2+ extraction) by XRD analysis. After discharging, a new phase caused by Ca2+ insertion reaction was observed in the XRD pattern that coexisted with the pristine orthorhombic V2O5 phase. Moreover, this new phase disappeared and the peaks returned to the parent structure after charging (Ca2+ extraction), suggesting the good structural reversibility during charge-discharge process. Compared with the pristine V2O5, the lengths of the b and c axes of CaxV2O5 decreased slightly, while the length of a axes obviously increased by about 10%, resulting in the lattice volume of the new phase increased by 3%. The same authors found that the amorphous V2O5 showed enhanced Ca2+ intercalation capacity as compared to the crystalline counterpart [439]. Bervas et al. [326] synthesized V2O5 xerogel/propylene carbonate nanocomposite by a sol-gel method. When Ca2+ was used as the guest cation, this nano composite showed a specific capacity as high as 310 mA h g-1 at a current density of 7.58 mA g-1, which was two to ten times better whan the case without propylene carbonate in the strucure. The authors attributed the superior Ca2+ storage performance to the propylene carbonate filled in mesoporous network of V2O5 xerogel, which ensures a very easy and fast transport of intercalating species to the inorganic network. Moreover, the significant 143

amount of structural water molecules in V2O5 xerogel shields the charge of the polyvalent cation. Based on the previous investigations, V2O5 may also be interesting as a host material for K+ [343, 441, 442], Ba2+ [343], and Y3+ [325] intercalation.

6. EFFECTS OF WATER AND SOLVENT CO-INTERCALATION For layered electrode materials, the solvents may co-intercalate with the working ion and alter the crystal structure, permanently remain in the structure or deintercalate with the working ion [443-445]. Moreover, water plays a key role in the intercalation process for V2O5 based cathode materials.Wang et al.108 investigated the effect of incorporpating crystalline water into V2O5 lattice via the intercalation of water molecules on the lithium interclation properties. They made V2O5 sol-gel derived films by drop casting vanadium pentoxide sol onto conductive ITO glass and found that the films with different crystal water content behave very differently as host material for lihtium ion intercalation. The as-fabricated film was confirmed to be V2O5·1.6H2O and after thermal treatment of increasing temperatures, the xerogel film started to lose crystal water (Figure 53a) and underwent a change in crystallinity. The crystal structure, interlayer distance and grain size was dependent on the water content as confirmed by XRD (Figure 53b,c). Electrochemical characterizations demostrated that the V2O5·1.6H2O film has the lowest capacity and the least-satifsying cycling performance, while V2O5·0.3H2O exhibits the best performance, with an initial capacity of ∼275 mA h g−1 that stabilized at 185 mA h g−1 after 50 cycles (Figure 53d). The authors believed that removing too little water limits the electrochemical performance because a 144

significant amount of water may react with lithium to form Li2O; removal of too much water deteriorates the intercalation performance as well, due to the shrinkage of the interlayer distance and the emergence of the crystalline phase. V2O5·0.3H2O was found to be an optimal composition, which has the least possible water amount and keeps the large interlayer distance and the amorphous phase.

Figure 53. (a) TGA curve for V2O5·nH2O xerogels, (b) dependence of interlayer spacing on the n value in V2O5·nH2O, (c) dependence of grain size on the n value in V2O5·nH2O and (d) cycling performance at a current density of 100 μA cm−2 for V2O5·nH2O films obtained at 25, 110, 250 and 300 °C. The voltage ranges from 0.4 to −1.6 V vs. Ag/Ag+. Reprinted with permission from ref [108]. Copyright 2005 American Chemical Society.

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Bai et al. [446] have synthesized CuV2O5 nanobelts and took CuV2O5 as an example to thoroughly investigate the co-intercalation of hydrated protons into layered vanadates during the charge processes in aqueous electrolyte. Careful investigation on the electrochemical properties at different pH values showed that the as-prepared sample exhibited the best cycle performance at pH ∼8.5. The time-dependent evolution study on the intermediate cycling processes indicated that the pH value greatly affected the amount of hydrated protons inserting into the layer space of CuV2O5 and the repeated intercalation/deintercalation of hydrated protons collapsed the structure framework of the electrode with some amorphization. As a result, co-intercalation of hydrated protons was confirmed to be the primary factor accounting for the fast capacity fading at lower pH. Reducing the co-intercalation of protons was proposed to be an effective way to improve the cycling stability of CuV2O5 nanobelts in aqueous lithium ion batteries. The insights gained from the investigation of CuV2O5 make it possible to improve the cycling stability of other similar layered vanadates by adjusting the pH to an appropriate value in aqueous electrolyte.

The incorporation of crystalline water in V2O5 was found to work for SIBs application as well [447]. The hydrated V2O5·nH2O (n = 0.55) cathode displays excellent sodium storage capability of 338 mA h g−1 at a current density of 50 mA g−1, which is much higher than orthorhombic V2O5. Verified by the ex situ XRD and FTIR analysis, the crystalline water in V2O5·nH2O is not totally exchanged during cycling. However, the layer of V2O5·nH2O is shrinking/expanding along with the Na+ ion intertion/extraction and it is reversible, which is different from the irreversible pase transition of α-V2O5. It was noted that the interlayer spacing of V2O5·nH2O did not totally recover to its initial state (with a very small reduction). The authors believed that the shrinkage and 146

collapse of the layer structure during long-term cycling is one of the main reasons for the capacity fading. To solve this issue, Wei et al. [295] developed a novel and facile strategy through iron preintercalation. By pre-inserting iron, the interlayer shrinking and expanding of VOx is largely reduced from 3.74 to 0.94 Å. It also leads to larger stabilized interlayer spacing for reversible Na+ insertion/extraction. The Fe-VOx benefits from the inhibited lattice shrinking/expanding and stabilized interlayer spacing, and consequently exhibits enhanced cycling and rate performance.

Kim et al. [448] synthesized lamellar-structured Ni0.1VOx nanotubes by a microwave-assisted hydrothermal method and cation exchange reaction. When used as a cathode material for SIBs, the Ni0.1VOx nanotubes delivers a high initial discharge capacity up to 140 mA h g-1 and 100% efficiency. The intercalation mechanism and capacity fading effect were investigated in detail by using X-ray diffraction (XRD), transmission electron microscopy (TEM), Fourier transform infrared spectroscopy (FT-IR) and X-ray photoelectron spectroscopy (XPS) analyses, combined with the ab initio simulations. They found that solvent molecules co-intercalates with Na+ during Na+ intercalating into Ni0.1VOx nanotubes,which leads to the expansion of the interlayer spacing and carbon and oxygen adsorption. Based on the simulation and experimental results, the authors proposed that the solvent molecules coordinated the Na-insertion mechanisms into the amine interlayer during discharging. This effect is considered one of the main reasons for fast capacity fading in Ni0.1VOx nanotubes electrodes applied to a Na-ion battery. These understandings of the Na intercalation mechanism in materials based on Ni0.1VOx nanotubes would be useful to design more stable and high-performance VOx-based electrodes for Na battery applications.

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Co-intercalation is a potential approach to influence the mobility and voltage with which cations insert into battery electrodes. Tepavcevic et al. [418] prepared nanostructured bilayered V2O5 by electrochemical deposition method. As-prepared electrochemically synthesized bilayered V2O5 incorporates structural water and hydroxyl groups (Figure 54a). This open-framework electrode shows reversible intercalation/deintercalation of Mg2+ ion in common electrolytes such as acetonitrile. The authors investigated the Mg2+-ion charge, role of water, and solvent co-intercalation by using XPS, HRTEM and X-ray fluorescence (XRF) combined with molecular dynamics (MD) simulations. They have shown using molecular dynamics simulations that the presence of structural water and ensuing hydration of Mg ions is necessary for reversible cycling of Mg ions. Upon prolonged cycling (10–50 cycles) with Mg ions, the host structure becomes more disordered and polycrystalline but improves capacity for Mg2+. A water solvation shell partially shields the charge of the Mg2+ cations, decreasing polarization effects and the strength of their interaction with apical oxygens and hydroxyl groups. The strongly bound structural hydroxyl groups remain in the structure during cycling (Figure 54b) and act as a lubricant for reversible (de)intercalation of solvated Mg2+ ions between enlarged galleries of electrochemically synthesized V2O5. In contrast to water in the electrolyte, which can enter and swell V2O5 but easily comes out of the structure upon cycling due to its weak interaction with the metal oxide framework, structural hydroxyl groups are covalently linked to V2O5 and therefore do not get depleted during cycling [405, 408]. On the contrary, strongly bound hydroxyl groups can participate in hydrogen bonding of water molecules that hydrate Mg2+ ions, facilitating intercalation of integrated solvated Mg2+. The authors found using molecular dynamics (MD) 148

simulation of a V2O5 model slab composed of three bilayers immersed in water that water swells to a structure with an interlayer spacing of ∼12 Å by taking up water to screen all exposed surfaces, reaching two layers of water per V2O5 slab (Figure 54c). MD simulations suggest that upon insertion of Mg ions into the V2O5 structure, the spacing between bilayers narrows to ∼11 Å due to the interaction of Mg ions with bilayer apical oxygens and structural hydroxyl groups (Figure 54d). This is very different from the behavior found for intercalation of large Na atoms, where the size of the crystalline unit cell and the bond lengths were enlarged in order to accommodate the intercalation of a large Na ion.271 Although the spacing decreases, MD simulations show that a significant amount of water remains in the structure, solvating inserted Mg ions. HRTEM and high-angle annular dark field (HAADF) images proved that the Mg ions are inserted in between the V layers (Figure 54e,f). MD simulations suggest that Mg ions are placed close to the oxygen atoms that terminate bilayers (Figure 54d). The insertion of Mg ions in the V2O5 bilayer structure does not remove the first coordination water shell that solvates the ions. Although accompanied by closing of the gap between the bilayers, intercalation does not result in the collapse of the bilayered structure. Upon intercalation, some of the coordination sites of the Mg ions are replaced with an interaction with terminal hydroxyl groups from the bilayer. The authors concluded that the interlayer water plays two important roles: it maintains a sufficient interlayered space to allow the physical diffusion of solvated Mg cations and assists in stabilizing intercalated Mg2+ through dipole interactions. Removal of excess water reduces these two functions and results in poor electrochemical performance of highly crystalline orthorhombic V2O5. As described in section 5.2.3, An et al. [410] reported graphene decorated hydrated 149

vanadium oxide nanocomposite as a cathode material for Mg storage. They also concluded the crystal water in hydrated vanadium oxide plays an important role in shielding the charge of Mg ions to reduce diffusion barrier and increase Mg intercalation capacity. Other than water and solvent molecules, the NH4+ [162] and formamide molecules [164] can also intercalate into the the interspace between the layers of bulk V2O5, which has been used to synthesize V2O5 nanosheets as described in section 4.1.3.

Figure 54. XPS spectra of bilayered V2O5 before (A) and after 10 cycles of charging with Mg2+ ions (B) in the O 1s core level region in conjunction with molecular dynamics simulation of the 150

V2O5 bilayer immersed in water (C) and in the presence of Mg ions with V in the +4 state in the presence of water (D). HRTEM image of Mg-enriched V2O5 (E) in conjunction with an HAADF image (F), in which the relative position of V and Mg atoms obtained with EDS is shown in the inset (V: red, Mg: green). The lattice model in the viewing direction [102] is also shown on the HAADF image, in which bright lines depict the V plane. Reprinted with permission from ref [418]. Copyright 2015 American Chemical Society.

Sai Gautam et al. [423] presented a detailed investigation revealing the important role of H2O during Mg2+ intercalation in nanocrystalline Xerogel-V2O5 with first-principles calculations. The analysis of the stable phases of Mg-Xerogel-V2O5 (Mgx(H2O)nV2O5) and voltages at different electrolytic conditions reveals a range of concentrations for Mg in the Xerogel and H2O in the electrolyte where there is no thermodynamic driving force for H2O to shuttle with Mg during electrochemical cycling (Figure 55). While initial Mg intercalation up to xMg = 0.25 pulls H2O into the structure for both wet and dry electrolytes, further co-intercalation of water with Mg depends more sensitively on the water content of the electrolyte. Interestingly, the presence of water in the electrolyte changes the phase behavior of the Mg-Xerogel system from that of a two-phase reaction at a single voltage (superdry) to one with a capacity over a range of voltages (wet and dry). They have established the significant impact of water (or solvent) co-intercalation on the voltages and voltage profiles obtained (Figure 55b). They demonstrated that H2O shuttling with the Mg2+ ions in wet electrolytes yields higher voltages than in dry electrolytes. Overall, H2O co-intercalation in Xerogel-V2O5 has three important technological consequences: (i) higher Mg insertion voltages, (ii) change in phase behavior from a two phase regime (superdry) to one with

151

intermediate stable Mg concentrations (wet, dry), and (iii) higher kinetic rate of Mg insertion originating from the electrostatic shielding effect of the coordinating water molecules in the cathode.

Figure 55. (a) Grand-potential phase diagram at 0 K of Mg-Xerogel V2O5 as a function of various electrolytic conditions and Mg chemical potentials is shown. Each colored region represents a single phase with the indicated Mg and water content. The dashed lines display different electrolytic regimes, with μMg = 0 corresponding to full magnesiation. (b) Average Mg insertion voltage for low (red line) and high (blue) Mg concentrations as a function of the electrolyte water content (aH2O). Equations on the curves indicate the change in H2O content in the Xerogel as Mg is inserted in each electrolytic regime. Reprinted with permission from ref [423]. Copyright 2016 American Chemical Society.

Recent experimental results from Sa et al. [449] corroborate computational studies to some extent. The electrochemical performance of the V2O5 xerogel was tested in coin cells using Mg metal as the anode, 1M Mg(TFSI)2 in diglyme as the electrolyte, and a charging rate of 20 µA cm-2 between 0 and 3.2 V. Both the capacity and voltage were low (~1 V during discharge and ~50 mA h g-1), which suggests that this particular combination of anode, cathode and electrolyte is not suitable for good cycling performance. However, a host of techniques were used to monitor 152

structural changes in the xerogel and observe co-intercalation of diglyme. Specifically, 1H NMR spectra demonstrated that as Mg coordination clusters are displaced in the cathode during cycling, diglyme reversibly fills the interplanar gap.

In summary, the intercalation process of alkali-ion in V2O5 is often complex. Water has a strong influence on the insertion of alkali-ions into V2O5. Incorporating a proper amount of crystalline water into V2O5 lattice could improve alkali-ions storage performance. These structural water molecules maintain sufficient interlayered space to allow the physical diffusion of alkali-cations. Moreover, for the intercalation of Mg2+ into V2O5, these water molecules assist in stabilizing intercalated Mg2+ through dipole interactions. The measured voltages are subjected to change if the co-intercalation of the solvent/electrolyte with the shuttled ion occurs, leading to a codependence on the solvent/electrolyte chemical potential. For the bilayered V2O5 with very large interlayer spacing, solvent molecules may also co-intercalate with the working ion and alters the structure of materials and significantly influence its electrochemical performances. The role of water and solvent co-intercalation should be considered for the structural design and synthesis of high performance V2O5 as host mateiral for alkali-ion intercalation. In addition, it has been reported that CTA+ (CTAB: N-cetyl-N,N,N,-trimethyl-ammonim bromide) [450, 451], organic cations ([(C6H5)4As]0.19+, [(C6H5)4P]0.20+, [(C4H9)4P]0.19+, [(C2H5)4N]0.58+) [452], Fe3+ [295], polypyrrole [453-455], polyvinyl pyrrolidone (PVP) [456], poly(3,4-ethylenedioxythiophene) (PEDOT) [457, 458], Polyaniline(PANI) [459], Poly(N-[5-(8-hydroxyquinoline)methyl]aniline) (PNQA) [460], ferrocene [461], poly(ethylene oxide) (PEO) [462], poly-o-methoxyaniline [463], poly(oxymethylene-oxyethylene)

(POMOE) 153

[464],

and

even

meso-(tetra-4-methylpyridinium)-porphyrin (TMPyP) [465, 466] can be intercalated into the V2O5·nH2O xerogel. Due to the pillar effect [467], the intercalation of these molecules and/or ions will change the interlayer space, valence state of vanadium, electronic and ionic conductivity, crystalline character, structural stability, and morphology of V2O5·nH2O, which could open interesting possibilities for the development of high performance vanadium oxides-based electrode materials for alkali-ion batteries. More generally, investigations of solvent co-intercalation properties in other layered materials will be useful and important in designing the next generation of rechargeable alkali-ion batteries.

7. FULL-CELL STUDIES OF V2O5 There have been numerous studies reporting V2O5-based nanostructures with improved electrochemical performance in half-cells, but only limited reports dedicated to the electrochemical

properties

when

coupled

with

anodes

in

full-cell

configurations

[129,153,255,468-474]. Full-cell assembly is very necessary to further evaluate the practical application of V2O5 for secondary batteries. Cheah et al. [129] evaluated the lithium storage properties of electrospun V2O5 nanofibers (VNF) in full-cell assembly with a spinel Li4Ti5O12 anode and cycled between 1 and 2.5 V at ambient temperature conditions. Prior to full-cell assembly, VNF is electrochemically lithiated using Swagelok fittings by discharging to 2.5 V vs Li. The full-cell Li-V2O5/Li4Ti5O12 exhibited an energy density of ~215 W h kg-1, which is comparable with the LiFePO4/Li4Ti5O12 system. Though capacity fading was noted during cycling and ~84% of the initial capacity is retained with an operating voltage of ~1.8 V, this work clearly

154

demonstrates the possibility of using LiV2O5/Li4Ti5O12 configuration for high power applications such as hybrid electric vehicles and electric vehicles in the near future. Kong et al. [153] fabricated flexible full-cell with V2O5 encapsulated into carbon nanotube (V2O5@G) as a cathode and lithiated SnOx-encapsulated carbon nanotube network (SnOx@G) as an anode. The flexible full-cell exhibited good cycling stability and excellent mechanical properties. The capacity calculated based on the entire weight of V2O5@G mass can reach ~200mA h g-1 in the first cycle; after a slight decay in the first two cycles, the capacity of flexible full-cell delivered over 90% capacity retention during the subsequent cycles. Moreover, the capacity of flexible full-cell remained nearly unchanged even after cycling five times at the bend state of 180 angles for 5 cycles. McNulty et al. [470] reported the first full inverse opal Li-ion battery, where both the cathode and anode binder-free electrodes were composed of 3D nanocrystal assemblies as inverse opal (IO) structures of intercalation-mode V2O5 IO cathodes and conversion-mode Co3O4 IO anodes. The full V2O5/Co3O4 IO cell provided a stable, high capacity Li-ion cell that retains 150 mA h g-1 over 175 cycles and excellent rate capability and capacity retention during symmetric and asymmetric (slow discharge, fast charge) conditions. The authors also demonstrated that an effective infiltration of electrolyte in nanoscale and porous materials is very critical in maintaining high capacity and its retention over long-term cycles. It should be mentioned that the full V2O5/Co3O4 IO cell was cycled in a potential window of 3.0–0.2 V, which is lower than the most commonly used Li-ion cell electrode pairing. For practical use, several V2O5/Co3O4 IO cell could be connected in series to raise the nominal voltage, depending on the power requirement. Xu et al. [471] constructed a full-cell with rGO/V2O5 as cathode and lithiated graphite as anode. The full 155

battery exhibited a high capacity of 213 mA h g-1 at a current density of 100 mA g-1 and an outstanding energy density of 106.5 W h kg-1. Zhou et al. [472] assembled a full-cell using V2O5 nanobelts as cathode and lithiated graphite as anode. When cycled in the voltage range of 1.5-3.8 V at 300 mA g-1, this V2O5/lithiated graphite full-cell exhibited a capacity of 241.4 mA h g-1 for the 1st discharge and 165.5 mA h g-1 during the 50th discharge. The above results suggest that V2O5 material could be a promising candidate as a cathode material for LIBs in the near future.

Despite the encouraging achievements in the configuration and electrochemical performance of full-cells with V2O5 as cathode material, there remain many issues to be overcome. For example, massive production of high-performance V2O5 electrode materials with economically viable cost remains a great challenge for the practical application of nanostructured V2O5 for conventional LIBs. Nanostructured V2O5 electrode materials generally have a low tap density, which is unfavorable for improving the volumetric capacity of the full-cell batteries. The long-term cycling stability of full cells with V2O5-cathodes is still far from satisfactory due to the capacity mismatch in the cathode/anode pairs and/or the poor electrode/electrolyte compatibility [129,470]. Further studies on V2O5-based cathodes should focus on the energy density and cycling performance of LIBs in practical applications instead of observed specific capacity and capacity retention in cion-type half-cells. In particular, it is of great importance to reveal the failure mechanisms for the full-cells with V2O5-based cathodes that is very different from that known for half-cell with lithium metal counter electrode. Matching appropriate anodes and selecting reasonable operating voltage windows are vital for fully utilizing the high theoretical capacity of V2O5-based cathodes and increasing the energy density of the full-cells [475]. In addition, exploration of other 156

significant components of full cells, such as electrolytes, additives, separators, and binders, will also help to improve the electrochemical performance of full-cells using V2O5-based cathode materials.

8. CONCLUSIONS AND PROSPECTS Demand for high-power and high-energy electronic devices and instruments is always increasing. Better batteries are required for better electric vehicles, grid-scale stationary storage, and our portable electronics, to name a few of the most salient applications. The development of this new generation of rechargeable batteries requires substantial improvements in energy density, rate capability, cycling stability, energy efficiency, and safety characteristics. As one of the critical limiting factors, high performance and low-cost cathode material holds the key in the field of new alkali-ion and alkaline-earth ion batteries. Among the potential cathode materials, V2O5 has advantages that include high lithium-ion storage capacity, low cost, abundant sources, and good safety properties when used as alkali-ion intercalation electrode, and therefore has been considered as one of the most attractive contenders to replace commercial LiCoO2, LiFePO4, and other cathodes in LIB applications. However, to make V2O5 cathode material practical, great challenges related to its poor cycling stabilities, intrinsically low ion diffusion coefficient, relatively low discharge voltage, and moderate electrical conductivity must be overcome. Extensive research published in open literature has demonstrated all the above shortcomings can be circumvented through the proper design of chemical composition, surface properties, crystal 157

structure and crystallinity, and nano and microstructures. The many structural, chemical, and electrochemical aspects of V2O5 summarized in this review highlight the influence of morphology and crystal structure on intercalation mechanisms.

Decreasing the particle size of V2O5 to the nanoscale regime has gone along way and is one of the most effective approaches due to the shortened transport lengths for both electrons and alkali-ions and increased surface energy and electrode/electrolyte contact area. In addition, the nanosized electrode materials sustain larger stresses before pulverization compared to the bulk counterparts, which could benefit the improvement of cycling stability of electrode materials. Compared to those low-dimensional nanostructures, the three-dimensional hierarchical nanostructures are more attractive due to their percolated ionically and electrically conductive networks, which can greatly enhance the transport kinetics, and maintain the structural stability. Moreover, the pores in the complex micron-sized architectures accommodate the large volume changes and strain generated during the intercalation/deintercalation process, which enhances the structural stability of the electrode material and thus improves the cycling stability. Numerous methods are reported for fabricating a variety of nanostructured V2O5 with peculiar electrochemical properties. However, some fabrication methods are generally time-consuming and the processes involve various sophisticated steps and expensive and/or toxic chemicals. To bridge fundamental research studies and commercial interests, developing facile, efficient, reliable, and environmentally friendly methods with economically viable cost for massive production of novel V2O5 nanostructures is still urgently demanded.

158

Nanotechnology combined with heterogeneous hybridization could further improve the electrochemical performance of V2O5 due to their synergic properties. This arises from simultaneously integrating multiple nanocomponents, each tailored to address a different demand, for example high energy density, high conductivity, and excellent mechanical stability. Implementation of these strategies addresses the problems that are related to the ionic and electronic conductivity as well as structural stability of V2O5. Theoretically, it is possible to rationally tune the electrical conductivity and/or ion diffusion coefficient in V2O5 as well as its structural stability by simultaneously doping with different cations due to the synergy effect. Doping of V2O5 by one species of metal cation has been well investigated; while there have been few investigations on the doping of V2O5 by two or three kinds of cations. The presences of structural defects (O vacancy and/or low valence V ions) and/or disorders on both atomic and nanometer length strongly influence the electrochemical performance of V2O5, such as electrochemical capacity, reactivity, electrical conductivity, and alkali-ion transport. In particular, these structural defects (vacancies and/or local disorder) may serve as possible nucleation centers for phase transformation during the intercalation/deintercalation process and therefore enhances cycling performance of V2O5. Therefore, defect-mediated higher storage capacities and better charge and mass transport may offer some routes towards better performance in cathode materials. Whereas how to accurately tune the defect sites and stabilize the structure for long-term cycle needs to be addressed. As a typical layered intercalation compound, the activation energies for alkali-ion diffusion in V2O5 are anisotropic. Therefore, design and synthesis of V2O5 materials with well-defined crystal plane staking offers the possibility of further improving the 159

electrochemical performance of V2O5 and will be the subject of further research in the years to come.

Na-ion batteries are attractive because sodium resources are seemingly inexhaustible as well as ubiquitous, and therefore cost considerably less; flexibility in choice of current collector also cuts the cell cost. However, the accommodation of sodium in traditional host materials is difficult because the ionic radius and reduction potential of sodium are much larger than that of lithium. Therefore, the reversible sodiation process induces large distortions in the host lattice that ultimately pulverize the electrode and lead to cell failure. To date this has been combatted with the inclusion of defects; techniques of note include cationic disordering, amorphization, doping, partial cation reduction, and manipulation of intrinsic defects. Defects can reduce the stress and the electrostatic repulsion between adjacent oxygen layers, which can directly alter the migration energy and diffusion barriers the alkali-ion must overcome during intercalation. Vanadium pentoxide and other vanadate derivatives have also been investigated as a prospective cathode material for magnesium-ion batteries. Its high theoretical capacity (590 mA h g-1 for insertion of two Mg-ion ions) and fairly high voltage (≥ 2.7 V vs. Mg/Mg2+) are very attractive, but it has been hampered by poor cyclic stability and slow transport kinetics. Most investigations have focused on overcoming these issues by tailoring the nanostructure of the material and screening the charge of Mg-ions during insertion. Moreover, water and solvent co-intercalation play a key role in the Mg2+ intercalation process, which needs to be considered for improving the cycling stability of vanadium pentoxide materials. Understanding how water co-intercalation improves intercalation may provide clues for searching other co-intercalants with a less detrimental effect on the metallic 160

anodes. Although the electrochemical performance of V2O5 in MIBs has been modest to date, massive interest and investment in recent years has brought more sophisticated spectroscopic techniques online, such as MAS NMR. These techniques allow researchers to probe the local bonding environment and electronic structure of materials in ways that have not been previously possible, and ask more valuable and insightful questions. In addition, computational simulations are also powerful tools and can be carried out along with experiments to rationalize the observed enhancement of performances.

So far research on V2O5 has been largely limited to the manipulation of crystallinity, nano- and microstructures, doping of cations and introducing oxygen vacancies to achieve a storage capacity close to theoretical limit with good rate performance and cycling stability. However, the research is scattered. Since some work focuses on one lithium-ion intercalation at a narrow voltage window and other work pushes the voltage window to allow the third lithium ion intercalation, the detailed comparison of all the work is difficult. However, the general trend and fundamental understanding are clear: a high lithium ion storage capacity close to theoretical limit is achievable with good rate performance and cycling stability. The lessons learned from the study of vanadium oxide and the fundamental understandings achieved in lithium-ion intercalation in vanadium oxide are valuable assets for the study of other electrode materials and for the design and development of new electrode materials.

Figure 56 summarized some important aspects need to be addressed in the years to come for V2O5 cathode material. In spite of the extensive research done so far, there are many issues

161

requiring more detailed work. (1) The stability of nanostructures during extended cycling is one such example. Will the surface oxygen vacancies and doping cations stay homogeneously dispersed after many cycles? (2) Is there a feasible and general way to enhance the discharge voltage? (3) The interlayer distance in V2O5 varies when guest species are inserted; the interlayer distance directly impacts the storage capacity in addition to the transport properties. (4) Co-intercalation seems to be an interesting approach and deserves more work to explore the full potential of V2O5 as a cathode, particularly for Na-, Mg-, and Ca-ion batteries. Crystalline water enlarges the interlayer distance, which favors the intercalation of lithium ions and appears to promote the intercalation of sodium and magnesium ions as well, but will the water molecules be stable during repeated insertion and extraction of guest ions? (5) Will there be other molecules or clusters more effective at enlarging the interlayer distance? Will the co-intercalation of alkaline ions together with inactive small molecules be an effective way to shield the electrostatic interaction between divalent magnesium ions with vanadium oxide crystal lattice? We present these and other open questions as challenges for the battery community to grapple with as we seek to address the critical issue of energy storage, one of the most urgent scientific challenges of our time.

162

Vanadium pentoxide

Developing facile, reliable, low-cost, and friendly methods for massive production of high performance V2O5 nanostructures. Fabricating flexible free-standing electrode for smart wearable electrics.

Investigations on performance of V2O5 in full Li-ion cells should be encouraged.  Li-ion batteries

 Na-ion batteries  Mg-ion batteries  Other metal (Zn2+, Ca2+, and Al3+) batteries

Discover new concepts in the fundamentals of electrochemistry using on-time spectroscopic techniques and computational simulations.

Structure modification in terms of pre-intercalation, tuning the interlayer, structure defect, and crystal plane staking. Intercalation and degradation mechanism of multi-valent ions (Mg2+, Zn2+, Ca2+, Al3+) in of V2O5.

Figure 56. Some important aspects need to be addressed in the years to come for V2O5 cathode material for lithium-ion batteries and beyond.

ACKNOWLEDGEMENTS This work was supported in part by the National Science Foundation of the U.S. (CMMI-1030048 and DMR-1505902), University of Washington TGIF, National Natural Science Foundation of China (51664012, 5120406, and 51464009), Guangxi Natural Science Foundation of China (2015GXNSFGA139006 and 2014GXNSFBA118238). EU and RCM would like to acknowledge support from the State of Washington through the University of Washington Clean Energy Institute.

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nanocomposite

materials:

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