Crafting Inorganic Materials for Use in Energy Capture and Storage

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Invited Feature Article

Crafting Inorganic Materials for Use in Energy Capture and Storage Yukti Arora, Charu Seth, and Deepa Khushalani Langmuir, Just Accepted Manuscript • DOI: 10.1021/acs.langmuir.8b02953 • Publication Date (Web): 26 Oct 2018 Downloaded from http://pubs.acs.org on October 28, 2018

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Crafting Inorganic Materials for Use in Energy Capture and Storage Yukti Arora#, Charu Seth# and Deepa Khushalani* Materials Chemistry Research Group, Department of Chemical Sciences, Tata Institute of Fundamental Research, Mumbai, India 400005. # *

Both authors contributed equally to this work [email protected]

ABSTRACT

Harnessing solar energy effectively by judicious use of photoactive inorganic/hybrid structures has become a pivotal requirement in the pursuit of environmentally benign technologies. The synthesis of new inorganic materials whose stoichiometry, structure and activity can be tuned, while maintaining a high level of architectural homogeneity, and to successfully evaluate each of the materials as viable components in specifically energy capture and storage based applications is being presented here. Two of our current projects are detailed involving (i) new 1D structured hybrid perovskite that is a more temporally and thermally stable analogue to the oft-cited methyl ammonium lead iodide and (ii) a new electroactive material that can function not only as a conventional electrode in a battery but also, because of the material’s inherent photoactivity, it can be employed in solar battery, and hence the concept that energy capture and energy storage can be coupled in a single device is also being detailed.

ACS Paragon Plus Environment

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INTRODUCTION

It is well accepted now that concurrent to the increase in population, the rapid economic growth of developing nations has forced us to recognize the challenge of harnessing renewable sources of energy supply. This needs to be done consistently over a prolonged period and importantly in addition to harnessing, it has also become important to devise schemes to store this energy efficiently. This task has been made even more relevant in order to mitigate the pervasive issues of air and water pollution, of which the non-renewable sources of energy are the largest culprits. Towards this aim, there are several research avenues that are being pursued under the large umbrella theme of “energy harvesting and storage”, of which solar cells and batteries/supercapacitors are some of the more popular devices being evaluated.1 All these devices are multicomponent entities and consist of several different ‘layers of materials’ that intimately interface with each other so as to maximize charge transport and minimize losses due to recombination or other forms of internal resistance.2 The chemical composition of each layer, the manner in which it is deposited, the nature of interaction with the neighboring electrodes etc. are just some of the many aspects of a device that have to be carefully considered in order to achieve maximum theoretical efficiency. As such, in the search for novel materials for use in energy based applications, a number of challenges have to be overcome depending on the exact device to be fabricated. Solid phase materials have thus far been preferred due to their mechanical, thermal and temporal stability along with ease of fabrication. In using these solid state structures, predominantly semiconductors are exploited due to the ease of modification of their optical and electrical properties by altering the stoichiometry or doping.3 In addition, aspects such as crystallinity, work function, charge carrier density, charge carrier mobility/lifetime, electronic structure (direct/indirect band), and band alignments for a heterojunction system are crucial in


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defining their eventual usage. For any researcher, the ability to manipulate the aforementioned parameters is pivotal and this is commonly done by carefully choosing the stoichiometry and moreover the overall architecture of the deposited phase is also given substantial consideration. Processing the envisaged materials so that there is ease of deposition in order to prepare devices in a facile manner is pivotal. Crystalline thin film formation, oriented nanowires/tubes or 2D and 3D spatially ordered nanostructures are some of the more relevant morphologies being commonly targeted.

Work in our lab, over the past few years, has revolved around the concept of synthesis of a range of inorganic materials which are then judiciously exploited in applications involving energy harvesting and storage. The aim has been to target a variety of homogenous nanostructured morphologies of semiconductors and ideally these are grown directly on a device based substrate such that interfacial conducive interactions are maximized. A diverse set of synthetic protocols are exploited to try and achieve such a structure so that there is purity in chemical composition, homogeneity in the nanostructure size and shape obtained, and also ideally a contiguous 2D growth on the substrate is achieved. Detailed below are two currently ongoing research areas in our lab. The first work deals with a new hybrid perovskite that has been recently synthesized exhibiting a 1D structure of the lead halide network, while the latter example showcases the versatility of an old oft-cited semiconductor, namely BiVO4, and its foray into a new area, that of energy storage. A. Energy Capture Among all renewable energy fuel sources, solar energy has emerged as one of the most practical alternatives to conventional non-renewable sources. This is mainly due to the fact that solar energy reaching the earth from sun is substantial, i.e. 3 × 1024 J per year or ca. is 104 times more than what


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the entire human population currently consumes annually.4 In terms of effectively exploiting solar energy, one option is to approach this problem as converting solar energy to an electrical impulse using devices referred to as solar cells.5 Solar cells are devices consisting of multiple chemical layers and the main functioning is based on the photovoltaic effect which is closely related to the photoelectric effect. The principle driving all the solar cells stems from absorption of a photon of energy to obtain an excited electron and a hole (charge carriers). The sole aim of the solar cell is to ensure that these excited species do not recombine but instead are separated effectively without loss of energy and are viably extracted from the device (via opposite electrodes) generating electricity. Over the past few decades, there has been an enormous research focused on developing new light absorbers that can effectively form the active layer in a solar cell. One of the latest contenders (that satisfies the many parameters needed for a viable light absorber and which was introduced over nine years ago) is a member of a family commonly referred to as hybrid perovskites.6 These are compounds with ABX3 stoichiometry and the most researched stoichiometry is the Methyl ammonium lead iodide (MAPI): CH3NH3PbI3. They have gained popularity because the solar cells that have been fabricated using this light absorbing material are found to routinely give efficiencies upwards of almost 20%. The current understanding of these materials is that they have very high absorption coefficient, the excited electron is purported to have high diffusion lengths ca. 1 m, as such the light absorbing layer in the device can be easily experimentally achievable (ca. 200300 nm), exciton lifetimes are very low and generation of charge carriers is easy. Moreover, one of the ways that the band-gap can be tuned is by the amount of halide concentration and can be varied from 1.7-2.2 eV.7–9


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From a structural point of view, these materials exhibit a correspondence between the packing of lattice and the Goldschmidt tolerance factor given as 𝑡=

𝑟𝐴 + 𝑟𝑋 2(𝑟𝐵 + 𝑟𝑋)

Where rA, rB, and rX are the ionic radii of cation A, cation B, and anion X respectively.10 A value of t ~ 1 implies a high symmetry, cubic structure of perovskite where A occupies the body centre, B the corners, and X the edge centers. Deviation of the value of t from unity on the lower side may result in distorted perovskite structures, and on the higher side may result in hexagonal or tetragonal structures, Table 1.

Table 1. List of a few organic lead iodides with the predicted structures based on Goldschmidt tolerance factor and the actual structures at room temperature (RT).

Compound

t

Predicted structure

Actual structure at RT

CH3NH3PbI3 (MAPI)

0.91

Cubic perovskite

Tetragonal

CH5N2PbI3 (FAPI)

0.99

Cubic perovskite

Hexagonal (δ phase)11,12

CsPbI3 (CsPI)

0.81

Distorted perovskite

Orthorhombic (γ phase)11,13,14

C2H5NH3PbI3 (EAPI)

1.05

Pseudocubic perovskite/Hexagonal

Hexagonal15

CH6NPbI3 (GuaPI)

1.04

Hexagonal

C3N2H5PbI3 (ImPI)

0.99

Cubic perovskite

Hexagonal (Orthorhombic when it forms 2D perovskite)16,17 Hexagonal (our work)18

The well-studied MAPI exhibits a tolerance factor value of 0.91, formamidinium lead iodide (FAPI) 0.99 and all-inorganic cesium lead iodide (CsPI) 0.81.15 The relatively less studied ethyl


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ammonium lead iodide (EAPI) shows a tolerance factor slightly greater than 1, and adopts a hexagonal structure.15 Guanidinium lead iodide (GuaPI), however forms a hexagonal 1D perovskite structure as predicted by the Goldschmidt tolerance factor, however is not always phase pure and consists of 2D perovskite stoichiometry Gua2PbI4 along with the 1D perovskite GuaPbI3. Moreover, it is not a photoactive material.16,17,19 While the Goldschmidt tolerance factor relation sheds light on the kind of packing a lattice adopts depending on the sizes of ions involved, thereby acting as a guiding principle in selecting stoichiometries for photovoltaic and optoelectronic applications, however this semi-empirical relation is rendered inconsequential when it comes to temperature dependent phase transitions. In hindsight, it is purported to employ this mathematical relation to comment on the packing efficiency of the lattice rather than specifically predicting whether a lattice will be perovskite or not. The low value of the above mentioned tolerance factor for MAPI insinuates the lattice instability of the material which has been vastly discussed in literature.20–24 MAPI not only undergoes temperature dependent phase transitions,25,26 but also irreversible chemical transformations under ambient conditions (light, moisture and oxygen) to form byproducts.20–22 Thus, the aspect that has frustrated researchers is the lack of temporal and thermal stability of this material, despite all the above mentioned impressive advantages. In order to address this issue, MAPI’s structure has been altered using various halides (iodide has been replaced with Cl or Br or SCN or mixed halides have been used) and the ‘B’ metal cation has been varied (Lead has been replaced with a variety of cations such as Sn, Ge etc.). However, fewer are the number of reports where the organic cation ‘A’ has been modified. Literature has surprisingly focused on only two examples: either formamidinium – FAPI analogue has been studied or predominantly Cesium based perovskites – CsPI have been evaluated.


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A subsequent altered direction therefore that has been undertaken is to form mixed cation based structures (vide infra). In principle, hybrid perovskite consists of 3-dimensional arrays of inorganic PbI64- anions surrounded by organic ammonium counter ions. In the conventional MAPI, the organic cation is a non-spherical methyl ammonium species that has a dipole moment and a strong polarizing capability. It is important to understand that the main 3D structure of this cubic perovskite with Pb, I and methyl ammonium components is a symmetric structure with respect to a, b, and c axes, Figure 1. It can be considered to contain infinite number of PbI64- octahedral chains in the c direction (infinite in the a and b directions).

Figure 1. Cubic perovskite structure of MAPI generated from the cif provided in ref. 25.25 (a) Cubic unit cell with methyl ammonium cation in the body center, lead on the corners and iodine at the edge centers, (b) a-b projection of the extended lattice showing PbI64- octahedra sharing corners (iodine), viewed along c axis, (c) b-c projection of the extended lattice showing PbI64octahedra sharing corners (iodine), viewed along a axis. The structures shown in (b) and (c) are identical; the corner shared PbI64- framework extends symmetrically and indefinitely in the three directions.

A closely related family of solids exist that are based on the same ionic species (lead halide octahedra and a small organic ammonium cation) which, however, are now arranged in 2dimensional layers. The layers consist of integer (n) increments of lead halide octahedra and they


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are delineated by a larger cation. Shown below in Figure 2 is a conventional representation of the family of 2D layered family of perovskites. Here, n-butyl ammonium is the long alkyl chain ammonium cation which yields compounds with different layer (n) values (e.g. n = 1, n = 2, n = 3...n = ∞). In this way, n = ∞ is a cubic 3D perovskite such as MAPI, whereas the other n values describe 2D perovskite structures. The main reason for this arrangement is that now the structure consists of mixed cations – some of the CH3NH3+ cations are now partially replaced by other preferably larger organic cations. The initial aim of 2D perovskites being introduced in the literature was to enhance the moisture stability of the conventional 3D hybrid perovskites through the long ammonium cations that are sandwiched between the inorganic layers. Subsequently, it has interestingly been observed that the stoichiometric ratios of the various components of 2D perovskites affect not only the thermal stability but also concomitantly the n value dramatically influences the optical and electronic properties. As such, the dimensionality of perovskites has been found to have a direct impact on the conductivity.27

Figure 2. A scheme showing the evolution of the perovskite dimensionality from n = 1, for a 2D perovskite structure, to n = ∞ for a 3D perovskite structure.


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It is interesting to note that Lead iodide (PbI2), the inorganic component of hybrid perovskites, as such is a layered material consisting of two-dimensional layers of PbI64- octahedra that are interconnected in a plane, Supporting Information Figure S1.28 Small molecules like ammonia, methyl amine, pyridine etc. can intercalate between the layers of PbI2, thereby influencing the optical and electronic properties.29,30 Nevertheless, the overall structure is still layered. However, addition of small ions like cesium, methyl ammonium, or formamidinium at appropriate reaction conditions results in a three dimensional cage-like network of PbI64- octahedra and the cavities within the cage are occupied by the small cation, thereby forming a cubic perovskite structure yielding the structure of CsPI or MAPI as discussed above or FAPI, respectively. Thus, positively charged cations cause PbI2 to lose its 2D lattice structure and dramatically change the optical and electronic properties. For instance, intercalation of methyl amine in PbI2 causes the colour to change from bright yellow (pristine PbI2) through pale yellow to white, depending on the concentration of the intercalated amine,30 whereas the methyl ammonium cation on reacting with PbI2 results in black coloured methyl ammonium lead iodide perovskite. The Raman signatures of methyl amine-intercalated PbI2 are entirely different from methyl ammonium lead iodide (MAPI); the former shows low wavenumber phonon modes corresponding to the lattice of PbI2, whereas the latter shows features of the organic cation motion.30,31 Similarly, addition of long ammonium cations (RNH3+ where R: n-butyl ammonium, n-hexyl ammonium, phenylethyl ammonium etc.) to PbI2 (at appropriate reaction conditions) results in a completely new structure and stoichiometry: (RNH3)2PbI4 (n = 1 2D perovskite being one of the extremes of the known 2D hybrid perovskite, vide supra).32,33 Although, in terms of the arrangement of PbI64- octahedra, the latter structure is still 2-dimensional like PbI2 is, however, the intercalation of RNH3+ alters the electronic structure and optical properties of the material dramatically.


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It has been shown so far only in the case of 3D and 2D hybrid perovskites that the electronic structure near the band edges is governed predominantly by the inorganic framework. For MAPI, the valence band maximum is composed of I 5p, Pb 6s, and I 5s states (in decreasing order of their contributions), whereas the conduction band minimum is dominated by Pb 6p states and I 5p.34 The frontier orbital/band states are exclusive of the organic component contributions.34,35 The conductivity of the perovskite material arises from the delocalization of charge in the inorganic framework. This means that structures such as dihydrated MAPI, (CH3NH3)4PbI6·2H2O, which consist of interestingly solitary PbI64- octahedra (that are completely un-connected) would not be conducting at all36–38 which is indeed the case.39 This dihydrated MAPI can, in principle, be called a 0D perovskite structure since the octahedra are distinct and isolated. 3D and 0D perovskites are therefore the two extremes of the octahedral connectivity; 2D and 1D should, in principle, lie between these two in terms of the coupling of the octahedra. This connectivity would then manifest itself ultimately in the conductivity of the material. Qian et al. have demonstrated, in a systematic theoretical study, a series of 0D, 1D, 2D and 3D hybrid perovskites and their optical and electronic properties specific to the direction of extension of inorganic framework. They have interestingly found that the conductivity is almost non-existent in the direction parallel to the organic component (i.e., orthogonal to the inorganic framework).40 The photon to current conversion efficiencies were also simulated for the chosen 0D, 1D, 2D and 3D model systems and the values are impressively similar for 1D and 2D systems. However, the experimental evidences of directional growth (parallel to the inorganic framework) enhancing the charge transport are available only for 2D perovskites and not for 1D.41–44 Furthermore, recently Xiao et al. have introduced the concept of electronic dimensionality (along with the aforementioned structural dimensionality) and this refers to the connectivity of atomic


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orbitals of the valence band maxima and conduction band minima.39 According to the model systems investigated by the authors, a material could have the two types of dimensionality going hand in hand (i.e. both the structural and electronic aspects are 0, 1, 2 or 3D) in which case the prediction of opto-electronic properties can be made easily while in other cases this may not necessarily be so. An important aspect that has been further highlighted is that even with the same structural dimensionality, the connectivity of the octahedra determines the band gap of the material. It has been suggested that band gap increases when the connectivity goes from corner sharing to edge sharing to face sharing.45 As such this has opened a new area and the reader is referred to a recent review by Shi et al. for further details.46 As a tangential approach to the formation of stable hybrid perovskites, we have recently synthesized directly 1-dimensional version of a hybrid material with lead iodide that shows remarkably improved stability with respect to methyl ammonium lead iodide.18 Since the organic cations have a tremendous role to play in determining the moisture resistance as well as the band gap of the perovskite material, we attempted the complete substitution of methyl ammonium cation with a symmetric, bulkier and aromatic cation, imidazolium. The resultant crystal structure was neither a 3D cubic perovskite nor a 2D perovskite structure, despite the stoichiometry of thus formed imidazolium lead iodide (as determined from the single crystal structure data) being C3N2H5PbI3 (or ABX3) which is the same as that of a 3D cubic perovskite. The stoichiometry of the hybrid perovskite material is determined by the charge on the organic cation connectivity of lead halide octahedra or the number of halide ions shared between adjacent octahedra. ImPI, for example, has face-shared octahedra and the stoichiometry ABX3. On the other hand, a 1D perovskite, C4N2H14PbBr4, reported by Z Yuan et al. has edge shared octahedra surrounded by divalent A cations, and the resulting stoichiometry of ABX4.47


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Since methyl ammonium cation is aliphatic, non-spherical and asymmetric with the positive charge localized only on the nitrogen atom and not on the entire moiety, it is known to wobble in the symmetric cubic lattice to maintain the charge neutrality,48,49 thereby making the structure fluxional and unstable. With this rationale, the replacement of methyl ammonium by imidazolium cation was pursued. Imidazolium lead iodide (ImPI) adopts a hexagonal unit cell and consists of one-dimensional chains of PbI64- octahedra running parallel to the c axis, Figure 3. Because of the unidirectional extension of the inorganic part (PbI64- octahedral framework) in ImPI, it can be considered to be a 1D perovskite. The inorganic octahedral chains possess negative charge which is neutralized by the imidazolium cations flanking on the three sides. The overall unit cell is hexagonal, as stated above with the following cell parameters: a = b = 15.62 Å, c = 7.99 Å, α = β = 90°, γ = 120° and P63/m space group, CCDC number:1822500.

Figure 3. The crystal structure of Imidazolium lead iodide (ImPI). (a) Hexagonal unit cell of ImPI viewed along the c-axis, Adapted with permission from Ref. 18. Copyright 2018, License number 4454770625727, John Wiley and Sons (b) The projection of extended lattice of ImPI viewed along the direction perpendicular to c-axis, showing chains of PbI64- octahedra running parallel to c-axis.


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The Goldschmidt tolerance factor value for ImPI, surprisingly is 0.99, thereby predicting a cubic perovskite structure,15 however, as purported above, this value rather be used as a qualitative measure to estimate the packing efficiency. From the high value of the tolerance factor it can be inferred that the lattice of ImPI is well packed as compared to the MAPI, and a quantitative measure to substantiate the compactness of lattice is through the analysis of atom-atom distances in the crystal structure. Figure 4 shows the planar projections of ImPI and MAPI unit cells with the atoms of interest labelled. The minimal variation in Pb-I bond lengths of ImPI tabulated in Figure 4d indicates higher symmetry of PbI64- octahedra in it as compared to those in MAPI. Moreover, due to three iodine ions being shared by two octahedra in ImPI in contrast with one shared iodine in MAPI illustrates better lattice stability of the inorganic network in ImPI since it is less prone to distortions such as octahedral tilting. The least distance between the sites bearing positive and negative charges, i.e., between nitrogen and iodine in both the materials as shown in Figure 4d, is lower in ImPI than MAPI, implying a tighter packing in ImPI.


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Figure 4. (a) Planar projection of ImPI unit cell, viewed along c-axis, (b) MAPI tetragonal unit cell viewed along c-axis, (c) MAPI tetragonal unit cell viewed along b axis, (c) tabulated atomatom distances in ImPI and MAPI. The ImPI and MAPI crystal structures were deposited by the authors in the CCDC database, the CCDC numbers being 1822500 and 1822501, respectively. Adapted with permission from Ref. 18. Copyright 2018, License number 4454770625727, John Wiley and Sons. How the robust lattice structure of ImPI in comparison with MAPI manifests itself in the phase stability is evaluated by the low temperature and high temperature XRD of both the materials. ImPI stays in the same hexagonal phase throughout the low temperature range (3 K to room temperature), as shown in the X-ray diffractograms in Figure 5. The only changes it undergoes is the contraction of lattice as the temperature is decreased, which is characterized by the lattice plane positions shifting to higher 2θ values or the reduction in d spacing. From RT to 3 K, the unit cell contracts by about 5.6 %. MAPI, on the other hand undergoes a phase transition at 160 K from the tetragonal phase to orthorhombic phase, as is evident from change in the X-ray diffractograms shown in Figure 5b, along with the tetragonal lattice contraction in the range of RT to 160 K and


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orthorhombic lattice contraction in the range of 160 K to 3 K. In the high temperature range, going from RT to 373 K, MAPI undergoes a phase transition from tetragonal to cubic perovskite phase at 333 K, whereas this threshold is higher for ImPI which goes from hexagonal phase to a higher symmetry phase at 373 K, which again proves the better lattice integrity of ImPI over MAPI, Figure 6.

Figure 5. (a), (b) Low temperature X-ray diffractograms of ImPI and MAPI, (c) Increase in 2θ peak position or decrease in the inter-planar spacing of (001) plane of ImPI on reducing the temperature, (d) Comparison of the X-ray diffractograms of MAPI at 180 K (tetragonal phase) and 160 K (orthorhombic phase). Adapted with permission from Ref. 18. Copyright 2018, License number 4454770625727, John Wiley and Sons.


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Figure 6. (a), (b) High temperature X-ray diffractograms of ImPI and MAPI. (c) Comparison of X-ray diffractogram of ImPI at 373 K and 353 K showing a clear phase change as the multiplet peaks turn to singlet on raising the temperature, (d) Comparison of X-ray diffractogram of MAPI at 333 K and 313 K showing an apparent phase transition from tetragonal to cubic as the multiplet peaks turn to singlet on raising the temperature. Adapted with permission from Ref. 18. Copyright 2018, License number 4454770625727, John Wiley and Sons.

After establishing the superior lattice integrity of ImPI over the well-known contender MAPI, we also have demonstrated the stability of ImPI under ambient conditions. The reader is referred to ref. 18 for full film characterization details.18 In all the comparisons made henceforth, MAPI has been used as the reference material since the structural instability is its Achilles’ heel which is well-highlighted in the literature (vide supra for references). The ImPI thin films and powder are


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much more stable than the MAPI counterparts (thin film as well as the crystal powder) as can be observed from Supporting Information Figure S2. The thin films of both the materials were prepared using similar optimized protocols under identical conditions. The intensities of (4 -2 1) plane of ImPI and (3 1 0) plane of MAPI have been plotted against the ageing time of the film samples under ambient conditions (unsealed sample kept at room temperature with 65 % humidity). The decay in the intensity of X-ray diffracted from the above mentioned planes signifies a decrease in the crystallinity of the sample with time which has been used a measure of lattice instability or the degradation of the material. The slopes of the profiles of diminishing intensity over time for ImPI and MAPI, assuming a linear decay, were computed and ImPI shows an enhanced stability which is three times more than that of MAPI. The X-ray diffractograms of the ImPI and MAPI powdered crystals were also recorded after every five-day interval of exposure to the same ambient conditions as for the films. The X-ray diffractogram in Supporting Information Figure S2a shows that ImPI crystals stay stable for the entire 35-day long duration of the experiment whereas MAPI starts showing extraneous lattice plane peaks 25th day onwards, apart from the native tetragonal phase. ImPI does not show any additional features in the diffractogram as the crystals age, while the additional features in the case of MAPI are a result of moisture adducts: CH3NH3PbI3.H2O and CH3NH3PbI3.2H2O. From the aforementioned data, it can therefore be summarized that the thermal and temporal stability can also now be accessed by forming the lower dimensional states of the lead halide structure. This has been done now not by slight modification of the organic cation, as has been the case for 2D hybrid perovskites, but instead by the complete replacement of the organic moiety. By the judicious choice of complete replacement of the organic cation we have been able to synthesize 1D version of the hybrid perovskites. Measurement of the electrical and optical properties of ImPI


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have been embarked upon and the photovoltaic performance of this material has been found to be marginally inferior to that of MAPI for devices prepared under fully ambient conditions. Moreover, a broad, red-shifted, near-white light luminescence of ImPI has been intriguingly observed and this is currently being explored for mechanistic detials.18 In terms of future directions, work is also currently underway to exploit the orientation dependent properties of this 1D perovskite and how this influences their opto-electronic properties. For conventional 2D hybrid perovskites, it is known that orientation of the unit cell with respect to the underlying substrate governs directly the electrical conductivity of the material and the eventual IV (current-voltage curves) and EQE (external quantum efficiency) characteristics of the ensuing solar cell. As such, these aspects are being further pursued for the 1D ImPI phase. B. Energy Storage Despite the effort being devoted to energy harvesting using renewable sources, the major hurdle that still exists in preventing the large scale usage of these novel devices is that those sources of energy are intermittent. A viable large scale usable form of an energy storage device is still to be achieved. Although batteries and super-capacitors have been referred to for some while now, there is still a void in terms of development of electrical energy storage devices which are able to meet the requirements of superior energy and power density. For this, the development of new electrode materials is being actively pursued. Batteries are by far the most commonly employed devices for storing electrical energy. The energy density of a rechargeable battery is mainly determined by the specific capacity and operating voltages of the electrode materials. As such, the chemical composition of the electrodes has been the main focus of research in recent years. The structure of a conventional battery consists of an


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anode, cathode, electrolyte and a separator, where cathode is typically Li-containing metal oxide with a layered structure (such as LiCoO2) or tunnel structure (such as LiMnO2) so as to facilitate the intercalation storage mechanism whereas anode materials are generally either insertion-type (e.g. graphite), or conversion-type (e.g. cobalt oxide) or alloying-type (e.g. Sn).50 Conceptually, batteries are inherently simple in terms of the mechanism of operation involved, however it is intriguing to note that their development has progressed much more slowly compared to other types of advanced functional devices.51 The rate of progress has been slow predominantly due to lack of suitable electrode materials, electrolytes and adequate understanding of the interfacial interactions involved.52 Present lithium ion battery technology works on intercalation based materials which have a relatively smaller number of charge storage sites, leading to limited energy densities, they suffer from slow rate of charging/discharging, ageing (poor cycling stability), and transportation is also restricted due to flammable electrolytes being employed in these devices.52 As such, there is a thrust to move away from lithium altogether, and systems based on Na+, Mg2+ and Al3+ are being considered. The monovalency of Na+ is advantageous as a large portion of knowledge accumulated for LIBs (lithium ion batteries) is being translated to these SIBs (sodium ion batteries) but unfortunately the larger size of this ion leads to significant differences in the intercalation mechanism and also the low melting point of sodium leads to safety concerns. The other two multivalent ions are also being exploited however their work is being impeded by intrinsic limitation of intercalation chemistries of Mg2+ and Al3+.53 In the search for new materials, a few aspects need to be considered when specifically evaluating stoichiometries for their potential to function as new electrode materials in batteries. The composition needs to consist of redox active ion(s) with potential that falls under a sufficiently large range, must be a good conductor, should ideally have an extended structure, and be inert in


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a variety of electrolytes. Moreover, the composition should ideally consist of earth abundant materials and be synthetically facile to achieve. Towards this objective, over the last few years, work in our lab has centered around the use of inorganic vanadates for energy storage. Specifically, Bismuth Vanadate, BiVO4, is an interesting transition metal oxide and thus far it has been predominantly evaluated for ferroelasticity, gas sensing, and more recently as a photoanode for solar water splitting owing to its small band gap (ca. 2.3 eV).54 For most of the electrochemical or photoelectrochemical applications of BiVO4, the operative potential window range is from 0.0 V to +1.2 V. However, for its application in the field of energy storage we have exploited BiVO4 in the potential window 0.0 V to -1.2 V, where its cyclic voltammogram (CV) shows quasi reversible faradaic behavior with redox peaks. Figure 7 shows some of the morphologies that have been successfully achieved with BiVO4 synthesis in our lab.55,56 Nanostructures consisting of 10-20 nm diameter sized particles that agglomerate to form large porous oval clusters with diameter in the range of 200-300 nm can be seen in Figure 7a, and Figure 7b showcases high aspect ratio rod-like structures which were obtained hydrothermally. More interestingly, aligned nanowires could also be grown directly on conducting substrates as shown in Figures 7c and 7d. SEM cross-sectional view indicates perpendicular growth of these 1D structures on FTO with diameter in the range of 200-300 nm and length in the range of 5-7 µm. TEM analysis along with selected area electron diffraction (SAED) confirm the crystalline nature of the nanowires. The chemical structure was confirmed with XRD and it is a monoclinic scheelite structure (JCPDS - 04-010-5713, Figure 8a) in which each V ion is coordinated to four O atoms in a tetrahedral site and each Bi ion is coordinated to eight O atoms each from a different VO4 tetrahedral units.57 Therefore, basic structural unit consists of VO4 tetrahedron and BiO8 dodecahedron. There are two types of V-O bonds viz. 1.692


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and 1.767 Å and there exist four types of Bi-O bonds ranging from 2.354 to 2.628 Å. This information is important as the polyhedra representation of BiVO4 structure provides an interesting insight into the material where it can be considered to consist of a layered structure of alternating layers if BiO8 dodecahedron and VO4 tetrahedron, Figure 8b. As reported by Zhou et al. BiVO4 has two different types of polyhedra. There exists a strong hybridization between V and O atoms resulting in covalent bond formation, V-O dipole, in VO4 tetrahedra, whereas there is a relatively weak hybridization between Bi and O atoms, resulting in ionic bond formation, Bi-O dipole, in the BiO8 dodecahedra.58

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Figure 7. SEM image of (a) BiVO4 oval nanostructures, (b) BiVO4 nanorods, inset shows the SAED pattern indicating high crystallinity, (c) and (d) show the top and cross-sectional view, respectively, of aligned BiVO4 nanowires on a conducting substrate, (e) TEM image of BiVO4 nanowire, inset shows the SAED pattern showing high crystallinity, and (f) HRTEM image of the tip of a nanowire showing defined lattice fringes.


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Figure 8. (a) XRD pattern of monoclinic phase of BiVO4, (b) polyhedra structure of BiVO4, where blue colored polyhedra are BiO8 units whereas grey colored polyhedra are VO4 units, (c) and (d) show CV and CD profiles of BiVO4 acquired at 20 mVs-1 and 7 Ag-1, respectively.

As a result of this structure, it is has become apparent from our work that the reversible redox activity of BiVO4 is feasible. As such, the cyclic voltammetry data presented in Figure 8c is for 20 mVs-1 scan rate under the potential window 0.0 V to -1.2 V. The conventional CV curve of BiVO4 shows two anodic (oxidation) peaks and one cathodic (reduction) peak, this electrochemical energy storage could be attributed to quasi reversible redox process Bi3+ ↔ Bi0. Two anodic peaks at 0.50 V and -0.38 V result from the two step oxidation of Bi0 to Bi3+ via Bi+ and single cathodic peak at -0.75 V can be assigned to direct reduction of Bi3+ to Bi0.56 As mentioned above, the type


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of bonding between Bi-O and V-O, which is more ionic in the former and more covalent in the latter, is of interest as it lends to the idea that Bi-ion can more easily diffuse in/out of the lattice, when subjected to charging/discharging, as opposed to V5+ which is covalently bound to the framework. As a consequence, faradaic behavior of solely Bi-ion is captured and the vanadate units are contributing to the pseudocapactive aspects of the CV curve. On increasing the scan rate an increase in the current response can be simultaneously observed which indicates faster interfacial redox kinetics.55 It should be highlighted that monoclinic BiVO4, despite having three dimensional crystal structure, unlike most of the two dimensional layered battery materials, shows impressive reversibility. It should, however, be noted that upon characterizing the working electrode after cyclic voltammetry (upwards of 50 cycles), signatures of Bi2O3, NaVO3 and V2O7 were recorded in the XRD pattern.59 It is anticipated that Bi-ion diffuse out of the lattice while charging, and it is the Na+ from the electrolyte (NaOH) that diffuses into the lattice to chargebalance the system and this results in the formation of deleterious side phase, NaVO3. The electrochemical capacitance of this vanadate was also evaluated by galvanostatic charge discharge (CD) measurements. The CD profiles in the potential window 0.0 V to -1.0 V appeared to be non-symmetric and this suggested pseudo-capacitive nature of material for charge storage applications, Figure 8d. CD profile, during charging, shows capacitive behavior with a gradual increase in voltage over time, however upon discharging it shows faradaic mechanism depicted by a plateau that persists over a long period of time (discharge time is also substantially longer than the charging time). There is however a steep voltage drop (upon the first few seconds of discharge) which is being ascribed to internal resistance.55 This internal resistance is predominantly due to the non-optimized electrolytic interfaces and this is currently the focus in our lab.


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In terms of the storage capacity, depending on the morphology employed, when subjected to cyclic charging and discharging, only ca. 40 % of the initial discharge capacity was retained after 200 cycles.55 It was determined that the poor cycling stemmed from either electrode contraction or expansion that accompanied pseudo reversible diffusion of ions in/out of the lattice and also there was degradation of the electroactive material (viz. formation of NaVO3 as detailed above). To circumvent the aforementioned problem of poor stability and to further augment the storage capacity of BiVO4, various composites have been formed of BiVO4 with different types of capacitive components.55,56 A capacitive component needs to have high specific surface area, good electronic conductivity, inherent chemical stability (under a wide temperature range and a variety of pH conditions) and propensity for charge storage.56 Considering aforementioned requirements we have employed carbon allotropes such as single walled carbon nanotubes and graphene which augmented the electrochemical behavior (at loading levels as low as 10 wt % of the carbon component) and storage capacity by a factor of 4 in comparison to pristine BiVO4. Also, columbic efficiency was found to be ca. 90% after 200 charge discharge cycles showing its long term stability. Specifically, BiVO4 in conjunction with graphene (152 Fg-1, 7 Ag-1) performed better than SWCNT/BiVO4 composite (110 Fg-1, 7 Ag-1) for the same amount of loading. Moreover, we have also incorporated 2D transition metal dichalcogenides (TMDs), such as molybdenum disulfide, and compared the composites’ behavior with graphene analogues.56 MoS2/BiVO4 nanocomposite demonstrated much larger values of charge storage, longer discharge times and improved cycling stability in comparison to pristine BiVO4, or even graphene/BiVO4 composites. It should be noted that there are reports in literature that point to the deleterious effects of addition of TMDs for electrode materials. This mainly stems from adverse ion interactions (due to large interfacial interactions with the electrolyte ions) that lead to parasitic reactions which contribute


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to electrolyte decomposition.60 With our work however, we were able to circumvent this problem by synthesizing explicitly small lateral dimensional particles of MoS2. AFM topographic image of MoS2 showed lateral dimension of the particles in the range of 50-200 nm with a multitude of heights inferring the presence of monolayer, bilayer and tri-layer regions which was also corroborated from Raman data. It was shown that as the lateral size of MoS2 was decreased, there was a higher preponderance of low coordination edges, step-edges as well as the corner atoms as compared to the basal plane atoms.61 The effect of these states appeared to dominate over those of the basal atoms and these contributed to higher charge storage sites. In addition, since the MoS2 nanoparticles are made up of only a few-layered sheets, it is being hypothesized that such loosely stacked sheets are able to accommodate structural changes in a better manner upon cycling when compared to bulk MoS2 leading to a more stable and higher energy density material.62 An interesting offshoot of our work with BiVO4 has been the realization that this material is also a photoactive n-type semiconductor and therefore there exists the possibility of this property to be coupled with its energy storage capability.59 It is useful to note that currently there are a few reports that cite merging of energy capture and storage in a single device with a single active material as opposed to devices that consist of independent batteries and solar cells which are subsequently coupled.63–65 Currently, there are reports on devices based on either electrostatic or electrochemical charge

storage

using

technologies

involving

photovoltaic

cells

(such

as

DSSCs),

photoelectrochemical or photochemical assisted generation of high energy charge carriers and these are being interfaced with energy storage electrodes.63–65 In general, light assisted energy storage systems are composed of two kinds of materials: light harvesting and energy storage component. Light harvesting component consists of materials that are capable of absorbing light and generating extractable electron-hole pairs, while energy storage component consists of


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materials that can trap the charges and store them under periods of illumination, and subsequently release the same during the discharge process. WO3/TiO2 and Ni(OH)2/TiO2 are the two most widely studied hybrid energy storage systems where TiO2 is the light absorber whereas WO3 and Ni(OH)2 are the charge storage components. In WO3 photo generated electrons are usually stored in the form of tungsten bronze (NaxWO3),66 and in Ni(OH)2 photo generated holes are stored in the form of chemical reaction where conversion of Ni2+ to Ni3+ takes place.64 Unfortunately, these tandem devices show poor efficiencies mainly owing to multiple interfaces being involved. From already existing work, it is clear that there is a high complexity involved in coupling energy capture and storage in a single device due to stringent requirements such as semiconductor band gap and alignment, charge transport kinetics, energy conversion efficiency, and material stability.64 Therefore, we adopted an alternate approach to couple energy capture and storage in a single device with the aim to minimize interfaces that in principle would lead to better performance and charge transport efficiency. We investigated the redox behavior of BiVO4 directly in the presence and absence of light in order to study if it is feasible for an electroactive component to be also photoactive, as a bi-functional material, in a single energy storage device.59 The advantage of this work was in reducing the number of interfaces by employing a single bi-functional material and the preliminary data has shown that in situ charge generation and storage can be made viable. To elaborate, the electrochemical measurements were performed at room temperature in a three electrode flat cell employed with Hg/HgO as the reference electrode, Pt foil as the counter electrode and 2 M NaOH as the electrolyte. The working electrode was a BiVO4 slurry that was dropcast on the graphite plate (area of coating, 1 cm2). Under dark conditions, conventional BiVO4 CV curve was obtained with two oxidation peaks and a single reduction peak that could be


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attributed to pseudo reversible faradaic process, Bi3+↔Bi0. As discussed above, despite the irreversible damage caused to the working electrode in the presence of 2M NaOH, under the current experimental setup, enough BiVO4 was deposited such that 70 CV cycles could be feasibly carried out. Subsequently, the working electrode was irradiated with 100 W tungsten lamp with a 400 nm long pass filter. As the working electrode was impinged with the light, CV of BiVO4 was found to be intriguingly perturbed, Figure 9a.

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Following were the two major perturbations observed under ‘light’ conditions: an augmentation in the area under the CV curve and even more important observation was the deflection in the reduction peak potential towards positive potential, however no such shift was observed for the oxidation peak potentials. To avoid any influence of the initial ‘activation process’,67 first 20 CV cycles were carried out in dark. Subsequently, 21st cycle onwards light was impinged on the working electrode for 5 cycles and the current was simultaneously measured, then the light was


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turned off for next 5 cycles. This way BiVO4 was impinged with visible light in a pulsating mode and its redox behavior was monitored. An average augmentation of ca. 288 mA.mV and 175 mA.mV was recorded in the area under the reduction and oxidation peaks, respectively, Figure

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Figure 10. (a) and (b) depicts the variation in the area under reduction and oxidation peak, respectively, as a function of cycle number on irradiating BiVO4 with visible light, (c) and (d) shows the variation in the reduction and oxidation peak potential of BiVO4 as a function of cycle number, respectively, in the presence and absence of light.

The data has been normalized with respect to cycle number 18. In principle, it can be inferred from this observation that a larger amount of BiVO4 was being probed electrochemically simply by impinging it with visible light. An even more intriguing observation was the asymmetric behavior


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with respect to the shift in the redox potentials. Figure 10c and 10d show the change in the reduction and oxidation peak potentials upon pulsing with light. It is clearly observed that reduction peak deflects towards positive potential indicating that Bi3+ → Bi0 is now a more feasible reaction. An average deflection of ca. 59 mV was recorded for reduction peak, however a small deflection of only ca. 4 mV was observed in the oxidation peak potential. This asymmetric variation directly indicated that in situ charge carriers (that are generated in the presence of light) play a pivotal role in affecting the faradaic process. It can be inferred that the load on the external device to perform the reduction step can be reduced simply by impinging the material with light. Also, this showed that material was charged at the same potential, however it discharged at a higher potential. In BiVO4 (n-type semiconductor) excited electrons that are formed on photo irradiation are more mobile and can reach the surface faster than the hole that are perhaps buried inside, and hence we observe its effect in the form of facile reduction i.e. shift in the reduction peak is significant as compared to the shift in the oxidation peak. Another set of experiments was performed on BiVO4, where electrons and hole scavengers were introduced into the system. The purpose of adding these scavengers was that they would nullify the extra concentration of photoinduced charge carriers. Removal of light induced holes (in the presence of Na2SO3, hole scavenger) did change the reduction potential however not substantially, there was a small shift on impinging the system with light. This behavior could be attributed to the fact that in case of BiVO4 (which is an n-type semiconductor) excited holes that are generated on photo-illumination are not mobile enough and hence, are not contributing to the redox behavior of BiVO4. Therefore, their presence or absence does not play a vital role. However, it should be highlighted that in the presence of O2 (electron scavenger) reduction peak did not deflect from its position in the presence of light. This clearly indicated that excited electrons are predominant contributors to charge storage


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and on quenching these carriers, effect of light is completely suppressed. Hence, this observation gives a clear insight to the fact that charge carriers that are generated on photo-illumination contribute to the faradaic process apart from augmenting the capacitive storage. Galvanostatic charge-discharge measurements were also performed on the system in the presence and absence of light, again in pulsed mode. Figure 9b shows CD curve of the working electrode, where alternate cycles were run in the presence and absence of light. A significant enhancement in the discharge duration was observed on impinging the working electrode with visible light, however the charging time was unaffected. The overall discharge duration got augmented by ca. 30 % (averaged over 3 set of cycles). Discharge time and specific capacitance of all the cycles in dark and light have been tabulated, Table 2. Table 2. Shows the discharge time and specific capacitance of BiVO4 in the presence and absence of light. Condition

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Specific Capacitance (F/g) 735 930 603 876 570 735

Moreover, we have performed additional experiments involving a p-type semi-conductor (Co3O4) to showcase how the behavior is similar but opposite for a photo and redox active semi-conductor where the holes are now the majority carriers, Supporting Information Figure S3. The reader is referred to published work for more details, ref 59.59


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As a further to the above work, it can be observed that one major concern with BiVO4 is the poor cycling stability that limits the number of charge discharge cycles that can evaluated. As such, to mitigate this stability issue, currently formation of viable Bi3+ electrolytes are being pursued along with formation of better interfaces so as to minimize charge carrier losses. Moreover, formation of other energy storage materials that are both photo- and redox-active are being considered such that that single devices (as opposed to tandem or coupled photovoltaic+electrochemical devices) can be considered feasible. Conclusion In summary, we have detailed two of our recent endeavors in pursing new materials for energy capture and storage. In part A we have showcased a study where we have introduced a new organic cation, imidazolium, into a hybrid perovskite structure and have successfully formed both single crystals and thin films of this hybrid ImPI structure. Imidazolium is a larger and more symmetrical cation and results in a structure that shows intriguingly a hexagonal unit cell. This structure has direct ramifications as it is a non-centrosymmetric unit cell and as such, ImPI shows improved thermal stability, lack of phase change at low temperature and three times increased stability under ambient conditions over MAPI. Interestingly due to its 1D character with respect to the PbI6 octahedra, this material has the potential to showcase interesting opto-electronic behavior. In part B we have introduced the use of a new electroactive material BiVO4 which has been synthesized in morphologies ranging from simple nanostructures to oriented nanowires directly grown on transparent electrodes. When the working electrode is formed in conjunction with either SWCNT or few layered, nanostructured MoS2 then these composites show much larger values of charge storage, longer discharge times and improved cycling stability in comparison to pristine BiVO4 or graphene/BiVO4 composites, and hence these are being actively considered as promising


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candidates for energy storage. Moreover, the work has now progressed in even more of a novel direction in that we have evaluated the capability of merging energy capture and storage. This is conceptually novel in that the idea involves directly storing solar radiation as opposed to forming devices that consist of independent batteries that are separately coupled with solar cells. This is being achieved simply because BiVO4 is also an n-type semiconductor with a band gap of 2.3 eV. Hence, it is electroactive and photoactive. Importantly, depending on the majority charge carriers that are generated, either of the faradaic processes (reduction in BiVO4 or oxidation in another complementary p-type semiconductor Co3O4) has been observed to be facilitated. On introducing scavengers into the system, these disproportionate perturbations in the CV have been found to be suppressed indicating the direct influence of excited charge carriers on the faradaic behavior. References (1)

Sun, H.; Zhang, Y.; Zhang, J.; Sun, X.; Peng, H. Energy Harvesting and Storage in 1D Devices. Nat. Rev. Mater. 2017, 2, 17023.

(2)

Sygletou, M.; Petridis, C.; Kymakis, E.; Stratakis, E. Advanced Photonic Processes for Photovoltaic and Energy Storage Systems. Adv. Mater. 2017, 29 (39), 1700335.

(3)

Li, Y.; Qian, F.; Xiang, J.; Lieber, C. M. Nanowire Electronic and Optoelectronic Devices. Mater. Today 2006, 9 (10), 18–27.

(4)

Arunachalam, V. S.; Fleischer, E. L. The Global Energy Landscape and Materials Innovation. MRS Bull. 2008, 33 (4), 264–288.

(5)

Goetzberger, A.; Hebling, C.; Schock, H.-W. Photovoltaic Materials, History, Status and Outlook. Mater. Sci. Eng. R Reports 2003, 40 (1), 1–46.


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(6)

Saparov, B.; Mitzi, D. B. Organic–Inorganic Perovskites: Structural Versatility for Functional Materials Design. Chem. Rev. 2016, 116 (7), 4558–4596.

(7)

Heo, J. H. et al. Efficient Inorganic–organic Hybrid Heterojunction Solar Cells Containing Perovskite Compound and Polymeric Hole Conductors. Nat. Photonics 2013, 7, 486.

(8)

Xiao, Z.; Bi, C.; Shao, Y.; Dong, Q.; Wang, Q.; Yuan, Y.; Wang, C.; Gao, Y.; Huang, J. Efficient, High Yield Perovskite Photovoltaic Devices Grown by Interdiffusion of SolutionProcessed Precursor Stacking Layers. Energy Environ. Sci. 2014, 7 (8), 2619–2623.

(9)

Stranks, S. D.; Eperon, G. E.; Grancini, G.; Menelaou, C.; Alcocer, M. J. P.; Leijtens, T.; Herz, L. M.; Petrozza, A.; Snaith, H. J. Electron-Hole Diffusion Lengths Exceeding 1 Micrometer in an Organometal Trihalide Perovskite Absorber. Science (80-. ). 2013, 342 (6156), 341 LP-344.

(10)

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Supporting Information. Crystal structure of PbI2 generated from its crystallographic information file, temporal evolution of X-ray diffractograms of ImPI and MAPI materials (both crystal powder and films grown on substrate), and electrochemical data of Co3O4 rods is presented. Corresponding Author *[email protected] Author Contributions The manuscript was written through equal contributions of all authors. All authors have given approval to the final version of the manuscript. #These authors contributed equally to work presented in this manuscript. Funding Sources The Department of Atomic Energy, India is gratefully acknowledged for providing plan funds to support this work.


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Langmuir

TOC.

Energy Capture

vs

vs

Energy Storage No shift

10

BiVO4

Current (mA)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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0

-10 In Dark In Light

-20

-30

Shifts towards +ve potential -1.2

-1.0

-0.8

-0.6

-0.4

-0.2

0.0

Potential (V) vs Hg/HgO


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Yukti Arora is currently pursuing her Ph.D. degree in the Materials Chemistry Group, Department of Chemical Sciences at TIFR. She received her Bachelor’s degree in Chemistry from the University of Delhi, India (2012). She was awarded gold medal in her Master’s degree with specialization in Organic Chemistry from the University of Delhi, India (2014). Her current research interests are synthesizing novel materials for solar energy storage and understanding the underlying mechanism.

Charu Seth is a Ph.D. student in Materials Chemistry Group, Department of Chemical Sciences at TIFR. She completed B.Sc. in Chemistry from St. Stephen’s College, University of Delhi, India in 2013 after which she joined the M.Sc.-Ph.D. program at TIFR. Her research interest lies in synthesis and study of opto-electronic properties of nanostructures for energy capture applications.

Deepa Khushalani received her Ph.D. from University of Toronto in 1998 and subsequently worked as a Postdoc in University of Bristol. From 2001-2004 she worked as a lecturer in University of Kent before returning to India in 2004 where she has been working at TIFR ever since. Her area of specialization involves materials chemistry with emphasis on exploiting alternate sources of clean energy involving photovoltaics, energy storage devices and catalysis.


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Crafting Inorganic Materials for Use in Energy Capture and Storage

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