Large-scale synthesis of hexagonal corundum-type In 2 O 3 by ball milling with enhanced lithium storage capabilities

Hexagonal corundum-type indium oxide (h-In2O3) is the structure that normally exists in a hightemperature and pressure environment. This structure has been realised from ambient environment stable cubic indium oxide (c-In2O3) using a high-energy ball milling approach at room temperature, in which the rearrangements of InO6 polyhedral units take place via plastic deformation and large defect creation during the milling process. More interestingly, the high-temperature h-In2O3 structure as anode materials in lithium-ion batteries exhibits lithium storage capabilities enhanced by up to 8 times compared to the c-In2O3 phase. This study demonstrates an effective ambient environmental approach for the production of high-pressure/temperature structures, h-In2O3, which may be extended to explore new phases and novel properties in other oxide systems.


Introduction
2][3] As an important functional material, indium oxide (In 2 O 3 ), an n-type semiconducting binary oxide with a wide band-gap (3.5-3.7 eV), has attracted much interest for several decades due to its high electrical conductivity and good optical transparency. 4,5here are three different crystal structures of In 2 O 3 : cubic bixbyite-type (c-In 2 O 3 ) stable under ambient conditions, hexagonal corundum-type (h-In 2 O 3 ) and orthorhombic Rh 2 O 3 (II)-type structures stable only under high pressures and temperatures. 68][9] With the recent report by Idota et al. that some tin-based oxides exhibit high capacities in lithium ion batteries (LIBs), new opportunities seem to be opening for exploring new anode materials based on metal oxide materials. 102][13][14] The reversible lithium storage performance of h-In 2 O 3 has not been reported so far.h-In 2 O 3 has been found to exhibit stable electric conductivity and high sensitivity to gases, compared to c-In 2 O 3 .6][17][18] There are two main methods for fabricating h-In 2 O 3 in the literature: the static or shock-induced phase transformation method involving high pressures (>15.3GPa) or high temperatures and high pressures (>1000 C and >6.5 GPa) 6,19,20 and special chemical methods under ambient pressure. 15,18,21,22These methods need not only an elevated temperature for the decomposition of metal salts, but also complex reaction conditions, which are disadvantageous from the manufacturing point of view.
In this letter, a large-scale and room-temperature production approach is demonstrated.The high-pressure phase h-In 2 O 3 can be produced from the c-In 2 O 3 phase via a high-energy ball milling process at room temperature.The low-environmental footprint synthesis method is based on a simple mechanical milling process with advantages of medium pressure range, low production temperature, and no need of any solvents and catalysts.More interestingly, the h-In 2 O 3 phase exhibits enhanced reversible storage capacity up to 8 times higher and good cycling performance in lithium ion batteries compared to the initial c-In 2 O 3 , indicating the advantages of high-pressure structures as electrode materials for energy storage.

Synthesis of h-In 2 O 3 powders
In a typical process, 9 g of c-In 2 O 3 powders (99.995%,Alfa Aesar) were loaded into a stainless steel container with four hardened steel balls (diameter of 25 mm), and then lled with N 2 -15%H 2 gas to a pressure of 300 kPa.A vertical rotating ball mill was used for the milling experiments under high-energy impact mode, which was realized by placing an external magnet under the container at a 45 position in relation to the vertical direction and the use of the mill rotation speed around 150 rpm. 23he temperature of the container surface was about 40 C during the milling process.The milling action was turned off every 2 days for 2 h for cooling purposes.

Characterization
The phase changes of milled samples were monitored with an X-ray powder diffraction (XRD) spectrometer (Panalytical X'Pert PRO system) using Cu Ka radiation at a wavelength of 1.5418 Å. Simulation and analysis of the XRD patterns of various phases were performed using the Material Studio program.Unit-cell parameters were obtained using DICVOL91, and structure renement was carried out using the Pawley method. 24Scanning electron microscopy (SEM) analysis was performed on a Zeiss Supra 55 VP instrument.Transmission electron microscopy (TEM) analysis was performed on a JEOL 2100 (operating at 110 kV) apparatus.High resolution TEM studies were carried out with a JEOL JEM 2011 LaB 6 (operating at 200 kV) instrument.Raman spectra were recorded on a Horiba Jobin Yvon Raman microscope in a backscattering geometry at the 532 nm line of an argon ion laser, provided with a CCD detector system.

Electrochemical measurements
The working electrodes of Li-ion battery cells were fabricated by mixing 70 wt% active material (h-In 2 O 3 powders), 15 wt% acetylene black (Super-P), and 15 wt% polyvinylidene uoride (PVDF) binder in N-methyl-2-pyrrolidone (NMP) to form homogeneous slurry.A thin layer of the slurry was coated on Cu foil to form the anode electrode; the electrode was dried overnight at 120 C in vacuum to remove the solvent.Coin cells (CR2032) were assembled using Li metal as the counter electrode in a dry glove box lled with argon.The electrolyte consisted of 1 M LiPF 6 in a non-aqueous solution of ethylene carbonate and dimethyl carbonate in a volume ratio of 1 : 1. Cyclic voltammetry (CV) measurements were carried out using a Solartron 1470E electrochemical workstation at a scanning rate of 0.1 mV s À1 .A galvanostatic cycling test of the assembled cells was performed using the voltage window of 0.01-3.0V versus Li/Li + .treatment.Three new diffraction peaks emerging around 32.8, 57.7, and 58.9 are indicative of a transformation to a new phase aer milling for 6 days.

Results and discussion
With further milling up to 8 days, the diffraction patterns of the sample show distinctive features indicating phase transition completion (Fig. 1).The Pawley renement for this new phase using an R 3c structural model yields a very good t with R wp ¼ 5.2% (Fig. 2b To further examine the different crystalline structures of the milled In 2 O 3 , Raman spectroscopy measurements were performed on the samples before and aer milling for 8 days.Group theory predicts the following representation for the center of Brillouin zone (q ¼ 0) optical vibrational modes of c-In 2 O 3 . 26The irreducible representation is given as where A g , E g , and T g represent Raman-active modes, A u and E u represent inactive modes, and T u represents infrared-active modes. 27Fig. 3 shows representative Raman spectra of the indium oxides in the frequency range of 110-900 cm À1 before milling and aer milling for 8 days.Five Raman active phonon modes of the initial sample are discernible and all the observed modes correspond well to the band positions of c-In 2 O 3 reported in the literature. 26Aer milling for 8 days, different features appear in the Raman spectrum.The positions of the six Raman peaks observed at 161, 218, 304, 377, 506 and 595 cm À1 are in good agreement with the previously described h-In 2 O 3 structure, predicting seven Raman active modes (2A g1 + 5E g ) by the factor group analysis. 22The Raman results actually provide additional evidence of the h-In 2 O 3 phase obtained from the ambient phase c-In 2 O 3 directly by a simple high-energy ball milling process.
The SEM image (Fig. 4a) shows that most h-In 2 O 3 particles are spherical and have diameters in the range of 0.1-1 mm.The histogram (Fig. 4b) shows that most particles have a diameter around 200 nm.The TEM image in Fig. 5a      It is well known that energy barriers oen exist in chemical reactions and phase transformations.The reaction or transformation becomes possible when sufficient energy is provided to overcome these barriers.A huge energy barrier should exist for the transformation from c to h-In 2 O 3 as it normally takes place at high temperatures and high pressures. 5In the case of the room-temperature milling process, a huge amount of energy is injected into the materials by repeated high-energy impacts and the energy is stored in the form of structural defects and large surface energy.Aer 6 days of milling, when the stored energy in the system is sufficiently high, c to h-In 2 O 3 transformation takes place even at room temperature.][30] The electrochemical performance of the h-In 2 O 3 in the cell conguration of Li/h-In 2 O 3 was evaluated.The CV measurement (Fig. 6a) was carried out to investigate the electrode reaction processes at a scan rate of 0.1 mV s À1 .In the rst cycle, two apparent peaks are observed at 0.58 V and 1.06 V in the discharge process.The peak at 0.58 V could be assigned to the formation of Li x In alloy proceeding in a multi-step process during lithium insertion. 12,13These two peaks disappear in further cycles, indicating the irreversible reduction of h-In 2 O 3 with a large irreversible capacity in the rst cycle.New reduction peaks at 0.78 and 0.46 V appear aer the rst cycle and shi to a lower voltage during further cycles.Two oxidation peaks at 0.51 and 0.70 V could be ascribed to the de-alloying of Li/In which form in the discharge process.Another two broad oxidation peaks at 1.18 and 1.78 V are observed, which may be ascribed to the reverse alloying reaction among Li x In alloys of different stoichiometry and the further formation of metallic phase indium.These two peaks fade away quickly aer the rst several cycles indicating the poor reversibility of this reaction.Fig. 6b shows the 1 st , 2 nd , 5 th , and 50 th cycles charge-discharge curves of the h-In 2 O 3 at a constant current density of 30 mA g À1 and a cutoff voltage window of 3.0 and 0.01 V.The curves show different plateaus due to the redox reactions associated with Li + insertion and extraction and are consistent with the CV results.The ramp above 0.87 V in the discharge curve possibly derives from the irreversible reduction of h-In 2 O 3 .The discharge plateaus around 0.42 V indicate the formation of Li x In alloy phases and the charge plateaus around 0.69 V are also in agreement with the cyclic voltammogram prole.
The h-In 2 O 3 electrode delivers initial discharge-charge capacities of 1256 and 729 mA h g À1 , respectively, which are higher than the corresponding capacities (1016 and 453 mA h g À1 ) from c-In 2 O 3 as shown in Fig. 6c.Subsequently, the charge and discharge capacities of the h-In 2 O 3 electrode approach a constant level at about 390 mA h g À1 aer 50 cycles with a coulombic efficiency of 98.8%.This reversible capacity is much higher than that of the c-In 2 O 3 (49 mA h g À1 ) electrode aer 30 cycles and some other oxides such as MnO 2 (about 200 mA h g À1 aer 10 cycles), 31 Co 3 O 4 (about 200 mA h g À1 aer 20 cycles) 32 and Mn 3 O 4 (below 200 mA h g À1 aer 10 cycles). 33In addition to the improved reversible capacity and excellent cycling behavior of the h-In 2 O 3 , the rate capabilities at various current rates from 30 to 300 mA g À1 are illustrated in Fig. 6d.A reversible capacity of 125 mA h g À1 can be sustained at the highest current rate of 300 mA g À1 .A capacity of 382 mA h g À1 at 30 mA g À1 is retained aer 25 cycles of charge and discharge at various current densities, demonstrating good reversibility and structural stability of the h-In 2 O 3 .
The better performance in terms of capacity and cyclability of h-In 2 O 3 compared to c-In 2 O 3 may be attributed to three aspects.First, the structure of c-In 2 O 3 is composed of an edgeconnected network of InO 6 octahedra and trigonal prisms, as shown in Fig. 7a, in which Li ions occupy the interstitial sites in this lattice framework.Consequently, the restricted Li diffusion in three-dimensional volume leads to poor electrochemical performance of c-In 2 O 3 as the LIB anode.In contrast, the h-In 2 O 3 structure consists of a three-dimensional array of edges and corners sharing InO 6 trigonal prisms, giving rise to channels and cavities perpendicular to the c-axis (Fig. 7b) and these channels provide elegant pathways for diffusion and storage of Li-ions.Second, the h-In 2 O 3 has a better electrical conductivity than c-In 2 O 3 which is advantageous for high capacity and fast

Paper
Journal of Materials Chemistry A charging and discharging processes. 34Third, the large quantities of defects and smaller grain size in the h-In 2 O 3 particles, as shown in Fig. 5d, not only provide additional sites for the insertion of Li ions, but also facilitate the diffusion of ions. 35All these properties suggest that h-In 2 O 3 is an outstanding candidate for anode materials of LIBs.

Conclusions
High-energy ball milling is a high yield, less hazardous and straightforward route to produce the h-In Typical XRD patterns of In 2 O 3 were collected and are plotted in Fig. 1 as a function of milling time.All diffraction peaks of the initial material can be indexed to a pure cubic bixbyite-type structure (Ia 3) In 2 O 3 (c-In 2 O 3 ) crystal with space group no.206, lattice constants a ¼ b ¼ c ¼ 10.11(1) Å, and unit cell volume V 0 ¼ 64.8(1) Å3 , in good agreement with those previously reported in the literature.25As shown in Fig.2a, it has two different types of In 3+ (In 1 and In 2 ) occupying both octahedral and trigonal prismatic interstices within the lattice of O 2À ions, and O atoms at the 8b, 24d, and 48e Wyckoff positions, respectively (inset of Fig.2a).No changes were observed in the XRD patterns of the samples aer milling for 2 and 4 days except peak intensity reduction and shape broadening.The signicant peak shape changes are caused by grain size reduction and structural distortion as the typical effect of ball milling

Fig. 1
Fig. 1 XRD patterns of the In 2 O 3 recorded at different milling times.

Fig. 2
Fig. 2 Pawley full-profile refinements of the diffraction patterns collected from the In 2 O 3 powders before milling (a) and after milling for 8 days (b).In the case of the high pressure phase, a good fitting with R wp ¼ 5.2% is obtained for the diffraction pattern in (b).Red, black, blue solid lines, and tick marks represent experimental, calculated, residual patterns, and the positions of calculated Bragg reflections, respectively.The insets show the respective crystal structure, and the different types of indium atoms are marked with different colors.
) corresponding to the high pressure and high temperature h-In 2 O 3 .The lattice parameters rened within this space group for the new phase are a ¼ 5.438(2) Å, c ¼ 14.474(3) Å, and unit cell volume V 0 ¼ 61.78(1) Å3 (Z ¼ 6).h-In 2 O 3 has a different atomic arrangement, containing only trigonal biprism coordinated In 3+ ions within the lattice of O 2À ions, with O atoms located at Wyckoff positions 12c and 18e, as shown in Fig. 2b.It can be explained that plastic deformation and increased defect concentrations created by ball milling play important roles in the occurrence of phase transformations in the In 2 O 3 system.This mechanical activation initiates the lattice change of O 2À ions and the shi of In 3+ ions from octahedral and trigonal prismatic to trigonal biprism sites (Fig. 2), resulting in distortion and redistribution of InO 6 octahedra units which are the building blocks of the h-In 2 O 3 structure.
shows typical morphology and sizes of the starting c-In 2 O 3 particles.The highresolution TEM image in Fig. 5b presents an interdistance of 0.407 nm between the parallel fringes corresponding to the dspacing of the (211) plane of c-In 2 O 3 .The TEM image presented in Fig. 5c reveals the smaller particle and grain sizes of h-In 2 O 3 particles.A typical selected area diffraction (SAD) pattern is shown in Fig. 5c.All electron diffraction rings can be indexed to the h-In 2 O 3 phase which is consistent with the XRD analysis

Fig. 3
Fig. 3 Micro-Raman spectra of In 2 O 3 powders before milling and after milling for 8 days in the 100-900 cm À1 frequency regions.

Fig. 4
Fig. 4 SEM image of h-In 2 O 3 particles (a) and histogram of diameter distribution of the particles (b).

Fig. 5
Fig. 5 TEM images of the In 2 O 3 powders (a) before milling and (c) after milling for 8 days; the SAD pattern corresponding to part c is shown in the inset.HRTEM images of the In 2 O 3 powders (b) before milling and (d) after milling for 8 days.
results.The HRTEM image (Fig.5d) of an h-In 2 O 3 particle shows the interdistance of the fringes of 0.283 nm, in agreement with the lattice distance of the (104) planes of the h-In 2 O 3 .These results further conrm the phase transition from c-In 2 O 3 to h-In 2 O 3 during the high-energy ball milling process.

Fig. 6
Fig. 6 Electrochemical properties of h-In 2 O 3 : (a) cyclic voltammogram of the h-In 2 O 3 electrode; (b) discharge-charge voltage curves under a current density of 30 mA g À1 ; (c) cycling performance and coulombic efficiency of h-In 2 O 3 and initial c-In 2 O 3 at 30 mA g À1 for 50 and 30 cycles, respectively; (d) rate performance of h-In 2 O 3 at rates from 30 to 300 mA g À1 .

Fig. 7
Fig. 7 Polyhedral representations of (a) c-In 2 O 3 and (b) h-In 2 O 3 , perpendicular to the c-axis.
2 O 3 from c-In 2 O 3 at room temperature.The structural transition is supported by XRD analysis, Raman measurements, and HRTEM examinations.The structural relationships and the possible transformation mechanisms can be explained in terms of the continuous topological variation of atom coordination or continuous displacement of the atoms induced by distortion and redistribution of InO 6 polyhedra units under high-energy milling impacts.As anode materials for lithium-ion batteries the h-In 2 O 3 exhibits an enhanced and stable capacity of 390 mA h g À1 (aer 50 cycles) compared to c-In 2 O 3 (49 mA h g À1 , aer 30 cycles).The high energy ball milling method can induce solidstate phase transformations and produce new metastable structures with enhanced or new properties.