Electrochemical activity of a-MoO3 nano-belts as lithium-ion batterycathode{

Uttam Kumar Sen and Sagar Mitra*

Received 6th July 2012, Accepted 18th September 2012

DOI: 10.1039/c2ra21373g

Few metal oxides have seen renewed interest because of their novel reactivity towards Li, leading to a

large storage capacity. However, apart from their large capacity gain, it suffers from cycling

instability, large polarization loss and poor rate performance. Herein, we report on the structural,

morphological and electrochemical properties of a-MoO3 nano belts prepared by a simple

hydrothermal method and used as a cathode for lithium-ion battery application. During the electrode

preparation, we observed that the MoO3 nano-belt composite sample cast on stainless steel (SS)

substrate leads to a better electrochemical performance towards Li compared to aluminium (Al) or

nickel (Ni) substrates. The reason behind the poor performance was considered here, due to surface

passivation on Al substrates. This report comprises experimental results depicting (i) a sustained

reversible capacity of 140 mA h g21 for over 50 cycles at a rate of 200 mA g21, (ii) outstanding rate

capabilities with reversible capacities as high as 320 mA h g21 at a rate of 50 mA g21and (iii)

electrochemical stability of a-MoO3 nano belts towards a stainless steel substrate. Being able to make

such highly oriented a-MoO3 nano belt-based electrodes, through the hydrothermal process and

providing the electrochemical results, together show another efficient way to use MoO3 electrodes as

a cathode in lithium-ion batteries.

Introduction

Rechargeable Li-ion batteries can offer high energy density,

flexibility, light weight and long cycle life; they are fast gaining

popularity as the technology of choice for portable computing

and telecommunication equipment for today’s requirements.1

Moreover, with the increasing awareness of side issues linked to

air pollution combined with the foreseen oil shortage, we are in a

new period where the use of renewable energy sources and

electric transportation becomes a must.

One of the major challenges of next generation Li-ion

technologies for high power applications such as hybrid vehicles

and clean energy storage are energy density, power density,

safety and cost.2 As the demand for performance exceeds the

capabilities of the existing Li-ion technology, new electrode

materials with superior electrochemical properties, performance

and low cost must be developed.

Orthorhombic molybdenum oxide (a-MoO3), which is the

most stable form of molybdenum oxides, shows a significant

interest because of its unique electrical and electrochemical

properties, especially as it undergoes the Li-ion intercalation–

deintercalation process in its two dimensional layered structures.

MoO3 also exists in two other metastable forms known as

b-MoO3 (monoclinic) and h-MoO3 (hexagonal) and these can be

easily converted to a-MoO3 upon heating. Due to the ability of

Li+ intercalation in the inner layers, a-MoO3 can act as a host

material for Li+ and can be used as a cathode material for Li-ion

batteries.3 Despite the high lithium mobility even at room

temperature, the main disadvantage in the electrochemical

application of pure MoO3 is poor electronic conductivity (s =

1 6 1023 S cm21).4,5 Therefore the electrochemical performance

of pure a-MoO3 is not well studied. As per the literature, the

electrochemical performance of any material can be improved by

increasing the utilization rate of active materials i.e. to have a

better electrode–current collector interface or by using nano

materials.6,7

So far, worldwide attention has been paid to the nanostruc-

tured materials as electrodes for Li-ion batteries, due to their

attractive properties like small particle sizes, large active surface

area and high surface energy.8 As we know, the properties of

metal oxides depend closely on the microstructure, including

crystal size, orientation, morphology and crystallographic

density. For example, enhanced electrochemical activity was

observed for LiCoO2 and V2O5 fibers and tubes compared to

bulk samples.9–11

Herein, we report on the simple hydrothermal based synthesis

of anisotropic a-MoO3 nano-belts. Structural and morphological

studies have been conducted as well as reported using different

Electrochemical Energy Laboratory, Department of Energy Science andEngineering, IIT Bombay, Mumbai-400076, India.E-mail: sagar.mitra@iitb.ac.in; Fax: +91 22 2576 4890;Tel: +91 22 2576 7849{ Electronic supplementary information (ESI) available. See DOI:10.1039/c2ra21373g

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characterization techniques. The electrochemical behaviour was

studied using cyclic voltammetry (CV), in situ electrochemical

impedance spectroscopy (EIS) and galvanostatic charge–dis-

charge methods. As per our knowledge, 1.7 Li-ions per formula

of a-MoO3 can be taken during the first discharge cycle, which is

close to the gravimetric capacity of 317 mA h g21, but a serious

issue of capacity fading has been observed in all the previous

reports.3,12,13 Here, we observed that the capacity became

stabilized within a few initial cycles, giving rise to a capacity of

150 mA h g21 (47% of the initial capacity) at the 25th cycle, and

at the end of 100 cycles the capacity was found to be 127 mA h

g21 (40% of the initial capacity). The results show that a-MoO3

nano-belts exhibit high capacity and excellent cyclic stability

compared to previous reports. As per our knowledge, this report

will be the first of its kind showing cyclic stability and the role of

the current collector on the performance of a-MoO3 nano-belts

as a cathode. The present results are encouraging and show an

opportunity for a-MoO3 nano-belts to be used as a moderate

potential cathode in lithium-ion battery applications.

Experimental section

Synthesis of a-MoO3 nano belts

Anisotropic a-MoO3 nano-belts were prepared by a simple

hydrothermal method, as reported by Li et al.14 In a typical

experimental procedure, 2.0 g sodium molybdate (Na2MoO4?

2H2O, Merck) was dissolved in 10 ml deionized water. Then 5 ml 4

M perchloric acid (HClO4), (Merck) was added dropwise to the

molybdate solution under constant stirring. At the first step a

colourless solution was obtained, which turned turbid after stirring

for an hour. This solution was then transferred into a 35 ml Teflon-

lined stainless steel autoclave, which was kept inside a muffle

furnace at 180 uC for 24 h. A thick white coloured precipitate was

obtained, which was washed several times with distilled water and

acetone and then dried over a hot air oven at 60 uC for 12 h. The

dried sample was calcined at 300 uC and 500 uC in a pure nitrogen

atmosphere for further physical characterizations.

Characterization

All solid samples were systematically examined by X-ray

diffraction (XRD) at room temperature (25 uC) using a Philips

X’pert diffractometer with Cu-Ka radiation (l = 1.5418 A) at 40

kV and 40 mA. Different metal–oxygen vibrational modes of

a-MoO3 were characterized using a Raman spectrometer (Jobin

Yvon HR800) equipped with a 514.5 nm laser at 10 mW power.

A field emission gun scanning electron microscope (FEG-SEM,

JEOL-7600F) with a resolution of about 1 nm was used to study

the surface morphology of the samples. Further investigations

were done by the use of a high resolution field emission

transmission electron microscope (HR-TEM, JEOL-2100F).

Cell fabrication and electrochemical measurements

All the electrochemical performances of the materials were

carried out in two electrode Swagelok type cells in a Li/

electrolyte/MoO3 cell configuration. The complete cell comprises

a-MoO3 as the cathode, borosilicate glass fiber sheet (Whatman

GF/D) soaked with a 1 M LiPF6 solution [in EC : DMC/1 : 1

mass ratio (LP-30, Merck, Germany)] as a separator and pure

Lithium (Alfa Aesar) as a counter as well as a reference

electrode. The cells were assembled in an argon-filled glove box

(Lab Star, Mbraun, Germany) with a water and oxygen

concentration level of y1 ppm. The electrode materials were

hand ground with 12% carbon black (Super C-65, Timcal,

Switzerland) and 8% PVDF (Sigma Aldrich) for half an hour. A

slurry of the composite was prepared from the mixed powder by

adding a few drops of NMP (Qualigens, India). This slurry was

then tape cast on Al foil, stainless steel (SS 304) plates and Ni

mesh. Then the electrodes were dried at 120 uC under vacuum for

12 h and pressed, unless otherwise mentioned.

Electrochemical Impedance Spectroscopy (EIS) experiments

were carried out at open circuit voltage (OCV) in frequency

ranges 1 MHz–0.01 Hz in Bio-logic VMP-3. The in situ EIS

experiment was performed in a Bio-logic VMP-3 instrument and

the impedance experiment was performed at nine different

cathodic polarizations, starting at OCV, 2.8 V, 2.76 V, 2.65 V,

2.45 V, 2.3 V, 2.18 V, 2.0 V, and 1.51 V vs. Li/Li+. At each point,

the EIS was taken within the frequency range 1 MHz–0.01 Hz

and with voltage perturbation DV = 5 mV. A selection of the

points and the other details are given later in the results and

discussion section, while the points and the discharge curve,

along with the potential points where EIS was performed, are

shown in Fig. S1, ESI.{ The cycling voltammetry (CV) profile

was obtained by measuring the I–V response at a scan rate of 0.2

mV s21, and the cut-off voltage was 1.5 V–3.5 V. Bio-logic VMP-

3 was employed for CV measurement. The electrochemical

charge–discharge experiments were performed on an Arbin

Instrument, USA (BT2000 model) with various current rates. All

the electrochemical measurements were done at a constant

temperature of 20 uC.

Results and discussions

B Structural characterization

In the literature, a-MoO3 was prepared from peroxomolybdic

acid, and various methods15–17 have been used to obtained

peroxomolybdic acid. In the first step, the molybdenum salt was

converted to molybdic acid (H2MoO4), which gets oxidized by

the use of a strong oxidizing agent like H2O2 or HNO3. In the

present study we have used HClO4 or HNO3 to form

peroxomolybdic acid from sodium salt of molybdate, and used

the hydrothermal synthesis process to get pure anisotropic

a-MoO3 powder. In this particular work, we have observed, at

180 uC, that the shape of the nano-belts was more uniform and

the yield was maximum. It was also observed that when

phosphoric acid was used instead of HClO4 or HNO3, a

colourless solution was obtained at the end of the hydrothermal

reaction. The mechanism of MoO3 formation could be as

follows,

MoO42{

ClO4ð Þ{

OxidationMo2O3 O2ð Þ4� �2{

Mo2O3 O2ð Þ4� �2{ Heat

MoO3zH2OzO2

The XRD pattern of a-MoO3 powder is shown in Fig. 1a,

which is indexed as the orthorhombic phase of MoO3 and a

space group of Pbnm (JCPDS card No. 35-0609). As explained

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earlier, the a phase of MoO3 is thermodynamically most stable

among the three possible phases. The other two meta-stable

phases of MoO3 can be readily converted to the a phase by

heating the sample at 450 uC.18,19 The XRD pattern of the as-

prepared sample was compared with annealed samples heated at

300 uC and 500 uC, to convert the hexagonal and monoclinic

forms, respectively, to the orthorhombic form. Fig. 1a shows

that the peak positions and their relative intensities were

unchanged during heating at higher temperature, which proves

that the as-prepared sample is phase pure a-MoO3. For this

reason all the electrochemical performance was done with the as-

prepared samples. The diffraction pattern (Fig. 1a) shows that,

except for (0k0), all other peaks have very low intensity as

compared to the standard data. The observed diffraction

patterns with the highest intensities are (020), (040), and (060),

which indicates the preferential orientation of the a-MoO3 nano-

belts in the [001] direction.12,20

The Raman spectroscopy result shown in Fig. 1b also supports

the XRD results. The band positions of the as-prepared sample

and the annealed samples are the same. The observed band

positions are assigned as 290 cm21(B2g, B3g), 337 cm21(B1g, Ag),

667 cm21(B2g, B3g), 819 cm21(Ag, B1g) and 995 cm21(Ag, B1g),

respectively, which is in good agreement with the result obtained

from pure a-MoO3.21,22

From the high end scanning electron micrographs, it has been

observed that the a-MoO3 has a belt like morphology (Fig. 2).

The FEG-SEM images shown in Fig. 2a and b illustrate that the

a-MoO3 belts were formed uniformly, having a width in the

range 100–500 nm and a length of several micrometers. The belt

thickness was observed to be varying from 20 nm to 100 nm, but

careful observation shows that each belt consists of several thin

layers of MoO3 (shown in Fig. 2c). Essentially, thin layers of

MoO3 are the building blocks that combined to form the nano-

belt morphology. The number of MoO3 layers assembled

determines the thickness of the individual belt. The number of

layers varies from belt to belt, and as a result their thickness also

varies. The FEG-TEM image shown in Fig. 2d also shows that

a-MoO3 has the morphology of nano-belts, which is in good

agreement with SEM analysis. Lattice fringes were observed at a

higher resolution in TEM studies, shown in Fig. 2e. The lattice

spacing was found to be 3.6 nm and 4.0 nm, which correspond to

the d-spacing of the (100) and (001) planes of a-MoO3,

respectively.3,13

The selected area electron diffraction (SAED) pattern

recorded perpendicular to the anisotropic growth axis of an

individual nano-belt is attributed to the [010] zone axis (shown in

Fig. 2f). Combining with the TEM and SAED patterns indicates

that the highly crystalline nano-belts of a-MoO3 have grown

along the [001] orientation. Because of the gravitational force,

the majority of the nano-belts grounded selectively on the (010)

base planes of the TEM grid under free sedimentation during

TEM sample preparation, and show excellent crystal growth

direction along the [001] direction.20

Electrochemical performance

Since the pure a-MoO3 nano-belts are poor electronic con-

ductors, all the samples were initially mixed with conductive

carbon (12 wt%) and binder (8 wt%) to perform the electro-

chemical studies. Before going into the details of electrochemical

studies, we have observed enormous capacity fading with the

sample cast on aluminium substrate compared to other

substrates. Stable capacity retention was observed when the

sample was cast on a stainless steel current collector.

Selection of current collector

Lithium-ion battery electrodes are composite materials, in which

the active mass is bound to a metal current collector with a

polymeric binder such as polyvinylidene di-fluoride (PVDF). In

addition, the composite electrode contains a conductive additive,

usually carbon, and the composite is pasted on the metallic

current collector. Current collectors play an important role in the

performance of electrode materials for Li-ion batteries, and

aluminium (Al) as a current collector is the preferred choice for a

cathode material. Apart from Al, chromium (2.5–4.5 V), nickel

(¡ 4.5 V) and stainless steel (SS-304) (¡ 5.0 V) have also been

used as current collectors for cathode materials.23 For the

potential range 1.5 V–3.5 V vs. Li/Li+ the possible current

Fig. 1 (a) XRD and (b) Raman spectra of a-MoO3.

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collectors are Al, Ni and SS. In this work, we have observed that

a-MoO3 is not compatible with the Al substrate and, as a result,

leads to poor electrochemical performance. Cyclic voltammo-

grams (CVs) of a-MoO3 on the Al substrate are shown in

Fig. 3a. It was found that during the cathodic process, prominent

peaks were observed at 2.7 V and 2.15 V vs. Li/Li+, but in the

reverse process (anodic) a weak peak at 2.5 V vs. Li/Li+ was

observed. The reduction peak at y2.7 V vs. Li/Li+ disappeared

after the 1st cycle for all the substrates used in the present study.

Previously, the reduction peak was assigned as irreversible

formation of the LixMoO3 phase.24 It was well studied by

researchers using ex situ XRD and TEM techniques that the

irreversible intercalation process occurred at 2.7 V vs. Li/Li+,

forming the compound of LixMoO3 (0 , x , 0.25).24,25 XRD

and TEM studies also reveal the coexistence of a-MoO3 and

Li0.25MoO3 during the 1st irreversible process.24,25 Similar

observations were found during the CV of a-MoO3 using Ni as

the current collector (shown in Fig. 3b). Two prominent cathodic

peaks at 2.44 V and 2.03 V and one anodic peak at 2.57 V vs. Li/

Li+ were found. From the 2nd cycle onwards, only one cathodic

peak at 2.3 V was observed, whereas the anodic peak position

gradually shifted from 2.57 V (in 1st cycle) to 2.52 V (2nd cycle)

to 2.46 V (5th cycle). The observed cathode peak positions were

at lower potentials compared to the Al substrate. Moreover, we

observed one reversible intercalation–deintercalation of Li+ in

the potential range 2.15 V–2.25 V vs. Li/Li+, not only in the 1st

cycle but also in consecutive cycles with decreasing intensity for

all the substrates used currently. The above results agree well

with the previous literature by Tsumura et al.24

The best electrochemical performance of the a-MoO3 nano-

belts was observed when the SS substrate was used as a current

collector. For this reason all the electrochemical performance

and cyclic stability of a-MoO3 was done on a SS current

collector. Fig. 3c shows the cyclic voltammograms (CVs) of the

cathode materials within the potential range 1.5 V–3.5 V vs. Li/

Li+. In the first cycle, two distinct peaks were observed at 2.55 V

and 2.15 V vs. Li/Li+ in the cathodic process (discharge process).

However, in the following anodic process (charge process) there

was only one distinct peak at 2.56 V, similar to the Al substrate.

In consecutive cycles the cathodic peak at 2.55 V disappeared,

whereas the 2.15 V peak shifted to 2.23 V. On the other hand,

there is not much change in the anodic peak positions, which is

consistent with earlier results.3,12,26 According to the litera-

ture12,23 the prominent set of peaks (2.23/2.56 V or 2.15/2.56 V)

in the first cycle is due to the intercalation of Li-ions into the

interlayer spacing of the [MoO6] octahedron layers, whereas

the 2.55 V peak in the cathodic process is due to intercalation of

the Li-ion into the irreversible sites of the [MoO6] octahedron

interlayers. As a result, once the Li-ion was inserted into the

spacing of the [MoO6] octahedron interlayers (causing the

development of 2.55 V peak), it was trapped inside the cavity,

which led to the unavailability of these sites for further Li-ion

insertion and as a result no further peak at 2.55 V was observed.

But on a closer look we can observe that from the 2nd cycle a

new cathodic hump at 2.8 V was observed, which signifies that

some unrecoverable sites were still available for Li+ insertion,

but on a higher potential side. This extra intensity of the hump/

peak gradually diminished and finally disappeared from the 5th

cycle. Similar CVs were also observed by Chen et al.12 It was also

noticed during the charge–discharge performance (Fig. 3c) that a

substantial amount of capacity fading was there in the initial 5

cycles.

Electrochemical impedance spectroscopy (EIS) is one of the

most commonly used techniques to elicit the electrochemical

processes occurring at the electrode–electrolyte interphase, and

has been widely applied to the layered transition metal

Fig. 2 FEG-SEM images are shown in (a), (b) and (c). FEG-TEM images are shown in (d) and (e). (f) SAED pattern of a-MoO3.

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oxides.27,28 Here, the EIS technique was used to find out the

electronic conductivity, electronic structure, phase transitions

and their effects on the electrochemical performances of the

a-MoO3 electrode on different substrates with approximately

equal electrode material loading and similar film thickness.

Fig. 3d shows the increase in diameter of the high frequency

semicircle when the same composition of MoO3 was cast over the

SS, Ni and Al current collectors, respectively. After careful

investigation of the EIS spectra, we observed that the semicircle

in the high frequency region is a summation of two semicircles

(one very small depressed semicircle observed in the very high

frequency region and a large semicircle in the high to middle

frequency region). These two semicircles are attributed to SEI

formation and charge-transfer resistance from the high to the

low frequency region, respectively.

It was observed that the charge transfer resistance was a

minimum in case of SS and a maximum for Al. Surface

passivation on the Al surface could be the reason behind the

higher charge transfer resistance. A thin passivating layer of

Al2O3 always forms on the Al foil, but during electrode

fabrication it might be possible that a thicker layer of oxide is

formed. There can be a probability of a conversion reaction

among MoO3 and Al during electrode drying at 120 uC, resulting

in Mo and Al2O3. The presence of a small amount of Mo

nanoparticles can catalyze the electrolyte decomposition, which

causes corrosion on the Al surface.23 Zhou et al.29 also report

that poor performance on the Al current collector over SS is

expected due to corrosion, but the actual reason is not yet

understood.

A detailed study of impedance spectroscopy was undertaken

for SS supported a-MoO3. As shown in Fig. 4, in situ EIS was

carried out for SS supported a-MoO3 during the 1st discharge

cycle at a current rate of 50 mA g21. As discussed earlier, two

types of Li-ion intercalation processes are there, the first one is

for irreversible Li-ion intercalation, occurring at 2.75 V, which

accounts for a Li-ion insertion of 0.25 moles of Li-ions per

formula unit of MoO3, while the other one is at 2.3 V vs. Li/Li+.

For a better understanding of the intercalation mechanism, the

entire discharge profile was divided into three zones, and for

each zone three sets of impedance measurements were carried

out at three different potentials (shown in Fig. 4a). Zone-I was

restricted to the first stable plateau i.e. first intercalation layer

(Fig. 4b) while Zone-II (shown in Fig. 4c) represents the second

discharge plateau and the sloppy part was represented by Zone-

III (Fig. 4d).

As per the previous literature, the Nyquist plot of electro-

chemical intercalation–deintercalation for LiCoO2-based elec-

trodes is commonly observed in three distinct parts: a semicircle

in the high frequency range, which is generally attributed to the

migration of Li+ across the surface film; another semicircle in the

Fig. 3 CVs of a-MoO3 on (a) Al, (b) Ni and (c) SS current collector. (d) EIS of a-MoO3 on different substrates at OCV (in this case OCVs are 3.3 V

for Al, 3.1 V for Ni and 3.0 V for SS).

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middle frequency range, which is ascribed to the charge transfer

process from the electrolyte to the layered structure; and in the

low frequency region the observed step inclined line is attributed

to the solid-state diffusion process of the Li-ion into the LiCoO2

lattice.30 However, the explanations and observations for the

layered cathodes based on EIS experiments are debatable in

many ways. According to Nobili et al.,27 the EIS spectra for the

layered materials are interpreted in terms of the following

physical processes: (1) at very high frequency, the presence of the

semicircle is due to the presence of the electrode surface

passivating layer, (2) at an intermediate frequency range, the

dispersion is due to the charge transfer process, (3) in the low

frequency region, the semicircle is associated with the electronic

properties of the material and finally, (4) at very low frequency,

the inclined line is due to ionic diffusion.31 Very few reports

could see four different frequency dispersions and describe the

reaction mechanism clearly. As per Fig. 4b, the Nyquist plot of

the MoO3 cathode at the open circuit potential (1st EIS at 3.0 V),

shows a depressed small semicircle in the high frequency range,

as well as a large semicircle in the high to middle frequency range

and a slightly inclined line in the low frequency range. With

increasing electrode polarization potentials, the depressed small

semicircle in the high frequency region does not change

significantly. However low frequency dispersion started appear-

ing from 2.8 V onwards, which is a starting point of the co-

existence of MoO3-LixMoO3. The middle to low frequency

semicircle could be attributed to a change in electronic properties

of the electrode materials and this was observed from point 2 (see

Fig. 4b). It was also observed that with increasing cathodic

electrode polarization, the extremely low frequency inclined line

showing resistive behavior (more inclined towards the real axis)

is due to an increase of grain boundaries, which reduces Li-ion

diffusion. The more resistive nature of ion diffusion is observed

from point 2 (see Fig. 4b) and again started inclining towards an

imaginary impedance axis from point 8 (see Fig. 4d). After the

complete conversion of MoO3, the Li-ion can easily diffuse

through the materials, due to nanometric nature of the

composite formation. One more interesting point to note is that

after point 7 (see Fig. 4d), the diameter of the high to middle

frequency semicircle corresponding to charge transfer resistance

decreases with more cathodic polarization, which was ascribed to

the formation of a more conductive lithiated MoO3 phase. It

seems that in the low frequency region the line is not a sloppy

one. It has been found that the impedance response in the very

low frequency region is relatively more sensitive towards

experimental conditions and reproducible 2–5% of the time,

even when using the same battery cell.32 So, the last few points

can be considered as non-reproducible or an experimental

limitation. Now if we neglect the last 3 to 4 points, which are

at a very low frequency range, then the impedance spectra

(Fig. 4b–d) can easily be explained. During the reduction, at

point 1, Li+ starts accumulating at the surface of the active

Fig. 4 (a) 1st discharge cycles of a-MoO3 on SS substrate at a current rate of 50 mA g21, (b) EIS at Zone-I (c) EIS at Zone-II (d) EIS at Zone-III at

three different cathodic potentials.

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material, due to which a sharp increase in the impedance line

(point 2) was observed in the low frequency regions. At 2.7 V,

Li+ starts intercalating on the MoO3 layers, which causes two

phase (MoO3 and Li0.25MoO3) formation, and as a result solid-

state diffusion becomes more sluggish (point 3) as the resistive

nature is increased due to an increase in the grain boundaries.

But as soon as the first plateau ends and LixMoO3 starts forming

(start of 2nd plateau) the small diffusion characteristics are

visible and the sloppy nature of the impedance curve becomes

more prominent from point 4 to point 6 (Fig. 4c). At the end of

the 2nd plateau (Fig. 4d) the lithium diffusion improves due to

the formation of lithiated MoO3, and as a result the slope at the

low frequency impedance line is improved from point 7 to point

9. After 2.2 V there was no Li intercalation, only accumulation

of Li on the surface of the electrode. So, in the very low

frequency range (for points 8 and 9) a small sharp increment in

the impedance line along the Z99-axis was observed, which is due

to lithium accumulation33 (illustrated in Fig. S2, ESI{).

Rate and cycling performance

The galvanostatic charge discharge properties of a-MoO3 on SS

are shown in Fig. 5a and b. Two prominent discharge plateaus

are observed in the first cycle, which is in good agreement with

the CV experiment. From Fig. 5a, it was shown that the

discharge plateau with a higher rate of cycling shifted to a lower

potential, and the charge plateau was shifted to a higher

potential, which means the polarization increased when the

current density increased. A similar observation was also

reported by Zhou et al.3 The first discharge capacity recorded

was 320 mA h g21 at a rate of 50 mA g21 and the capacity was

315 mA h g21 for 100 mA g21, corresponding to the y1.7 Li+

insertion per mole of MoO3, but the capacity was decreased to

280 mA h g21 (equivalent to 1.5 Li+ per mole) for the first cycle

with a discharge rate of 200 mA g21. Almost 107 mA h g21

(y0.57 Li+/MoO3) capacity loss was observed in the initial 5

cycles, which was due to Li+ insertion in unrecoverable sites of

the [MoO6] octahedron, and the observations are consistent with

the CV results. In the remaining 95 cycles the capacity fading was

smaller (shown in Fig. 5b and c). A total 74% capacity (with

respect to the 5th cycle) was retained at the end of 100 cycles,

which is the best among all the reported results with any current

rate. Initially, charge–discharge performances for a-MoO3 on

Al, Ni and SS substrates are shown in Fig. 5c, and the result

reflects a similar conclusion as obtained from CV. After the

initial few cycles a huge capacity fading was observed for

a-MoO3–metal, except for the SS substrate, which means the Li+

reversibility was lost completely due to the complete loss of the

Fig. 5 (a) Charge–discharge curves as a function of rate for the first two cycles, (b) charge–discharge curve on SS substrate at 200 mA g21, (c) charge–

discharge cycles of a-MoO3 on different substrates at a current rate of 200 mA g21, (d) power cycle performance at different current rates.

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MoO3 structure and the formation of a large resistive interface.

The reason for this phenomenon is not yet fully understood. A

deep investigation is required to reveal the inside chemistry using

in situ techniques. When the Al substrate was used as the current

collector (Fig. 5c), the reversible capacity reaches almost zero

after the initial few cycles. By using the Ni current collector the

performance was improved a little in comparison to the Al

current collector. The observed capacity decreased gradually in

the initial 10 cycles and then stabilized for the remaining cycles

with a constant capacity of 52 mA h g21.

To know the electrochemical stability and the quality of

capacity retention, a power plot (shown in Fig. 5d) has been

made with different current densities starting from 100 mA g21,

200 mA g21, 300 mA g21, 400 mA g21 and 500 mA g21, and then

reverts back to 100 mA g21 then 200 mA g21 for the Li/

electrolyte/a-MoO3 half cell.

Fig. 5d shows that at high rate of about 500 mA g21 (y2.7C)

the material exhibits a discharge capacity of about 111 mA h g21.

The C rate is calculated for the capacity obtained due to 1 Li+

insertion in a formula unit of MoO3 in 1 h. Here, a rate of 1C is

equivalent to 186 mA h g21 capacity gain/loss in 1 h. After

reversing the rate to 100 mA g21, it shows a stable capacity of

around 170 mA h g21 and at the end of the 100th cycle it shows a

stable capacity of 128 mA h g21 with a rate of 200 mA g21

(which equals to a rate of y1.1 C). The above results show the

excellent capacity retention and robustness of the electrode

material at high as well as low rate performance.

In summary, this kind of work falls within the scope of

improving the electrochemical performance of the MoO3

material as a cathode in a lithium-ion battery. In addition to

previous findings in the reported literature i.e. (1) electrochemi-

cal performance enhancement by using lithiated MoO3 nanos-

tructures,34 (2) by making composites with conductive

polymers35 and (3) by imposing oxygen deficiency,5 the present

study provides a simple way to improve the electrochemical

performance by improving the electrode–current collector inter-

face. A common approach alluded to above is focusing on the

realization of better 2D nanostructured materials and a better

current collector interface for MoO3 in order to achieve the best

electrochemical performance. A real weakness of these materials

towards applications is irreversible capacity loss during the

initial stage of charge–discharge, poor capacity retention and

large polarization. Such limitations were suggested to be due to

irreversible structural change during reduction, low electronic

conductivity of phase pure a-MoO3 materials and kinetic

limitation related either to some intrinsic properties of materials

or non-compatibility with current collectors.

Conclusions

The present work emphasizes a few important points. It

highlights the advantages of simple preparation of anisotropic

a-MoO3-based cathode materials at low temperature using a one

step hydrothermal process. Very importantly, the electrochemi-

cal studies show that a-MoO3 material is more stable with a

stainless steel substrate in terms of capacity retention and rate

capabilities, while the reason behind the poor cycling perfor-

mance of aluminium current collectors was the formation of the

resistive interface. Here, we have improved the discharge

capacity as much as possible (140 mA h g21) and the capacity

is stable up to 100 cycles at a rate of 200 mA g21. During

electrochemical cycling against Li, we have also observed a two

phase reaction, and at the end of the reduction process loss of the

MoO3 structure is observed as in previous reports. Furthermore,

we have tried to explain the same phenomenon by electro-

chemical impedance spectroscopy. All the above results show

that the a-MoO3 material could be reconsidered as a suitable low

potential cathode for lithium-ion battery applications.

Acknowledgements

We thank the National Centre for Photovoltaic Research and

Education (NCPRE)–Ministry of New and Renewable Energy,

Govt. of India and IRCC-IIT Bombay for support. The authors

are thankful to the staff members of SAIF, IITB for their

assistance with electron diffraction and FEG-SEM analysis.

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