Superstructure ZrV2O7 nanofibres: Thermal expansion, electronic and lithium storage properties

Article abstract of DOI:10.1039/c6cp05694f

ZrV₂O₇ has attracted much attention as a negative thermal expansion (NTE) material due to its isotropic negative structure. However, rarely has investigation of the lithium storage behaviors been carried out except our first report on it. Meanwhile, the electrochemical behaviors and energy storage characteristics have not been studied in depth and will be explored in this article. Herein, we report on the synthesis, characterization and lithium intercalation mechanism of superstructure ZrV₂O₇ nanofibres that were prepared through a facile solution-based method with a subsequent annealing process. The thermal in situ XRD technique combined with the Rietveld refinement method is adopted to analyze the change in the temperature-dependent crystal structure. Benefiting from the nanostructured morphology and relatively high electronic conductivity, it presents acceptable cyclic stability and rate capability. According to the operando evolution of the XRD patterns obtained from electrochemical in situ measurements, the Li intercalation mechanism of the solid solution process with a subsequent conversion reaction can be concluded. Finally, the amorphous state of the electrodes after the initial fully discharged state can effectively enhance the electrochemical performances.

Full text of DOI:10.1039/c6cp05694f

View Article Online
View Journal
PCCP
Accepted Manuscript
This article can be cited before page numbers have been issued, to do this please use: Q. Li, Y. Zhao, Q.
Kuang, Q. Fan, Y. Dong and X. Liu, Phys. Chem. Chem. Phys., 2016, DOI: 10.1039/C6CP05694F.
This is an Accepted Manuscript, which has been through the
Royal Society of Chemistry peer review process and has been
accepted for publication.
Accepted Manuscripts are published online shortly after
acceptance, before technical editing, formatting and proof reading.
Using this free service, authors can make their results available
to the community, in citable form, before we publish the edited
article. We will replace this Accepted Manuscript with the edited
and formatted Advance Article as soon as it is available.
You can find more information about Accepted Manuscripts in the
Information for Authors.
Please note that technical editing may introduce minor changes
to the text and/or graphics, which may alter content. The journal’s
standard Terms & Conditions and the Ethical guidelines still
apply. In no event shall the Royal Society of Chemistry be held
responsible for any errors or omissions in this Accepted Manuscript
or any consequences arising from the use of any information it
contains.
www.rsc.org/pccp

Please do not adjust margins
Physical Chemistry Chemical Physics
Page 1 of 9
Page 2 of 14
View Article Online
DOI: 10.1039/C6CP05694F
Journal Name
ARTICLE
Superstructure ZrV O nanofibres: thermal expansion, electronic
2
7
and lithium storage properties
Received 00th January 20xx,
Accepted 00th January 20xx
a
b, c
b
b
b
a
Qidong Li, Yanming Zhao, Quan Kuang,* Qinghua Fan, Youzhong Dong and Xudong Liu
DOI: 10.1039/x0xx00000x
2 7
ZrV O has attracted much attention as a negative thermal expansion (NTE) material due to its isotropic negative structure.
However, rare investigation of the lithium storage behaviors has been carried out except our first report on it. Meanwhile,
the electrochemical behaviors and energy storage characteristics have not been in-depth studied and will be explored in
this project. Herein, we report on the synthesis, characterization and lithium intercalation mechanism of superstructure
www.rsc.org/
2 7
ZrV O nanofibres that were prepared through a facile solution-based method with subsequent annealing process.
Thermal in-situ XRD technique combined with Rietveld refinement method are adopted to analyze the change of
temperature-dependent crystal structure. Benefiting from the nanostructured morphology and relatively high electronic
conductivity, it presents acceptable cyclic stability and rate capability. According to the operando evolution of the XRD
patterns obtained from electrochemical in-situ measurement, the Li intercalation mechanism of solid solution process with
subsequent conversion reaction can be concluded. Finally, the amorphous state of the electrodes after initial fully
discharged state can effectively enhance the electrochemical performances.
1
2-14
Li-ion batteries (LIBs).
Among them the binary TMOs
Introduction
always suffered from poor conductivity and huge volume
change during electrochemical discharge/charge process for
Rechargeable lithium-ion batteries (abbreviated as LIBs, and
similarly hereinafter) with high capacity, long lifespan, no
memory effect, and environmental benignity have been
1
5-17
LIBs,
while the ternary MTMOs showed their potential in
synergistically enhancing the electronic/ionic conductivity,
1
reversible capacity and mechanical stability.
8-20
1
-7
extensively applied in portable electronic devices. However,
As one of the MTMOs family, transition metal vanadates
TMVs) with abundant sources and relatively high ionic
both the commercial used LiCoO cathode and graphite anode
2
(
encounter certain problem such as low-capacity which cannot
meet the growing demand for high-energy/power densities
applications like electric vehicles, plug-in hybrid electric
vehicles, and further large-scale grid storage. So, it is an
essential task to seek new materials with higher capacities to
meet the rapidly growing requirements of LIBs with higher
conductivity seem to be a superior group and have attracted
1
8, 21-35
many researchers’ interests in recent years.
the excellent performed anode materials such as Co
Apart from
1
2
2 2 7
V O ,
1
8, 22
23
24
Co
versatile potentials in cathode materials. Layered structured α-
CuV nanofibres were applied as a cathode material for LIB
by Chen et al., with high discharge capacities (447-514 mAh g
3
V
2
O
8
,
FeVO
4
,
Zn
3
V
2
O
8
,
etc., TMVs also show their
8
-9
2 6
O
energy/power densities.
Transition metal oxides (TMOs)
-1
including binary and ternary (also called the mixed transition
metal oxides (MTMOs)) ones with their unique morphologies
have attracted considerable interests in recent years because
of their promising applications in many fields such as
-1
at 20 mA g and 37 °C) and excellent high-rate capability
6
2
displayed.
3 8
Alkali metal vanadates Li1+xV O has been
researched for almost 40 years as a cathode material for its
2
7-30
1
0-11 higher capacity and better safety features.
Many lithiated
electrocatalysts for the oxygen reduction reaction,
x 2 5
Li V O phases occurring during the electrochemical reaction
electrode materials for electrochemical capacitors (ECs), and
+
-
+ xLi + xe = Li
V
2
O
5
x 2 5
V O indicate the multi-electron
transportation process, leading to the widespread study of
3
1-33
different Li
x
V
2
O
5
phases as cathode materials.
Zhang et al.
reported ε-Cu0.95V O
2 5
nanoribbons as a high-capacity cathode.
He found that the obtained materials could host 2.64 Li in a
reversible way, leading to a reversible capacity up to 292 mAh
-
1 34
g . Recently, the sol-gel synthesized ZrV
2 7
O micro-particles
had been evaluated as a Li-absent cathode material for
3
5
rechargeable Li batteries. The relatively high capacity and
acceptable cyclic ability reveal its potential use in the future.
While according to the reported TMVs electrodes, the
This journal is © The Royal Society of Chemistry 20xx
J. Name., 2013, 00, 1-3 | 1
Please do not adjust margins

Please do not adjust margins
Physical Chemistry Chemical Physics
Page 3 of 14
ARTICLE
Page 2 of 9
View Article Online
DOI: 10.1039/C6CP05694F
Journal Name
electrochemical performances have been found to be Materials characterization
influenced by the morphologies of the synthesized
8, 21-24
X-ray diffraction (XRD) data of ZrV
2
O
7
NFs for comparison and
1
materials].
As is known to us, different kinds of
Rietveld refinement analysis were collected on a Bruker D8
Advance diffractometer with Cu Kα radiation (λ= 1.5406 Å).
The XRD data acquired at temperature varied from room
temperature (RT) to 400 °C were obtained from a PANalytical
X’Pert Pro MPD diffractometer (Cu Kα radiation: λ= 1.5406 Å)
equipped with an auto-thermophilic sample stage.
Morphology and composition were analyzed by field-emission
scanning electron microscopy (FESEM, Nova NanoSEM430),
and energy dispersive spectrometer (EDS, Oxford instruments)
attached to the SEM.
nanostructures have been fabricated and investigated
chronically to relieve the strain and deformation of those
materials for flexible usages, mechanical engineering, or
battery electrodes with deformation before/after Li
insertion/extraction. Also, the importance of nanoscale
architecture had been proved in especially MTMOs on
enhancing their electronic and ionic transportation properties.
Downsizing materials to nanoscale level is an effective
approach to enhance the electrochemical performance due to
the increased surface areas and the shortened Li-ion diffusion
Electronic conductivity measurements
3
6-38
paths of the active material particles.
In this article, the traditional negative thermal expansion
NTE) material ZrV nanofibres (NFs) have been synthesized
2 7
Superstructure ZrV O NFs were pressed into pellets (10 mm in
diameter and ~1 mm in thickness) and annealed at 800 °C for 2
h under oxygen atmosphere to measure the direct current (dc)
electronic conductivity. The two sides of the pellets were
deposited with conductive silver epoxy to strengthen the
(
2 7
O
successfully by using a solution-based method under ambient
conditions for the first time. XRD data together with the
2 7
Rietveld refinement results indicates that the obtained ZrV O
2 7
contact between the ZrV O pellets and electrical silver wires.
nanofibres are in its superstructure form. The rate capability
and cyclic performance are also enhanced by using the
An AUTOLAB PGSTAT302N (Metrohm, Netherlands) was used
to acquire the I-V curves based on a cyclic voltammetry (CV)
obtained superstructure ZrV
2
O
7
nanofibres as cathode
-1
process with the scan rate of 0.1 V s in the voltage range of -
.0 V~1.0 V.
materials for Li-ion batteries. Furthermore, The Li intercalation
mechanism was explored and discussed.
1
Electrochemical measurements
LIR2032-type coin cells were assembled in a high purity argon-
filled glove box. Lithium metal was used as counter and
Experimental
2 7
reference electrode. Superstructure ZrV O NFs, conductive
Raw materials
acetylene black and polyvinylidene fluoride (PVDF) with the
weight ratio of 70:20:10 were blended in N-methyl-2-
pyrrolidone (NMP) solution to obtain a homogeneous slurry,
and then coated onto clean aluminum foil which served as a
current collector. After drying in an oven at 90 °C for 24 h, the
electrode was cut into small discs to serve as working
All chemicals, zirconium nitrate pentahydrate (Zr(NO
vanadium pentoxide (V ), oxalic acid (H ) and urea
CO(NH ) were used as received without further purification
3 4 2
) •5H O),
2
O
5
2 2 4
C O
(
2 2
)
to prepare the precursor. All aqueous solutions were freshly
prepared using double distilled water as solvent.
Synthesis of nanofibres-composed ZrV
A sol-gel method was used to synthesize the target product. In carbonate (DMC) (1:1, v/v) was used as the electrolyte and
a typical process, 0.4042 g V and 0.8362 g H in a Celgard 2320 microporous membrane as the separator.
stoichiometric ratio of 1:3 were firstly dissolved in 10 mL Galvanostatic discharge/charge tests were performed with the
2 7
O superstructure
6
electrodes. 1 M LiPF in ethylene carbonate (EC)-dimethyl
2
O
5
2 2 4
C O
-1
actively stirred deionized water at 60 °C until a clear blue current density of 10 mA g in the voltage range of 1.5- 4.0 V
+
solution was formed. The initial purpose to use oxalic acid was (vs. Li/Li ) using a LAND-CT2001A cell test system. The specific
to dissolve V
process can be described as the following formula:
+ 3H = VOC + 3H O + 2CO
Next, 0.8079 g Zr(NO •5H O in stoichiometric amount was electrochemical impedance spectroscopy (EIS) at the
2 5
O , namely the reducing agent. The reaction capacity was evaluated for the active material only. Cyclic
30
voltammetry (CV) potentiostatic in the voltage range of 1.5- 4
+
−1
V
2
O
5
2
C
2
O
4
2
O
4
2
2
↑
V (vs. Li/Li ) at a scanning rate of 0.1 mV s and
3
)
4
2
dissolved in 5 mL deionized water and then added dropwise to corresponding cell voltage in the frequency range from 100
the aforementioned solution. The V/Zr ratio was set as 2.35 in kHz to 0.01 Hz was carried out by using an AUTOLAB
which the V content was excess due to the easily vaporization PGSTAT302N (Metrohm, Netherlands) workstation. In-situ X-
of V in high temperature during annealing process. The ray diffraction experiments were performed using a specifically
detailed V/Zr ratio adjusting information can be found in the developed electrochemical cell formed with a beryllium
supplementary information (Fig. S1). 0.5 g of urea was used as window as current collector. The electrochemical
surfactant agent, and the mixture was vigorously stirred at 60 measurements were controlled by a LAND-CT2001A test
-
1
°
C until a sol formed. After drying at 90 °C in an air oven system. The cell was cycled at a current density of 10 mA g
-1
overnight, the gel was calcined at 800 °C for 2 h in a during discharging process and 20 mA g in following charging
continuously flux of oxygen. The calcination process can be procedure. All the above measurements were conducted at
expressed as follows:
VOC + 2Zr(NO
The whole process were conducted at atmospheric pressure.
room temperature.
4
2
O
4
3 4 2 2 7 2 2
) + O ↑= 2ZrV O + 8CO ↑+ 8NO ↑
2
| J. Name., 2012, 00, 1-3
This journal is © The Royal Society of Chemistry 20xx
Please do not adjust margins

Page 3 of 9
Physi cP al el aC s he ed mo ins ot tr ya dC j uh set mm i ac rag li nP shysics
Page 4 of 14
View Article Online
DOI: 10.1039/C6CP05694F
Journal Name
ARTICLE
2 7
Fig. 1 (a) (b) (c) SEM images of the as-synthesized superstructure ZrV O
nanofibres. (d) Point EDS result of the individual nanofibre in (c). The
quantitative analysis is shown inset.
Results and discussion
Typical SEM images of the as-synthesized superstructure
ZrV O nanofibres were shown in Fig. 1. The overview image of
2
7
Fig. 1(a) shows that the nanofibres have nearly uniform
diameter along their entire length, which stretch along several
micrometers. Moreover, these nanofibres are interconnected
into a porous network architecture. Closer observations (Fig.
Fig. 2 (a) XRD patterns of the as-synthesized superstructure ZrV
2 7
O nanofibres.
(b) Rietveld refinement of the samples. (c) The crystal structure of ZrV O
2 7
superstructure. (d)(e) Crystal structure of normal structure of ZrV
2 7
O with
different view directions.
1(b)&(c)) indicate the uniform smooth surfaces of the
nanofibres with width of 200- 400 nm. A representative X-ray
energy dispersive spectrometer (EDS) image of the point
marked with cross fork in Fig. 1(c) that taken from the FE-SEM
attachment is demonstrated in Fig. 1(d), in which the signals of
O, Si, V, Zr, Al and K can be observed. Zr/V/O signals reveal the
existence of desired products. The excrescent element signals
intensity of the peaks. Fig. 2(b) shows the Rietveld refinement
for ZrV O NFs carried out by using GSAS program via the
2
7
_
EXPGUI interface, and a space group of Pa 3 in cubic
superstructure was chosen as the refinement model. The
computational XRD pattern fits the experimental data points
very well as apparently indicated by the relatively small Rwp
(K, Si and Al) were from the mica plate used as loading
(
9.06%) and Rp (6.64%), which reveals single phase
substrates in SEM testing. Quantitative analysis of the EDS
result is presented as well (inset of Fig. 1(d)). The nearly same
Zr/V/O atomic proportion (8.37/18.73/64.26) to the
superstructure ZrV O NFs with no impurity phases detected
7
2
were obtained under the fabrication conditions. The crystal
structure of ZrV O superstructure consists of corner-sharing
2
7
2 7
stoichiometric ratio further prove the formation of ZrV O
VO (light blue) tetrahedral and ZrO (gray) octahedral has
4
6
sample. In order to further confirm the element proportion of
the synthesized sample, we did another experiment using
Elemental Analysis technique to serve as an extra evidence.
The results show that the weight percentage of Zr, V and O are
been shown in Fig. 2(c), from which VO double layers along
4
the b-direction can be clearly seen. The Zr cations lie in the
interlayer space, coordinating with nearest six oxygen atoms
to build ZrO octahedral chains, which are separated from each
6
30.03 %, 32.01 % and 37.96 %, which correspond to the atomic
other by VO tetrahedral double chains. Each oxygen of a ZrO6
4
percentage of 9.89 %, 18.85 % and 71.26 %. The atomic
percentage also coincide with the stoichiometric ratio of
octahedron is shared with a VO4 tetrahedron. Each VO4
tetrahedron shares three of its four O atoms with a ZrO6
octahedron, whilst the fourth is shared with another VO4
tetrahedron leading to a V O pyrovanadate group. Obviously,
2 7
ZrV O . Moreover, apart from the point EDS analysis of the
obtained sample, a large selected area of the fibrous region
was also applied for EDS quantitative analysis (See Fig. S2 in
supplementary information). The much rational results
compared to point EDS results indicate the homogenous and
2
7
the crystal structure of the ZrV O superstructure form would
7
2
be visually regarded as a 3×3×3 superposition of the cubic unit
cell of normal ZrV O as displayed in Fig. 2(c), (d)&(e).
2
7
2 7
successful synthesis of ZrV O .
However, the superstructure form shows much weaker
reflections according to the refinement results from single-
The phase composition and purity of the as-prepared
3
9
crystal synchrotron X-ray data of ZrV Specifically, two of
samples were investigated by using XRD technique. Both
_
2
O
7
.
2 7
_
standard XRD patterns of normal ZrV O (space group: Pa3,
the six V groups in superstructure ZrV O are constrained
2 7
2
O
7
ICSD #59396) and its superstructure form (space group: Pa3,
ICSD #84883) were presented for comparison with the
experimental result (shown in Fig. 2(a)). It can be clearly
distinguished that all the diffraction peaks of the obtained
by symmetry to contain linear V-O-V linkages, while the
remaining four are free to bend away from 180°. This is
significantly different from the normal ZrV
2 7
O structure of
2 7
which the six V O groups are constrained to be 180° on
_
account of the bridging O atom lies on a 3 site. The six linear V-
2 7
sample correspond to that of the ZrV O superstructure rather
than the normal form according to the position and relative
O-V linkages endow it a much higher structure symmetry than
This journal is © The Royal Society of Chemistry 20xx
J. Name., 2013, 00, 1-3 | 3
Please do not adjust margins

Please do not adjust margins
Physical Chemistry Chemical Physics
Page 5 of 14
Page 4 of 9
View Article Online
DOI: 10.1039/C6CP05694F
ARTICLE
Journal Name
3
.
9
superstructure ZrV
2
O
7
The lower symmetry of ZrV
2
O
7
thermal expansion phenomenon. All the patterns were
superstructure would affect the electron and Li-ion successfully indexed with a cubic lattice using the program
passageways along a, b, and c axes during discharge/charge Dicvol embedded in FullProf Suite, and the lattice parameters
process when applied as cathode materials for Li-ion batteries. were further least-square refined by the program PIRUM. The
Most obviously, the nonlinear V-O-V linkages of ZrV
2
O
7
obtained cell parameters versus temperature variation is
superstructure would result in zigzag Li-ion channels, thus plotted in Fig. 3(b), from which the positive thermal expansion
reducing the ionic diffusion rate. This phenomenon can be (PTE) (RT~120 °C) and NTE (above 120 °C) stages are obviously
clearly seen from the schematic V-O-V linkages of both normal presented. And the transformation temperature of about 120
and superstructure ZrV
2
O
7
illustrated in Fig. S3 (Supplementary °C is referred from the result, which differs to that of the
(about 102 °C) to some extent. The
Information). Nevertheless, the stability of the superstructure reported normal ZrV
2 7
O
would be better to buffer the volume change during Li ions morphology and structure may be responsible for the small
insertion/extraction because of the static disorder of oxygen difference. Except for the safety issue concerned in LIBs at
4
0-41
atoms in bent V-O-V linkages and the lower free energy.
higher temperature or self-heating during discharge/charge
Generally, ZrV O was highly noted as a compound that process, the nanofibres-composed ZrV O are also believed to
2
7
2 7
showed a bulk contraction in volume on heating in the past be significantly useful for thermal expansion areas used in
few decades, so-called negative thermal expansion (NTE) cooperation with positive thermal expansion materials due to
material. Attributing mainly to its relatively high-symmetry their tougher and stronger interfacial contact caused by larger
structure originated from NaCl structure type (Space group contact area of nanoscale size. As can be seen from Fig. S4 in
4
_
Fm3m) among which the cation is Zr and (V O ) radical acts the supplementary information, the linear thermal expansion
+
4-
2
7
3
as the anion, the ZrV O type structure is also unique in of superstructure ZrV O nanofibres is presented. Above 120
9
2
7
2 7
providing for isotropic negative structure over a very broad °C in NTE area, a mean coefficient of linear expansion of -
-6
-1
temperature range rather than the commonly anisotropic NTE 5.57×10 °C can be obtained.
4
1-42
materials.
present the claimed feature.
ZrV O7 was studied to be above 102 °C to
4
The electronic properties of the obtained superstructure
2
3-44
On the other hand, the ZrV O NFs would be of great importance to the application in
2
7
thermal-related structure change is also seemed as helpful for electrode materials for Li-ion batteries, in which the high
safety issue in Li-ion batteries because the shrinking Li-ion internal electronic conductivity can reduce the voltage
passageways derived from lattice contraction of NTE materials polarization and heat release. Furthermore, cathode materials
is expected to obstruct the diffusion of Li ions just like the with good conductivity can easily increase the electrochemical
function of separator in the situation of thermal runaway. The activity and benefit the charge transfer among the cathode
4
6-47
thermal expansion investigation of superstructure ZrV O NFs materials.
2
Hence, the electronic conductivity of
7
was therefore carried out from RT to 400 °C, attempting to superstructure ZrV O NFs pellet at temperatures ranging from
2
7
find the morphology-structure-property relationship of this 100 to 220 °C were studied. Fig. 4(a) illustrates the schematic
compound. One can see distinct peaks shift from Fig. 3(a), with diagram of the electronic conductivity measurement set-up.
their corresponding temperature labeled aside. The strongest The fast Ohmic response of the CV test can be shown as a
peak is magnified to demonstrate the change clearly. As the
temperature increases, the peaks move to lower angles,
indicating the augment of d-spacings according to Bragg
5
equation:
4
nλ ꢀ 2dꢁsinθ
Hence, the positive thermal expansion of the unit cell can be
referred at the beginning due to the cubic structure of ZrV O .
2
7
The peaks move backward to higher angles with the
continuously increasing temperature, revealing the negative
4
| J. Name., 2012, 00, 1-3
This journal is © The Royal Society of Chemistry 20xx
Please do not adjust margins

Page 5 of 9
Physi cP al el aC s he ed mo ins ot tr ya dC j uh set mm i ac rag li nP shysics
Page 6 of 14
View Article Online
DOI: 10.1039/C6CP05694F
Journal Name
ARTICLE
5
0-51
galvanostatic current-voltage process. I–V curves of the same conductivity.
In detail, in PTE period, during which the
2 7
ZrV O pellet in different temperature ranges are shown in Fig. scattering of phonon to transporting electrons become
4(b) and Fig. 4(c), respectively. All the curves seem to be linear, stronger with the increasing temperature, the electronic
indicating favourable Ohmic contacts formed between the conductivity is apparently influenced to reduce. When the
metal electrodes and the pellet. The linear fitting result shows temperature increases continuously over 120 °C, mean oxygen
-
5
-1
that the conductivity of ZrV
2 7
O at RT is 2.8×10 S cm , which is angles in the structure may open up, causing the large
much higher than that of many traditional cathode materials, amplitude transverse vibrations of the oxygens through
-
9
-1 48
(10 S cm ) and LiMn
-6 -1 49
(10 S cm ).
such as LiFePO
4
2
O
4
coupled librations of the tetrahedra and octahedra until a
4
0
From the I–V curves, it is seen that the electronic conduction higher symmetry state is reached. The higher symmetry with
of the pellet progressively decayed with increasing better periodic lattice field would lower the scattering
temperature from 100 to 120 °C, and was back to a rising probability of phonon to electrons. Meanwhile, the intrinsic
2 7
trend when the temperature was higher than 150 °C, which excitation of ZrV O at higher temperature can enhance the
comply a similar tendency with the structural evolution versus electronic conductivity as well. An analogous abrupt tendency
temperature (indicated in Fig. 3). It’s not difficult to connect of conductivity and lattice constant has also been observed in
52
the same trend with some causal relationship. In concrete other NTE materials such as Mn Ga0.75Si0.25N, Mn xSn1
−
xC (x
3 y 1
−
≤ 0.1) and Mn Zn Sn yC (y= 0.4, 0.5) compounds, in which
3
3+
5
3
terms, the ZrV O superstructure started to inflate with the
2
7
positive thermal expansion process beneath 120 °C, in which the magnetic transition are primarily responsible for their
the length of bonds were enlarged, resulting in the increasing unusual properties. Further exploration need to be conducted
atom distance and transport distance. Electronic transport in the future whether the same reason can be used to explain
along lager distance generates a diminished electronic the abnormal phenomenon. The conductivity, σ, of the
4
7
conductivity on this occasion. When the temperature was superstructure ZrV O NFs pellet is plotted in Fig. 4(d) as ln(σT)
2
7
increased to the NTE period, an opposite trend was shown in versus reciprocal temperature in the range of 100~220 °C. This
the same way. As a quasi-rigid unit modes (QRUMs) termed plot is used to analyze the data in the frame of the general
mechanism that display phonon-driven contraction material, formula for electronic conductivity of transition metal oxides
5
4
the vibrations of V-O-V linkages are mainly responsible for its proposed by Mott, where the conductivity is given by:
claimed NTE effect and the abrupt change of electronic
This journal is © The Royal Society of Chemistry 20xx
J. Name., 2013, 00, 1-3 | 5
Please do not adjust margins

Please do not adjust margins
Physical Chemistry Chemical Physics
Page 7 of 14
Page 6 of 9
View Article Online
DOI: 10.1039/C6CP05694F
ARTICLE
Journal Name
ꢇ
ꢃꢄ ꢆ ꢈꢉꢃ1 ꢊ ꢈꢉ
ꢏ
ꢋꢍ
storage performances. The rate capability at different current
densities of the superstructure ZrV O NFs electrodes were
ꢅ
σ ꢀ ꢂ
ꢎexp ꢁꢃ ꢊ2αRꢉexp ꢁꢃ ꢊ
ꢏ
ꢉ
ꢋꢌꢍ
2
7
or
carried out in ambient temperature to further investigate their
application possibility. As shown in Fig. 5(b), the test was set at
lnσT ꢀ lnA ꢊ
ꢋꢍ
is a phonon frequency, C the concentration, R the
-1
current densities of 10 mA g , 20 mA g , 50 mA g , 100 mA g
-1
-1
-1
where υ
0
-1
-1
and 200 mA g every 5 cycles, and turned back to 10 mA g in
average hopping distance, W the activation energy, and α the
5
rate of wave function decay. The two periods can be linearly
fitted with different slope, indicating their conformation to the
5
succession. The rate performance seemed to be much
35
impressive than that of the normal structured ZrV
2
O
7
.
The
better rate capability should be ascribed to the nanoscale
architecture of ZrV , which can shorten the Li diffusion
2 2 7
typical Mott law. As can be analogized to the reported Ca V O
5
6
2 7
O
single crystal,
the conduction mechanism should be
distance in the bulk phase, increasing the electrode-electrolyte
interface and availing the electrolyte penetration into the
inter-connected electrode. The larger specific surface area
makes more active material utilized at a high discharge/charge
rate, thus a high specific capacity can be expected. Besides, the
better electronic conductivity (Fig. 4) contributes to the rate
performance as well. Cyclic performance was explored at the
attributed only to the electron conduction that stems from the
vanadium ions at lower temperatures (< 102 °C). And the
4
addition of predominantly ionic conduction by mobile Zr or
+
2
O
-
ions above 102 °C is the reason for the steeper slope
shown in Fig. 4(d) in the corresponding temperature range.
-1
Representative discharge/charge curves at 10 mA g for
the 1st, 2nd and 10th cycles are depicted in Fig. 5(a) within a
-1
+
voltage window of 1.5~4 V vs. Li/Li . The initial specific
current density of 10 mA g . As shown in Fig. 5(c), both charge
-1
-1 and discharge capacities can retain at about 135 mAh g
-1
discharge and charge capacity of 279 mAh g and 208 mAh g
can be achieved, which are much higher than that of most
state-of-the-art Li-contained cathode materials. The resulting
columbic efficiency of 74.6% indicates the irreversible
conversion reaction and inert reaction products formed similar
(
64.5% of the reversible capacity of 1st charge) after 30 cycles.
Although the cycle performance are not so impressive
compared to some typical cathode materials on account of the
serious voltage polarization, the properties are clearly
improved by the nanostructured technique when compared
5
7
to most conversion electrode materials. However, the values
of charge and discharge capacity are slightly lower than the
normal structured ZrV O reported by our group recently
3
5
with that of the micron-sized ZrV O . The columbic efficiency
2
7
(CE), calculated by dividing the charge capacity by discharge
2
7
-1
-1
capacity, increased rapidly from 74.6% for the first cycle to
99.1% for the second one. During the subsequent cycles, it
maintains at a high level of 98%, demonstrating a good
reversibility between the charge and discharge processes. It
can be seen that some CE data points are above 100% which
can be attributed to the slight variation of temperature during
(
282.6 mAh g and 227.7 mAh g for discharge and charge
3
5
respectively). The reason would be related to the lower
symmetry of ZrV O superstructure, which has nonlinear V-O-V
2
7
linkages induced zigzag Li-ion channels. CV curves in Fig. 5(a)-
inset reveal similar redox peaks to the dQ/dV results in Fig. S5.
As detailed presented, the three main reduction peaks located
at 2.37 V, 2.26 V and 1.94 V in first cathodic process entirely
correspond to the differential capacity versus potential
6
0
st
nd
cycling. The Nyquist plots at open circuit state, after 1 , 2 ,
th
th
and 20 cycles (Fig. 5(d)) present similar semicircles and
1
0
quasi-straight lines. Fig. 5(d)-inset is the corresponding
2
(
dQ/dV) curves, in which the main plateau around 2.5 V vs.
6
+
Li/Li , namely, the sharp peak centered at
equivalent circuit model for the impedance response.
Following the usual practice, R stands for the combination of
∼2.43 V in the
5
dQ/dV results, is attributed to the reduction of V to V and
+
4+
s
4
+
3+
electrolyte resistance and ohmic resistance of cell
components. The semicircle could be attributed to the charge
partial reduction of V to V through typical conversion
reactions according to the previously reported α-CuV O
2
6
2
nanofibres. The other small peaks should be ascribed to the
6
4
+
3+
fully reduction of V to V together with the intercalation
process of lithium. The splitting of the redox peaks was
originated from the different lithium sites with small energy
5
8-59
difference for holding the inserted Li-ions.
The CV and
dQ/dV results consistently affirm the plateau assignments of
the initial discharge curve. The small difference of the peaks
locations between dQ/dV and CV result comes from the
voltage polarization at different of galvanostatic
discharge/charge
and
potentiostatic
CV
testing
35
conditions. There are no obvious oxidation peaks can be
found in the anodic process of differential capacity curves, the
relatively broad swell should be caused by the extraction of Li-
x 2 7
ions from the Li ZrV O matrix. The subsequent discharge and
charge process show the similar profiles in view of the
discharge/charge curves and differential capacity of the 2nd to
10th cycles based on Fig. 5(a), revealing good reversibility of Li
6
| J. Name., 2012, 00, 1-3
This journal is © The Royal Society of Chemistry 20xx
Please do not adjust margins

Page 7 of 9
Physi cP al el aC s he ed mo ins ot tr ya dC j uh set mm i ac rag li nP shysics
Page 8 of 14
View Article Online
DOI: 10.1039/C6CP05694F
Journal Name
ARTICLE
transfer resistance of the electrode reaction Rct and constant lattice parameter is consistent with its volume change. When
+
phase element (CPE) of the electrode/electrolyte interface, the solid-solution limit of Li in ZrV
2
O
7
superstructure was
reached with further insertion, no prominent peaks were
, Warburg detected in the as-defined stage-2 except the diffraction peaks
impedance). The specific values (Table S1) can be obtained by of BeO under the resolution of the diffractometer, revealing
data fitting according to the equivalent circuit model in Fig. that the Li ZrV solid solution in stage-1 totally decomposed
(d)-inset. As can be seen from Table S1, for the fresh into an amorphous state via conversion reaction. Even after
electrode, the charge transfer resistance value is 502.5 Ω. recharged to 4.0 V, the superstructure or normal form of
+
and the inclined line in the low frequency region to Li
diffusion processes within the active material (Z
w
x
2 7
O
5
st
nd
After 1 and 2 cycle, the values of Rct decrease to 235.3 Ω ZrV
th
2
O
7
were no longer recovered from the non-crystalline
th
and 166.8 Ω. And to 140.53, 87.39 after 10 cycle and 20
composite. Namely, despite modification of the coordination
cycle, respectively. Change of charge transfer resistance was environment of the local vanadium or zirconium ligands during
reported as the important factor to capacity fading and rate the following discharge/charge process, the electrodes of
6
1-62
performance.
2 7
Specifically, lower Rct can upgrade the rate superstructure ZrV O NFs still stay in an amorphous state
performance and cycling stability. The decreasing Rct upon throughout the whole discharge/charge process in subsequent
cycling reflect a better electronic contact of the electrode cycles. The in-situ XRD detection results unfortunately negates
superior to the initial electrode, indicating a decrease in inner the aforementioned thermal safety advantage in virtue of its
6
3
resistance after cycling. With lithiation and de-lithiation, the NTE characteristic under higher temperature. The excellent
nanofibres constitution of the active material would turn into electronic conductivity of ZrV is also beyond hope after the
2 7
O
increasingly inter-connected and densely packed structure, collapse of the initial structure. However, from the EIS results
which can enhance the inter-particle contact and electrical we can see the well-retained conductivity by the decomposed
6
4
s x x
contact. While the plot and values of R did not change much, VO and ZrO , which is still helpful for the electrochemical
revealing the electronic conductivity and ion diffusion performance. Much deeper research for detailed reason why
pathways in the composite electrodes are well maintained the final amorphous state of the electrode materials exist had
after several cycles, leading to a good cycling stability during already been demonstrated by Tarascon et al. They
the chemical process. The values of R
s
and Rct are also similar investigated copper vanadium oxides (CVOs) with different
nanofibres and stoichiometric constitution (Cu
indicating the suitability of TMVs in Cu2.33 11) as cathode materials for LIBs.
application of electrode materials. The above results Though with entirely same anions and cations, different Li
demonstrated that the superstructure ZrV nanofibres can intercalation mechanisms are presented by them on account
serve as a promising cathode material for lithium-ion battery of the diverse vanadium sites derived from crystal structures.
after optimization. And the further optimization strategy to In detail, ε-Cu0.95 11 can react
enhance the electrochemical performance may lie in the with Li through a reversible combination of displacement and
better morphology construction and electronic conductivity intercalation process. While for Cu 10 and Cu , the
promotion such as carbon coating. initial phase would convert into amorphous state upon
To better understand the reaction mechanism and the electrochemical discharge. Structural considerations involving
structural change of ZrV during electrochemical Li cation local stability and long-range crystal stability are
insertion/extraction between 1.5 and 4.0 V vs. Li/Li , the initial responsible for the significant difference. They found the
discharge and charge processes were investigated by in-situ vanadium and oxygen in Cu 10 and Cu bond as VO
XRD detection (Fig. 6). At the beginning of the reaction, only tetrahedrons. Because the V and V ions are not stable in a
to other TMVs electrodes such as α-CuV
2
O
6
5 2 2 2 7 2 5
V O10, Cu V O , ε-Cu0.95V O ,
26, 65
34, 66-67
CoV
2
O
6
nanosheet,
4 4
V O11 and Cu1.1V O
2
O
7
2 5 4 4
V O , Cu2.33V O11 and Cu1.1V O
5
V
2
O
2 2 7
V O
2 7
O
+
5
V
2
O
2
V
2
O
7
4
4
+
3+
4+ 3+
5
+
2 7
the well-defined peaks of crystalline ZrV O are visible in the tetrahedral environment, the V /V /V reduction process
diffraction pattern at the open-circuit voltage (OCV) of 3.12 V would involve a change in vanadium coordination, leading to
+
vs. Li/Li . We mark the whole discharging process into two the noncrystallization of the initial structure. We know that the
2 7
main stages, of which the stage-1 is from OCV to 1.65 V vs. vanadium coordination in ZrV O also lie in tetrahedral sites
+
+
Li/Li and 1.65~1.5 V vs. Li/Li belongs to stage-2. During the from Fig. 2 (c), which is in consistent with the final amorphous
stage-1 of the initial discharge, the intensities of these peaks phase after electrochemical discharge. However, from the in-
2 7
decrease progressively without observing any new peaks of situ XRD detection of ZrV O during discharge, we found that
second phase in the diffraction patterns. Note that the peaks before the conversion reaction to ultimate amorphous state,
positions shift to lower angles and gradually become broader the electrode undergo a solid-solution process during which
and weaker (clearly shown in the magnification of the highest the Li ions are inserted into ZrV
peak of ZrV superstructure), which indicates the start of a collapse. This phenomenon is different from that of the
Li ZrV solid-solution reaction with subsequent conversion reported ones. In addition, the amorphous nature of cycled
reaction which gradually degraded the crystallinity and/or electrode that indicates the structural consistency has been
reduced the crystalline size of ZrV . The shift of the peaks, proved to be in favor of effectively preventing the lattice stress
which should be attributed to the insertion process of Li ions as well as buffering the huge volumetric change during the Li
intercalation into the interstitial sites of ZrV matrix, implies ions insertion/extraction and providing continuous Li ions
the expansion of the ZrV structure. On account of the cubic diffusion pathways and insertion sites, which greatly improve
nature of ZrV superstructure, the linear evolution of the the reversible capacity and sustain the structural stability.
2 7
O matrix without structural
2 7
O
x
2 7
O
2
O
7
2 7
O
2 7
O
6
8-69
2 7
O
This journal is © The Royal Society of Chemistry 20xx
J. Name., 2013, 00, 1-3 | 7
Please do not adjust margins

Please do not adjust margins
Physical Chemistry Chemical Physics
Page 9 of 14
Page 8 of 9
View Article Online
DOI: 10.1039/C6CP05694F
ARTICLE
Journal Name
1
8 G. Z. Yang, H. Cui, G. W. Yang and C. X. Wang, ACS Nano,
014, , 4474.
Conclusions
2
8
In summary, we have fabricated ZrV
2
O
7
nanofibres via a facile 19 B. Liu, J. Zhang, X. F. Wang, G. Chen, D. Chen, C. W. Zhou and
G. Z. Shen, Nano Lett., 2012, 12, 3005.
solution-based route. The homogeneous 1D fibrous
morphology after calcination can widen and upgrade their NTE
and electrochemical performances. The XRD pattern combined
2
0 H. S. Jadhav, R. S. Kalubarme, C. N. Park, J. Kim and C. J. Park,
Nanoscale, 2014, , 10071.
1 F. F. Wu, C. H. Yu, W. X. Liu, T. Wang, J. K. Feng and S. L.
Xiong, J. Mater. Chem. A, 2015, , 16728.
6
2
2 7
with Rietveld refinement reveals that the obtained ZrV O NFs
3
are crystallized in a 3×3×3 superstructure form. As a 22 F. F. Wu, S. L. Xiong, Y. T. Qian and S. H. Yu, Angew. Chem.
Int. Ed., 2015, 54, 10787.
conventional negative thermal expansion material, the
2
2
3 D. H. Sim, X. Rui, J. Chen, H. Tan, T. M. Lim, R. Yazami, H. H.
Hng and Q. Yan, RSC Adv., 2012, , 3630.
4 L. H. Gan, D. Deng, Y. Zhang, G. Li, X. Wang, L. Jiang and C. R.
Wang, J. Mater. Chem. A, 2014, , 2461.
nanoscale form would highly widen its use in corporation with
other positive thermal expansion materials on account of the
tougher interfacial contaction between them. After being
2
2
evaluated as a cathode material for Li-ion batteries, the 25 L. X. Zeng, X. X. Huang, C. Zheng, Q. R. Qian, Q. H. Chen and
M. D. Wei, Dalton Trans., 2015, 44, 7967.
enhanced rate capability and cyclic performance are
2
2
6 H. Ma, S. Y. Zhang, W. Q. Ji, Z. L. Tao and J. Chen, J. Am.
Chem. Soc., 2008, 130, 5361.
7 G. Pistoia, M. Pasquali, M. Tocci, V Maney and R. V. Moshtev,
J. Power Sources, 1985, 15, 13.
impressing on account of the nanoscale effect. In-situ XRD
analysis preliminarily demonstrates the Li intercalation
mechanism of this compound should be a solid solution
process with subsequent conversion reaction to an ultimate 28 S. Panero, M. Pasquali and G. Pistoia, J. Electrochem. Soc.,
983, 130, 1225.
9 J. Dai, S. F. Y. Li, Z. Q. Gao and K. S. Siow, J. Electrochem. Soc.,
998, 145, 3057.
1
amorphous state. Combining major merits of this
superstructure form ZrV NFs such as the high capacity and
2
3
2 7
O
1
electronic conductivity, this functional material possesses
bright prospects in developing advanced materials for the next
generation NTE application and LIBs through rational and
delicate design.
0 A. Q. Pan, J. Liu, J. G. Zhang, G. Z. Cao, W. Xu, Z. M. Nie, X. Jie,
D. Choi, B. W. Arey, C. M. Wang and S. Q. Liang, J. Mater.
Chem., 2011, 21, 1153.
1 W. J. Wang, H. Y. Wang, S. Q. Liu and J. H. Huang, J Solid
State Electrochem., 2012, 16, 2555.
2 M. B. Smirnov, E. M. Roginskii, V. Y. Kazimirov, K. S. Smirnov,
R. Baddour-Hadjean, J. P. Pereira-Ramos and V. S. Zhandun,
J. Phys. Chem. C, 2015, 119, 20801.
3
3
Acknowledgements
3
3
3
3
3
3
3
4
4
4
4
4
3 J. C. Wang, G. H. Gao, X. W. Zhou, J. D. Wu, H. Y. Yang, Q. Li
and G. M. Wu, J. Solid State Electrochem., 2014, 18, 2459.
4 W. Hu, X. B. Zhang, Y. L. Cheng, Y. M. Wu and L. M. Wang,
Chem. Commun., 2011, 47, 250.
5 Q. Kuang, Y. M. Zhao, Y. Z. Dong and Q. H. Fan, Electrochim.
Acta, 2015, 170, 229.
This work was funded by the NSFC Grant (Nos. 51402107 and
51372089) supported from the NSFC Committee of China.
Notes and references
6 Q. F. Zhang, E. Uchaker, S. L. Candelaria and G. Z. Cao, Chem.
Soc. Rev., 2013, 42, 3127.
1
J. M. Tarascon and M. Armand, Nature, 2001, 414, 359.
M. Armand and J. M. Tarascon, Nature, 2008, 451, 652.
R. V. Noorden, Nature, 2014, 507, 26.
2
3
4
7 C. Liu, F. Li, L.P. Ma and H. M. Cheng, Adv. Mater., 2010, 22
E28.
,
B. Dunn, H. Kamath and J. M. Tarascon, Science, 2011, 334
,
8 P. G. Bruce, B. Scrosati and J. M. Tarascon, Angew. Chem. Int.
Ed., 2008, 47, 2930.
9 J. S. O. Evans, J. C. Hanson and A. W. Sleight, Acta Crystallogr.
B, 1998, 54, 705.
0 G.D. Barrera, J. A. O. Bruno, T. H. K. Barron and N. L. Allan, J.
Phys.: Condens. Matter., 2005, 17, R217.
1 N. Khosrovani, A. W. Sleight and T. Vogt, J. Solid State Chem.,
928.
H. Wu, G. Chan, J. W. Choi, Y. Yao, M. T. McDowell, S. W. Lee,
A. Jackson, Y. Yang, L. Hu and Y. Cui, Nat. Nanotechnol.,
5
6
2
012,
L. Q. Mai, X. C. Tian, X. Xu, L. Chang and L. Xu, Chem. Rev.,
014, 114, 11828.
7, 310.
2
7
8
T. Sasaki, Y. Ukyo and P. Novák, Nat. Mater., 2013, 12, 569.
Y. Yao, M. T. McDowell, I. Ryu, H. Wu, N. Liu, L. Hu, W. D. Nix
and Y. Cui, Nano Lett., 2011, 11, 2949.
1
997, 132, 355.
2 V. Korthuis, N. Khosrovani and A.W. Sleight, Chem. Mater.,
995, , 412.
1
7
9
1
1
B. Jang, M. Park, O. B. Chae, S. Park, Y. Kim, S. M. Oh, Y. Piao
and T. Hyeon, J. Am. Chem. Soc., 2012, 134, 15010.
0 F. Y. Cheng, J. Shen, B. Peng, Y. D. Pan, Z. L. Tao and J. Chen,
3 Q. Q. Liu, X. N. Cheng, X. J. Sun, J. Yang and H. H. Li, J. Sol−Gel
Sci. Technol., 2014, 72, 502.
4 H. L. Yuan, B. H. Yuan, F. Li and E. J. Liang, Acta Phys Sin.,
2012, 61, 226502.
5 C. G. Pope, J. Chem. Educ., 1997, 74, 129.
Nat. Chem., 2011, 3, 79.
1 Y. Y. Liang, H. L. Wang, J. G. Zhou, Y. G. Li, J. Wang, T. Regier
and H. J. Dai, J. Am. Chem. Soc., 2012, 134, 3517.
2 Y. M. Chiang, Science, 2010, 330, 1485.
4
4
6 Q. Kuang, Y. M. Zhao and J. T. Xu, J. Phys. Chem. C, 2011,
15, 8422.
1
1
1
1
3 J. R. Miller, P. Simon, Science, 2008, 321, 651.
4
4
7 M. Park, X. C. Zhang, M. Chung, G. B. Less and A. M. Sastry, J.
Power Sources, 2010, 195, 7904.
4 G. Q. Zhang and X. W. Lou, Angew. Chem. Int. Ed., 2014, 126
187.
5 C. X. Peng, B. D. Chen, Y. Qin, S. H. Yang, C. Z. Li, Y. H. Zuo, S.
Y. Liu and J. H. Yang, ACS Nano, 2012, , 1074.
6 X. H. Wang, Y. Fan, R. A. Susantyoko, Q. Z. Xiao, L. M. Sun, D.
Y. He and Q. Zhang, Nano Energy, 2014, , 91.
7 H. Wu, M. Xu, Y. C. Wang and G. F. Zheng, Nano Res., 2013,
, 167.
,
9
8 S. Q. Shi, L. J. Liu, C. Y. Ouyang, D. S. Wang, Z. X. Wang, L. Q.
Chen and X. J. Huang, Phys. Rev. B, 2003, 68, 195108.
9 J. Molenda and W. Kucza, Solid State Ionics, 1999, 117, 41.
0 A. P. Giddy, M. T. Dove, G. S. Pawley and V. Heine, Acta
Cryst. A, 1993, 49, 697.
1
1
1
6
4
5
5
6
8
| J. Name., 2012, 00, 1-3
This journal is © The Royal Society of Chemistry 20xx
Please do not adjust margins

Page 9 of 9
Physi cP al el aC s he ed mo ins ot tr ya dC j uh set mm i ac rag li nP shysics
Page 10 of 14
View Article Online
DOI: 10.1039/C6CP05694F
Journal Name
ARTICLE
5
1 A. K. A. Pryde, K. D. Hammonds, M. T. Dove, V. Heine, J. D.
Galex and M. C. Warren, J. Phys.: Condens. Matter, 1996,
0973.
8,
1
5
5
2 Y. Sun, C. Wang, Y. C. Wen, L. H. Chu, M. Nie and F. S. Liu, J.
Am. Ceram. Soc., 2010, 93, 650.
3 Y. B. Li, W. F. Li, W. J. Feng, Y. Q. Zhang and Z. D. Zhang, Phys.
Rev. B, 2005, 72, 024411.
4 N. F. Mott, J. Non−Cryst. Solids, 1968,
1, 1.
5 J. Muster, G. T. Kim, V. Krstic, J. G. Park, Y. W. Park, S. Roth
and M. Burghard, Adv. Mater., 2000, 12, 420.
6 A. F. G. de I'Eprevier and R. C. Buchanan, J. Electrochem. Soc.,
5
5
5
5
5
5
6
6
6
6
1
982, 129, 2562.
7 M. J. Aragón, B. León, C. P. Vicente and J. L. Tirado, J. Power
Sources, 2011, 196, 2863.
8 M. V. Reddy, B. L. W. Wen, K. P. Loh and B. V. R. Chowdari,
ACS Appl. Mater. Interfaces, 2013, 5, 7777.
9 A. M. Cao, J. S. Hu, H. P. Liang and L. J. Wan, Angew. Chem.
Int. Ed., 2005, 44, 4391.
0 L. F. Zhang, J. Zhou and C. Y. Zhang, J. Mater. Chem. A, 2014,
2, 14903.
1 D. C. Li, Y. K. Sasaki, K. Kobayakawa and Y. Sato, J. Power
Sources, 2006, 157, 488.
2 Y. K. Sun, J. M. Han, S. T. Myung, S. W. Lee and K. Amine,
Electrochem. Commun., 2006, 8, 821.
3 Y. Yu, C. J. Niu, C. H. Han, K. N. Zhao, J. S. Meng, X. M. Xu, P. F.
Zhang, L. Wang, Y. Z. Wu and L. Q. Mai, Ind. Eng. Chem. Res.,
2
016, 55, 2992.
4 Y. R. Zhong, M. Yang, X. L. Zhou, Y. T. Luo, J. P. Wei and Z.
Zhou, Adv. Mater., 2015, 27, 806
5 L. Zhang, K. N. Zhao, Y. Z. Luo, Y. F. Dong, W. W. Xu, M. Y.
Yan, W. H. Ren, L. Zhou, L. B. Qu and L. Q. Mai, ACS Appl.
6
6
Mater. Interfaces, 2016,
6 M. Morcrette, P. Rozier, L. Dupont, E. Mugnier, L. Sannier, J.
Galy and J. M. Tarascon, Nat. Mater., 2003, , 755.
8, 7139.
6
2
6
6
7 F. Y. Cheng and J. Chen, J. Mater. Chem., 2011, 21, 9841.
8 D. Zhao, L. R. Zheng, Y. Xiao, X. Wang and M. H. Cao,
ChemSusChem, 2015,
9 C. Li, X. Miao, W. Chu, P. Wu and D. G. Tong, J. Mater. Chem.
A, 2015, , 8265.
8, 2212.
6
3
This journal is © The Royal Society of Chemistry 20xx
J. Name., 2013, 00, 1-3 | 9
Please do not adjust margins