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Cite this: J. Mater. Chem. A, 2015, 3,
509
Received 2nd October 2014
Accepted 5th November 2014
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Reaction pathway and wiring network dependent
Li/Na storage of micro-sized conversion anode
with mesoporosity and metallic conductivity†
Zhonghui Cui,‡a Chilin Li,‡*a Pengfei Yu,a Minghui Yang,*b Xiangxin Guo*a
and Congling Yinc
DOI: 10.1039/c4ta05241b
www.rsc.org/MaterialsA
Micro-sized or monolithic electrode materials with sufficient mesoporosity and a high intrinsic conductivity are highly desired for highenergy batteries without the trade-off of electrolyte infiltration and
accommodation of volume expansion. Here metallic nitrides consisting of mesoporous microparticles were prepared based on a
mechanism of solid–solid phase separation and used as conversion
anodes for Li and Na storage. Their superior capacity and rate
performance during thousands of cycles benefit from the preservation
or self-reconstruction of hierarchically conductive wiring networks.
The conversion efficiency is also highly dependent on the reaction
pathway and product. Exploring more conductive and percolating
mass/charge transport networks particularly in a deep sodiation state
is a potential solution for activation of Na-driven conversion
electrochemistry.
In the pursuit of high performance anodes for Li-ion batteries
(LIBs), numerous material prototypes and related nanotechnologies have been attempted to construct a desired (micro)
structure in these years.1 Their electrochemistry and electroactivity signicantly depend on the corresponding reaction
mechanisms. Among them, carbonaceous (e.g. graphite) and Tibased (e.g. TiO2 and Li4Ti5O12) materials usually undergo
insertion reaction sometimes concomitant with interfacial
storage or (pseudo)capacitive charging.2,3 These materials
enable a high-rate energy storage but with a limited reversible
capacity (<300 mA h g1) in view of their structural integrity
during cycling. In order to substantially improve the capacity
a
State Key Laboratory of High Performance Ceramics and Superne Microstructure,
Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050,
China. E-mail: [email protected]; [email protected]
b
Dalian National Laboratory for Clean Energy, Dalian Institute of Chemical Physics,
Chinese Academy of Sciences, Dalian 116023, China. E-mail: [email protected]
c
College of Material Science and Engineering, Jiangxi University of Science and
Engineering, Ganzhou 341000, China
† Electronic supplementary information (ESI) available: Experimental section and
Fig. S1 to S10 are given. See DOI: 10.1039/c4ta05241b
‡ These authors contributed equally to this work.
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performance, it requires to break certain bonding in the original structure and create new bonds with Li, as already shown in
alloying (e.g. for Si) and conversion (e.g. for FeOx) reactions.4,5
Multi-electron transfer oen occurs in these phase transformation processes (e.g., by utilizing all the oxidation states of
redox metal cation during conversion reaction), therefore
resulting in considerably higher reversible capacities near or
more than 1000 mA h g1. However, in most of the cases,
remarkable volume expansion and morphology evolution
appear to be inevitable due to repeated bonding/debonding
with a volume of Li atoms. Comparing the huge volume
expansion ratio (up to 400%) for a typical alloying reaction of
Si,4 for the conversion reaction the ratio is considerably smaller
(<200%) and therefore could alleviate electrode cracking and
delamination.6 A main challenge for converting oxides or uorides lies in the preservation of the electron/ion mixed
conductive network, because less conductive Li2O or LiF
nanodomains are always produced when lithiation starts and
then they probably aggregate to form a high-area insulating
matrix aer deep discharging.5,7
Since the pioneering work by Poizot et al. in 2000,8 nanosizing of transition metal oxides (MOx) or uorides (MFy)
simultaneously combining with suitable electron wiring (e.g., in
situ carbon decoration) has been widely investigated to improve
their conversion efficiency. This strategy is benecial to mass
and charge transports of the discharged active species (metal M
and Li2O or LiF), which are conned in small scale, as shown in
previous reports on nanoparticle or thin lm samples.8,9 Transition metal nitrides (MNz) as another promising candidate of
conversion materials display two extra advantages over their
oxide or uoride counterparts: (1) discharged product Li3N is a
lithium superionic conductor (6 103 S cm1 at room
temperature) and can serve as additional ion wires;10 (2) nitrides
themselves are considerably more electronically conductive
than most of oxides and uorides.11 As a consequence, nitrides
are expected to require a less amount of carbon additives
without the cost of serious degradation of mixed conductive
networks during the conversion reaction. Apart from enhancing
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electron conductivity, a nitridation strategy is also useful to
increase ion conductivity by inducing lattice defects or modifying structure moieties as shown in LiPON solid electrolyte or
the recent Na3TiP3O9N electrode, respectively.12 However,
nitridation technologies used to construct well-dened nanostructures are still very limited. Previous works on conversion of
nitrides mainly focused on the samples either in the form of
thin lms or consisting of discrete nanoparticles, unfavorably
leading to mass lack or a low density of active species.13
Furthermore, nanoparticle powders are easy to diffuse and thus
are unsafe for physical health and industrial production.
Building mesoporosity in high-density micro-sized or monolithic materials is a desired solution to these aforementioned
issues without the trade-off of electrolyte inltration and
accommodation of volume expansion.
Most recently, room temperature Na-ion batteries (NIBs)
have shown competitive potential in the application elds of
large-scale transportation and stationary energy storage, owing
to the abundant reserve of sodium in nature and its low cost.14
However, many reaction mechanisms and their corresponding
structure prototypes, which have been successfully applied in
LIBs, cannot be extended to NIBs because of the 34% larger ion
radius of Na+ (1.02 Å) than that of Li+ (0.76 Å).15 Apart from the
strategy of open framework to discover new Na-storage materials,16 an exquisite electrode design is universally required to
alleviate the possibly more signicant volume expansion and
electric contact loss during sodiation (i.e., to achieve a robust
ion/electron wiring network). The latter method is crucial when
one resorts to a phase transformation or conversion reaction
characterized by successive generation of new volumes and
interfaces. It enables the acceleration of Na transport at the
electrode–electrolyte or interparticle (interphase) interfaces. To
date, Na-driven conversion processes in oxides and uorides are
extremely irreversible compared with the corresponding Lidriven processes.17,18 To the best of our knowledge, Hector et al.
reported the only example of Na-storage in nitride (Ni3N in their
work) with a reversible capacity as low as 100 mA h g1 at low
rates (0.1 C or 0.5 C) within 20 cycles.19 Most of the previous
reports on conversion anodes assumed the similar reaction
pathway and products for Na- and Li-storage, and therefore
intuitively ascribed the poor conversion efficiency to the sluggish mass transport.17a,19
In this work, for the rst time we propose that Na-driven
conversion electrochemistry could introduce a different reaction pathway and wiring network state when investigating
conversion nitride. We use a template-free method based on
solid–solid phase separation to prepare mesoporous vanadium
nitride (VN) by ammonolyzing bulk Zn-based ternary oxide
Zn3V2O8 (Experimental section in ESI†).20 In terms of constructing porous structures, the ex situ introduction or in situ
generation of a sacricial template (particularly in nano-scale)
evenly in the bulk phase is still a great challenge. Here at a
preparation temperature above 500 C, Zn as a volatile species
sublimes out of the product to create mesopores and in addition three O2 anions are displaced by two N3 anions
(Fig. S1†). VN is chosen as a model material because of its metal
conductivity (bulk conductivity of 1.23 104 S cm1), higher
510 | J. Mater. Chem. A, 2015, 3, 509–514
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than that of many other nitrides.21 By simply mixing with a
small amount of conductive carbon (10 wt%), micro-sized VN as
Li-storage conversion anode displays a highly reversible
capacity as large as 600 or 400 mA h g1 under high rates of 2 C
or 10 C for even thousands of cycles. For Na-storage, the
reversible capacity is still preserved at 300 mA h g1 at 0.1 C or
100 mA h g1 at 2 C even aer 1000 cycles. The rangeability of
the vanadium valence state and permeability of the conductive
network are responsible for the capacity discrepancy between
Li- and Na-storage.
Fig. 1a shows the X-ray diffraction (XRD) pattern of VN
prepared by ammonolysis of Zn3V2O8 at 700 C for 6 h and
conrms the phase purity of the ammonolyzed product. The VN
m with a rened lattice
crystallizes in the space group Fm3
parameter of a ¼ 4.1432(2) Å, but with relatively broad diffraction peaks due to the small crystalline domain sizes. The
specic surface area of mesoporous VN, measured by the Brunauer, Emmett and Teller (BET) method (Fig. 1b), was 54.4 1
m2 g1. The average pore sizes ranged from 20 nm to 40 nm
(Fig. S2†). There was some microporosity (pore diameter # 2
nm) that accounts for 3.6 m2 g1 of the surface area, and a total
micropore volume of 1.2 103 m3 g1. Scanning electron
microscope (SEM) images in Fig. 1c and d further conrm the
well-dened mesoporosity of micro-sized nitride particles with
pores on the scale of 20–40 nm. Because micro-sized particles
are too thick to be analyzed by high-resolution transmission
electron microscopy (HRTEM), we have to choose a thin region
peeled from the sample edge for TEM characterization. The
mesoporosity and polycrystallinity of VN are also indicated from
the HRTEM and selected area electron diffraction (SAED) in
Fig. S3.†
Fig. 2 compares the galvanostatic performance of microsized mesoporous VN as LIB and NIB anodes. In the Li-storage
case (Fig. 2a), the rst discharge capacity of 1350 mA h g1 at
0.1 C slightly exceeds the theoretical value (1237 mA h g1 based
on three electron transfer by completely reducing V3+ to V0). The
excess capacity should be associated with the electrochemical
Fig. 1 (a) PXRD pattern and its refinement profile of VN prepared by
ammonolyzing Zn3V2O8. (b) Plots of nitrogen sorption isotherms at
196 C for mesoporous VN powder. (c) and (d) SEM images of VN
microparticles with internal mesoporosity at different scales.
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Fig. 2 Voltage vs. capacity profiles of micro-sized mesoporous VN as
(a) LIB and (b) NIB conversion anodes at 0.1 C during the first ten cycles
with a voltage range of 0.01–3.0 V. Capacity and coulombic efficiency
of VN anodes as a function of cycle number under high-rate and longterm cycling: (c) at 2 C for Li-storage up to 1000 cycles, (d) at 10 C for
Li-storage up to 3000 cycles and (e) at 2 C for Na-storage up to 1000
cycles.
formation of solid electrolyte interphase (SEI) on electrode
surface, which is oen seen in phase transformation reactions,
particularly when the voltage is close to 0 V vs. Li+/Li.5 The rst
discharge prole displays a distinct two-stage electrochemical
process (sloped from 2.5 V to 0.75 V and less sloped from 0.75 V
to 0.01 V). This two-stage characteristic is well held during the
following charge and discharge with transitions located at 1.75
V and 0.75 V, respectively. The main difference between the rst
and following cycles lies in the coulombic efficiency (CE) or in
the discharge capacity of the lower voltage region, which is
decreased by 44% during the following cycles in view of
considerably less Li consumption for SEI formation. The
following voltage proles are highly reversible with a capacity of
800 mA h g1 and a CE stabilized to 99.5% aer 10 cycles. The
transition voltages are almost in accordance with the ones
where half the capacity is achieved. Their gap between
discharge and charge is as large as 1 V, indicating a different
conversion pathway or sequence between Li uptake and
extraction processes.5 This is also hinted from the opposed
slope tendency of discharge and charge proles, i.e., in contrast
to discharge, the charge prole is sloped from 0.01 V to 1.75 V
and is less sloped from 1.75 V to 3 V.
The micro-sized VN as an LIB anode displays excellent rate
and cycling performance as shown in Fig. S4a,† 2c and d. The
reversible capacity at 0.1 C is stabilized at 700 mA h g1 even
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Journal of Materials Chemistry A
aer long-term (100 cycles) and high-rate (10 C) cycling. Under
considerable higher rates of 1 C, 2 C, 5 C and 10 C, the capacities
of 600, 550, 500 and 400 mA h g1, respectively, are still retained
with high reversibility and CE (100%). More surprisingly, the
capacity at 2 C is increased gradually from 600 mA h g1 to 700
mA h g1 within an extremely long-term cycling of 1000 cycles
(Fig. 2c). A considerable longer cycling of 3000 cycles is
successfully achieved at a rate as high as 10 C, despite a gradual
decrease in reversible capacity from 400 mA h g1 to 300 mA h
g1 aer the rst 700 cycles (Fig. 2d). Note that such an excellent
Li-storage electrochemistry can be achieved as long as a small
amount of conductive carbon (e.g. 10 wt%) is added. Considering that the conductive additive merely wires the outer surface
of the nitride, this indicates that the inner surface of microsized nitride is always sufficiently conductive during the
conversion reaction.
The electrochemical behavior of VN as an NIB anode is
overall different from an LIB anode, mainly in terms of voltage
prole and capacity. Firstly, the two-stage feature of the rst
discharge prole becomes less discernable (Fig. 2b). The
following voltage proles are sloped without evident ‘slope
transition points’ between 0.01 V and 3 V. The voltages where
half the capacity is achieved are located at 1.5 V and 0.75 V for
the charge and discharge processes, respectively. The voltage
gap (0.75 V) at ‘mid-capacity’ appears to be smaller than for the
Li-storage case (1 V), unexpectedly denoting a smaller polarization for Na-storage than for Li-storage. Because it was stated
that Na-storage voltages of compounds are usually 0.18–0.57 V
lower than the corresponding Li-storage voltages,22 this
abnormal phenomenon can be interpreted by the difference of
reaction pathways, products (intermediates), and thus their
thermodynamic potentials. Secondly, the releasable capacity for
Na-storage is lower than for Li-storage. The rst discharge
capacity at 0.1 C is close to 800 mA h g1, and the following
charge capacity is stabilized at 320 mA h g1 with a CE of 98%
aer 10 cycles. Similar to the Li-storage case, the irreversibility
in the rst Na-storage cycle could also be ascribed to the accumulation of SEI, which should mainly occur below 0.7 V during
discharge (Fig. 2b). The cyclic voltammetry (CV, Fig. S5†) is used
to further compare the electrochemistry evolution of VN for LIB
and NIB. At rst glance, sodiation/desodiation of VN displays a
more pseudocapacitive behavior without remarkable characteristic redox peaks apart from a gradually increased cathodic
current towards 0 V. This current increase becomes considerably sharper for the lithiation/delithiation process, where some
additional redox peaks are discernable, particularly those that
are reversible around 2 V. For both the cases, the rst cathodic
process shows a larger current area than the following cycles of
high reversibility, agreeing with the rst irreversibility, as
shown in the galvanostatic measurement (Fig. 2a and b).
With considerably higher rates of 0.5 C, 1 C, 2 C, and 5 C, the
highly reversible capacities are achieved at 250, 200, 150 and
100 mA h g1, with a CE close to 100% (Fig. S4b†). Furthermore,
the capacity at 2 C is gradually decreased to 100 mA h g1 within
the rst 300 cycles, and then is stabilized therein up to at least
1000 cycles (Fig. 2e). Note that such a stable cyclability even at
high rates is obtained without any use of specialized binder
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additive. The Na-storage performance of our micro-sized VN is
considerably superior to most previous reports on conversion
anodes, for which state-of-the-art nanostructure and nanowiring usually were pursued.17,19
In order to get insight into the origins of the outstanding
cyclability and capacity discrepancy, HRTEM (Fig. 3a and b) and
SAED (Fig. S6†) were used to investigate the evolution of reaction products and mixed conductive networks. As expected,
dense (rock salt) structured VN is converted into Li3N and V
aer sufficient Li injection.13a The crystallized Li3N phase can be
detected from XRD, whereas the peaks for the other phases (V or
residual VN) are almost invisible owing to their nanodomain
characteristics (Fig. S7†). Note that the generated metallic V
nanocrystallites are as small as 5 nm, and moreover, they are
welded with each other through grain boundaries to form
chain-like electron channels (scheme in Fig. 3c). These electron
wires are probably percolating in the Li3N matrix. Superionic
Li3N can serve as Li-ion wires and appears to display relatively
better crystallinity near the surface than in bulk. The existence
of unreacted VN nanodomains cannot be ruled out in view of
the difficult propagation of reaction frontiers or conversion
interfaces in some thick regions. In contrast, sodiation of VN
results in a discrete distribution of N-decient VN0.35 (JCPDS 060624) nanodomains (<5 nm) as shown in Fig. 3b and S4b.†
Therefore, the electron transport in V-based products is
conned and interrupted by the surrounding Na-contained
nitrides, which are difficult to identify due to their amorphous
Fig. 3 HRTEM images of (a) lithiated and (b) sodiated VN samples after
discharging to 0.01 V. (c) Scheme of Li- and Na-driven conversion
reactions and the corresponding ion/electron mixed conductive
networks.
512 | J. Mater. Chem. A, 2015, 3, 509–514
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state. Regarding chemical stability, the matrix nitride is probably NaN3 rather than Na3N, which belongs to anti-ReO3
structure and is quite unstable in the ambience.23 The additional corrugated lattice stripes can be assigned to those of the
carbon additive. One should note that the real distribution may
more or less deviate from the present scenario due to metastability of Na-driven reaction products susceptible to e-beam
irradiation.
The difference of conversion efficiency and product evolution between lithiation and sodiation is also indicated by
comparing the X-ray absorption near-edge structure (XANES)
spectra of V K-edge in pristine, discharged and recharged VN
electrodes (Fig. 4). For the pristine sample, a pronounced preedge peak (denoted as A) is found around 5470 eV, indicating a
dipole-allowed transition from V 1s to 3d orbitals, which likely
stems from the distortion of VN6 octahedral units as a consequence of the decrease of ligand symmetry and the break of the
inversion center.24 This is in accordance with the nanocrystalline structure of VN with disorder to a certain extent. In addition, the poor-resolved peaks around 5490 eV and 5505 eV
(denoted as B and C respectively) should correspond to the 1sto-np (n $ 4) transitions, shape resonances and/or multiple
scattering.24a When lithiating to 0.1 V, the absorption edge
shis to the lower energy and the pre-edge peak almost disappears. It conrms the reduction of the vanadium oxidation state
from V3+ to V0 with the change of original coordination symmetry.24b However, the absorption edge of discharge state does
not exactly reach the position of that for metal V due to the
possible existence of unreacted VN nanodomains, which is also
implied by the trace residual of the pre-edge peak. When
recharging to 3 V, the absorption edge shis back to the higher
energy with the reappearance of a remarkable pre-edge peak,
indicating the electrochemical debonding of Li–N and renitridation of V. However the spectrum is difficult to return to the
original position of VN due to conversion insufficiency during
the rst cycle.
In the Na-storage case, the XANES spectra show the similar
shi tendency as for Li-storage. The absorption edges also lie
between those of metal V and VN but with less shiing as well as
Fig. 4 Normalized K-edge XANES spectra of pristine VN and the
samples after discharge to 0.1 V or recharge to 3 V for Li- and Nastorage. The K-edge of V foil is also measured as a reference.
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with a less suppressed pre-edge peak for sodiation than for
lithiation. The shallower conversion depth for sodiation leads
to the formation of N-decient VN0.35 (as shown by TEM) rather
than N-free V in the completely discharged state. The V–N
coordinations in VN0.35 should be still distorted and favorable
for dipole-allowed 1s-to-3d transition. The incompleteness of
Na-driven conversion may be associated with the sluggish Nacontained mass transport at interfaces, large volume repulsion
effect or weaker Na–N bonding than Li–N. These probably
trigger a different reaction pathway with more difficult Nextraction from V–N to form Na–N, although the exact reason is
still unclear.
From the aforementioned results, it is concluded that the
pathway, capacity and product of the electrochemical conversion reaction are signicantly dependent on the size of alkali
cation.18 Although nano-engineering usually brings about
enrichment of defect structures beneting Na insertion or
adsorption for example in the case of FeOx,17b we propose that
another potential factor for electrochemistry upgrade is the
self-reconstruction or maintenance of both internal and
external mixed conductive networks during the conversion
reaction. In this work, the superior Li-storage electrochemistry
of VN is attributed to the real-time self-reconstruction of
conductive network (Fig. 3c). The pristine particle itself has a
well-dened internal conductive network, where the entire VN
framework serves as electron wire, whereas the concomitant
mesopores facilitate electrolyte inltration and serve as a
reservoir for storage of Li-ion wires. Aer complete lithiation,
the mesoporous reservoir is still well maintained owing to the
integrity of micro-sized framework (the overview SEM in
Fig. S8a†), although the evolution of the pore size and geometry is unavoidable due to conversion-driven volume expansion and SEI formation. As far as the local framework is
concerned, therein electron chains consisting of welded V/VN
nanodomains intersect with each other and percolate over the
Li-ion conductive Li3N matrix. Such a self-reconstructed local
framework of mixed conductivity appears to be uniform in
component distribution from the color contrast of the overview TEM in Fig. S9.† Similarly, a Na-driven conversion
process also enables the preservation of integrity of mesoporous structure (Fig. S8b†), although its pore spaces further
shrink due to larger volume expansion and/or thicker SEI
formation. The lower capacity for Na-storage (roughly half of
that for Li-storage) is mainly caused by incomplete vanadium
reduction and conned ion/electron transports. Firstly, the in
situ formed VN0.35 nanodomains are isolated and do not
interconnect to form an electron chain. Secondly, the Nabased matrix (e.g. NaN3) is not so conductive as the superionic
conductor Li3N. To the best of our knowledge, the viewpoint of
reconstruction of a conductive network and reaction pathway
has not been given enough attention to interpret the performance discrepancy between Na- and Li-based conversion
anodes. Some previous reports about oxides assumed a similar
conversion mechanism between them involving the generation of Na2O and Li2O, but did not clarify the intrinsic reason
for Na-driven conversion insufficiency.17 For nitrides, the
complexity of N-bonding chemistry would be more likely to
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Journal of Materials Chemistry A
trigger deviation of conversion mechanism, e.g., generating
Li3N and NaN3. The enrichment of nitride (electro)chemistry
would also bring about new opportunity for the application of
nitride conversion electrodes in NIBs. This result inspires us
to consider a new viewpoint to design future Na-based
conversion electrodes, i.e., exploring more conductive and
percolating mass/charge transport networks, particularly in a
deep sodiation state.
Another advantage for our material lies in the low content
(10 wt%) of carbon additive. Theoretically, the metallic
microparticles merely require point-to-point contact to maintain the external conductive network. Therefore, a lower
content of additive is sufficient to guarantee interparticle
electron transport than for the isolated nanoparticles with a
higher surface area. In our case such electrical contact is not
easy to lose due to the integrity of the micro-sized framework,
which merely suffers from a small volume change due to
internal mesoporosity. Otherwise, under the carbon-free case,
the electrochemical activity degrades signicantly, although a
gradual capacity increase during long-term cycling is observed
for Li-storage (Fig. S10†). This abnormality further indicates
the self-optimization or self-activation capability of the
conductive network.
Conclusions
In summary, we report a micro-sized nitride characterized by
high mesoporosity and conductivity as potential conversion
anodes for LIBs and NIBs. The superior rate and cycling
performances (e.g., 600 mA h g1 and 400 mA h g1 at 2 C and 10
C for Li-storage, and 300 mA h g1 and 100 mA h g1 at 0.1 C
and 2 C for Na-storage, respectively) benet from the selfreconstruction and preservation of intraparticle and interparticle mixed conductive networks. The capacity discrepancy is
closely associated with the differences of reaction pathway (or
degree of V reduction) and wiring network conductivity. This
synthesis methodology based on solid–solid phase separation
can be extended to more monolithic materials with internal
mesoporosity, which have advantages in terms of improving
energy density, decreasing particle-diffusion risk, tolerating
volume change and protecting conductive network.
Acknowledgements
The beamline of BL14W1 in the Shanghai Synchrotron Radiation Facility are acknowledged for hard X-ray absorption ne
structure (XAFS) measurements. We acknowledge the project
supported by the National Natural Science Foundation of China
under Grant no. 51372263, no. 21471147 and no. U1232111. C.
L. Li would like to thank the support from the “Hundred
Talents” program of Chinese Academy of Sciences and the
Science Foundation for Young Researchers of State Key Laboratory of High Performance Ceramics and Superne Microstructures. M. H. Yang would like to thank the National
“Thousand Youth Talents” program of China.
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