electrical-properties-defect-structures-and-ionic-conducting-mechanisms-in-alkali-tungstate-li2w2o7

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Electrical Properties, Defect Structures, and Ionic Conducting
Mechanisms in Alkali Tungstate Li2W2O7
Jungu Xu,* Xiangyu Xu, Huaibo Yi, Yun Lv, Nansheng Xu, Lunhua He, Jie Chen, Xiaojun Kuang,*
and Kevin Huang*
Cite This: Inorg. Chem. 2021, 60, 8631−8639 Read Online
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ABSTRACT: Discovery of new high-conductivity solid-state ionic conductors has been a long-
lasting interest in the field of solid-state ionics for their important applications in solid-state
electrochemical devices. Here, we report the mixed oxide-ion and Li-ion conductions, together
with their conducting mechanisms in the Li2W2O7 material with triclinic symmetry. The process
for the ionic identity is supported by several electrochemical measurements including
electrochemical impedance spectroscopy, DC polarization, oxygen concentration cell, and
theoretical analysis of neutron diffraction data and bond-valence-based energy landscape
calculations. We show from electrochemical measurements strong evidences of the predominating
oxide-ion conducting and minor Li-ion chemistry in Li2W2O7 at high temperatures, while the
bond-valence-based energy landscape analysis reveals possible multidimensional ionic migration pathways for both oxide-ions and
Li-ions. Thus, the presented results provide fundamental insights into new mixed ionic conduction mechanisms in low-symmetry
materials and have implications for discoveries of new ionic conductors in years to come.
■ INTRODUCTION
Solid-state oxide-ion conductors are key components in clean
and efficient solid oxide cells (SOCs) for power generation and
chemical production.1−4 The magnitude of their ionic
conductivity directly determines the performance of SOCs
and, more importantly, the temperature at which SOCs are
operated. Too high an operating temperature would invoke
significant material deterioration and unwanted reactions, thus
leading to cell performance degradation. The high operating
temperature also requires the use of expensive exotic materials,
which adds costs to SOCs and is an impediment to
commercialization. Therefore, discovery of high-conductivity
solid-state oxide-ion conductors and understanding the
conduction mechanisms are of paramount importance to the
development of low-cost and reliable SOCs.
Conventional solid-state oxide-ion conductors are primarily
found in crystalline oxides with high crystallographic
symmetries and oxygen vacancies or interstitials as mobile
lattice defects, representatives of which with oxygen vacancies
are fluorite-structured yttria-stabilized zirconia (YSZ)5 and
perovskite-structured strontium and magnesium co-doped
LaGaO3 (LSGM).6−8 Those with oxygen interstitials have
been found in apatite-structured materials,9−13 melilite-
structured materials,14−18 and, more recently, reported
Ba7Nb4MoO20-based hexagonal perovskite-related oxides.19
In addition to solid crystalline materials, molten binary
oxides, such as Bi2O3,
20 V2O5,
21 and TeO2,
22 have also been
found with high oxide-ion conductivity. Very recently, we have
reported a ternary oxide Na2W2O7 with a high oxide-ion
conductivity above its melting point.23 The solid crystallized
Na2W2O7 has orthorhombic symmetry with the space group of
Cmca, in which W has mixed tetrahedral and octahedral
coordination with oxygen, forming infinite W2O7 chains along
the a axis. Each WO6 octahedron is corner-sharing with two
other octahedra, while each WO4 tetrahedron connects two
corner-shared WO6 octahedra via corner sharing. Below its
melting point, Na2W2O7 did not show any oxide-ion
conductivity, above which, however, Na2W2O7 exhibits an
appreciably high oxide-ion conductivity that closely relates to
its disordered coordination-number-variable WO4+x mixed
polyhedral network.
In this work, the study focused on another solid-state alkali
tungstate Li2W2O7 that has the same general formula
(A2W2O7, A = alkali metal) as Na2W2O7 but with a totally
different crystal structure. Figure S1 shows the crystal structure
of Li2W2O7 in triclinic symmetry with space group P1̅, in which
all the W atoms have 6-coordination with O, forming distorted
edge-shared octahedra and one-dimensional chains along the c
axis. These chains are interlinked by distorted LiO4 tetrahedra
by corner-sharing two of the four O atoms in the LiO4
tetrahedra. At the same time, these LiO4 tetrahedra are
corner-shared with each other, also forming 1D chains along
the c-axis direction. Pralong et al. had previously cursorily
Received: February 28, 2021
Published: June 2, 2021
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© 2021 American Chemical Society
8631
https://doi.org/10.1021/acs.inorgchem.1c00609
Inorg. Chem. 2021, 60, 8631−8639
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investigated the lithium-ion conduction in Li2W2O7 from room
temperature to 500 °C in air24 and without oxide-ion
conduction being studied. Here, the electrical properties,
defect structures, and ionic conducting mechanisms of the
Li2W2O7 material were thoroughly studied and mixed lithium-
ion and oxide-ion conductions, which are dominant at high
temperature, were revealed.
■ EXPERIMENTAL SECTION
Stoichiometric amounts of Li2CO3 and WO3 were first weighed and
intimately mixed by ball milling in ethanol. The dried mixture was
then pressed into pellets and calcinedat 550 °C for 8 h. The pellets
were then broken up and ground followed by sintering at 660 °C for 6
h. The final product has a density greater than 95% of its theoretical
value.
The phase composition of the samples was examined by X-ray
diffraction (XRD, Panalytical X’pert Pro X-ray diffractometer) with
Cu Kα radiation over a 2θ range of 10−80° with an interval of
0.01313° and a scan speed of 10° min−1. Compositional analysis was
carried out by inductively coupled plasma (ICP) elemental analysis,
for which the sample was dissolved in a mixed HNO3/HF solution.
Time-of-flight (TOF) neutron powder diffraction (NPD) data were
collected from a general purpose powder diffractometer (GPPD)25
(90° bank) at the China Spallation Neutron Source (CSNS).
Refinements were performed using Topas Academic software.26
The total electrical conductivity of Li2W2O7 was measured by
electrochemical impedance spectroscopy (EIS) using a Solartron
1260 frequency response analyzer with a 10 mV AC stimulus over a
frequency range of 107 to 10−1 Hz. Prior to EIS measurements, the
sample was coated with gold paste on its two surfaces and cured at
600 °C for 3 h. For each EIS measurement, the sample was exposed to
one specific atmosphere of interest and temperature and was
equilibrated for 1 h. To determine the oxide-ionic transport number
of Li2W2O7, an oxygen concentration cell was fabricated, and the
electromotive force (EMF) of the cell was measured. To make the
oxygen concentration cell, a Li2W2O7 pellet was sealed on an alumina
tube and its two surfaces were exposed to two gases with a fixed but
different partial pressure of oxygen (pO2). The theoretical EMF values
of the above two oxygen concentration cells were calculated using the
Nernst equation. In parallel, potentiostatic DC polarization of
chronoamperometry was also performed in oxygen to separate
oxide-ion and possible Li-ion conductivity. Like the EIS measurement,
the sample was coated with platinum paste on its two surfaces and
cured at 700 °C for 3 h.
A solid oxide fuel cell (SOFC) with Li2W2O7 as the electrolyte was
also tested to verify the oxide-ion conducting chemistry. In this cell,
the mixture of NiO and Gd-doped CeO2 (GDC) with a volume ratio
of 3:2 was used as the anode, and La0.6Sr0.4CoO3 was used as the
cathode. To fabricate the cell, separate layers of the anode and
Li2W2O7 were first co-pressed into a pellet with a diameter of 15 mm
and then fired at 650 °C for 4 h to achieve a good electrolyte/anode
bonding and a dense electrolyte. The thicknesses of the anode and
electrolyte layers are 0.85 and 0.65 mm, respectively. Last, the
La0.6Sr0.4CoO3 cathode was screen-printed onto the surface of the
Li2W2O7 side followed by firing it at 650 °C for 3 h. The effective
surface area of the cell was 0.5 cm2. A silver mesh was used as the
current collector. A flow of 50 mL min−1 air and 30 mL min−1 dry H2
was supplied to the cathode and anode sides as the oxidant and fuel,
respectively. The measurements were only carried out at 700 °C. The
characteristic OCV−t, V−I, and P−I curves were collected using an
electrochemical workstation (Bio-Logic VSP, France).
The Li-ion and oxide-ion migration pathways were systematically
investigated using the program SoftBV27 by calculating their bond-
valence-based energy landscapes (BVELs) that can provide
information on the connecting local minima and saddle points
(identified by fractional coordinate values). Diffusion pathways were
determined with regions of low bond-valence site energy by direct
visualization of the connectivity of the isosurfaces and by examination
of the calculated pathway segments. The structure from Rietveld
refinement against NPD data was used as the input for the BVEL
calculations. The spatial resolution was set to 0.1 Å for both Li-ion
and oxide-ion calculations.
■ RESULTS AND DISCUSSION
Phase Analysis. Shown in Figure 1a is the XRD pattern of
synthesized Li2W2O7. All the peaks can be well indexed into
the triclinic Li2W2O7 phase (JPDCS #70-0869), and no
secondary phase is detected. Figure 1b showing Le Bail
refinement based on the XRD data using the triclinic cell
parameters further confirms the phase purity of the synthesized
Li2W2O7.
Electrical Properties. Preliminary electrical properties of
Li2W2O7 were then studied by AC impedance spectroscopy.
Figure 2a shows the Arrhenius plots of the total conductivities
of Li2W2O7 under O2, air, and Ar atmospheres. The total
conductivity is calculated based on the bulk resistance (Rb, the
total of grain and grain boundary responses) using eq 1 as
follows:
σ =
×
L
S Rb
t
(1)
where L and S are the thick and conducting cross-sectional
areas of the sample pellet. The insensitivity of the total
conductivity to pO2 in oxidizing and inert atmospheres is
indicative of the ionic conduction, and the electronic
conduction, if any, is negligible. As the melting point of
Li2W2O7 is ∼730 °C, EIS measurements were carried out only
at ≤700 °C. At 700 °C, the conductivity of Li2W2O7 is 6.6 ×
10−3 S cm−1. The EIS spectra recorded under different
atmospheres at 700 °C were then further analyzed to
determine the nature of charge carriers. Figure 2b shows that
Figure 1. (a) XRD data of the synthesized Li2W2O7; (b) Le Bail
refinement against the experimental data.
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all these impedance spectra exhibit only an electrode response
with a nonzero and almost identical intercept at high
frequency. Since this intercept represents the ohmic con-
tribution from the grain and grain boundary of the electrolyte,
its independence on pO2 suggests the nature of pure ionic
conduction14,28 from 10−4 atm (pure Ar) to 1 atm (pure O2
gas). The electrode response in the low-frequency range is,
however, very sensitive to pO2, i.e., the higher the pO2, the
lower the electrode resistance. This strong dependence on pO2
infers the nature of active oxygen electrocatalysis (O2 + 2e− ↔
2O2−) at the electrolyte/electrode interface, which is routinely
observed in the oxide-ion conductors (e.g., the recently
reported ferroelectric Na0.5Bi0.5TiO3-based oxide-ion conduc-
tors28).
Further cell electromotive force (EMF) measurements based
on oxygen gas concentration cell of air|Li2W2O7|O2 resulted in
an EMF of ≈31.21 mV at 700 °C, more than 95% of the
theoretical EMF calculated using the Nernst equation,
confirming again the oxide-ion conduction in Li2W2O7 under
a high pO2 range, with negligible electronic conduction.
As aforementioned, the mobility of the lithium-ion in
Li2W2O7 had also been identified by Pralong et al.24 Thus, the
contribution of lithium-ion conduction to the total con-
ductivity should be evaluated. One way to separate mixed ionic
conductivity is to combine EIS and DC polarization measure-
ments, which has been well documented in the literature.29−32
In this method, the electrolytes with mixed mobile ions are
usually sandwiched between two identical electrodes (sym-
metric); the symmetrical electrodes are blocked for one ion but
reversible for the other. For instance, Watanabe et al.32 used
the combined AC/DC method to estimate the Li+ transport
number in a polymer electrolyte consisting of poly(ethylene
oxide) (PEO) and a lithium salt, in which both the cation and
anion are mobile. In their work, a lithium disk was used as the
electrode, which was reversible to the Li+-ion but blocked for
anions. Similarly, Anantha and Hariharan29 investigated the
Na+ transport number in a PEO−NaNO3 polymer electrolyte
with metallic Na as the reversible electrode for Na+-ions. In thepresent work, since it is impossible to use pure Li metal as the
reversible electrode as in the above method to directly estimate
the Li+ transport number in Li2W2O7 because of the high
chemical reactivity and low melting point of lithium metal
(∼180 °C), we instead used porous Pt as reversible oxygen
electrodes under a pure oxygen atmosphere to block off Li+.
For this method, the AC impedance plot requires the full-range
impedance responses that include impedances from the bulk
electrolyte (the total of grain and grain boundary responses)
and electrode response resistance (including charge transfer
and diffusion resistances).31 For the DC measurements, the
final steady current (Is) in the chronoamperometric curve
results from the migration of the reversible ions only, i.e.,
oxide-ions in the present work. Meanwhile, for the blocked
ions, Li-ions here, their initial migrations are driven by the
applied electric field and would eventually be balanced out by
diffusion due to the concentration gradient. The total
resistance calculated using the applied DC bias V and steady
current Is includes not only the electrolyte resistance but also
the polarization resistance due to the electrode reaction (Re),
the value of which can be obtained from the full-range AC
impedance response plot. Therefore, to calculate the oxide-ion
conductivity from the DC measurement, the resistance of the
oxide-ion transporting in the electrolyte should be deduced by
stripping the polarization resistance from the total resistance,
i.e., (V/Is − Re). Therefore, oxide-ion conductivity in the DC
measurement can be calculated using the following equation:
σ =
× [ − ]
L
S V I R( / )o
s e (2)
As the total conductivity of Li2W2O7 can be obtained from
eq 1, the transport number of the oxide-ion in this work can
thus be calculated as
σ
σ
=t
t
o
o
(3)
From eqs 1−3, the to is finally deduced as
=
−
t
R
V I R( / )o
b
s e (4)
and the transport number of the Li+-ion can be simply
calculated by tLi = 1 − to, with a negligible electronic
conduction assumption.
To ensure the reliability of this method, a standard 8YSZ
pellet sample (∼14.85 mm in diameter, ∼1.0 mm in thickness)
with two identical Pt electrodes was first measured to
determine its oxide-ion transport number, which is known to
be close to 1.33−35 Figure S2a shows the EIS spectrum
recorded at 750 °C under a pure oxygen atmosphere. The bulk
resistance Rb of this 8YSZ pellet at this temperature is ∼4.58
Ω, while Re including charge-transfer and oxygen-diffusion
resistances is ∼1.12 Ω. The subsequent DC measurements at
different DC biases resulted in an almost identical value of V/
Is, ∼5.75 Ω (see Figure S2b). Thus, the oxide-ion transport
Figure 2. (a) Arrhenius plots of the conductivities of Li2W2O7 under different pO2; (b) EIS spectrum at 700 °C under different pO2.
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number in this YSZ pellet at 750 °C is calculated to be ∼0.99,
consistent with that reported in the literature33−35 and
confirming the validity of this combined AC/DC technique
with porous Pt under high pO2 as a reversible electrode to
determine the oxide-ion transport number in ionic conducting
materials.
We then performed combined AC/DC measurements on
the Li2W2O7 pellet (∼10.44 mm in diameter, ∼1.0 mm in
thickness) under a pure oxygen atmosphere at 700 °C, with a
porous oxygen-reversible Pt electrode on the opposite sides of
the pellet. Figure 3a showing the EIS spectrum gives Rb =
17.29 Ω and Re = 4.32 Ω. For the DC measurements, Figure
3b shows the current vs time profiles under fixed potentials. As
reported,30,32 the Is used for calculating the transport number
must be taken from the region in which the Is vs V plot is
linear. Here, as shown in Figure 3c, the Is is ohmic up until the
potential reaches 50 mV, at and above which it levels off. A
similar behavior where the plot was linear only when the
applied voltage was less than 50 mV was also observed in
Hashmi and Chandra’s work.30 Accordingly, on the basis of the
data in the linear part of the Is vs V plot in the DC
measurements and the data from the EIS spectrum, using eq 4,
the oxide-ion transport number (to) was calculated to be
∼0.78, and the transport number for the Li-ion is ∼0.22 = 1 −
to. It is worth noting that, although the oxide-ion and Li+-ion
transport numbers derived here using this combined AC/DC
method may not be extremely accurate, we can conclude that
the Li2W2O7 material shows dominant oxide conduction.
To evaluate the potential of Li2W2O7 as an electrolyte for
SOFCs and investigate its phase stability under the operational
conditions of SOFCs, a solid oxide fuel cell using Li2W2O7 as
the electrolyte was then fabricated (details are described in the
Experimental Section). Figure 4a shows the V−I and P−I
curves, which were acquired within the first 4 h. The close-to-
theoretical open-circuit voltage (OCV) confirms the oxide-ion
conduction chemistry. The Li-ion conduction is not expected
Figure 3. (a) EIS spectrum for Li2W2O7 at 700 °C using porous Pt as a reversible electrode; (b) obtained currents vs DC potentials; (c) steady-
state currents (Is) vs applied DC voltages.
Figure 4. (a) Power generating characteristic of an SOFC using
Li2W2O7 as the electrolyte obtained in the first 4 h at 700 °C; (b)
delayed electronic conduction behavior of Li2W2O7 in reducing
atmospheres.
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to be prevalent in SOFC conditions because of the lack of
chemical potential gradient of Li. A low peak power density of
22.5 mW cm−2 was clearly due to the thick electrolyte and
unoptimized electrodes. While Figure 4a shows that the SOFC
exhibits a reasonable OCV at the beginning, Figure 4b (black
plot) indicates that this initial OCV experiences a sudden drop
from 1.13 to 0.85 V after ∼4 h, implying the occurrence of
electronic conduction, likely associated with the reduction of
W6+ by H2. The delayed electronic conduction is further
supported by the conductivity measurements of Li2W2O7
under a reducing atmosphere (10% H2−90% N2) at 700 °C.
The result shown in Figure 4b (blue plot) indicates a 5 h
delayed but enhanced conductivity, which can be ascribed to
the partial reduction of W6+ into lower oxidation states and
therefore introducing electronic conduction. This is verified by
the post-mortem XRD examination of the reduced sample,
which reveals two impurity phases (Figure S3): Li0.36WO3
(PDF #76-1497) and Li2WO4 (PDF #12-0760), and the
former carries lower oxidation states of W ions.
Conducting Mechanisms. To achieve a good fundamen-
tal understanding on how ionsmove in Li2W2O7 and establish
a relationship between the defect structure and conductivity,
room-temperature powder neutron diffraction (PND) was
performed. The collected data were then analyzed by Rietveld
refinement to determine structural parameters. Prior to
Rietveld refinement, ICP analysis was carried out to determine
the real composition of the as-made Li2W2O7 sample. The
result turned out to be ∼Li1.92W2 for cations. The deficiency of
lithium in the sample may be due to the volatilization during
the preparation process. The refinement was carried out with
the background, lattice parameters, and peak shape parameters
refined in sequence as the first stage. This was followed by the
positional parameters, atomic site occupancies, and finally the
isotropic atomic displacement parameters being refined. For
the atomic site occupancies, they were fixed to be unity for the
W sites, while for the Li sites, their occupancies were
constrained to reasonable values to avoid obvious deviation
from the lithium content derived from ICP analysis. Figure 5
showing the refinement indicates a good agreement between
the measured and calculated diffraction patterns, and Table 1
lists the obtained structural parameters. The refinement results
revealed that ∼97% and ∼93% are occupied for the Li1 and
Li2 sites, respectively. The refined lithium content in the
structure is therefore a little bit lower than that obtained from
the ICP experiment. Accordingly, oxygen vacancies are
expected to form in the structure and act as charge balancing
defects to Li vacancies. The occupancies on the oxygen sites
were then refined subject to the charge neutrality constraint.
The considerably low atomic displacement parameters for all
the atomic sites infer a high precision for these refined atomic
positions. Thus, the high Li-ion and oxide-ion conductivities in
Li2W2O7 may originate from the intrinsic lithium and oxygen
vacancies. Similar cases were also reported in the well-known
solid Li-ion conductor Li7La3Zr2O12 garnets.
36 Selected bond
lengths and bond valence sum in the refined structure are listed
in Table 2.
To investigate how exactly Li-ions and oxide-ions migrate in
Li2W2O7, the bond-valence-based method was applied in this
work. This method has been well documented in the literature
about how it is used for analyzing ion transport pathways
statistically from crystal structure data37,38 and has been
extensively employed for the investigation of conduction
pathways in various oxide-,7,39−44 sodium-, and lithium-ion
conductors.22,45−49 For calculations using the bond-valence
method, the ion to be tested was placed sequentially at all
points of a three-dimensional grid covering a unit cell.
Positions with a low bond-valence mismatch ((Vi − ∑Sij),
where Vi corresponds to the formal oxidation state of ion i and
Sij is the bond valence for the pair of ions i and j), which is
correlated to a low energy site and connecting equilibrium
sites, form an infinite network for a potential migrating
pathway.46
Here, we calculated the bond-valence-based energy land-
scapes (BVELs) of both Li+ and O2− ions using the refined
lattice and positional parameters from the PND data, with a
spatial resolution of 0.1 A. Figure 6a shows the iso-surface (in
orange) of the BV-based energy for a Li-ion at 0.675 eV in the
crystal structure of Li2W2O7 shown in Figure S4a, where the
most stable position of a Li-ion is set to 0 eV. The one-
dimensional (1D) zigzag pathway along the ⟨010⟩ direction is
evident, and the calculated energy barrier for this 1D diffusion
is ∼0.673 eV. Therefore, the pathway for the Li-ion follows (1)
Li2a−Li2b−Li1a−Li2a and (2) Li2a−Li2b−Li1b−Li2a.
Here, “a” and “b” are used to specify the atoms with the
same crystallographic coordinate in a unit cell but different
positions on the pathway. In detail, for pathway (1), a Li-ion
diffuses first from Li2a to an interstitial site, i5 (0, 0.556,
0.431), via a local saddle point, and then to Li2b. A continual
migration from Li2b to Li1a requires the passage of an
interstitial position at i6a (0.719, 0.056, 0.403) before finally
reaching Li2a by passing through another interstitial position
at i2a (0.073, 0.111, 0.778). Thus, pathway (1) for Li-ion
migration follows specifically Li2a−(i5)−Li2b−(i6a)−Li1a−
(i2a)−Li2a. Similarly, for pathway (2), there are also three
essential interstitial positions, i.e., i5, i2b, and i6b, and the
crystallographic coordinates for i2b and i6b are the same as
those for i2a and i6a, respectively. Thus, pathway (2) follows
the moving sequence of Li2a−(i5)−Li2b−(i2b)−Li1b−
(i6b)−Li2a. The interstitial positions are marked in red
crosses in Figure 6a. From Figure 6a and Figure S4b, one can
see that no connection is formed along the ⟨100⟩ and ⟨001⟩
directions when the iso-surface energy is set to 0.675 eV.
However, when the iso-surface energy is set to 0.685 eV,
Figure 6b shows that connections along both ⟨100⟩ and ⟨001⟩
directions are also formed besides the ⟨010⟩ direction. ThisFigure 5. Rietveld fitting plots for ND data of Li2W2O7.
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insight suggests a three-dimensional (3D) diffusion behavior
for the Li-ion in Li2W2O7. The calculated energy barrier for
such 3D diffusion is 0.679 eV. Along the ⟨100⟩ direction, the
pathway follows Li1a−(i8)−Li2a−(i2a)−Li1a or Li2a−
(i2a)−Li1a−(i8)−Li2a, whereas along the ⟨001⟩ direction,
the pathway follows Li2b−(i8)−Li1a−(i6a)−Li2b or Li2a−
(i6b)−Li1b−(i8)−Li2a. One can see that the i8 interstitial
site (0.448, 0.708, and 0.319) is essential for the connection in
both ⟨100⟩ and ⟨001⟩ directions.
For oxide-ion migration, the calculated results suggest that a
1D diffusion pathway is formed along the ⟨111⟩ direction with
an energy barrier of ∼0.593 eV. Figure S5a depicts the BVEL
with an iso-surface at 0.6 eV. The iso-surface was connected
along the ⟨111⟩ direction, and the connected areas cover
oxygen sites O1, O2, O3, O4, O6, and O7, inferring that the
oxide-ion diffusion will take place within these oxygen sites. As
revealed by Rietveld refinement on the PND data, the O5 site
is isolated due to its full occupancy. A likely oxide-ion pathway
can be viewed as O4−O2−O6−O2−O4, which is indicated by
red arrows in Figure S5b. Two-dimensional (2D) diffusion for
the oxide-ion on the (1, −1, 0) plane is favorable due to its low
energy barrier (∼0.62 eV); a BVEL with an iso-surface at 0.625
eV viewed from different directions is depicted in Figure S5c,e.
For this 2D diffusion, a possible diffusion pathway along the
⟨110⟩ direction is O2−O4−O1−O6−O2, as shown in Figure
S5d, while another diffusion pathway along the ⟨001⟩
direction, i.e., O4−O2−O6−O1−O2−O4 shown in Figure
S5f, is also possible. The calculations suggest that only when an
oxide-ion overcomes an energy barrier (or energy saddle
point) of ∼0.76 eV can it migrate through the O5 site. The 3D
migration of the oxide-ion in Li2W2O7 by interconnecting the
(1, −1, 0) plane requires a higher energy barrier (∼0.91 eV)
(see Figure 7a). This predicted energy barrier is also close to
the experimental value of ∼1.1 eV shown in Figure 1a. An
oxide-ion needs to hop between the (1, −1, 0) plane from the
O6 site in one plane to a nearest O5 or O7 site in another (1,
−1, 0) plane, as indicated by redarrows in Figure 7b. The
distances for O6 to hop to O5 and O7 are 2.872(6) and
Table 1. Refined Structural Parameters of Li2W2O7 at Room Temperaturea
atoms sites x y z occupancy Biso (Å
2)
W1 2i 0.1785(2) 0.2523(1) 0.3045(2) 1 0.13(1)
W2 2i 0.6667(2) 0.4511(2) 0.188(2) 1 0.13(1)
Li1 2i 0.287(1) 0.023(1) 0.859(3) 0.97(2) 0.41(1)
Li2 2i 0.905(1) 0.266(3) 0.706(2) 0.93(1) 0.39(3)
O1 2i 0.8711(3) 0.6037(2) 0.0497(3) 0.98(1) 0.16(1)
O2 2i 0.7018(2) 0.7481(1) 0.3468(3) 0.99(1) 0.14(1)
O3 2i 0.8228(2) 0.9751(3) 0.8231(1) 1.0 0.22(1)
O4 2i 0.5782(1) 0.5851(1) 0.8211(3) 0.99(1) 0.16(1)
O5 2i 0.6287(1) 0.2083(2) 0.0738(4) 1.0 0.15(2)
O6 2i 0.9648(2) 0.1969(3) 0.3552(3) 0.99(1) 0.13(1)
O7 2i 0.7687(1) 0.4359(3) 0.536(2) 1.0 0.23(1)
aLattice parameters: a = 8.244(3) Å, b = 7.005(2) Å, c = 5.026(4) Å, α = 85.324(2)°, β = 102.212(1)°, γ = 110.356(1)°; space group, P1̅. The
refined composition is ∼Li1.9W2O6.95.
Table 2. Selected Bond Lengths and Ion Bond Valence Sum
(BVS) for Li2W2O7 Obtained from Rietveld Refinement of
PND Data at Room Temperature
bond length (Å) bond length (Å) bond length (Å)
W1−O6 1.735(2) W2−O5 1.741(2) Li1−O2 1.833(3)
W1−O3 1.763(5) W2−O7 1.793(1) Li1−O5 1.924(4)
W1−O2 1.828(4) W2−O1 1.888(3) Li1−O3 2.015(4)
W1−O1 2.011(4) W2−O4 1.934(2) Li1−O6 2.191(3)
W1−O4 2.135(4) W2−O4 2.101(1) BVS for
Li1
0.99
W1−O7 2.254(4) W2−O2 2.191(1) Li2−O7 1.907(3)
BVS for
W1
5.97 BVS for
W2
5.86 Li2−O1 1.982(1)
BVS for
O1
2.04 BVS for
O5
1.86 Li2−O3 2.015(3)
BVS for
O2
2.02 BVS for
O6
1.83 Li2−O6 2.092(3)
BVS for
O3
1.78 BVS for
O7
1.97 BVS for
Li2
0.97
BVS for
O4
2.15
Figure 6. (a) BV-based energy landscape (yellow areas) for an oxide-
ion in Li2W2O7 with an iso-surface at 0.678 eV viewed along the
⟨001⟩ direction, and WO6 polyhedra are omitted for clarity; (b)
BVELs at 0.685 eV viewed along the ⟨010⟩ direction. The pictures are
plotted in the range of 0.0 ≤ x ≤ 2.0, 0.0 ≤ y ≤ 2.0, and 0.0 ≤ z ≤ 2.0,
and only one unit cell line is depicted.
Figure 7. (a) BVEL with an iso-surface at 0.91 eV; (b) possible
pathways of an oxide-ion to hop from O6 in one W2O7 layer to O5 or
O7 in another W2O7 layer.
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3.055(8) Å, respectively. In fact, the O6−O7 line is the edge of
a Li2O4 tetrahedron, while the O6−O5 line is the edge of a
Li1O4 tetrahedron. Thus, we can draw a conclusion that Li-
ions play an essential role in promoting 3D oxide-ion migration
in Li2W2O7.
■ CONCLUSIONS
In this work, we have identified Li2W2O7 as a new reducible
oxide-ion and Li-ion conductor through experimental measure-
ments and theoretical analysis. The Rietveld refinement using
PND data reveals that Li2W2O7 possesses intrinsic vacancies
on both Li sites and O sites. These vacancies provide the basic
mobile defects for Li2W2O7 to conduct the Li-ion and oxide-
ion. The total ionic conductivity in an oxidizing-to-inert
environment is ∼6.6 × 10−3 S cm−1 at 700 °C and dominated
by oxide-ion conduction. The theoretical study on Li-ion and
oxide-ion conducting mechanisms by a bond-valence-based
method reveals energy barriers of ∼0.68 and 0.91 eV for 3D
migration of Li+ and O2−, respectively; the latter is close to the
experimental value. Overall, this work provides fundamental
insights into the electrical properties and ionic conducting
mechanisms in the lowest-symmetry Li2W2O7 as a solid-state
ionic conductor and has important implications for discoveries
of new ion conductors in years to come.
■ ASSOCIATED CONTENT
*sı Supporting Information
The Supporting Information is available free of charge at
https://pubs.acs.org/doi/10.1021/acs.inorgchem.1c00609.
Crystal structure of Li2W2O7, EIS spectrum and
potentiostatic DC measurements of chronoamperomet-
ric curves of the 8YSZ sample recorded at 750 °C under
pure oxygen, XRD pattern of Li2W2O7 after reduction,
crystal structure of refined Li2W2O7 based on PND data,
BVELs at 0.678 eV for an oxide-ion in Li2W2O7, and
BVELs for the 1D and 2D oxide-ion migration pathways
(PDF)
■ AUTHOR INFORMATION
Corresponding Authors
Jungu Xu − MOE Key Laboratory of New Processing
Technology for Nonferrous Metals and Materials, Guangxi
Universities Key Laboratory of Non-ferrous Metal Oxide
Electronic Functional Materials and Devices, College of
Materials Science and Engineering, Guilin University of
Technology, Guilin 541004, China; orcid.org/0000-
0003-4034-3772; Email: xujungu@glut.edu.cn
Xiaojun Kuang − MOE Key Laboratory of New Processing
Technology for Nonferrous Metals and Materials, Guangxi
Universities Key Laboratory of Non-ferrous Metal Oxide
Electronic Functional Materials and Devices, College of
Materials Science and Engineering, Guilin University of
Technology, Guilin 541004, China; orcid.org/0000-
0003-2975-9355; Email: kuangxj@glut.edu.cn
Kevin Huang − Department of Mechanical Engineering,
University of South Carolina, Columbia, South Carolina
29201, United States; orcid.org/0000-0002-1232-4593;
Email: huang46@cec.sc.edu
Authors
Xiangyu Xu − MOE Key Laboratory of New Processing
Technology for Nonferrous Metals and Materials, Guangxi
Universities Key Laboratory of Non-ferrous Metal Oxide
Electronic Functional Materials and Devices, College of
Materials Science and Engineering, Guilin University of
Technology, Guilin 541004, China
Huaibo Yi − MOE Key Laboratory of New Processing
Technology for Nonferrous Metals and Materials, Guangxi
Universities Key Laboratory of Non-ferrous Metal Oxide
Electronic Functional Materials and Devices, College of
Materials Science and Engineering, Guilin University of
Technology, Guilin 541004, China
Yun Lv − MOE Key Laboratory of New Processing Technology
for Nonferrous Metals and Materials, Guangxi Universities
Key Laboratory of Non-ferrous Metal Oxide Electronic
Functional Materials and Devices, College of Materials
Science and Engineering, Guilin University of Technology,
Guilin 541004, China
Nansheng Xu − Department of Mechanical Engineering,
University of South Carolina, Columbia, South Carolina
29201, United States
Lunhua He − Beijing National Laboratory for Condensed
Matter Physics, Institute of Physics, Chinese Academic of
Sciences, Beijing 100190, China; Songshan Lake Materials
Laboratory, Dongguan 523808, China; Spallation Neutron
Source Science Center, Dongguan 523803, China
Jie Chen − Spallation Neutron Source Science Center,
Dongguan 523803, China
Complete contact information is available at:
https://pubs.acs.org/10.1021/acs.inorgchem.1c00609
Notes
The authors declare no competing financial interest.
■ ACKNOWLEDGMENTS
This work was supported by the National Natural Science
Foundation of China (no.21965008), Guangxi Natural
Sc ience Foundation (2017GXNSFAA198203 and
2019GXNSFGA245006), the Open Fund of Key Laboratory
of New Processing Technology for Nonferrous Metal &
Materials, Ministry of Education/Guangxi Key Laboratory of
Optical and Electronic Materials and Devices (no. 20AA-4),
and the “High-level innovation team and outstanding scholar
program of Guangxi institutes.”
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