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From solid electrolyte to zinc cathode: vanadium substitution in ZnPS
3
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2021
J. Phys. Mater.
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024005
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J. Phys. Mater.
4
(2021) 024005
https://doi.org/10.1088/2515-7639/abe365
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PAPER
From solid electrolyte to zinc cathode: vanadium substitution
in ZnPS
3
Andrew J Martinolich
, Skyler D Ware
, Brian C Lee
and Kimberly A See
Division of Chemistry and Chemical Engineering, California Institute of Technology, Pasadena, CA 91125, United States of America
E-mail:
ksee@caltech.edu
Keywords:
zinc battery
,
zinc cathode
,
multivalent ionics
,
multivalent battery
Supplementary material for this article is available
online
Abstract
Development of next generation batteries is predicated on the design and discovery of new,
functional materials. Divalent cations are promising options that go beyond the canonical Li-based
systems, but the development of new materials for divalent ion batteries is hindered due to
difficulties in promoting divalent ion conduction. We have developed a family of cathode materials
based on the divalent cation conductor ZnPS
3
. Substitution of V for Zn in the lattice concomitant
with vacancy introduction yields isostructural but redox-active materials that can reversibly store
Zn
2
+
in the vacancies. A range of voltammetry and galvanostatic cycling experiments along with
x-ray photoelectron spectroscopy support that redox is indeed centered on V and that capacity is
dependent on the V content. The voltage of the materials is limited by the irreversible
decomposition of the
[
P
2
S
6
]
4
polyanion above 1.4 V vs. Zn/Zn
2
+
. The reversible capacity before
anion decomposition is limited to half the vacancies and is due to the relative ratios of oxidized and
reduced V centers. Such observations provide useful design rules for cathode materials for divalent
cation based battery technologies, and highlight the necessity for a holistic interpretation of
physical and electronic structural changes upon cycling.
1.Introduction
Energy storage technologies are a ubiquitous part of modern life. Devices based on Li
+
intercalation at the
anode and cathode dominate the market of rechargeable devices, largely centered around Li intercalated
graphite anodes and LiCoO
2
cathodes [
1
]. Li-ion batteries exhibit highly reversible electrochemistry due to
the intercalation based behavior but are nearing theoretical limits [
2
]. Beyond performance, concerns
regarding the sustainability of Li-ion batteries and their dependence on threatened resources, namely Li and
Co, drive research into new chemistries [
3
]. Divalent working ions present attractive opportunities to help
relieve the environmental and economic strain imposed by current technologies [
4
]. Despite being
conceptualized as early as 1840 [
5
], development of batteries using divalent working ions such as Mg and Zn
has been limited compared to the Li counterparts.
The development of divalent batteries necessitates the design and discovery of new cathode materials that
can accommodate the reversible intercalation and deintercalation of the divalent cation via the redox of the
framework. However, the design principles surrounding divalent ion mobility in solid state materials are
limited. A few materials have been reported to function as cathodes for batteries based on divalent metal
intercalation. The canonical example for Mg
2
+
cations is the Chevrel phase Mo
6
S
8
[
6
], which has also
exhibited reversible Zn intercalation [
7
]. Computational results indicate that Mg
2
+
intercalation into the
Chevrel phase is aided by electronic screening of the high charge density of the divalent cation by the
surrounding host lattice, enabled by the metallic nature of the Chevrel material. The importance of a metallic
cathode was supported by the observation of a much higher Mg
2
+
mobility in Ti
2
S
4
than Zr
2
S
4
, which was
attributed to the higher electronic conductivity in the former [
8
].
A range of intercalation cathodes for Zn batteries have been studied [
9
,
10
], including vanadium oxides
[
11
13
], manganese oxides [
14
,
15
], Prussian blue analogues [
16
,
17
], and beyond. Direct comparisons to
© 2021 The Author(s). Published by IOP Publishing Ltd
J. Phys. Mater.
4
(2021) 024005
A J Martinolich
et al
Figure1.
Schematic of V substituted ZnPS
3
. Substitution of V
3
+
for Zn
2
+
is charge compensated by vacancies on the transition
metal sites. The layers above and below are not shown.
the traditional Li intercalation material LiCoO
2
have also been made. The material ZnCo
2
x
Al
x
O
4
exhibits
reversible cycling in nonaqueous electrolytes [
18
]. Al was substituted into the spinel lattice to stabilize the
material at high states of charge (low Zn content). However, at
x
>
0
.
67
no electrochemical activity is
observed, which is attributed to poor electronic conductivity of the material with greater Al substitution,
despite the possible low energy conduction pathways between tetrahedral sites in a spinel lattice [
19
].
In aqueous electrolytes, reversible Zn
2
+
intercalation is facilitated by the coinsertion of water molecules
into the lattice [
12
,
20
]. The presence of water in the V
2
O
5
cathode material, for example, is suggested to
impart structural stability and promote Zn
2
+
mobility. However, such attributes are limited to aqueous
electrolytes and the kinetics of divalent ion conductivity and charge transfer are negatively impacted when
measured in nonaqueous systems [
21
]. Proton intercalation can also occur in aqueous based electrolytes,
which convolutes the charge storage mechanisms in prospective batteries aiming to use Zn
2
+
working ions
[
22
,
23
]. Thus, broadening the design rules for divalent cation cathode materials is necessary to promote the
development of next generation battery technologies.
Despite the suggestions that divalent ionic conductivity requires either high electronic conductivity or
co-intercalated solvent, we recently reported the mobility of Zn
2
+
in the electronically insulating material
ZnPS
3
with a low activation barrier of ionic conduction of 350 meV [
24
]. ZnPS
3
crystallizes in 2D
honeycomb layers of edge sharing Zn octahedra surrounding the P
2
S
4
6
polyanion, with the layers separated
by a van der Waals gap of approx. 3.3 Å. Theoretical data suggest that the flexibility of the polyanion and the
low electron density in the van der Waals gap facilitate Zn
2
+
conduction with a low activation energy. The
material is amenable to a range of cation substitutions [
25
], including aliovalent substitution.
Here, we report the preparation and electrochemical behavior of vanadium substituted ZnPS
3
.
Substitution of V
3
+
for Zn
2
+
is charge compensated by vacancies on the metal site (figure
1
) [
26
,
27
] and
provides us with the ability to tune the electronic properties of a lattice without altering the favorable ionic
properties provided by the structure or anionic chemistry. Introduction of V into the lattice directly impacts
the electronic structure of the material and enables reversible redox as evidenced by cyclic voltammetry (CV)
and galvanostatic cycling data; greater capacity is observed at higher levels of vanadium substitution. At
higher potentials (
>
1.4 V vs. Zn/Zn
2
+
), irreversible oxidation occurs at the polyanion, as indicated by x-ray
photoelectron spectroscopy (XPS). Our results provide a promising new avenue toward the discovery of
functional divalent cathodes, wherein the installation of redox active cations on lattices that allow divalent
cation mobility enables the desired functionality.
2.Experimentalmethods
2.1.Materialpreparation
All manipulations were performed in an Ar-filled glovebox with H
2
O and O
2
concentrations below 1 ppm.
Zn
1
1.5
x
V
x
PS
3
was prepared from Zn (Alfa Aesar, 99.9%), V (Alfa Aesar, 99.5%), P
2
S
5
(Acros Organics,
98
+
%) and S (Acros Organics, 99.5
+
%). The reactants were combined in stoichiometric quantities
corresponding to
x
values of 0.067, 0.20 0.33, 0.40, and 0.50, ground to homogeneity in a mortar and pestle,
compressed into 0.25
′′
cylindrical pellets, and sealed in an evacuated vitreous silica ampoule. The reaction
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A J Martinolich
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mixture was then heated at 2
C min
1
to 400
C and annealed for 24 h. The product was then collected in
the glove box, reground, repelleted, sealed in an evacuated silica ampoule, heated to 600
C at 2
C min
1
and annealed for 24 h. The
x
=
0 end member does not require the second annealing step [
24
]. The resultant
powders were collected in the glove box and stored under inert conditions.
2.2.Powderx-raydiffraction
Powder x-ray diffraction measurements were collected on a Panalytical X’Pert Pro Diffractometer equipped
with a CuK
α
x-ray source. The diffraction patterns were fit by the Rietveld method using GSAS-II [
28
].
Crystal structures were visualized using Vesta [
29
].
2.3.Electrochemicalmeasurements
CV and galvanostatic charge/discharge measurements were collected on composite electrodes of
60 wt% Zn
1
1.5
x
V
x
PS
3
, 20 wt% PTFE binder (Sigma Aldrich), and 20 wt% Super P conductive carbon (Alfa
Aesar,
99%). The composite mixture was ground to homogeneity using a mortar and pestle and pressed
into 0.25
′′
diameter discs. The mass of active material in each electrode was between 5 and 15 mg. The
electrodes were assembled into 2032 coin cells with a spring, stainless steel spacer, 18 mm diameter Whatman
GF-D glass fiber separator, and Zn foil counter/reference electrode (Sigma Aldrich, 99.9%). The Zn foil was
cleaned with a razor blade prior to assembly to remove any oxidized surface species. All measurements used
500 mM Zn(TFSI)
2
(Sigma Aldrich, 95%) in propylene carbonate (Sigma Aldrich,
>
99%) electrolyte
(8 drops, approx. 100 mg). Measurements were collected on a Biologic BCS 805 battery cycler or VMP3
potentiostat.
2.4.X-rayphotoelectronspectroscopy
Samples were oxidized or reduced to the specified voltages using linear sweep voltammetry in preparation for
XPS measurements. The composite electrodes were dried under dynamic vacuum for
>
24 h prior to data
collection. The samples were ground in a mortar and pestle and mounted on carbon tape in the glove box
before transferring to the instrument, with brief air exposure. XPS data were collected using a Surface
Science Instruments MProbe ESCA controlled by Hawk Data Collection software. Element-specific scans
were collected with a resolution of 0.065 eV and a spot size of 500
×
1200
μ
m. The sample chamber was
maintained at
<
2
×
10
8
Torr. The XPS data were analyzed using CasaXPS analysis software. All spectra were
recorded with a neutralizer on to minimize charging effects and referenced to adventitious carbon at 285 eV.
3.Resultsanddiscussion
V substitution into ZnPS
3
is achieved via stoichiometric substitution during initial preparation of the
material. A composition of Zn
1
1.5
x
V
x
PS
3
was targeted to introduce both redox-active V
3
+
centers and
vacancies on the metal sites (figure
1
). The stoichiometric mixture of Zn, V, P
2
S
5
, and S was combined in an
Ar-filled glove box and annealed at 400
C and 600
C to yield a black powder. The powder x-ray diffraction
pattern and quantitative Rietveld refinement are shown in figure
2
(a). The fit corresponds to a single phase
that is isostructural to ZnPS
3
, indicating that partial substitution of V for Zn does not induce a structural
change to the parent material. Mixed valency and partial vacancy substitution in the
M
PS
3
family of
materials is precedented, specifically in the case of V
0.78
PS
3
[
26
]. The color change from white to black upon
partial substitution of V for Zn is a clear indicator of the effect of metal substitution on the electronic
properties of the material.
The refined lattice parameters of Zn
1
1.5
x
V
x
PS
3
(0
x
0
.
50) are plotted as a function of V content in
figure
2
(b). The lattice contracts in all directions upon V substitution, which is expected due to the smaller
ionic radius of V
3
+
than Zn
2
+
, 0.64 Å and 0.74 Å respectively [
30
,
31
], along with the introduction of
vacancies into the lattice. The lattice contracts by 1.1% and 0.9% along
a
and
b
, respectively, but only 0.6%
along
c
. The lesser contraction along
c
is expected due to layered nature of the material, in that the van der
Waals gap will be more affected by the electron density around the sulfur anions than partial substitution in
the metal layer. The contraction of the metal layer may increase the activation barrier of intralayer ionic
conduction due to restricted conduction pathways for the Zn
2
+
. While the structural contraction may
increase the activation barrier, we previously hypothesized that the ionic conduction in ZnPS
3
is mediated by
entropic vacancy defects on the metal site. Thus, increasing the vacancy concentration through aliovalent
substitution likely increases the bulk conductivity of the material. Additionally, the presence of electronic
carriers in the material may also promote divalent ionic conductivity through the lattice, as previously
hypothesized for Mg
2
+
in Mo
6
S
8
and Ti
2
S
4
[
8
,
32
,
33
].
CV was used to understand the redox properties of Zn
1
1.5
x
V
x
PS
3
. CVs with varied positive switching
potentials are shown in figure
3
for the
x
=
0.33 material. The first scan was initially swept positive to 1 V
3
J. Phys. Mater.
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A J Martinolich
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Figure2.
(a) Powder x-ray diffraction of Zn
0.50
V
0.33
PS
3
and single phase quantitative Rietveld refinement. No crystalline
impurities are observed. (b) Lattice parameters of the solid solution Zn
1
1.5
x
V
x
PS
3
. The linear decrease in the lattice parameter
supports the quantitative substitution of V for Zn.
Figure3.
Cyclic voltammetry of Zn
0.50
V
0.33
PS
3
using 500 mM Zn(TFSI)
2
in propylene carbonate electrolyte and a Zn metal
reference/counter electrode. The voltammetry was collected at 0.1 mV s
1
and initially swept to positive potentials. After two
cycles, the upper voltage cutoff was successively increased by 200 mV. The cathodic wave observed at approx. 500 mV shifts to
lower potentials once the second anodic wave is observed.
then negative to 0.1 V. A cathodic wave is observed at approx. 500 mV on the negative sweep, with no
correlated anodic wave observed on the reverse sweep. After reaching the upper voltage cutoff for the first
cycle (1 V), a small wave is observed on the second negative sweep at a similar voltage to the cathodic wave in
4
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scan 1, but with smaller peak current. The smaller peak current indicates that most of the redox-active
material was reduced on the first sweep and very little remains in the material by the second sweep.
After two scans with a switching potential of 1 V, the switching potential was increased to 1.2 V. Sweeping
to more positive potentials above 1 V yields an incomplete anodic wave, which appears to be coupled to the
cathodic wave observed in the initial CVs. On the negative sweep, the cathodic wave shifts to slightly higher
potentials than originally observed to approx. 700 mV. The next CV scan overlays the preceding scan,
indicating that the oxidation and reduction processes are reversible. The peak splitting is wide (530 mV),
however, suggesting that the processes are not Nernstian, which is likely due to overpotentials associated with
ionic conduction or charge transfer resistance. Increasing the anodic voltage cutoff to 1.4 V maintains the
reversibility and yields a more symmetrical wave, with no changes observed between the first and second
cycle.
Sweeping to higher potentials affects the reversibility and redox properties of Zn
0.50
V
0.33
PS
3
. The
reversibility of the redox processes is determined by both the peak splitting and evaluating the charge passed
on oxidation vs. the charge passed on reduction with the capacitive background subtracted. A less sensitive
method is to evaluate the peak current of the anodic and cathodic waves and this can give a general idea of
reversibility. If a process were reversible, the peak current would be the same for both waves. Deviations from
such behavior indicate some degree of irreversibility. After twice sweeping to a switching potential of 1.4 V,
two CVs each were swept to 1.6, 1.8, 2.0, and 2.2 V (figure
3
). The CVs to 1.6 V begin to show an increase in
the current at the highest voltages on both cycles. On the reverse sweep, the cathodic wave changes shape and
begins to broaden, indicating alteration to the kinetics of the redox properties of the material, which could be
due to either a structural transformation or distortion yielding poorer Zn
2
+
mobility through the lattice.
Upon sweeping to a positive switching potential of 1.8 V, a second oxidative wave is clearly observed. The
second oxidation precedes a broad reduction wave that is shifted to lower potentials than originally observed,
centered at approx. 450 mV. The second oxidation also affects the oxidation processes in later cycles. The
original anodic wave observed below 1.4 V broadens and shifts to higher potentials, mirroring the behavior
of the cathodic wave. The redox behavior for the
x
=
0.33 material is also observed for the
x
=
0.20 and
x
=
0.40 materials (see the SI available online at
stacks.iop.org/JPMATER/4/024005/mmedia
). Slightly higher
current densities are observed for the first oxidative wave for the
x
=
0.40 material, suggesting that the first
wave is related to vanadium redox. For
x
=
0.067 and
x
=
0.50, much smaller oxidative waves are observed
before the large, irreversible wave above 1.4 V, suggesting that a threshold amount of vanadium is required to
enable redox activity, while too much changes the properties of the material and disallows reversible redox.
Voltammetry for all materials are shown in the SI.
XPS spectra in the Zn, V, P, and S regions were collected on the
x
=
0.33 material at various oxidation
states after voltammetry to determine the redox mechanism of Zn
1
1.5
x
V
x
PS
3
(figure
4
). The material was
either reduced to 0.1 V or reduced to 0.1 V then subsequently oxidized to 1.4 or 2.0 V at 0.1 mV s
1
. The Zn
binding energy remains unchanged in all cases (figure
4
(a)). The V binding energies shift as a function of
oxidation and reduction, indicating the direct involvement of the metal in the redox (figure
4
(b)). The V
binding energy of the pristine material is centered at 514 eV, which then shifts to 513 eV upon reduction to
0.1 V. Subsequent oxidation to 1.4 V shifts the V binding energy back to 514 eV and is nearly identical to the
pristine sample, indicating reversible V redox below 1.4 V. Subsequent oxidation to 2 V results in the
observation of a second peak at higher binding energy (517.5 eV), indicating further oxidation of the V at 2
V [
34
]. The low signal-to-noise ratio from V in the XPS spectra is likely a combination of low V content in
the sample and the comparatively small relative sensitivity factor.
XPS of the P and S in Zn
0.50
V
0.33
PS
3
were collected to understand the changes in anion oxidation state
when the material is oxidized or reduced (figures
4
(c) and (d)). The new peak denoted with an asterisk in the
S region of the XPS spectra after oxidation or reduction is attributed to adsorbed TFSI
anions, which
remain after exposure to the electrolyte. Upon reduction to 0.1 V or oxidation to 1.4 V, no major changes are
observed in either the P or the S regions of the XPS spectra, indicating that the polyanion is stable below
1.4 V. However, upon oxidation to 2 V, both spectra exhibit a shoulder at higher binding energies suggesting
polyanion oxidation. Based on the XPS data, general assignments to the different redox properties observed
in the CV (figure
3
) can be made. V in reduced during the initial reduction process at 500–700 meV and is
then reversibly oxidized at approx. 1.2 V. Sweeping to voltages
>
1.4 V yields a second anodic wave that we
assign to polyanion oxidation.
To confirm the contribution of the redox from the [P
2
S
6
]
4
polyanion, CV was executed on
unsubstituted ZnPS
3
. No cathodic wave is observed on the initial negative sweep. However, a single oxidative
wave is observed at approx. 1.7 V (figure S3). The oxidative wave is in the same position as the irreversible
wave observed in the V substituted material. The observation of an irreversible wave in both materials
suggests that the [P
2
S
6
]
4
polyanion is oxidized in both cases. Irreversible polyanion oxidation is likely to
5
J. Phys. Mater.
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A J Martinolich
et al
Figure4.
(a) Zn, (b) V, (c) P, and (d) S regions of the XPS spectra of Zn
0.50
V
0.33
PS
3
at various states of oxidation. Reduction to 0.1
V only affects the XPS spectrum in the V region, while oxidation to 2 V yields changes in the V, P, and S regions. The peak at
168 eV in the S 2p region of the spectra is due to the TFSI
counteranion present in the electrolyte, denoted with (
).
significantly alter the structure and shift the subsequent redox features, as observed in figure
3
and in
electrochemical cycling experiments at high potentials. Various oxidation states of [P
x
S
y
]
n
polyanions are
accessible beyond [P
2
S
6
]
4
, including the [P
2
S
6
]
2
in the mixed polyanionic compound Zn
4
(P
2
S
6
)
3
and
[PS
4
]
3
in Zn
3
(PS
4
)
2
[
35
,
36
]. Zn is tetrahedrally coordinated in both materials, which would require
significant structural rearrangement from the structure of ZnPS
3
, as suggested by the changes in the
electrochemistry after oxidation beyond 1.4 V.
The reversibility of the V centered redox is supported by rate dependent voltammetry between 1.4 and
0.1 V (figure
5
). Upon increasing the scan rate from 0.1 to 10 mV s
1
, the peak current of the reversible
anodic and cathodic waves increases as well. The current is linearly correlated to the square root of scan rate,
indicating electrochemical reversibility of the low voltage redox couple. Above 1 mV s
1
, the voltage of the
waves begins to shift, indicating some irreversibility which may be related to slower kinetics of Zn diffusion
through the Zn
1
1.5
x
V
x
PS
3
lattice rather than in the nonaqueous electrolyte. However, changing the
concentration of the electrolyte does lead to a slight decrease in the maximum observed current (figure
5
(b)),
suggesting that the kinetics of Zn mobility in the solid state lattice are not completely resolved from kinetic
limitations in the liquid electrolyte.
The presence of Zn
2
+
cations, vacancies, and redox active V centers in the pristine material suggests that
the material can be either oxidized or reduced as synthesized. Thus, CVs were measured on Zn
0.50
V
0.33
PS
3
initially sweeping to either positive or negative potentials (figure
6
). Indeed, we observe that the material can
be either oxidized or reduced from its pristine state, indicating that Zn incorporation is possible without the
necessity of charging first. When sweeping to more positive potentials first, Zn must be removed from the
octahedral sites within the metal layer as V is oxidized. When sweeping to more negative potentials initially,
it is possible that Zn could occupy either the vacancies in the metal layer, introduced due to aliovalent
substitution, or vacancies in the van der Waals gap as the V is reduced. On the return sweep in both cases,
the correlated wave increases in current density. The increase in current density on the return wave provides
further support that the redox features at 0.6 and 1.1 V are related to the same electrochemical processes and
are reversible. The correlation between the waves also provides indirect evidence that Zn is being introduced
into the metal layer rather than into the van der Waals gap, as it may only originate from the transition metal
when sweeping to more positive potentials initially, and no changes in potential are observed on subsequent
cycles.
6
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Figure5.
(a) Scan rate dependent cyclic voltammetry of Zn
0.50
V
0.33
PS
3
. (b) The linear relationship between the square root of the
scan rate and the maximum observed current indicates the reversibility of the reduction and first oxidation observed. Using a less
concentrated electrolyte slightly decreases the observed maximum current, indicating that the observed kinetics cannot be
isolated from the kinetics of the electrolyte.
Figure6.
CVs of Zn
0.50
V
0.33
PS
3
swept to either positive or negative potentials from open circuit potential, indicating the
intermediate state of charge of the pristine material.
Based on the observation of the reversible, low voltage redox wave in the voltammetry, Zn
1
1.5
x
V
x
PS
3
was cycled galvanostatically with a current density of 1 mA g
1
between 1.4 and 0.1 V, starting first with the
charge cycle (figure
7
). The first cycle exhibits a Coulombic efficiency much greater than 100% (defined here
as
Q
out
/
Q
in
or
Q
discharge
/
Q
charge
) because Zn is able to occupy the vacancies present in the pristine material
upon discharge. A
>
100% Coulombic efficiency on cycle 1 is observed for
x
=
0.20,
x
=
0.33, and
x
=
0.40. The second cycle exhibits much higher capacity on the charge, and a comparable capacity on discharge,
indicating that the extra Zn
2
+
introduced during discharge can be removed. The materials exhibit reversible
capacity that is dependent on the V content, with greater V content yielding higher capacity. The correlation
between capacity and V content further supports that the reversible charge storage is enabled by V redox.
7
J. Phys. Mater.
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A J Martinolich
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Figure7.
Galvanostatic cycling data of Zn
1
1.5
x
V
x
PS
3
. (a) The discharge and charge curves of the first two cycles of the
x
=
0.20,
0.33, and 0.40 materials with a current density of 1 mA g
1
. The materials are first charged then discharged. (b) The capacity
normalized to the number of vacancies in the materials plotted as a function of cycle number. All materials exhibit a maximum
reversible capacity of half of the available vacancies.
The capacity over several cycles is normalized to the vacancy content and plotted in figure
7
(b). The
materials show gradual capacity fade followed by more stable cycling. When normalized to the number of
vacancies in the material, we observe that all materials are reversibly accessing approximately half of the
vacancies present in the vanadium substituted materials. Such an observation elucidates the nature of
divalent cation intercalation and mobility in the lattice. Filling half the vacancies on discharge correlates to
reduction of half of the V
3
+
to V
2
+
. Once half of the V
3
+
is reduced, Zn incorporation ceases, even if the
vacancy concentration remains relatively high. For example, filling half the vacancies in the
x
=
0.40
material leaves 0.1 equivalents of unfilled vacancies per formula unit. The same number of vacancies are
present in the pristine
x
=
0.20 material, however, the presence of 0.1 equivalents of vacancies alone does
not provide the same electrochemical activity. Instead, it appears that the mixed valency on the substituted
vanadium centers (dis)allows the reversible redox. Such behavior highlights the delicate balance between
electronic and crystal structure that may influence divalent ion based energy storage materials. Capacity
retention data for the
x
=
0.20, 0.33, and
x
=
0.40 materials with replicate cells are shown in the SI.
Zn
0.50
V
0.33
PS
3
was cycled at 60
C in an attempt to increase the experimental capacity of the material.
While the hysteresis between the charge and discharge curves decreases, no significant increase in capacity is
observed (figure S7). Additionally, the Coulombic efficiency begins to decrease drastically at later cycles,
eventually giving way to parasitic current on the charge curve and cell failure. We hypothesize that the poor
Coulombic efficiencies and cell failure are due to slow oxidation of the [P
2
S
6
]
4
polyanion, which is accessed
at lower potentials when higher temperatures are applied. The inability to access greater capacities at elevated
temperatures suggests that the capacity is thermodynamically limited to filling half of the vacancies (approx.
26 mAh g
1
for
x
=
0.33) and it is not an issue of kinetics.
While the materials with
x
=
0.20, 0.33, and
x
=
0.40 all exhibit reversible Zn storage upon galvanostatic
cycling, no appreciable reversible oxidation is observed for the
x
=
0.067 or
x
=
0.50 materials at room
temperature or at 60
C (figures S4–S6). For Zn
0.90
V
0.067
PS
3
, it is likely that the low vanadium content
and even lower vacancy concentration in the material leads to no appreciable charge storage
(
Q
theoretical
=
9
.
5
mAh g
1
if 100% of the vacancies are filled). However, the material with the most V also
exhibits no reversible charge storage. Multiple hypotheses can be provided to account for this unexpected
behavior. It is possible that when a majority of the metal centers in the material are V instead of Zn, the
electronic structure of the material changes greatly and disallows vanadium oxidation and reduction. For
example, V
0.9
PS
3
is a Mott insulator, where both the highest occupied valence band and lowest unoccupied
conduction band are composed of V
d
states [
37
,
38
]. At high V content it is possible the materials become
poorer electronic conductors which leads to electrochemical inactivity. An alternative hypothesis is that there
is a threshold amount of Zn necessary in the lattice to support divalent ion mobility. When the amount of V
8
J. Phys. Mater.
4
(2021) 024005
A J Martinolich
et al
is greater than the amount of Zn, it is likely that the percolative pathways for Zn
2
+
mobility and conduction
are disrupted and thus the ions are immobilized [
39
]. However, all materials exhibit irreversible oxidation
above 1.4 V, further supporting that the second anodic wave is due to the oxidation of the [P
2
S
6
]
4
anion but
also confirming that Zn
2
+
can be removed from the material. Therefore, it is unlikely that the high V
content material is inactive due to sluggish Zn
2
+
conduction.
4.Conclusion
Here, we have shown that V substitution into the Zn
2
+
cation conductor ZnPS
3
yields reversible redox
activity. The introduction of vacancies allows for either electrochemical oxidation or reduction from the
pristine state. The storage capacity is directly correlated to the V content of the material, suggesting that the
V centers are redox active, which is supported by voltammetry in concert with XPS measurements at various
redox states. The charge storage capacity of the materials is limited to half of the synthetically-introduced
vacancies, which is hypothesized to be due to the ratio of oxidized and reduced V centers and their influence
on the electronic structure of the material. These insights provide a useful new design principle for
prospective cathode materials for divalent working ions, and highlight the importance of both electronic
structure as well as a suitable host lattice for divalent cation conduction and mobility in the pursuit of next
generation energy storage materials.
Acknowledgments
This research was supported by the Arnold and Mabel Beckman Foundation through the Beckman Young
Investigator Award. The authors thank Dr FAL Laskowski for many enlightening discussions. AJM was
partially supported through a postdoctoral fellowship from the Resnick Sustainability Institute at Caltech.
XPS data were collected at the Molecular Materials Research Center in the Beckman Institute at Caltech.
ORCIDiDs
Andrew J Martinolich
https://orcid.org/0000-0002-7866-9594
Skyler D Ware
https://orcid.org/0000-0002-3249-1946
Brian C Lee
https://orcid.org/0000-0002-0898-0838
Kimberly A See
https://orcid.org/0000-0002-0133-9693
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