Article
pubs.acs.org/JPCC
Potential-Resolved In Situ X‑ray Absorption Spectroscopy Study of
Sn and SnO2 Nanomaterial Anodes for Lithium-Ion Batteries
Christopher J. Pelliccione,*,† Elena V. Timofeeva,‡ and Carlo U. Segre*,†
†
Department of Physics & CSRRI, Illinois Institute of Technology, 3101 S. Dearborn St., Chicago, Illinois 60616, United States
Energy Systems Division, Argonne National Laboratory, Argonne, Illinois 60439, United States
‡
S Supporting Information
*
ABSTRACT: This work provides detailed analysis of processes occurring in metallic Sn
and SnO2 anode materials for lithium ion batteries during first lithiation, studied in situ
with rapid continuous X-ray absorption spectroscopy (XAS). The X-ray absorption near
edge structure (XANES) and extended X-ray absorption fine structure (EXAFS) spectra
provide information on dynamic changes in the Sn atomic environment, including type
and number of neighboring atoms and interatomic distances. A unique methodology was
used to model insertion of Li atoms into the electrode material structure and to analyze
the formation of SnLi phases within the electrodes. Additionally, analysis of fully lithiated
and delithiated states of Sn and SnO2 electrodes in the first two cycles provides insight
into the reasons for poor electrochemical performance and rapid capacity decline. Results
indicate that use of SnO2 is more promising than metallic Sn as an anode material, but
more effort in nanoscale and atomic engineering of anodes is required for commercially
feasible use of Sn-based materials.
■
INTRODUCTION
Lithium-ion batteries (LIBs) are the standard power source for
portable electronic devices such as cell phones, laptops, tablets,
etc. Graphite is a common anode material in LIBs because of its
long cycle life and adequate specific capacity (theoretical
capacity 372 mAh/g).1,2 However, for LIBs to progress and
become an efficient and cost-effective option for electric
vehicles, significant improvements need to be made in the
energy and power densities of current electrode materials.3−5
Metallic tin is an attractive alternative to graphite because of the
nearly three times higher theoretical capacity (994 mAh/g)
enabled by the formation of Li−Sn alloys. However, due the
large number of lithium atoms involved in the discharging
process (4.4 Li for every Sn atom), the initial crystal structure
of metallic tin undergoes significant alterations, including a ca.
260% volumetric expansion that severely degrades the anode’s
crystallinity, resulting in a rapidly declining capacity.6 For tin to
successfully supplant carbon-based anodes, these catastrophic
volumetric changes must be controlled to achieve extended
cycle life with enhanced specific capacities.
Using oxidized tin phases has been shown to significantly
improve the longevity of Sn-based anodes through the
conversion of SnO2 to metallic Sn and Li2O during the first
lithiation/delithiation cycle. Although SnO2 exhibits a lower
theoretical capacity than metallic Sn (790 and 994 mAh/g,
respectively),7,8 the cycle life is significantly better, with the
reversible capacity on the 20th cycle showing ca. 70%
improvement over metallic Sn.9,10 These structural changes
during Li alloying with SnO2 result in a primarily amorphous
environment in which the Li2O network provides rigid
structural support for metallic tin clusters embedded within,
© 2016 American Chemical Society
controlling the stress of the volumetric expansions during the
lithiation process.11−13 Common techniques to study battery
materials such as X-ray diffraction (XRD) require long-range
crystalline order to acquire representative spectra, thus are
generally ineffective to study greatly amorphized systems.14−16
A technique that can adequately account for both amorphous
and crystalline contributions is needed to have a complete
understanding of the mechanism of lithium insertion/removal
and the effect it has on the observed capacity.
X-ray absorption spectroscopy (XAS) is an ideal technique to
study the structural changes of battery electrodes because it is
element specific so it can probe only elements of interest but
also sufficiently accounts for both the crystalline and
amorphous phases present within the sample.17−24 XAS,
specifically X-ray absorption near-edge structure (XANES)
spectroscopy and extended X-ray absorption fine structure
(EXAFS) spectroscopy, provide detailed information about the
local electronic and atomic environment (<6 Å) around each
absorbing atom. The XANES region is primarily responsive to
changes in the electronic structure of the absorbing atom (i.e.,
changes in oxidation state), while the EXAFS region indicates
the neighboring atomic structure. Theoretical models used to
interpret the EXAFS region of the spectra help deduce the
coordination numbers, interatomic distances, and element
species surrounding the absorbing atom as the material is
lithiated and delithiated.
Received: December 15, 2015
Revised: February 10, 2016
Published: February 17, 2016
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Photon Source at Argonne National Laboratory. Both the Sn
and SnO2 electrodes were placed in pouch-type electrochemical
cells with Li metal as the reference/counter electrode and
measured in transmission mode at the Sn K−edge (29.2 keV)
with a reference Sn foil measured simultaneously for proper
alignment of multiple scans. For in situ tests, a custom sample
chamber was used to ensure the pouch cell was under inert
atmosphere by flowing He through a sealed chamber.28 Pouch
cell preparation and mounting within the sample chamber were
conducted inside an Ar atmosphere glovebox (VAC). Since this
study is based around Sn materials in a half-cell configuration
with the more negative metallic Li as the counter electrode,
designations of “discharge” for lithiation and “charge” for
delithiation will be used throughout the text. In the context of a
full cell, where Sn-based materials would be used as anodes, the
nomenclature would be reversed. During the in situ XAS
experiment, both the Sn and SnO2 electrodes were lithiated/
delithiated galvanostatically at a rate of 124 mA/g (C/8)
between 0.01−1.50 V versus Li/Li+ using an EzStat Pro
potentiostat/galvanostat system (Nuvant Systems, Inc.). Before
the first lithiation, the electrodes were held at open circuit
voltage (OCV, ca. 3.0 V) to measure the spectra of the pristine
electrodes. During the following lithiation and delithiation
process, XAS measurements were continuously acquired with
each individual XAS scan lasting less than 2 min. Since each
XAS scan was significantly shorter than the total duration of the
discharge, a potential-resolved XAS analysis was conducted on
the first lithiation for both Sn and SnO2 electrodes. The fully
discharged and charged states of the first two cycles were also
analyzed. In those tests, the potential was held at 0.01 V for the
discharged states and 1.50 V for the charged states and multiple
scans (3−4) were taken at each discharged/charged state in
order to reduce the noise in the XAS measurements.
The EXAFS spectra were aligned, merged, and normalized
using Athena.29,30 The built-in AUTOBK algorithm was used to
minimize background below Rbkg = 1.0 Å. Each spectrum was fit
using Artemis with theoretical models of β-Sn metal,31 SnO2,32
and a three Sn−Li path (“short”, “medium”, and “long”) model
derived from several Li−Sn phases, including the fully lithiated
Li22Sn5 phase, previously developed by our group11 to obtain
quantitative details about the neighboring lithium atoms.33−37
As the tin-based electrode is converted from metallic tin to the
fully lithiated Li22Sn5, it goes through several Li−Sn phases
such as LiSn, Li2Sn5, etc. When in these initial Li−Sn phases,
the lithium atoms are exclusively observed at medium and long
distances from the tin atoms. As the electrode gradually
converts to the Li22Sn5 phase when fully discharged, the short
Sn−Li distances are also observed. Using this distinctive
signature in the location of the lithium atoms for the initial and
final states of the electrode, we are able to track the state of
discharge by comparing the relative ratios of lithium atoms at
short (ca. 2.76 Å), medium (ca. 2.90 Å), and long (ca. 3.30 Å)
distances away from the tin atoms.
All the theoretical EXAFS models were constructed starting
with the simplest interpretation and adding complexity only if
needed to achieve a reasonable fit to the experimental data. For
both the Sn and SnO2 series of potential-resolved EXAFS
modeling, all spectra were initially fit with the pristine crystal
structure. When that model could no longer adequately
describe the experimental data, the Sn−Li model was included.
If there was a path that resulted in parameters with no physical
significance (unusually high number of near neighbors, negative
Debye−Waller factor, etc.) or large estimated standard
XAS has not been used to investigate lithiation of pure
metallic Sn anodes; however, EXAFS studies have been
previously conducted on SnO2-based anodes to gain insight
on the specifics of the structural changes the electrode
undergoes during initial discharge/charge cycles. Kim et al.25
were able to directly observe the evolution of SnO2 into a
metallic tin phase during the first lithiation. Kisu et al.26 also
conducted a similar EXAFS study on 2−4 nm SnO 2
nanoparticles at various points of the first lithiation and
delithiation processes with similar results; through changes
observed around the Sn atoms, it is deduced that the oxygen is
converted to Li2O early in the charging process with the
simultaneous evolution of a metallic tin phase. In the fully
lithiated state, the metallic tin phase disappears from the spectra
and then returns upon delithiation.
These studies provide insight into the general changes
occurring within SnO2 electrodes during the lithiation/
delithiation process; however, they do not report any
theoretical modeling of the EXAFS nor do they discuss direct
observation of lithium atoms in the lithiated electrode material.
Pelliccione et al. have previously shown that detailed XAS
modeling results can be achieved including direct observation
of neighboring lithium atoms in anode materials.11,27 In this
work, we report a comparative study of capacity fading
mechanisms in metallic Sn and SnO2 electrodes using in situ,
continuous, rapid scanning XAS. This approach permits the
collection of XAS spectra directly correlated to a specific
electrode potential to investigate the structural changes as
lithiation progresses. Theoretical models fitted to the resulting
spectra allow the potential-specific determination of coordination number and interatomic distances. The direct observation
of lithium insertion into the electrodes is also included in the
theoretical modeling. Potential-resolved in situ EXAFS
measurements were conducted on both Sn and SnO2 electrodes
during the first lithium insertion, as well as in the fully
discharged and charged states of the first two cycles. Our results
provide comparative, quantitative details of the mechanism of
lithium insertion and resulting structural changes in Sn and
SnO2 nanoparticle electrodes which are consistent with
previous qualitative studies.25,26
■
METHODS
Electrode Preparation. Electrodes of Sn (Aldrich, 576883;
less than 150 nm) and SnO2 nanoparticles (MTI Corporation,
NP-SnO2; 50 nm) for electrochemical characterization were
casted onto 0.03 mm thick copper foil current collectors. The
suspension used for electrode casting was prepared from a
80:10:10 wt % mixture of active material (Sn or SnO2),
poly(vinylidene fluoride) binder (Aldrich, 24937-79-9) and
acetylene carbon black (STREM Chemicals, 138-86-4),
respectively, dispersed in 3 mL of 1-methyl-2-pyrrolidinone
(Aldrich, 872-50-4) per gram of solid. The suspension was
sonicated and mixed for over 24 h then applied to copper foil
with a doctor blade film casting knife (MTI Corporation, EQSe-KTQ-50) to create a uniform deposition thickness. The
resulting deposition was dried in an oven at 80°−100 °C for 3 h
and then calendered to improve the electrode performance.
The deposition density of active material was determined
through measuring the total change in mass after deposition
and normalizing per surface area of deposition.
X-ray Absorption Spectroscopy. XAS spectra were
acquired at the Materials Research Collaborative Access
Team (MRCAT) beamline, Sector 10-ID, of the Advanced
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deviation, it was excluded from further fitting of that spectrum.
All possible combinations of Sn−O, Sn−Sn, and Sn−Li paths
were explored to ensure the correct model was being used. All
spectra were fit using k, k2, and k3 weightings simultaneously to
confirm the fitting model was complete in a k-range of 2−10
Å−1. The R-space window was defined as 1.0−3.4 Å to fully
encompass the first and second atomic shells around the
absorbing atom. If no oxygen was present, the R-range was
reduced to 1.7−3.4 Å for analysis of the second shell.
■
RESULTS AND DISCUSSION
Figure 1 shows the discharge/charge capacity curves of the first
two cycles from both the Sn and SnO2 in situ electrodes. Both
Figure 2. Three-dimensional (3D) representation of selected |χ(R)|
plots for the metallic Sn electrode on the first lithiation cycle as a
function of electrode potential.
The XANES region of the XAS spectra indicate a change in
the average oxidation state of the Sn atoms (Figure 3). At 1.0 V,
Figure 1. Discharge and charge capacity curves for the first (solid line)
and second (dashed line) cycles for Sn (black) and SnO2 (red)
electrodes, during in situ XAS measurements.
the Sn and SnO2 exceed their respective theoretical capacities
on the first lithiation (discharge) most likely because of
formation of the solid-electrolyte interface (SEI). On the
second cycle, the metallic Sn electrode has a low Coulombic
efficiency of ca. 36%, while the SnO2 electrode has a Coulombic
efficiency of ca. 80%. This discrepancy in reversible capacity is
an indication of the state of the electrode, a highly reversible
capacity indicates lithium atoms are inserted and removed with
almost the same efficiency. A low Coulombic efficiency
indicates the lithium atoms that are inserted during discharging
are not removed during charge, suggesting the electrode has
started to degrade.
XAS Results of the First Lithiation of Metallic Sn
Nanoparticles. From the data collected during the first
lithiation of the metallic Sn electrode, 20 EXAFS spectra
corresponding to various states of lithiation (discharge) were
analyzed, and a selected number of these are shown in Figure 2
as a 3D representation of |χ(R)| [Fourier transform of k2χ(k)]
as a function of the electrode potential. The initial state of the
electrode clearly corresponds to the crystal structure of metallic
Sn. As the electrode lithiation progresses, the amplitude of the
metallic Sn−Sn peak at ca. 2.9 Å is quickly reduced, indicating
an amorphization of the crystal structure. When the electrode
potential reaches ca. 0.40 V the original sharp metallic Sn−Sn
peak broadens, and when the electrode is at 0.01 V, the peak
transitions into two broad peaks spanning the range of 1.5−3.0
Å. These peaks indicate large structural disorder in the
electrode and particularly the appearance of new neighboring
atoms between ca. 2−3 Å.
Figure 3. XANES of the metallic Sn electrode as a function of
potential in 0.20 V steps. Inset depicts shifts in the edge position.
the electrode is in the initial metallic Sn state as the edge energy
[defined as the maximum of the first derivative of xμ(E)] is
exactly at 29200 eV. As the material is being lithiated, the edge
position gradually shifts to lower energy, reaching ca. 29198 eV
at 0.40 V and ca. 29196 eV at 0.01 V. Since the edge position at
0.01 V is lower than metallic Sn, this indicates that Sn is now in
a Sn−-type oxidation state (i.e., Li+ has alloyed with Sn).
To quantitatively determine the structural changes during
first lithiation, the EXAFS spectra shown in Figure 2 were
modeled using theoretically calculated contributions from Sn−
O, Sn−Sn, and Sn−Li neighbors and the results are shown in
Figure 4. In the initial OCV state for metallic Sn nanoparticles,
there is a small contribution due to oxygen, which is assumed to
be surface oxidation of the nanoparticles as no oxidized phases
were present in the XRD patterns of the electrode starting
materials (Figure S1). This Sn−O contribution is only present
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Figure 5. Number of neighboring lithium atoms at the short, medium,
and long Sn−Li distances determined from EXAFS fits of the first
lithiation of the metallic Sn electrode as a function of potential.
modeling as 2.71 ± 0.03, 2.96 ± 0.03, and 3.36 ± 0.02 Å for the
short, medium, and long Sn−Li paths, respectively. The
number of near neighbors was allowed to vary freely for each
individual Sn−Li path. In the potential range from 0.6 to 0.45
V, only medium and long Sn−Li paths could be included in the
fits. Once the electrode potential reaches 0.45 V, Sn−Sn
distances are stabilized in an expanded state and a contribution
from the short Sn−Li path is first observed. As the electrode
approaches 0.05 V, a transition is observed in the relative ratio
of Sn−Li neighbors as a reduction in the number of lithium
atoms located at the longer distance and an increase in the
number of lithium atoms located at the shorter distance while
the total number of Sn−Li neighbors remains the same. This is
also consistent with the decrease in Sn−Sn near neighbors, as
lithiation progresses packing more Li into the crystalline
structure. Once fully discharged, the number of lithium atoms
at each distance is 3.1 ± 0.4, 5.0 ± 0.5, and 3.0 ± 1.6 for the
short, medium and long Sn−Li paths, respectively. If all the tin
had been fully converted to the Li22Sn5 crystalline phase, those
numbers should be 2, 6.5, and 5.5 Li atoms, respectively. Thus,
our experimental results indicate a significant degree of disorder
within the LiSn alloy in the most lithiated state.
XAS Results of the First Lithiation of SnO2 Nanoparticles. Analysis of the SnO2 electrode is carried out in much
the same way as for the pure Sn described above. Selected
|χ(R)| of the SnO2 electrode during the first lithiation are
shown in the 3D plot in Figure 6. At 1.00 V, the initial SnO2
crystal structure with both a Sn−O peak at ca. 1.4 Å and a Sn−
Sn peak at ca. 3.8 Å are intact. As the electrode potential is
reduced to 0.80 V, there is a significant reduction in the
intensity of both these peaks. By 0.60 V, a new Sn−Sn peak,
characteristic of metallic Sn, appears at ca. 2.8 Å with low
intensity. As lithiation continues, the Sn−O and metallic Sn−
Sn peaks continue to lose intensity, and once in the fully
discharged state, there are only two broad peaks observed
between 1.5 and 3 Å. These trends are in agreement with
previous XAS studies on tin oxide anodes.25,26
Figure 7 shows the XANES region of the spectra as a
function of electrode potential for SnO2. The Sn oxidation state
changes correlate to changes observed in the |χ(R)|. At 1.00 V,
the edge is at 29203 eV, and by 0.70 V, the edge has shifted to
Figure 4. Number of near neighbors and corresponding interatomic
distances for Sn−Sn (black), Sn−O (blue), and combined Sn−Li
(red) contributions determined from EXAFS fits of the first lithiation
of the metallic Sn electrode as a function of potential.
during first lithiation with very low number of near neighbors
(0.5 ± 0.1 oxygen atoms) and is not observed in later cycles.
The Sn−Sn contribution was modeled with the first two Sn−
Sn paths (3.01 and 3.16 Å) of the I41/amd metallic Sn crystal
structure.31 The pristine metallic Sn structure remains
unchanged until 0.65 V, corresponding to the first plateau
observed in the first discharge capacity curve (Figure 1).
Further discharging to 0.60 V results in a significant expansion
in both the short and long Sn−Sn interatomic distances (from
3.01 ± 0.01 Å to 3.05 ± 0.01 Å in the shorter and from 3.16 ±
0.01 Å to 3.20 ± 0.01 Å in the longer Sn−Sn paths). Between
0.60 and 0.40 V, the Sn−Sn distances are stable at these
expanded values, but once the electrode reaches 0.40 V, both
Sn−Sn distances begin to contract from 3.05 ± 0.01 to 2.91 ±
0.01 Å and from 3.20 ± 0.01 to 3.10 ± 0.01 Å for short and
long distances, respectively, at 0.01 V. The number of
neighboring Sn atoms continuously drops from 5.0 ± 0.4
atoms in the initial OCV state to 1.0 ± 0.3 at 0.01 V.
Sn−Li contributions begin to appear at 0.60 V with a total of
3.6 ± 2.3 neighboring lithium atoms from all the relevant Sn−
Li paths. The emergence of lithium at 0.60 V is in agreement
with previous in situ XRD studies that observe the evolution of
a Li−Sn crystal phase at around the same point in the lithiation
process.16 The total number of lithium atoms increases until
reaching a stable value of ca. 12 near neighbors at 0.40 V. In a
fully lithiated (0.01 V) state, we observe 11.2 ± 1.7 lithium
neighboring atoms. If the system was fully converted to Li22Sn5,
a total of 14 lithium near neighbors would be expected.
Figure 5 shows the number of near neighbor fitting results
for each of the three Sn−Li paths used to fit all spectra. The
distances for each path were determined by fitting the most
lithiated (0.01 V) state and setting those distances for all further
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Figure 6. 3D representation of |χ(R)| plots of SnO2 nanoparticle
electrode during the first lithiation.
Figure 8. EXAFS fitting results showing near neighbor and distances
for Sn−O (blue ■), Sn−Sn (black ●) and Sn−Li (red ▲) paths on
the first lithiation of the SnO2 electrode as a function of potential.
V, no Sn−Sn neighboring atoms are observed within the fitting
window of 3.5 Å. Once the electrode approaches 0.70 V, Sn−
Sn neighbors with distances characteristic of the metallic tin
phase emerge. This metallic tin is in the form of very small
clusters as deduced from the reduced number of neighboring
atoms (1.0 ± 0.6 Sn atoms opposed to 6 for bulk metallic Sn)
and a shifted interatomic distance (2.91 ± 0.02 Å) compared to
bulk metallic Sn (3.02 Å). As the lithiation continues, the Sn−
Sn distances and number of near neighbors are very stable with
no appreciable variation. At the end of the first lithiation,
beyond 0.05 V, there are no significant contributions either
from Sn−Sn or Sn−O neighbors in the spectra.
Significant contributions from Sn−Li bonds are first
observed in the SnO2 electrode at 0.40 V, much further along
in the lithiation process compared to the first observation of
Sn−Li contributions at 0.80 V in the metallic Sn electrode
(Figure 4). The number of Sn−Li neighbors linearly increases
with lithiation from 0.35 to 0.10 V, and at 0.05 V, a significant
jump in the total number of Sn−Li neighbors from 5.2 ± 0.9 to
8.0 ± 0.8 is observed, which also coincides with the
disappearance of residual Sn−O and Sn−Sn contributions.
When SnO2 is in the most lithiated state, 8.2 ± 0.9 lithium
atoms are observed around the tin atoms. Figure 9 shows the
number of neighboring lithium atoms at the short (2.71 ± 0.03
Å), medium (2.93 ± 0.03 Å), and long (3.43 ± 0.03 Å) Sn−Li
distances. In this SnO2 electrode, lithium first appears at short
and medium distances, as opposed to the medium and long
distances in the metallic Sn electrodes (Figure 5). The long
distance Sn−Li neighbors are observed starting from 0.10 V,
when the electrode is close to the most lithiated state. This
divergence in the process of lithiation is likely due to the
smaller sizes of the metallic tin clusters that form in the SnO2
system and also formation of a Li2O network surrounding the
tin sites.
Figure 7. XANES of SnO2 particles as a function of potential in 0.2 V
steps. Inset depicts edge position shifts.
29200 eV corresponding to the metallic-like state. This is in
agreement with the evolution of the Sn−Sn distances seen in
Figure 6. When the electrode is in the most lithiated state (0.01
V), the edge moves to 29198 eV similar to the results for the
metallic Sn electrode, indicating the formation of a LiSn alloy.
Figure 8 displays the EXAFS fitting results for near neighbors
and distances of the Sn−O and Sn−Sn paths along with the
total number of Sn−Li near neighbors. The number of
neighboring oxygen atoms drops very quickly in the early
stages of lithiation, dropping from 5.9 ± 0.5 oxygen atoms at
1.00 V to 2.8 ± 0.3 neighboring oxygen atoms at 0.80 V. The
number of oxygens continues to decrease until 0.05 V, where
Sn−O paths are no longer observed. Throughout the first
lithiation process, the distance between tin and oxygen is quite
stable, beginning at 2.05 ± 0.01 Å and decreasing to 2.02 ±
0.03 Å at 0.05 V.
The Sn−Sn distances of pristine tetragonal SnO2 crystal
structure32 remain unchanged until 0.85 V; however, the
number of Sn−Sn neighbors continuously decreases from ca. 4
to ca. 1 in this segment. Below this potential, the long-range
Sn−Sn neighbors (ca. 3.70 Å) are no longer observed in the
system. For a short period of lithiation between 0.85 and 0.75
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Figure 11. Comparison of the total number of neighboring lithium
atoms in the metallic Sn and SnO2 electrodes during 1st and 2nd
cycles.
Figure 9. Fitting results for individual Sn−Li paths: short (black ■),
medium (blue ●), and long (red ▲) for in situ SnO2 electrode as a
function of electrode potential on the first lithiation.
phase, indicating that some Sn atoms were not fully lithiated.
This is likely due to the inhomogeneity and nonequilibrium
nature of lithiation process, where during first lithiation, large
volumetric expansions of the outer shell of the Sn nanoparticle
forms a Li22Sn5-like phase, with reduced electrical conductivity
and slower Li ion diffusion. In this situation, the core of the
nanoparticles (ca. 1 Sn−Sn near neighbor) could get insulated
from the rest of the electrode and remain in the metallic Sn
phase, while the nanoparticle shell is in a highly lithiated
Li22Sn5-like phase.
In the SnO2 electrode, however, the first most lithiated state
exhibits no statistically significant contributions from Sn−Sn
neighbors, indicating complete conversion of newly formed Sn
clusters to LiSn alloy. In the first delithiated (charged) state,
there is a small contribution from Sn−Sn with 0.9 ± 0.2
neighbors at the compressed distances of 2.92 ± 0.01 Å and
3.08 ± 0.01 Å, similar to the environment in the metallic Sn
electrode.
In the second discharged state, both Sn and SnO 2
nanoparticles have neighboring metallic tin atoms, indicating
incomplete lithiation of Sn atoms as supported by reduced
capacity of both electrodes (Figure 1). The metallic Sn
electrode shows no appreciable change in the local tin atomic
structure from the first discharged to charged state, indicating a
limited amount of reversible lithiation. As the Sn electrode is
cycled further, independent of discharged or charged state, the
number of neighboring Sn atoms and the Sn−Sn distance
increase toward values typical of metallic Sn (2.99 ± 0.02 Å
after the second charge compared to 3.01 Å for Sn metal). This
suggests that larger clusters of metallic Sn are forming but not
participating in the lithiation reaction and most likely
electrically insulated. The SnO2 electrode shows a slight
decrease in the number of neighboring Sn atoms in the second
discharged state (0.6 ± 0.2 Sn atoms in the second discharged
opposed to 0.9 ± 0.2 in the first charge) with no statistically
significant change in the distance between them. After the 2nd
delithiation, the local Sn environment returns to what was
observed in the first delithiated (charged) state.
Figure 11 displays the fitting results for the number of
lithium near neighbors in the first two cycles for both Sn and
SnO2 electrodes. In the first lithiated (discharged) state, the
metallic Sn electrode has 11.7 ± 1.7 neighboring lithium atoms
XAS of Discharged and Charged States. In addition to
the potential-resolved study of the initial lithiation processes of
both Sn and SnO2 electrodes, XAS spectra were acquired at
each fully discharged and fully charged state of the first two
discharge/charge cycles. The fits to the spectra for each
electrode and state were conducted using the previously
described procedure and presented together in Figures 10 and
11 for comparison.
Figure 10. Comparison of changes in the distance and number of
neighboring Sn atoms for OCV, and fully discharged (lithiated) and
charged (delithiated) states in the 1st and 2nd cycles for both metallic
Sn and SnO2 electrodes.
Figure 10 shows the fit results for coordination number and
interatomic distances for the Sn−Sn neighbors. In the metallic
Sn electrode, the transition from OCV to the first lithiated
(discharged) state is accompanied by a large reduction in both
the number of neighboring tin atoms along with a contraction
in the distance between them. The presence of neighboring
Sn−Sn atoms in the fully discharged state of the metallic Sn
electrode is not consistent with complete conversion to Li22Sn5
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while the SnO2 electrode has a total of 8.2 ± 0.9 lithium
neighbors. The discrepancy between the number of Sn−Li
neighbors in Sn and SnO2 electrodes is likely due to the
difference in the nanoscale and atomic arrangements around
the Sn atoms. The average size of nanoparticles in the metallic
Sn electrode is ca. 100 nm, while the metallic tin clusters that
form during the first lithiation of SnO2 electrodes are very
small, likely on the order of a few atoms. When particles are at
this size, the ratio of surface to bulk atoms is high, and a
reduction in the number of Sn−Sn neighbors is observed
compared to a bulk Sn metal phase. Neither system reaches the
expected 14 neighboring lithium atoms typical of the Li22Sn5
crystal structure, indicating structural disorder and that the
entire electrode does not fully convert to this phase. This
conclusion is also supported by the low capacity observed past
the first lithiation in both electrodes.
In the first delithiated (charged) state, there is a small
reduction in the number of lithium atoms in the metallic Sn
electrode (from 11.7 ± 1.7 to 9.8 ± 1.7), while in the ideal case,
no Sn−Li neighbors would be observed in the charged state
corresponding to complete delithiation. Upon continued
cycling there is a monotomic reduction in the number of
lithium neighbor atoms and a corresponding increase in the
number of Sn−Sn neighbors, in both discharged and charged
states, indicating very poor reversibility of lithium insertion,
most likely due to a reduction of electrical conductivity within
the active material. This lack of reversibility in lithium insertion
and removal is also supported by the poor electrochemical
performance of the metallic Sn electrode (Figure 1). It may be
suggested that Li diffusion within the electrode assists in the
segregation of Sn atoms into larger clusters, which also become
electrically insulated.
For the SnO2 electrode, there is an appreciable reduction
from 8.2 ± 0.9 Li atoms to 5.9 ± 1.2 Li during the first
delithiation, and on the second lithiation, the total number of
lithium atoms returns close to the amount observed in the first
lithiated state. Upon subsequent delithiation, the number of
lithium atoms again approaches what was observed for the first
delithiated state, indicating better reversibility of lithiation in
SnO2 electrodes versus the metallic Sn material but still only
partial removal of Li. Interestingly, the Sn clusters in the SnO2
electrode do not grow during subsequent discharge/charge
cycles.
These local structural changes in both Sn and SnO 2
electrodes aligns well with the electrochemical performance
displayed in Figure 1. The initial lithiation of both Sn and SnO2
materials is accompanied by a large contribution of irreversible
capacity from SEI formation. This process is not reflected in
EXAFS data as only a minor fraction of Sn surface atoms are
directly exposed to SEI. The latter makes it difficult to directly
compare the local structural changes determined through
EXAFS modeling to the electrochemical capacities on the first
lithiation cycle, where majority of the SEI is formed. For that
reason, the EXAFS results from the first lithiation are correlated
to the electrode/cell potentials, not to the experimental
capacities. Once SEI is mostly stabilized, on the second
lithiation, both Sn and SnO2 electrodes show similar capacities
(ca. 600 mAh/g), which correspond well to the similarities in
the local structures at this state (Figures 10 and 11). On second
delithiatiation, the SnO2 electrode exhibits ca. 80% Coulombic
efficiency while Sn electrode only ca. 36%. These electrochemical results are supported by the EXAFS, where the SnO2
electrode shows a decrease in Li neighbors upon delithiation
(Figure 11). In contrast, the metallic Sn electrode has no
significant change in the number of Li neighbors for subsequent
lithiation cycles, suggesting limited process reversibility and,
thus, low Coulombic efficiency.
On the basis of this comparison of the electrochemical
cycling results and corresponding EXAFS data on the changes
in the local atomic environment around tin atoms in both
metallic Sn and SnO2 nanoparticle electrodes, the positive effect
of using SnO2 as the starting electrode material becomes clear.
During the first lithiation, metallic Sn nanoparticles undergo
large structural changes, evidenced by a continual decrease in
the number of Sn−Sn near neighbors throughout the lithiation
process, which is also accompanied by expansion and then
contraction of the Sn−Sn distances. The loss of the Sn crystal
structure during the first lithiation process and amorphization
of the electrode results in poor electrical conductivity within the
particles, which is manifested by the presence of metallic Sn
atoms in the most lithiated state and poor reversibility of
lithium removal upon charge (Figures 4 and 11), indicating that
these structural changes irreversibly passivate the electrode
material.
The SnO2 nanoparticles undergo different structural changes
during the first lithiation. In the initial stages (1.00 to 0.40 V),
there is clear conversion of SnO2 to metallic Sn and Li2O.
Further lithiation results in full conversion of metallic Sn to
LiSn alloy with little variation in the number of tin atoms and
the interatomic distances (Figure 8). It is suggested that Li2O
formed from SnO2 serves as a buffer for mitigating the
segregation of metallic Sn into large Sn particles, thereby
providing structural stability and maintaining sufficient
electrical conductivity within the active material as evidenced
by the better reversibility of lithium insertion/removal in the
first two cycles.
It is also worth noting the onset of observable Sn−Li
neighbors is quite different in the metallic Sn and SnO2
systems. Sn−Li paths are first observed at ca. 0.80 V in
metallic Sn and at ca. 0.40 V in SnO2. This drastic difference is
likely due to the process of converting SnO2 to Sn clusters
embedded in a Li2O matrix before lithiation of tin may begin.
Limited electrical conductivity and slow Li diffusion rates
through the Li2O matrix to the metallic Sn domains may also
contribute to the limited capacity and reversibility of Li
insertion in SnO2 particles.
Both electrodes show dramatic capacity fading within the first
few cycles. In the case of the metallic Sn electrode, the main
issue is the mobility of Sn atoms within the solid phase that
results in unsuppressed volumetric expansion to accommodate
lithium atoms into the crystal structure. Removal of some Li
atoms from expanded SnLi structures on the first delithiation
results in the loss of electical conductivity of the electrode
material with large fractions of Li remaining in the form of
SnLi, regardless of the electrode potential. In the case of the
SnO2 electrode, segregation of metallic Sn atomic clusters
within the Li2O matrix during the first lithiation limits the
mobility of Sn atoms but also the amount of lithium that can
efficiently diffuse to and from the Sn atomic clusters. To
achieve better cycle life performance of these tin-based
electrodes, careful and intricate atomic and nanoscale engineering of electrode material is required to mitigate the capacity
fading mechanisms identified above. In particular, use of Sn
compounds that would result in a better conducting network
upon Sn reduction to metal, as well as nanoscale confinement
of such materials, is suggested.
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DOI: 10.1021/acs.jpcc.5b12279
J. Phys. Chem. C 2016, 120, 5331−5339
The Journal of Physical Chemistry C
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CONCLUSIONS
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ASSOCIATED CONTENT
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S Supporting Information
*
The Supporting Information is available free of charge on the
ACS Publications website at DOI: 10.1021/acs.jpcc.5b12279.
Representative EXAFS fits in k2χ(k), Re[χ(R)], and
|χ(R)| along with a table of detailed fitting results for
both SnO2 and Sn EXAFS fits. XRD of starting Sn and
SnO2 nanoparticles is also presented (PDF)
■
Article
AUTHOR INFORMATION
Corresponding Authors
*E-mail: [email protected]. Tel: 312-567-3498.
*E-mail: [email protected].
Notes
The authors declare no competing financial interest.
■
ACKNOWLEDGMENTS
C.J.P. was supported by a Department of Education GAANN
Fellowship, award no. P200A090137. The project is supported
by the U.S. Department of Energy, Office of Basic Energy
Science and the Advanced Research Project Agency−Energy
(ARPA−E) under Award no. AR-000387. MRCAT operations
are supported by the Department of Energy and the MRCAT
member institutions. Use of the Argonne National Laboratory
Advanced Photon Source is supported by the U.S. Department
of Energy, under Contract no. DE-AC02-06CH11357.
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