Pidaparthy 2021 j._electrochem._soc._168_100509

Journal of The Electrochemical
Society
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Increased Disorder at Graphite Particle Edges Revealed by Multi-length
Scale Characterization of Anodes from Fast-Charged Lithium-Ion Cells
To cite this article: Saran Pidaparthy et al 2021 J. Electrochem. Soc. 168 100509
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Increased Disorder at Graphite Particle Edges Revealed by Multi-
length Scale Characterization of Anodes from Fast-Charged
Lithium-Ion Cells
Saran Pidaparthy,1,2
Marco-Tulio F. Rodrigues,1,* Jian-Min Zuo,2,3
and Daniel
P. Abraham1,*,z
1
Chemical Sciences and Engineering Division, Argonne National Laboratory, Lemont, Illinois, 60439, United States of
America
2
Department of Materials Science and Engineering, University of Illinois at Urbana-Champaign, Urbana, Illinois, 61801,
United States of America
3
Materials Research Laboratory, University of Illinois at Urbana-Champaign, Urbana, Illinois, 61801, United States of
America
Fast charging of lithium-ion cells increases voltage polarization of the electrodes and creates conditions that are favorable for Li-
deposition at the graphite anode. Repeated fast charging induces changes in the capacity-voltage profiles and increases the
probability of lithium-plating on the electrode. This higher probability results from structural, morphological and chemical
modifications that are revealed by multi-length scale characterization of graphite anodes extracted from discharged lithium-ion
cells, previously charged at rates up to 6 C. The distinct differences between anodes with lithium-plating and as-prepared electrodes
are clearly seen in analytical electron microscopy data. Scanning electrode microscopy (SEM) images show that the fast-charged
anode is significantly thicker, apparently because of the electrolyte reduction/hydrolysis products that accumulate in electrode
pores. High resolution electron microscopy (HREM) images reveal wavy graphite fringes near the particle edges. Analysis of
scanning electron nanodiffraction (SEND) data reveal higher d-spacings and greater lattice rotations, indicating disorder in the
graphite near the particle edges that extend about 20 nm into the bulk. The extent of this disorder is greater near larger internal
pores, highlighting nanoscale heterogeneities within particles. As graphite lithiation occurs primarily through edge planes, this
permanent disorder would hinder Li+
intercalation kinetics and favor Li0
plating during repeated cycling.
© 2021 The Author(s). Published on behalf of The Electrochemical Society by IOP Publishing Limited. This is an open access
article distributed under the terms of the Creative Commons Attribution 4.0 License (CC BY, http://creativecommons.org/licenses/
by/4.0/), which permits unrestricted reuse of the work in any medium, provided the original work is properly cited. [DOI: 10.1149/
1945-7111/ac2a7f]
Manuscript submitted August 25, 2021; revised manuscript received September 17, 2021. Published October 8, 2021.
Supplementary material for this article is available online
Advances in lithium-ion battery (LIB) technology have fueled the
rise of the electric vehicle (EV) industry.1
Most high-energy LIBs
contain a positive electrode (cathode) comprised of a layered oxide
with Ni, Co and Mn, a negative electrode (anode) comprised of
graphite, and an electrolyte with LiPF6 salt in a mixture of organic
solvents. These cell components are being continually optimized so
that LIB technology can deliver high energy densities safely and
reliably during a battery’s 10–15-year lifetime. State-of-the-art LIBs
can be fully charged in about 1 h with minimal degradation in cell
performance. However, full charging over shorter durations causes
irreversible damage to the battery via losses in the inventory of
mobile Li+
ions (i.e., capacity fade) and resistance increases (i.e.,
power fade) in the cells.2,3
The U.S. Department of Energy (DOE) is
funding research to identify the underlying mechanisms that limit
fast-charge performance of LIB cells and to develop technological
solutions that would enable full battery charge in less than 10 min.4–7
Solutions to these limitations will allay range anxiety concerns and
accelerate customer acceptance of EVs.
During cell charging, lithium ions are extracted from the
transition metal oxide cathode and intercalated between atomic
planes of the graphite particles. Under fast charge conditions, there is
salt enrichment in the oxide cathode and salt depletion in the
graphite anode, leading to non-uniform distribution of Li+
ions
across the electrode thickness.8,9
These conditions lower Li+
ion
conductivity in the electrolyte within the electrode pores and
increase impedance, causing polarization in both electrodes.
Cathode polarization causes the cell to reach its upper cut-off
voltage (UCV) before Li+
ions are fully extracted from the oxide,
while anode polarization causes the cell to dip below the Li-plating
potential creating conditions that are favorable for lithium nucleation
and growth on the graphite electrode.10–12
The situation is further
complicated by non-uniformities in local reactivity resulting from
factors that include uneven mechanical compression of the separator
that affects its porosity/tortuosity, and irregular distribution of active
particles that affects distribution of local currents within the
electrodes.13–16
These inhomogeneities cause large lithium concen-
tration gradients in the electrodes, both across the surface and along
its thickness, promoting non-uniform aging of the cell materials. The
high charge currents can also enhance structural stresses in the
electrode matrix causing the coating to peel off the current collector
and within the active material particles causing their fracture and
disconnection from the electron conduction network.17
In our recent series of articles we have used operando X-ray
techniques to quantify phase evolution in the graphite and lithium
concentration gradients that develop in the electrode during fast
charging.18,19
In addition, we have described techniques to identify
lithium deposition on the graphite electrode and detailed electro-
chemical charging protocols that can mitigate this plating.20–22
We
observed that the likelihood of Li-plating on the graphite anode
increases during repeated fast-charge cycling.10
We note here that
commercial graphite typically contains a mixture of the hexagonal
(2H) and rhombohedral (3 R) phases: the stacking sequence is
ABAB in the 2H structure and ABCABC in the 3 R structure.23
The 2H phase is more thermodynamically stable than the 3 R phase
at room temperature and pressure. The Li intercalation mechanism is
reported to be similar for both phases: during Li+
insertion, the
adjacent graphene sheets glide relative to each other to lower the
overall energy of the system and attain AAAA layer stacking.24
Any
material feature that hinders the movement of graphene sheets will
prevent Li+
intercalation. During fast charging, Li+
intercalation
into the graphite competes with Li0
nucleation on the particles.
Greater layer-to-layer stacking disorder in the graphite in the form of
added strain, local misfits, and misorientation angles in the stack,
would make Li+
intercalation more difficult and expedite Li0
z
E-mail: abraham@anl.gov
*Electrochemical Society Member.
Journal of The Electrochemical Society, 2021 168 100509

plating. Hence, an in-depth understanding of changes in the graphite
particles subjected to fast charging is crucial for the development of
materials that enable intercalation and mitigate deterioration of cell
performance during the electrochemical cycling.
In this work, we examine the morphological, chemical, and
structural changes to the graphite materials and electrodes, using a
suite of electrochemical and physicochemical techniques. The
characterization was performed across multiple length-scales on
as-prepared electrodes and on numerous electrodes harvested from
cycled cells; only representative data from the various techniques
are shown to highlight key changes that result from the cycling. The
data in this manuscript are from an as-prepared graphite (Gr-AP)
electrode that was never exposed to the electrolyte, graphite
electrode with no plating (Gr-NP) harvested from a cell after a
few low-rate cycles and graphite electrode with lithium plating (Gr-
LP) harvested from a fast-charged cell. For purposes of clarity and
contrast, we show only the Gr-AP and Gr-LP data in the main text;
data from the Gr-NP sample are included in the Supporting
Information (SI). Note that all Tables and Figures in the SI are
marked with the designator S, as in Fig. S1 (available online at
stacks.iop.org/JES/168/100509/mmedia).
We start by describing electrochemical tests, conducted on full
cells with LiNi0.5Co0.2Mn0.3O2 (NCM523) cathodes and graphite
(Gr) anodes, which include charging rates up to 6 C (here, 1 C is the
current needed to fully discharge the cell in 1 h). Following this, we
compare electrochemical behavior of the Gr-AP electrode with that
of the Gr-LP electrode. Then, after presenting high-resolution X-ray
diffraction (XRD) data from the materials, we show microscopy and
spectroscopy results obtained on electrode cross-sections fabricated
via ion milling and on sections prepared with a focused ion beam
(FIB). At the bulk-scale, we use scanning electron microscopy
(SEM) and energy dispersive X-ray spectroscopy (EDS) to probe the
morphological and chemical changes in cross-sections of the
graphite anode. At the nano- to atomic-scale, we rely on both
conventional transmission electron microscopy (TEM) and scanning
TEM (STEM) to reveal local structural disorder and chemical
variation within individual graphite particle sections through a
combination of high-resolution imaging, EDS, and scanning electron
nanodiffraction (SEND). Our investigations reveal that anodes from
the fast-charged cells have lower capacity, present a sloping voltage
profile, display an increase in thickness and show greater lattice
disorder in localized areas near the graphite particle edges.
Experimental Section
Electrode materials, cell assembly and electrochemical
testing.—In the NCM523//Gr cells, the positive electrode has a
71 μm thick coating, with 90 wt% Li(Ni0.5Co0.2Mn0.3)O2 (Toda),
5 wt% C45 carbon (Timcal) and 5 wt% PVDF binder (KF-9300,
Kureha), on a 20 μm thick Al current collector.10
The negative
electrode has a 70 μm thick coating, with ∼92 wt% SLC1506T
graphite (Superior Graphite), 2 wt% C45 carbon and 6 wt% PVDF
binder, on a 10 μm thick Cu current collector. The electrolyte (aka
Gen 2, purchased from Tomiyama) is a solution of 1.2 M lithium
hexafluorophosphate (LiPF6) salt dissolved in a 3:7 w/w solvent
mixture of ethylene carbonate (EC) and ethyl methyl carbonate
(EMC). Celgard 2320 is the separator between the electrodes. Before
cell assembly, the electrodes and separator are dried in vacuum
ovens set at 120 °C and 70 °C, respectively. The 3-electrode
measurements are conducted in custom-made stainless-steel cells,
which are assembled and tested inside an Ar-atmosphere glove box.
The reference electrode (RE) probe is placed between two separa-
tors, sandwiched between 20.3 cm2
disks of the oxide-positive and
graphite-negative electrodes (Fig. 1): the RE consists of a copper
wire (25 μm diameter) with a thin layer of Li metal plated on the
exposed tip.9
Electrochemical cycling is typically conducted at 30 °C with a
Maccor Model 2400 cycler. The Gr-NP electrode is from a cell that
experienced two C/10 and five C/2 cycles between 3 and 4.2 V at 30
°C. The Gr-LP electrode was harvested from a cell that experienced
multiple sequences of an “escalating C-rate” protocol (Fig. 1). To
accelerate cell aging, three of the sequences were at 30 °C, one at 40
°C and one at 50 °C. Including the initial cycles and periodic
reference performance tests (C/25 rate), the cell experienced a total
of 50 cycles; of these, 25 charge segments were at rates of 2 C and
above.
After cycling, cells were discharged to (and held for 10 h at) 3 V
before disassembly in an argon-atmosphere glove box. Various
diagnostic tests were conducted on the harvested cell materials to
determine the effect of cycling on the graphite material. For the
electrochemical tests, 2032-type coin cells were assembled with 1.6
cm2
area harvested-electrode samples, lithium metal counter elec-
trode, fresh Celgard 2320 separator, and fresh Gen2 electrolyte.
Cross-section polishing.—Cross-sections of the graphite elec-
trode samples were prepared via argon ion milling using the Gatan
Model 685.O Precision Etching Coating System II (PECS II). The
electrode was gently handled with metal tweezers and cut with sharp
metal scissors to leave at least one flat-side of ∼5 mm; all metal
tools were cleaned using an ethanol and acetone rinse prior to
handling the electrode. The cut electrode was mounted onto the
PECS II sample blade (Gatan, Product No. 693.08500) using fast-
drying silver paint (Ted Pella, Product No. 16040–30) and adjusted
with a micrometer to ensure ∼200 μm overhang of the flat-side of
the electrode from the top of the sample blade. Once the electrode
was secure on the blade, the electrode/blade polishing sample was
placed onto the cross-section specimen mount (Gatan, Product No.
Figure 1. In this escalating C-rate experiment, the NCM523//Gr cell is
charged with increasing currents to a fixed capacity of 87 mAh g−1
. Each
charge is followed by 1 h rest and C/5 discharge to 3 V and a hold. The cell
voltage changes during the cycles are shown in (a), while the graphite-
electrode potential changes are shown in (b). Lithium plating on the graphite
particles becomes likely when the electrode potential dips below 0 V vs
Li/Li+
; this Li-plating condition (LPC) is indicated by the dotted line in (b).
In the above experiment, Li-plating on the graphite electrode becomes likely
when the charge rate exceeds ∼2 C.

685.08540) and loaded into the PECS II instrument. After the sample
was loaded, the instrument was set to “Etch by Time” with a “Single
Modulation” etching condition, a Gun Tilt angle of 0° for both ion
guns, and a rotation rate of 0.5 RPM. The ion milling was performed
in two stages: first, a high-voltage milling at 8 kV (with current
readings in the range 20–40 μA) that was run for 4 h to create a
rough cross-section; second, a low-voltage milling at 4 kV (with
current readings of 7–15 μA) that ran for 30 min to apply a low-kV
fine polishing.
Synchrotron powder diffraction.—Powder diffraction data were
collected at beamline 11-BM of the Advanced Photon Source,
Argonne National Laboratory. Coatings from the as-prepared
electrodes and cycled electrodes, scraped of the copper current
collectors, were loaded into sealed capillaries for examination. An
X-ray wavelength of 0.41283 Å was used to collect the diffraction
data, over an angular range from –6.0° to 28.0° with a step size of
0.001° and a time per step of 0.1 s.
Scanning electron microscopy and energy-dispersive X-ray
spectroscopy.—Both imaging and elemental examinations were
performed using a JEOL 7000 F analytical SEM. The electrode
samples were prepared in the following ways: (1) for top-down
imaging, a small portion of the electrode is mounted onto a flat
sample holder via carbon tape; (2) for cross-sectional imaging, the
samples prepared by the PECS II ion milling are loaded onto a
standard cross-section specimen mount designed for the SEM. For
the SEM imaging, the instrument was set to 7 kV and operated under
secondary electron imaging mode with a medium probe current of
∼0.3 nA. For EDS acquisition, the sample was aligned to an optimal
working distance of 10 mm and the probe current was adjusted to a
setting of ∼15 nA. The mapping data were acquired using the
ThermoFisher NSS 2.0 software using its “Spectral Imaging” mode.
The datasets were summed for 150 frames, each acquired for 10 s
per frame with a 50 μs dwell time. The mapping data are reported as
counts per element and the quantification was performed by the
software automatically using a built-in database and algorithm based
on the standard phi-rho-z method. The line profile data for the cross-
section samples are acquired using the “X-ray Linescans” mode with
a single scan with 50 points from the bottom to the top of the anode
with a dwell time of 10 s per point. Line profile data are further
binned for every 5 μm step.
Focused-ion beam milling.—Samples for nano- to atomic-scale
electron imaging and diffraction were prepared with a ThermoFisher
Scientific Scios 2 FIB-SEM. The electrodes were cut with sharp
metal scissors and mounted onto a flat sample holder via carbon tape
and loaded into the instrument. The cross-section samples were
selected from areas near the electrode surface and prepared via
standard focused-ion beam lift-out procedure provided by
ThermoFisher Scientific onboard user guidance (included with the
instrument software). The samples were mounted onto one of four
posts of a PELCO FIB lift-out TEM copper grid and thinned to a
final thickness around 100–200 nm.
Transmission electron microscopy.—Transmission electron mi-
croscopy (TEM) was performed on a Hitachi 9500 TEM equipped
with a LaB6 emitter and operated at 300 kV. Image acquisition was
done using the K2 direct electron detector operated in counting
mode.
Scanning transmission electron microscopy, energy dispersive
X-ray spectroscopy, and scanning electron nanodiffraction.—
Nano- and atomic-scale imaging, chemical analysis, diffraction
data were acquired using the ThermoFisher Scientific Titan
Themis Z STEM equipped with a Schottky field emission gun
operated at 300 kV accelerating voltage. The electrode sample was
loaded into a standard double-tilt holder and, during operation, tilted
close to the 100
〈 〉 zone axis or the 210
〈 〉 zone axis (with respect to
the Graphite-2H phase), as both orientations allowed us to track the
graphite interplanar spacing in both imaging and diffraction modes.
Imaging was performed at 15–30 pA screen current with a
convergence angle of 18.0 mrad and acquired using a high-angle
annular dark field (HAADF) detector. The EDS acquisition was
performed using the above mentioned imaging conditions via the
Velox software (ThermoFisher Scientific) with an acquisition time
of ∼1 h using a continuous scan with a frame time of ∼1 s. SEND
data were collected in microprobe STEM (μP STEM) mode with a
convergence angle of 0.46 mrad. The SEND probe size is 1.7–-
1.8 nm at full-width half maximum. The diffraction datasets were
acquired at a camera length of 145 mm over a real-space scan area of
Figure 2. (a) Capacity-potential profiles of the graphite electrode from a NCM532//Gr cell with a reference electrode. The lithiation profiles are at the indicated
rates; the 0.2 C rate delithiation profile shown is after lithiation at the 2 C rate followed by a 1 h rest. (b) Profiles for as-prepared (black) and harvested electrodes
(red); the latter are from a cell subjected to multiple “escalating C-rate” cycling. The as-prepared electrode data are at C/100 rate, whereas the harvested electrode
plots are at both C/100 (solid) and C/10 (dashed) rates. In both (a) and (b), the specific capacities are obtained by dividing the cell capacity (mAh) by the graphite
weight (gGr
−1
) in the anode.

150 nm × 150 nm using 75 × 75 total steps with drift correction
performed after each line acquisition.
Results and Discussion
Electrochemical cycling.—Capacity-potential profiles of the
graphite electrode from a NCM523//Gr cell are shown in Fig. 2a.
The graphite is lithiated and delithiated during cell charge and
discharge, respectively. At low lithiation rates (<C/5), the inter-
calation progresses in distinctive stages seen as plateaus in the plots;
the Li+
ions fill spaces between the carbon sheets forming distinct
LixC6 phases (x = 0.22, 0.33, 0.5, 1). At higher lithiation rates
(⩾1 C), the plateaus are not distinct as intercalation occurs under far-
from-equilibrium conditions. Multiple mixed and strained phases are
probably formed so that the profile appears continuous, rather than
staged; note that LixC6 phases of various compositions have been
reported in the research literature.25–29
During the C/5 delithiation,
staging plateaus are always observed, regardless of the behavior
during the preceding lithiation cycle (Fig. 2a). We observed this
difference between the lithiation and delithiation behavior during
operando XRD studies; therein, we attributed the blurring of phase
boundaries during lithiation to lattice strain and the formation of
distinct LixC6 phases during delithiation to abrupt separation from
oversaturated solid solutions.18
As indicated earlier, the cells in this study were subjected to
multiple applications of the escalating C-rate protocol. Capacity-
voltage profiles from a representative cell, before and after the
protocol, are shown in Fig. S1. The electrochemical cycling causes a
significant loss in cell capacity. The Li+
ions are lost to the solid
electrolyte interphase (SEI) formed during electrolyte reduction on
the graphite electrode and to the dead lithium that accumulates in the
electrode because of repeated excursions below the Li-plating
potential (Fig. 1b).22,30
Figure 2b compares plots from a lithium-plated graphite elec-
trode (Gr-LP) harvested from the cell in Fig. S1, with plots from an
as-prepared graphite (Gr-AP) electrode. The lithiation profile of the
Gr-AP electrode (at C/100 rate) displays distinctive plateaus around
210, 117 and 83 mV vs Li/Li+
. These plateaus are also very distinct
in the C/10 rate lithiation data (not shown), though they occur at
slightly lower potentials because of polarization arising from
electrode impedance. In contrast, the plateaus are not distinct in
the lithiation profile of the Gr-LP electrode, even at a C/100 rate.
The capacity of this aged-harvested electrode (326 mAh g−1
) is also
smaller than that of the as-prepared electrode (357 mAh g−1
). The C/
10 lithiation data of this Gr-LP electrode resembles the 2 C graphite-
lithiation data in Fig. 2a; the sloped profile, the barely visible staging
plateaus and the lower capacity all suggest changes in the graphite
particles that alter lithium insertion into the material. Interestingly,
the staging plateaus are visible both in the C/100 and C/10
delithiation profiles of the harvested material, which is consistent
with the distinct differences between the lithium insertion and
extraction processes mentioned earlier.
The theoretical capacity of graphite is 372 mAh g−1
but the Gr-
AP electrode yields only 357 mAh g−1
even at very low rates. The
underlying reason for this smaller capacity is the inherent turbos-
tratic stacking disorder in the commercial graphite, which limits Li+
ion intercalation into the material.31,32
Fast-charging appears to
create additional disorder, further decreasing capacity. Moreover, the
increased structure-disorder increases the number of electronically
and geometrically non-equivalent sites for Li-intercalation causing
the sloping lithiation profile of the Gr-LP even at extremely low
rates.33
In contrast, most of the sites in Gr-AP are thermodynami-
cally equivalent for Li+
intercalation, thereby producing the
observed staging behavior at low lithiation rates.
Ensemble phase analysis via powder X-ray diffraction.—We
evaluated the bulk-scale structure and ensemble lattice spacing of the
graphite anode coatings using synchrotron X-ray diffraction.
Diffraction patterns obtained on the Gr-AP and Gr-LP electrodes
are shown in Figs 3a and 3b, respectively. Figure 3a shows that the
as-prepared electrode contains two distinct graphite phases: the 2H
phase (marked as #) and the 3 R phase (marked as *). The 3 R phase
is not observed in the Gr-LP electrode pattern (Fig. 3b); we postulate
that this phase transforms into the more stable 2H form during
repeated Li+
insertion into and extraction from the graphite, which
requires gliding of the graphene planes as mentioned in the
Introduction.
We calculated the graphite interplanar d-spacing from the strong
diffraction peak at ∼1.87 Å–1
, which corresponds to the (002)
reflection for the Graphite-2H phase or the (003) reflection for the
Graphite-3R phase. For both Gr-AP and Gr-LP electrode samples we
obtained d 3.356 Å,
interplanar
〈 〉 ≅ which indicates that the latter
sample is fully delithiated and that the amount of residual Li+
ions, within the graphite of the fully discharged cell, is negligible.
However, diffraction pattern of the Gr-LP sample (Fig. 3b) shows
new peaks at ∼2.70 Å–1
and ∼4.40 Å–1
, which arise from lithium
fluoride (LiF) (marked as ♢). The presence of LiF was confirmed by
selected area electron diffraction (SAED) data on a FIB cross-
section prepared from the top surface of the Gr-LP sample (Fig. S2).
LiF is known to be a component of the graphite SEI and forms
through decomposition reactions of the LiPF6 salt.34,35
There was no
evidence of other known SEI phases in the Gr-LP pattern, which
indicates that these compounds are either fully amorphous or lack
the long-range order needed to coherently scatter the X-ray beam.
SEM imaging and compositional analysis.—Figure 4 reveals the
dramatic change to the graphite anode surface after being subjected
to cycling under the fast-charge, lithium-plating condition. Figure 4a
shows a photograph of the Gr-AP sample (cut to a 14.2 mm
diameter); the EDS data in Figs. 4b and S3 indicate that the coating
contains only C and F, reflecting the graphite, C45 and PVDF
Figure 3. Synchrotron powder diffraction data from the following graphite
electrodes: (a) as-prepared (Gr-AP) and (b) cycled with lithium plating (Gr-
LP). The Bragg reflections are indexed according to JCPDS cards 01–071-
–3739 for Graphite-2H, 01–075–2078 for Graphite-3R, and 01–071–3743 for
LiF. The phases are marked using the following symbols: # for Graphite-2H,
* for Graphite-3R, and ♢ for LiF.

materials used for electrode preparation. Figure 4c is an SEM image
of the Gr-AP sample surface at 2k × magnification to illustrate that
the electrode is comprised of densely packed, potato-shaped graphite
particles up to 10 μm in size.
SEM and EDS data from the Gr-NP electrode are displayed in
Fig. S4. The graphite particles appear unchanged by the low-rate
cycles; the elemental analyses reveal the presence of C, O, F and P,
which arise from electrolyte decomposition products and residues on
the electrode. By contrast, Fig. 4d shows a photograph of the Gr-LP
sample (cut to 14.2 mm diameter) with an uneven white-coloration
across the surface; the average elemental composition (in wt%) of
the electrode surface (Fig. 4e) is 21% O, 49% F, 27% P and 3% C.
Figure 4f shows the surface morphology of the Gr-LP sample at 2k
× magnification; in this image, there is complete coverage of the
graphite particles by products formed by reactions of the plated
lithium with the electrolyte (also see Fig. S5).
To further determine extent of changes, we examined the
electrodes in cross-section to study morphological and elemental
characteristics throughout their depth (see Figs. 5, S6, S7 and S8).
Details of the coating thickness calculations are provided in Section
4 of the SI and Fig. S9. Figure 5a shows that the Gr-AP sample has
an average coating thickness of 76.6 μm. Figure 5b shows a FIB
cross-section of a typical graphite particle, revealing the presence of
internal pores. The elemental composition across the electrode
thickness was measured via EDS line-scans. For the Gr-AP sample,
the carbon signal, which arises predominantly from the graphite,
measures an average of 97 wt% across the depth of the coating layer
above the copper current collector, while the fluorine signal, which
arises from PVDF binder, measures an average of 3 wt% (Figs. 5c
S6). Cross-sections of the Gr-NP electrode (Fig. S7) reveal an
average coating thickness of 78.1 μm, which is about 2% greater
than that of the Gr-AP electrode. Irreversible increases in graphite
electrode thickness have been noted previously in electrochemical
dilatometry studies; these typically range from 2%–5% after a few
cycles.36–38
EDS scans reveal carbon, oxygen, fluorine and phos-
phorous signals, distributed across the electrode. These elements are
from the electrolyte decomposition products that accumulate in the
electrode.
By contrast, the Gr-LP electrode has an average coating thickness
of 91.3 μm, which does not include the ∼15–20 μm thick layer
resulting from lithium plating on the anode surface (see Fig. 5d).
This coating thickness is 19.2% and 16.9% greater than that of the
Gr-AP and Gr-NP electrodes, respectively. We attribute this
electrode thickness increase to SEI buildup in the electrode pores
and rearrangement of graphite particles during the electrochemical
cycling. The process is exacerbated by Li plating and subsequent
electrolyte reduction, which increases the amount of SEI. This SEI
buildup is also seen on individual graphite particles, on the exterior
surfaces and in the internal pores (Fig. 5e). Moreover, EDS line-
scans show that the elemental composition is non-uniform across the
electrode thickness (Figs. 5f and S8). The topmost portion, which
has no graphite, contains mainly oxygen, fluorine and phosphorus
indicating that the products of lithium plating are mainly inorganic—
for example, LiF and lithium (oxy)fluoro-phosphate compounds.
Below this portion, in the electrode coating, carbon is the dominant
element but we also observe an abundance of fluorine and
phosphorus, possibly (oxy)fluoro-phosphate compounds resulting
from hydrolysis reactions of the electrolyte salt.39,40
STEM-EDS compositional analysis.—Using STEM-EDS, we
further show that the (oxy)fluoro-phosphate compounds only coat
the surface of the graphite particle and do not penetrate the bulk.
Figure 4. Comparison of the anode surface morphology and composition. The as-prepared graphite (Gr-AP) anode surface details are elucidated via: (a)
photograph at 1x magnification, (b) EDS elemental composition, in weight percent, and (c) SEM image at 2000 × magnification. Similarly, the Gr-LP anode
surface details are elucidated via: (d) photograph at 1x magnification, (e) EDS elemental composition in weight percent, and (f) SEM image at 2000 ×
magnification. Scale bars in (a) and (d): 5 mm.

Figure 6a shows a survey of the Gr-LP particle section (prepared by
FIB) with EDS elemental analysis performed in the marked region,
which is magnified in Fig. 6b.
We reveal compositional heterogeneity by comparing the local
EDS spectra from graphite edges (marked as Area #1 and Area #3)
to the graphite center (marked as Area #2) shown in Fig. 6c(i). The
spectra from Area #1 and Area #3 show signals for O, F, and P in
addition to the signal from C. However, Area #2 only shows C
signature, which indicates a lack of (oxy)fluoro-phosphate com-
pounds in the graphite interior. Figure 6c(ii) shows the spectra
corresponding to the whole area in Fig. 6b. The elemental maps
(Fig. 6d) further confirm the spatial heterogeneity between carbon
and the (oxy)fluoro-phosphates; the C K signal appears only in the
graphite bulk, while the O, F, and P signals appear on the graphite
band edges. Moreover, the O, F, and P signals appear in the same
Figure 5. SEM imaging and elemental mapping of the electrode cross-sections. The as-prepared graphite (Gr-AP) sample details are elucidated via (a) SEM
image of the electrode cross-section, (b) SEM image of an individual particle section, and (c) EDS elemental data acquired across the thickness from bottom of
the current collector to the top of the coating surface—see white arrow in (a). Similarly, the lithium-plated graphite electrode (Gr-LP) sample details are
elucidated via (d) SEM image of the electrode cross-section, (e) SEM image of an individual particle section, and (f) EDS elemental data acquired across the
thickness depth from bottom of the current collector to the top surface—see white arrow in (d). Scale bars: (a), (d) are 25 μm; (b), (e) are 2 μm.

locations across the EDS maps. Additional EDS maps for different
regions of the Gr-LP electrodes are provided in Figs. S10 and S11.
The above observations indicate that the SEI only accumulates on
the graphite particle edges and does not penetrate the bulk material.
Electrolyte solvent co-intercalation would cause excessive strain
and/or damage to the host graphite. We do not see such damage in
the bulk graphite and, therefore, do not anticipate extensive
electrolyte penetration. We also observe that smaller pores within
the graphite yield weaker electrolyte residue signatures compared to
the larger pores. The electrolyte is known to penetrate the intra-
particle pores during electrochemical cycling.41
Larger pores pro-
vide easier access to the electrolyte; hence, more SEI would be
expected to accumulate in the larger pores than in the smaller pores.
Revealing localized disorder via high-resolution TEM.—High-
resolution TEM (HRTEM) data reveal the nature of the turbostratic
disorder that manifests in graphite particles at the lattice-level
(Fig. 7). Figure 7a, spanning a 50 nm × 50 nm field-of-view, shows
that the Gr-AP sample has straight and uniform graphite lattice
fringes corresponding to the constituent graphene planes. From this
image, we can qualitatively assess the degree of uniformity by
studying Region 1 (Fig. 7b) and Region 2 (Fig. 7c). The lattice-
fringe spacing, measured directly from the images of both regions, is
∼0.33 nm, which is consistent with the interplanar d-spacing of
graphite determined by XRD.
Fast-Fourier Transform (FFT) analysis is useful for quantifying
geometric and periodic patterns in electron microscopy images.42,43
The localized Fast-Fourier Transform (FFT) of the Region 1 and
Region 2 images are shown in Figs. 7d and 7e, respectively;
information to explain features observed in the HRTEM images is
contained in Section 6 and Fig. S12 of the SI. Briefly, the FFT of a
real-space intensity image (with distances x in nm) yields a
representative image in the corresponding spatial frequency domain
(with distances x
1/ in nm–1
).42,43
By applying FFT to the image in
Fig. 7b, the periodic stripe pattern of the graphite lattice (with
measured spacing 0.33 nm
∼ ) is mapped to points in Fig. 7d
corresponding to the frequency of the lattice spacing (i.e.,
1 0.33 nm 1
∼ / − from the center). Moreover, because the intensity
variation of the fringes lies along the direction normal to the fringes
in Fig. 7b (specified as direction 1 on the shown axis), the
corresponding spots in the frequency domain of Fig. 7d lie along
this same direction. The higher frequency spots in Fig. 7d that lie
along the same line (direction 1) correspond to additional harmonics
(i.e., n 1 0.33 nm 1
× ∼ / − where n is an integer) of the real-space
lattice; observation of such harmonics implies that the lattice is
relatively straight and uniform (i.e., ordered) within Region 1. In the
specified direction 2 (normal to direction 1), spots are occasionally
observed within the lattice fringes (Fig. 7b); these spots likely arise
from atomic columns within the graphite and the distance between
them correspond to the spacing between such columns. Similar
observations can be made for the Region 2 image (Fig. 7c) and its
corresponding FFT (Fig. 7e), thereby illustrating spatial uniformity
of the Gr-AP lattice.
Figure 6. STEM-EDS of a graphite particle cross-section from the Gr-LP sample. HAADF-STEM images of (a) graphite particle section and (b) a close-up of
the area where the EDS maps were acquired. (c) EDS spectra represented by the following regions: (i) from selected areas 1 (top-edge of graphite edge), 2 (center
of graphite band), and 3 (bottom-edge of graphite band) as marked in (b); (ii) whole-area EDS spectrum. (d) EDS maps of (i) C, (ii) O, (iii) F, and (iv) P signals.
Scale bars: (a) 2 μm, (b) 100 nm, (d) (i)–(iv) 50 nm.

By contrast, HRTEM images of the Gr-LP sample (Fig. 7f), show
significant nonuniformity. This nonuniformity is illustrated by
example from the magnified Region 1 (Fig. 7g) and Region 2
(Fig. 7h); localized FFTs of Regions 1 and 2 are shown in Figs. 7i
and 7j, respectively. Region 1, which is taken near the surface of an
internal pore within a graphite particle, is different from both
Regions 1 and 2 of the Gr-AP sample. From Fig. 7g, we find that
the Gr-LP lattice of Region 1 appears wavy; based on Fig. 7f, we
estimate that this waviness extends about 20 nm into the graphite
bulk. The FFT of Region 1 (Fig. 7i) confirms lattice disorder in the
following ways: (i) arcs are found instead of spots along direction 1;
(ii) the high frequency harmonic spot (marked by the white arrow) is
relatively faint compared to those in the Gr-AP sample. Furthermore,
the interplanar d-spacing from Fig. 7g measures up to ∼0.37 nm,
indicating that the lattice in this region does not return to its original
delithiated state. This nanoscopic increase in interplanar spacing is
not captured in the powder diffraction data (Fig. 3) because the latter
technique provides an average from all portions of the sample. On
the other hand, Region 2, which is in the bulk of the Gr-LP sample,
shows straight, uniform lattice fringes (Fig. 7h) similar to those
observed in the Gr-AP sample (Figs. 7b, 7c). The corresponding FFT
of Region 2 (Fig. 7j) has spots along direction 1 and a high-
frequency harmonic, which suggests greater ordering within the
graphite bulk. Taken together, the HRTEM images and FFT results
indicate there is greater disorder in the graphite lattice near the
Figure 7. Direct electron imaging of the graphite lattice via HRTEM. The as-prepared graphite (Gr-AP) electrode lattice details are elucidated via (a) TEM
imaging with highlighted Regions 1 and 2; real-space magnification of (b) Region 1 and (c) Region 2; localized Fast-Fourier Transforms of (d) Region 1 and (e)
Region 2. Correspondingly, the Gr-LP lattice details are elucidated via (f) TEM imaging with highlighted Regions 1 and 2; real-space magnification of (g)
Region 1 and (h) Region 2; and localized Fast-Fourier Transforms of (i) Region 1 and (j) Region 2. Scale bars: (a), (f) are 10 nm; (b), (c), (g), (h) are 2 nm; (d),
(e), (i), (j) are 5 nm–1
.

surface of its internal pores and that this disorder can extend up to
∼20 nm into the particle bulk.
Quantification of nanoscopic turbostratic disorder via SEND.—
In Fig. 8 and 9, we quantify the nanoscale heterogeneity of the
graphite by measuring the interplanar d-spacing and local lattice
rotation using SEND.44
In this technique, a ∼2 nm probe (at full-
width half-maximum) is rastered across a 150 nm × 150 nm real-
space area generating an electron diffraction pattern for each probe
position. SEND provides position-resolved structural information
regarding turbostratic disorder in the materials. Details of the SEND
data acquisition are briefly described in the experimental section;
data analysis information is provided in Section 7 of the SI, which
includes Fig. S13. In the SEND experiments, we quantify both the
graphite interplanar spacing and the lattice rotation by tracking
changes to the Ginterplanar vector (with respect to the Graphite-2H
phase) for each real-space sample coordinate of the SEND scan. The
quantification is done by determining (1) changes to the length of the
vector G d
1
interplanar interplanar
∣ ∣ = to measure position-to-position
nanoscopic interplanar d-spacing, which is then represented as a
dinterplanar map to highlight local d-spacing changes; (2) changes to
the in-plane orientation θ of each vector G ,
interplanar which is then
represented as a map of θ
∣∇ ∣ to identify locations where rotational
changes in the lattice occur.
Figure 8a(i) contains HAADF-STEM image from a Gr-AP
particle cross-section, with marked Region 1 (magnified in
Fig. 8a(ii)) and Region 2 (magnified in Fig. 8a(iii)) where SEND
data were acquired. First, we discuss the nanoscale structure of
Region 1 in Fig. 8b. The data indicate that the lattice rotation shows
negligible variation ( 0.09 nm 1
∼ ° − ) across the sample (Fig. 8b(i)) and
that the interplanar d-spacing map across Region 1 is relatively
uniform (Fig. 8b(ii)). Figure 8b(iii) is a plot of the overall d-spacing
from the map as a histogram; the data show a distribution of d-
spacings in the sample with a mean of 0.334 nm and a standard
deviation of 0.006 nm. Figure 8c illustrates that Region 2 is
comparable to Region 1 in terms of both lattice rotation
(Fig. 8c(i)) and interplanar spacings (Figs. 8c(ii) and (iii)).
However, we find a slightly elevated value for rotational variation
in the particle bulk of Fig. 8c(i). This observation points to an
important characteristic of graphite, which is that the graphene
planes tend to group as bands of 10–20 nm thickness (Fig. S14).
Within each graphite band, the magnitude of the rotational variation
is negligible; between band-to-band, the in-plane lattice angle can
exhibit sudden jumps. Overall, though, the results for the Gr-AP
sample are consistent with the HRTEM observations, which also
indicated a uniform lattice. Furthermore, SEND data from the Gr-NP
electrode (Fig. S15) are very similar to those from the Gr-AP
electrode, which indicate that the low-rate cycles have negligible
impact on the graphite particle structure.
By contrast, a non-uniform lattice and an increase in lattice
disorder at the particle edges are observed in particles from the Gr-
LP electrode. Figure 9a(i) contains a HAADF-STEM image with
areas marked Region 1 (magnified in Fig. 9a(ii)) and Region 2
(magnified in Fig. 9a(iii)) where SEND data were collected. In
Fig. 9b(i), we show position-to-position changes of the in-plane
rotation angle of Ginterplanar in Region 1, which is near a relatively
thin pore region. The graphite bulk in Region 1 has a uniform lattice
rotation ( 0.14 nm 1
∼ ° − ), while the graphite edge has greater disorder
(with rotational changes 2 nm 1
> ° − ). Figure 9b(ii) shows the corre-
sponding interplanar d-spacing reconstruction map; the d-spacing is
elevated ( 0.37 nm
> ) only within ∼2 nm of the particle edge while
the bulk has d-spacings similar to those in the Gr-AP sample. The d-
spacing distribution (Fig. 9b(iii)) shows a mean around 0.337 nm
and a standard deviation of 0.008 nm, only slightly larger than those
for the Gr-AP sample.
On the other hand, Region 2 (acquired near the edge of a larger
graphite pore) of the Gr-LP sample (Fig. 9c) shows greater lattice
disorder. The average rotation gradient in this region is 0.30 nm,
̃ °/
while the particle edges have values 4 nm
> °/ (Fig. 9c(i)). Moreover,
the frequency of lattice rotation (i.e., the “wavy” nature of the
lattice) occurs up to ∼20 nm into the graphite bulk from the pore
edge. A similar behavior is observed in Fig. 9c(ii), which shows
elevated d-spacings near the pore edge and extending into the
particle bulk. Figure 9c(iii) shows that the mean d-spacing of Region
2 is 0.35 nm, which is larger than that calculated from the X-ray
data. The d-spacing distribution also shows a tail that extends from
∼0.36 nm to 0.40 nm, which indicates that the graphite edges in
Region 2 have sustained greater damage than those in Region 1. This
result also points to the spatial heterogeneity in lattice disorder
within each graphite particle—the edges with greater exposure to the
electrolyte will experience more Li+
ion (de)intercalation reactions
and apparently greater damage, especially during high-rate charging.
Damage localized at the graphite particle edges is not entirely
unexpected.45
Prior studies have pointed out that Li+
ion intercala-
tion into graphite occurs primarily through the edge planes.46,47
At
high charge rates, significant build-up of Li+
ions is expected at the
edges planes; the concentration gradients between the surface and
bulk sites could induce stresses that are significant enough to deform
the graphite lattice near the particle edges.
Mitigating damage at the graphite particle edges is crucial to
prolonging the life of LIB cells subjected to fast charge. One
solution would be to raise the lower cutoff voltage (LCV) for cell
cycling (e.g. from 2.8 to 3.3 V) to prevent complete delithiation of
the graphite. The residual dilute LixC6 phases have higher d-spacings
than the pristine graphite,29
which would facilitate Li+
ion inter-
calation and (partially) mitigate the stresses induced during the
process. Alternative solutions could include the use of graphite with
higher d-spacings in the anode, as has been suggested for both Li+
and Na+
ion cells.48,49
However, such “expanded” graphite materials
typically have lower electronic conductivity, which would limit the
overall rate of intercalation. In addition, these materials have higher
surface areas and a proclivity to trap more lithium in the SEI formed
during electrolyte reduction reactions. Other solutions proposed to
facilitate rapid cycling include the use of functionalized-graphite that
have higher d-spacings at the particle edges,50
graphite with particle
edges aligned perpendicular to the current collector,51
graphite with
carbon coatings that have larger interlamellar spacing,52
and anodes
containing a mix of graphite and hard carbons.7
These efforts, along
with research on cathode and electrolyte materials, electrode
architectures and cycling protocols, are important steps towards
making fast-charging of LIB cells, a commercial reality.
Conclusions
Fast charging of LIB cells can degrade their energy and power
performance characteristics. Here we show representative data from
as-prepared graphite electrodes (Gr-AP) and from electrodes ex-
tracted from (discharged) cells that were previously charged at rates
up to 6 C (Gr-LP). Electrochemical measurements show that the
harvested electrodes have lower capacities than the Gr-AP electrode
and display a sloping lithiation profile, rather than a staged profile,
even at low rates. Analysis of X-ray diffraction data from the Gr-AP
and Gr-LP electrodes yield similar interplanar d-spacings, indicating
that the graphite bulk is not altered by the electrochemical cycling.
Similar conclusions about the graphite particle bulk can be drawn
from the analysis of HREM and HAADF-STEM images and SEND
data. However, these nanoscale data clearly indicate damage at the
graphite edges, which extend up to ∼20 nm into the particle bulk for
the Gr-LP sample. This damage is seen as arcs (rather than spots) in
the FFT data of HREM images and as increased rotations and higher
d-spacings in the SEND data; both datasets suggest increased
turbostratic disorder and buckling of the graphene planes at particle
edges. In addition to changes in graphite-edge structure, the Gr-LP
electrode shows a significant thickness increase from the build-up of
SEI and particle rearrangements that occur during the cycling. This
SEI contains a mix of compounds, including LiF (seen in the XRD

patterns) and (oxy)fluoro-phosphates suggested by the EDS data.
Larger quantities of this SEI are present adjacent to the larger
internal pores of the graphite, which likely have easier access to the
electrolyte than the smaller pores. Furthermore, the SEI is localized
Figure 8. Quantitative lattice rotation and interplanar d-spacing analysis of the as-prepared graphite (Gr-AP) sample. (a) HAADF-STEM images (i) at low-
magnification, showing Regions 1 and 2 where SEND acquisition was performed; (ii), (iii) magnified images of Regions 1 and 2, respectively. Diffractive image
reconstructions and analysis of (b) Region 1 and (c) Region 2 containing the following: (i) in-plane rotation gradient map, (ii) interplanar spacing map, (iii)
interplanar spacing distribution. Scale bars: (a) (i) is 2 μm; (a) (ii), (a) (iii), (b) (i), (b) (ii), (c) (i), (c) (ii) are 50 nm.

at the particle edges and does not penetrate the graphite bulk, suggesting that only Li+
ions (no electrolyte components) intercalate
into the particles (effective desolvation of ions).
Figure 9. Quantitative lattice rotation and interplanar d-spacing analysis of the fast-charged graphite (Gr-LP) electrode. (a) HAADF-STEM images (i) at low-
magniﬁcation showing Regions 1 and 2 where SEND data were acquired; (ii), (iii) magniﬁed images of Regions 1 and Region 2. Diffractive image
reconstructions and analysis of (b) Region 1 and (c) Region 2 containing the following: (i) in-plane rotation gradient map, (ii) interplanar spacing map, (iii)
interplanar spacing distribution. Scale bars: (a) (i) is 2 μm; (a) (ii), (a) (iii), (b) (i), (b) (ii), (c) (i), (c) (ii) are 50 nm.

The implication of our study is that even a small number of high-
rate cycles appears to be sufficient to induce significant and
permanent disorder in graphite domains that are closest to the
electrolyte, be it at particle surfaces or the inner pores. Such disorder
could affect the lithiation kinetics, preventing the graphite particles
from accepting charge and favoring the occurrence of Li plating.
Moreover, we show that these structural changes are highly non-
uniform across the various pores within a same particle, which could
either be a cause or consequence of reaction heterogeneities
commonly observed at high rates. Awareness and mitigation of the
structural evolution at the graphite/electrolyte interfaces is crucial to
the development of LIBs that can endure repeated fast-charging and
yet reliably deliver high-performance over the 10 + year lifespan of
electric vehicles.
Acknowledgments
SP acknowledges support from the U.S. Department of Energy
(DOE) Graduate Student Research (SCGSR) program. The SCGSR
program is administered by the Oak Ridge Institute for Science and
Education for DOE under contract number DE‐SC0014664. This
work was carried out in part at the Materials Research Laboratory
Central Research Facilities, University of Illinois. DA and MTFR
acknowledge support from DOE’s Vehicle Technologies Office
(VTO). The electrodes used in this article were fabricated at
Argonne’s Cell Analysis, Modeling and Prototyping (CAMP)
Facility, which is fully supported by the VTO. We are grateful to
our many colleagues, especially Stephen Trask and Andrew Jansen
at Argonne. Use of the Advanced Photon Source (APS) at Argonne
National Laboratory was supported by the U.S. Department of
Energy, Office of Science, Office of Basic Energy Sciences, under
Contract No. DE-AC02–06CH11357. We are grateful to Saul
Lapidus at the APS for help with the high-resolution powder
diffraction experiments. The submitted manuscript has been created
by UChicago Argonne, LLC, Operator of Argonne National
Laboratory (“Argonne”). Argonne, a U.S. Department of Energy
Office of Science laboratory, is operated under Contract No. DE-
AC02–06CH11357. The U.S. Government retains for itself, and
others acting on its behalf, a paid-up nonexclusive, irrevocable
worldwide license in said article to reproduce, prepare derivative
works, distribute copies to the public, and perform publicly and
display publicly, by or on behalf of the Government.
ORCID
Saran Pidaparthy https://orcid.org/0000-0002-0783-3094
Marco-Tulio F. Rodrigues https://orcid.org/0000-0003-0833-6556
Jian-Min Zuo https://orcid.org/0000-0002-5151-3370
Daniel P. Abraham https://orcid.org/0000-0003-0402-9620
References
1. A. Masias, J. Marcicki, and W. A. Paxton, ACS Energy Lett., 6, 621 (2021).
2. A. Raj, M.-T. F. Rodrigues, and D. P. Abraham, J. Electrochem. Soc., 167, 120517
(2020).
3. A. Tomaszewska et al., eTransportation, 1, 100011 (2019).
4. S. Ahmed et al., J. Power Sources, 367, 250 (2017).
5. A. M. Colclasure et al., Electrochim. Acta, 337, 135854 (2020).
6. X.-G. Yang, T. Liu, Y. Gao, S. Ge, Y. Leng, D. Wang, and C.-Y. Wang, Joule, 3,
3002 (2019).
7. K.-H. Chen, V. Goel, M. J. Namkoong, M. Wied, S. Müller, V. Wood, J. Sakamoto,
K. Thornton, and N. P. Dasgupta, Adv. Energy Mater., 11, 2003336 (2021).
8. D. W. Dees, M.-T. F. Rodrigues, K. Kalaga, S. E. Trask, I. A. Shkrob, D.
P. Abraham, and A. N. Jansen, J. Electrochem. Soc., 167, 120528 (2020).
9. M.-T. F. Rodrigues, K. Kalaga, S. E. Trask, D. W. Dees, I. A. Shkrob, and D.
P. Abraham, J. Electrochem. Soc., 166, A996 (2019).
10. I. A. Shkrob, M.-T. F. Rodrigues, D. W. Dees, and D. P. Abraham, J. Electrochem.
Soc., 166, A3305 (2019).
11. M.-T. F. Rodrigues, I. A. Shkrob, A. M. Colclasure, and D. P. Abraham,
J. Electrochem. Soc., 167, 130508 (2020).
12. T. Waldmann, B.-I. Hogg, and M. Wohlfahrt-Mehrens, J. Power Sources, 384, 107
(2018).
13. L. Somerville, J. Bareño, S. Trask, P. Jennings, A. McGordon, C. Lyness, and
I. Bloom, J. Power Sources, 335, 189 (2016).
14. J. Cannarella and C. B. Arnold, J. Electrochem. Soc., 162, A1365 (2015).
15. J. S. Okasinski, I. A. Shkrob, A. Chuang, M.-T. F. Rodrigues, A. Raj, D. W. Dees,
and D. P. Abraham, Phys. Chem. Chem. Phys., 22, 21977 (2020).
16. S. Müller, J. Eller, M. Ebner, C. Burns, J. Dahn, and V. Wood, J. Electrochem. Soc.,
165, A339 (2018).
17. A. Mukhopadhyay and B. W. Sheldon, Prog. Mater Sci., 63, 58 (2014).
18. A. Raj, I. A. Shkrob, J. S. Okasinski, M.-T. Fonseca Rodrigues, A. C. Chuang,
X. Huang, and D. P. Abraham, J. Power Sources, 484, 229247 (2021).
19. K. P. C. Yao, J. S. Okasinski, K. Kalaga, I. A. Shkrob, and D. P. Abraham, Energy
Environ. Sci., 12, 656 (2019).
20. M. T. F. Rodrigues, S.-B. Son, A. M. Colclasure, I. A. Shkrob, S. E. Trask, I.
D. Bloom, and D. P. Abraham, ACS Appl. Energy Mater., 4, 1063 (2021).
21. M.-T. Fonseca Rodrigues, V. A. Maroni, D. J. Gosztola, K. P. C. Yao, K. Kalaga, I.
A. Shkrob, and D. P. Abraham, ACS Appl. Energy Mater., 2, 873 (2019).
22. P. P. Paul et al., , Adv. Energy Mater., 11, 2100372 (2021).
23. H. Shi, J. Barker, M. Y. Saïdi, R. Koksbang, and L. Morris, J. Power Sources, 68,
291 (1997).
24. A. Van der Ven, Z. Deng, S. Banerjee, and S. P. Ong, Chem. Rev., 120, 6977
(2020).
25. H. Fujimoto, H. Kiuchi, S. Takagi, K. Shimoda, K. Okazaki, Z. Ogumi, and T. Abe,
26. H. He, C. Huang, C.-W. Luo, J.-J. Liu, and Z.-S. Chao, Electrochim. Acta, 92, 148
(2013).
27. S. Schweidler, L. de Biasi, A. Schiele, P. Hartmann, T. Brezesinski, and J. Janek,
J. Phys. Chem. C, 122, 8829 (2018).
28. J. Wilhelm, S. Seidlmayer, S. Erhard, M. Hofmann, R. Gilles, and A. Jossen,
J. Electrochem. Soc., 165, A1846 (2018).
29. E. Hazrati, G. A. de Wijs, and G. Brocks, Phys. Rev. B, 90, 155448 (2014).
30. S.-B. Son, D. Robertson, Z. Yang, Y. Tsai, S. Lopykinski, and I. Bloom,
31. R. E. Franklin, Acta Crystallogr., 4, 253 (1951).
32. J. R. Dahn, A. K. Sleigh, H. Shi, J. N. Reimers, Q. Zhong, and B. M. Way,
Electrochim. Acta, 38, 1179 (1993).
33. J. R. Dahn, R. Fong, and M. J. Spoon, Phys. Rev. B, 42, 6424 (1990).
34. P. Verma, P. Maire, and P. Novák, Electrochim. Acta, 55, 6332 (2010).
35. A. M. Andersson and K. Edström, J. Electrochem. Soc., 148, A1100 (2001).
36. D. Sauerteig, S. Ivanov, H. Reinshagen, and A. Bund, J. Power Sources, 342, 939
(2017).
37. B. Rieger, S. Schlueter, S. V. Erhard, J. Schmalz, G. Reinhart, and A. Jossen,
J. Energy Storage, 6, 213 (2016).
38. A. Y. R. Prado, M.-T. F. Rodrigues, S. E. Trask, L. Shaw, and D. P. Abraham,
39. R. Wagner, M. Korth, B. Streipert, J. Kasnatscheew, D. R. Gallus, S. Brox,
M. Amereller, I. Cekic-Laskovic, and M. Winter, ACS Appl. Mater. Interfaces, 8,
30871 (2016).
40. M. Stich, M. Göttlinger, M. Kurniawan, U. Schmidt, and A. Bund, J. Phys. Chem.
C, 122, 8836 (2018).
41. D. J. Miller, C. Proff, J. G. Wen, D. P. Abraham, and J. Bareño, Adv. Energy
Mater., 3, 1098 (2013).
42. A. Marion, , ed. A. Marion (Springer, US, Boston, MA) p. 255 (1991).
43. R. C. Gonzalez and R. E. Richard E. Woods, Digital image processing 2nd ed.
(Prentice Hall, Upper Saddle River, N.J) (2002).
44. J. M. Zuo and J. C. H. Spence, Advanced Transmission Electron Microscopy
(Springer, New York, NY) (2017).
45. V. A. Sethuraman, L. J. Hardwick, V. Srinivasan, and R. Kostecki, J. Power
Sources, 195, 3655 (2010).
46. R. J. Rice and R. L. McCreery, Anal. Chem., 61, 1637 (1989).
47. K. Persson, V. A. Sethuraman, L. J. Hardwick, Y. Hinuma, Y. S. Meng, A. van der
Ven, V. Srinivasan, R. Kostecki, and G. Ceder, J. Phys. Chem. Lett., 1, 1176 (2010).
48. T.-H. Kim, E. K. Jeon, Y. Ko, B. Y. Jang, B.-S. Kim, and H.-K. Song, J. Mater.
Chem. A, 2, 7600 (2014).
49. Y. Wen, K. He, Y. Zhu, F. Han, Y. Xu, I. Matsuda, Y. Ishii, J. Cumings, and
C. Wang, Nat. Commun., 5, 4033 (2014).
50. J.-S. Park, M.-H. Lee, I.-Y. Jeon, H.-S. Park, J.-B. Baek, and H.-K. Song, ACS
Nano, 6, 10770 (2012).
51. J. Billaud, F. Bouville, T. Magrini, C. Villevieille, and A. R. Studart, Nat. Energy,
1, 16097 (2016).
52. W. Cai, C. Yan, Y.-X. Yao, L. Xu, R. Xu, L.-L. Jiang, J.-Q. Huang, and Q. Zhang,
Small Struct., 1, 2000010 (2020).

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