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A 5V-class Cobalt-free Battery Cathode with High Loading Enabled by
Dry Coating
Weiliang Yao1, Mehdi Chouchane2, Weikang Li3, Shuang Bai1, Zhao Liu4, Letian Li4,
Alexander X. Chen3, Baharak Sayahpour1, Ryosuke Shimizu3, Ganesh Raghavendran3,
Yu-Ting Chen1, Darren H.S. Tan3, Bhagath Sreenarayanan3, Crystal K. Waters5, Allison
Sichler5, Benjamin Gould5, Dennis J. Kountz5, Darren J. Lipomi3, Minghao Zhang3,*,
Ying Shirley Meng2,3,*
1Materials Science and Engineering, University of California San Diego, La Jolla, CA
92093, USA
2Pritzker School of Molecular Engineering, University of Chicago, Chicago, IL, 60637,
USA
3Department of NanoEngineering, University of California San Diego, La Jolla, CA
92093, USA
4Materials and Structural Analysis, Thermo Fisher Scientific, 5350 NE Dawson Creek
Drive. Hillsboro, Oregon 97124, USA
5Advanced Performance Materials Chemours Discovery Hub, The Chemours Company,
Newark, DE 19713, USA
Co-correspondence: miz016@eng.ucsd.edu, shirleymeng@uchicago.edu
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Abstract
Transitioning toward more sustainable materials and manufacturing methods will be
critical to continue supporting the rapidly expanding market for lithium-ion batteries.
Meanwhile, energy storage applications are demanding higher power and energy densities
than ever before, with aggressive performance targets like fast charging and greatly
extended operating ranges and durations. Due to its high operating voltage and cobalt-free
chemistry, the spinel-type LiNi0.5Mn1.5O4 (LNMO) cathode material has attracted great
interest as one of the few next-generation candidates capable of addressing this
combination of challenges. However, severe capacity degradation and poor interphase
stability have thus far impeded the practical application of LNMO. In this study, by
leveraging a dry electrode coating process, we demonstrate LNMO electrodes with stable
full cell operation (up to 68% after 1000 cycles) and ultra-high loading (up to 9.5 mAh/cm2
in half cells). This excellent cycling stability is ascribed to a stable cathode-electrolyte
interphase, a highly distributed and interconnected electronic percolation network, and
robust mechanical properties. High-quality images collected using plasma focused ion
beam scanning electron microscopy (PFIB-SEM) provide additional insight into this
behavior, with a complementary 2-D model illustrating how the electronic percolation
network in the dry-coated electrodes more efficiently supports homogeneous
electrochemical reaction pathways. These results strongly motivate that LNMO as a high
voltage cobalt-free cathode chemistry combined with an energy-efficient dry electrode
coating process opens the possibility for sustainable electrode manufacturing of cost-
effective and high-energy-density cathode materials.
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Introduction
Reducing the content of cobalt in cathode materials is becoming a critical
requirement for the next generation of Li-ion batteries (LIBs). The cost of the battery
industry’s reliance on cobalt is not only highly volatile from a commercial perspective
(doubled to $80,000 USD per ton in one year1) due to strategic international stockpiling
and the monopoly in supply from the Democratic Republic of Congo, but also raises severe
ethical concerns with its mining and child labor.2 Therefore, reducing or completely
removing cobalt from cathode materials has become an urgent challenge to achieve
environmentally benign and sustainable battery manufacturing. Among the candidates,
cobalt-free LiNi0.5Mn1.5O4 (LNMO) with its high operating voltage (~4.7 V) can reduce
the number of cells for the battery pack system, thus providing higher volumetric energy
density. Despite its high energy density and low cost, LNMO is still facing various
commercialization challenges such as poor cycling stability and low electronic
conductivity (~ 10-6 S/cm).3-7
Among the efforts to improve the performance of LNMO, developing novel
electrolyte additives is the most common strategy to stabilize the interphase of both the
cathode and the anode.8-15 However, in full cells, most studies are limited to 200 cycles or
use cathode loading lower than 20 mg/cm², making them incompatible for practical
applications. Materials doping is another approach to stabilize the cathode electrolyte
interphase (CEI) while mitigating the electrolyte corrosion.16 Nevertheless, this approach
has only achieved limited success and the addition of expensive transition metals will raise
the cost. Another approach involves surface coatings, which aim to slow down the cathode
surface degradation and prolong cell cycling17-18 by providing a more robust CEI and
preventing transition metal dissolution. Unfortunately, the cost and equipment required for
scaling up this sophisticated synthesis process likely precludes high-throughput
manufacturing.19-20 Despite the wealth of research focused on improving the performance
of LNMO, few studies have considered a thick electrode approach to achieve practical
usage, i.e., at least 3.0 mAh/cm2 (~21 mg/cm2) per side to achieve around 300 Wh/kg
(Table S1). In these cases, the works targeting practical loadings were limited by either
low cycle number (less than 300 cycles) or poor capacity utilization.21-23
4
To realize the full potential of LNMO, high loading must be achieved
simultaneously with other modifications. Many fabrication strategies have been explored,
including repeated coextrusion/assembly to create artificial channels to reduce tortuosity
and improve the ionic flow,24 dispersing single-wall carbon nanotube (SWCNT) to
fabricate 800 𝜇m electrodes,25 utilizing novel binder such as polyacrylonitrile (PAN) to
enable high loading,21 and adjusting solid content in water based slurry with carbon micro
fibers (CMF).26 However, these methods are all slurry-based approaches and either have
very complex procedures or are limited to lab scale processing. Conventionally, N-Methyl-
2-pyrrolidone (NMP) is widely used as the mixing solvent due to its excellent chemical
and thermal stability as well as its ability to dissolve polyvinylidene fluoride (PVDF)
binder, which offers high mechanical and electrochemical stability in cathode operation.27-
28 However, NMP’s notorious toxicity and requirement of expensive solvent recycling
equipment make the slurry-based fabrication process costly (> $5M for NMP solvent
recovery equipment), more energy demanding, and less sustainable.29
Unlike the abovementioned slurry-based methods, fabrication using binder
fibrillation is a dry process, where polytetrafluoroethylene (PTFE) is the widely used
binder, with the first use in dry electrodes for LIBs reported in 1979.30 In this process,
PTFE particles are shear mixed to form adhesive fibrils which can closely bind both
conductive carbon and active materials. The dry coating process has recently drawn interest
through Maxwell Technologies reporting high long-term cycling performance in 10 Ah
pouch cells31 and Tesla’s announcement in 2020 of using dry electrodes in EV batteries in
their future vehicles.32 Compared to the slurry-based method, this dry process can easily
fabricate roll-to-roll electrodes with extreme thicknesses and no cracks in the electrode.33-
34 More importantly, with no drying process, the dry electrode method saves 45% to 47%
energy consumption and ~1% to 2% total battery cost compared to slurry-based method.20,
29 To reach TWh level energy storage, low cost, less energy waste and reduced
environmental pollution are key factors to achieve sustainable manufacturing. Figure 1
summarizes both conventional slurry-based and novel dry electrode fabrication methods.
In this work, we utilize the binder fibrillation process to fabricate LNMO electrode
at high loadings (>3.0 mAh/cm2 level) and demonstrate the performance improvement of
long-term cycling in the high voltage (> 4.7 V) LIBs’ application. With the combined
5
experimental and modeling efforts, we pinpoint the underlying mechanism of performance
improvement by using dry-coated LNMO electrodes in terms of reduced parasitic reaction,
a highly distributed and interconnected electronic percolation network, and robust
mechanical properties.
Figure 1. Schematic of dry electrode and slurry-based cathode fabrication procedures.
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Results & Discussion
Fabrication of dry electrodes at high loadings
Before electrode fabrication, we evaluated the inactive components such as the
current collector, binder and conductive carbon were stable under high voltage. Linear
sweep voltammetry (LSV) and cyclic voltammetry (CV) testing on various components
were carried out and the results are shown in Figure S1. Interestingly the Super C65 (SC65)
triggered more electrolyte decomposition than vapor-grown carbon fiber (VGCF), likely
because of its higher specific surface area.
To investigate the quality of the thick electrodes, plasma focused ion beam (PFIB)
was adopted to mill electrode cross-sections (Figure S2). The Xe+ based PFIB allows much
faster milling on large volumes (350
𝜇m × 100
𝜇m × 30
𝜇m) compared to conventional
Ga+ based FIB (~1 hour milling versus >12 hours milling). These cross sections revealed
that the slurry-based LNMO using VGCF exhibits severe carbon agglomeration, indicating
the conventional slurry mixing failed to disperse the carbon fibers as uniformly as the SC65.
Similar observations were also found on the top surface of these electrodes (Figure S3),
indicating that the VGCF powder tends to aggregate into a group of fibers. Consequently,
in slurry-based electrodes, most of LNMO particles are not connected by VGCF, which
severely hinders the electron flow from active materials to the current collector. This
behavior is reflected by the poor cycling performance in the full cell for the slurry-based
LNMO using VGCF (Figure S4). In contrast, the shear force applied during the dry process
mixing straightens the fibers without breaking them to form a ‘network’. The dry-LNMO
full cell, which has a more homogeneous carbon distribution (Figure S2A), shows much
more stable performance at both C/10 and C/3 rate. Interestingly, due to its nano-
agglomerate morphology, the SC65 carbon in the dry electrode is unable to construct an
effective electronic percolation network compared to VGCF, which has a similar fiber
morphology as the fibrillated PTFE binder (Figure S5). Thus, this work will focus on the
comparison between dry electrodes with VGCF and slurry-based electrodes with SC65.
Half cells were then fabricated to investigate the electrochemical performance. The
dry electrodes, even with areal loadings as high as 9.5 mAh/cm² (~240 𝜇m), still delivered
similar performance to the baseline 3.0 mAh/cm2 level electrode (Figure 2A and 2C). The
slurry-based LNMO, however, starts to show obvious performance degradation at 4.0
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mAh/cm2, as seen from the increased polarization and capacity degradation in the voltage
profiles (Figure 2B and 2D). This impact is even more pronounced at higher areal loadings,
as the 6.0 mAh/cm2 level slurry-based LNMO experienced severe electrode cracking after
drying. (Figure S6). In slurry-based electrodes at high areal loading, conductive carbon
and binder will easily float near the electrode surface and agglomerate due to capillary
action and diffusion.35 This inhomogeneity in the conductive carbon distribution leads to a
poor electronic percolation network, lowering the effective electronic conductivity further
with electrode thickness. In contrast, for dry electrodes, both in-plane and out-of-plane
electronic conductivities remain in the same order of magnitude as the areal loading is
increased from 3 to 9.5 mAh/cm² (Figure S7). This suggests that electron flow in both
directions is well maintained even at an ultra-high loading, so the VGCF is likely well
distributed. As a result, even at 6.0 mAh/cm2 level, the dry-LNMO half cells can still
deliver >110 mAh/g at a C/3 rate (Figure S8), outperforming the 4.0 mAh/cm2 level slurry-
based LNMO.
8
Figure 2. Electrochemical and mechanical evaluation of high loading LNMO using
both slurry-based and dry electrode method. (A) Half-cell performance of dry-LNMO
and (B) slurry-based LNMO at various areal loadings. SEM cross-section image of (C) 9.5
mAh/cm2 level dry-LNMO and (D) 4.0 mAh/cm2 level slurry-based LNMO. (E)
Normalized peel-off forces and thickness of peeled-off electrodes from both dry-LNMO
and slurry-based LNMO at areal loading of 3.0 mAh/cm2. (F) Representative stress-strain
curves of dry electrodes with various areal loadings.
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The mechanical behaviors of slurry-based and dry electrodes were then investigated
using 90° peel and tensile (pull) tests (Figure 2E and 2F). 90° peel tests (using scotch tape)
were used to compare the debonding behavior of dry and slurry-based LNMO electrodes
with the same areal loading (3.0 mAh/cm2) and thicknesses (~90 𝜇m). Figure 2E compares
the average thickness of each electrode removed during the peel test (i.e., delaminated from
the current collector and remaining on the adhesive tape) relative to the normalized peel-
off force (normalized with the width of the tape). For slurry-based LNMO electrodes, on
average, over half of the electrode was peeled off with less than 5.25 N/m force. In contrast,
the dry-LNMO electrodes required more force (~4.20 to 14.70 N/m) to delaminate while
removing less than half of the electrode thickness. Thus, more force was needed to
delaminate less material for the dry-LNMO electrode in comparison to the slurry-based
electrode. Likewise, the two types of electrodes differed in the manner by which they failed.
Qualitatively, from Figure S9A and S9B, we observed that the dry-LNMO electrodes
primarily experienced cohesive failure (i.e., delamination within the electrode layer), while
the slurry-based electrode experienced a mixture of both cohesive and adhesive (i.e.,
delamination at the Al-electrode interface) failure. Together, these observations suggest
that the dry-LNMO electrodes have better adhesion to the current collector, as well as
greater cohesive strength within the electrode. This difference in adhesive and cohesive
strength is likely due to the morphology of the electrode stemming from the dry fabrication
process. Dry-LNMO electrodes formed homogenously distributed PTFE fibrils, which
help maintain the electrode structure (e.g., dissipate mechanical energy) during mechanical
failure. Likewise, the PTFE binder can take multiple forms, including nanofibrils
(thickness ~20 nm, Figure S10) that can more effectively bind the LNMO and carbon
fibers, further increasing the cohesive strength of the electrode. In comparison, the drying
process of the slurry-based LNMO electrode results in phase segregation of the binder,
carbon, and active materials, leading to uneven binder distribution. This phase segregation
results in a concentration gradient where electrode layers closer to the current collector
contain a lower concentration of PVDF binder and layers further from the current collector
contain a higher concentration.
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The improved uniformity from the dry fabrication process is evident from the
tensile behaviors of dry-LNMO electrodes with various areal loadings (Figure 2F).
Despite the significant increase in thickness, the fracture behavior (Figure S11) and tensile
properties of all electrodes remained similar, with tensile strengths ranging between 0.25
– 0.30 MPa and fracture strains between 18 – 26% (Figure 2F and Table S3). The vertical
uniformity and robust mechanical properties of the dry electrodes help maintain a stable
electrode structure from cell fabrication to long-term electrochemical cycling, thus
providing fast electron transfer and reducing the cell impedance in the long run. More
importantly, the removal of the solvent drying process significantly reduces energy
consumption and eliminates the cost required to recover NMP, making the LIB
manufacturing process sustainable and more environmentally friendly.36
Modeling of the 2D electronic percolation network
To gain further insights from the PFIB-SEM cross-section images, a 2-D modeling
approach has been developed, using a sub-volume of the extracted electrode as an input for
the geometry. The two conditions that will be compared are the dry-LNMO with VGCF
(Figure 3A) and the slurry-based LNMO with SC65 (Figure 3B), through a single
discharge at a C-rate of C/3. The discussion over the fair comparison between these two
slices and more details on the modeling effort such as the model description, and the
parameters used can be found in the Supplementary Information.
From the simulated discharge curves in Figure 3C, the dry-LNMO displays a
specific capacity of 125.2 mAh.g-1 versus 101.6 mAh.g-1 for the slurry-based LNMO. This
difference is due to the higher overpotential in the slurry-LNMO, which will lead to a
shorter discharge time before reaching the cut-off voltage of 3.5 V. On average, a potential
difference of 0.24 V is observed for a given time of discharge between the slurry-based
and dry-LNMO. To investigate the cause of this higher overpotential, the distribution of
the log of the current density in the solid phase is plotted at the end of discharge in Figure
3D and 3E. In the two electrodes, most of the current is flowing through the carbon network,
which is the optimal scenario to prevent a significant ohmic drop due to the poor electronic
conductivity of LNMO (~ 10-6 S/cm).5 In the dry-LNMO on Figure 3D, 99.62% of the
current is flowing through the VGCF, while only 97.26% in the case of the slurry-based
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LNMO on Figure 3E. This indicates a better connectivity of VGCF network thanks to its
fiber morphology compared to the SC65. In the slurry-based LNMO, the higher amount of
current flowing through the LNMO will be detrimental to the performance of the battery,
because it will induce a higher electronic resistivity, hence a higher voltage drop, entailing
that the cut-off voltage will be reached sooner than previously observed.
This difference in electronic transport translates into the state of lithiation
([Li]LNMO/[Li]LNMO,max), with a more uniform and higher state of lithiation, i.e., a better
utilization, of the dry-LNMO (Figure 3F) than the slurry-based LNMO with SC65 (Figure
3G). It is noteworthy that the utilization is higher near the current collector than near the
separator. This gradient is characteristic of a system limited by electronic transport, with
poorly connected active material particles far from the current collector (the source of the
electrons) having a lower state of lithiation. The absence of strong intraparticle gradients
also demonstrates that the intercalation process is not limited by solid diffusion at this rate
of discharge. Overall, through this modeling effort relying on high quality PFIB-SEM
images, the efficiency of VGCF to form an electronic percolating network has been
highlighted versus the reference case of SC65. Even though the model does not capture the
long-term benefits from having VGCF (low surface carbon), that will hinder the side
reactions when compared to SC65 (high surface carbon), it is still able to show its merits
in terms of electronic conductivity.
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Figure 3. 2-D Modeling results based on high-quality PFIB-SEM images. PFIB-SEM
cross-sections of (A) dry coated and (B) slurry-based LNMO with SC65. (C) The simulated
discharge curves for dry coated and slurry-based LNMO. Finally, the 2-D modeling results
at the end of a discharge at C/3 are displayed with the current density in the solid phase
and the state of lithiation reported for the dry coated LNMO respectively in (D,F), and for
13
the slurry-based LNMO respectively in (E,G). The bottom of the electrodes corresponds
to the current collector location and the top to the separator location in the simulations.
Electrochemical performance and interfacial analysis
Long-term cycling performance using 3.0 mAh/cm2 LNMO electrodes was carried
out with graphite as the counter electrode. The dry-LNMO full cell achieved 80% capacity
retention after 300 cycles while slurry-based LNMO full cell only delivered 67% (Figure
4A). Notably, the dry-LNMO full cell enables high Coulombic efficiency (CE%)
throughout the testing. The CE% increases to 99.9% after 115 cycles while slurry-based
LNMO full cell can only reach maximum ~99.8% throughout the cycling. The slurry-based
LNMO full cell shows a continuously growing overpotential while the overpotential of
dry-LNMO full cell is stabilized after 300 cycles (Figure 4B). In addition, the oxidative
peaks from the slurry-based LNMO full cell are found to shift more towards the higher
voltage end, with a significant increasing trend in both redox peaks around 4.52 V
(Ni2+/Ni3+) and 4.7 V (Ni3+/Ni4+) (Figure S15A and B). Also, the reductive peaks move
towards the lower voltage end. In contrast, both oxidative and reductive peak positions
from the dry-LNMO full cell show limited or nearly no shift. These results suggest that a
drastic impedance rise and severe Li inventory loss occurred in the slurry-based LNMO
full cell. This analysis is enhanced by electrochemical impedance spectroscopy (EIS)
measurements (Figure S15C and Figure S16). Moreover, it was also found that after cell
disassembling, the adhesion between dry-LNMO and current collector was well
maintained while slurry-based LNMO had delaminated from the Al foil (Figure S17). Such
contact loss, possibly due to the corrosion of parasitic reaction products on the current
collector or to the weaker adhesion force illustrated in Figure 2E, will further exacerbate
the cell impedance growth.
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Figure 4. Electrochemical performance and interfacial analysis between the dry-
LNMO and slurry-based LNMO thick electrode full cells (A) cycling performance with
Coulombic efficiencies, (B) corresponding average charge and discharge voltages; (C)
XPS spectra of O 1s region of cycled LNMO; TEM images of the surface and/or CEI region
of (D) dry-LNMO and (E) slurry-based LNMO after cycling; (F) XPS spectra of F 1s
region of cycled graphite. TEM images of the surface and/or SEI region of graphite cycled
in (G) dry-LNMO and (H) slurry-based LNMO full cells.
To better understand the impact of interphase properties, we characterized the
CEI/SEI products from electrodes after 300 cycles using X-ray photoelectron spectroscopy
(XPS). From the O 1s spectra (Figure 4C), the intensity of lattice oxygen peak (529.9 eV)
in the cycled slurry-based LNMO is much higher than that in the cycled dry-LNMO. This
implies that the CEI of slurry-based LNMO is not well formed to fully cover the LNMO
surface, as illustrated in Figure 4D, with large variations in thickness including some areas
where no CEI is observed. In contrast, the dry-LNMO particle surface is protected with a
15
conformal 2-nm thick CEI layer. Notably different CEI properties in slurry-based LNMO
could be attributed to more significant attack by HF37 triggered by the higher specific area
conductive carbon electrolyte decomposition. The trace amount of H2O from the
electrolyte and from the carbonate solvent decomposition will react with PF5, which is the
major salt decomposition product, to form strongly acidic HF which will further corrode
the CEI and SEI. The absence of CEI layer on the slurry-based LNMO surface will lead to
more HF corrosion on the particle surface, followed by the increasing dissolution of
transition metal (TM) from the cathode and its redeposition on the graphite (Figure S18).
The dissolution of TM cations to the electrolyte and deposition on the graphite will cause
graphite poisoning, which ultimately lead to fast capacity decay.38 XPS results of cycled
graphite anodes (Figure 4F) further demonstrates the impact of interphase. In the F 1s
spectra, the Li-F peak intensity from graphite cycled in slurry-based LNMO full cell is
significantly higher than that of the dry-LNMO full cell. LiF is well-known as the
decomposition product of the LiPF6 salt.39 This indicates more salt decomposition is
triggered in the slurry-based LNMO full cell during cycling. Moreover, HRTEM results of
cycled graphite anode (Figure 4G and 4H) also show that a much thicker layer of SEI is
formed on the surface of graphite cycled in the slurry-based LNMO full cell. This
observation could suggest that more Li inventory is continuously consumed in the
accelerated side reactions triggered by conductive carbon with higher specific surface area.
Despite the surface TM dissolution observed, no bulk phase change is found in both
types of cycled LNMO electrodes based on the capillary XRD results (Figure S19). The
clear (1 1 1) peak shift indicates the loss of lithium inventories in the bulk structure. Note
that the loss of Li in bulk structure is only part of the lithium inventory loss in the full cell.
Even though there is ~12% capacity retention difference in cycling performance between
dry-coated and slurry-based LNMO full cells, the right shift of (1 1 1) peak is similar in
both cathodes. Thus, the major lithium inventory loss does not happen in the bulk structure,
but more in the interphase formation triggered by the continuous HF corrosion on both CEI
and SEI. In addition, it is also found that PTFE is coated on the carbon fiber during
electrode formation. Such insulating coating further prevents the electrolyte decomposition
on the carbon active sites while allowing fast electron flows within the electrode structure
(Figure S20).
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Enhanced Long-term Cycling Performance
Figure 5. Long-term cycling performance of dry-LNMO full cells with different
electrolyte (A) Specific capacity and (B) Coulombic efficiency of dry-LNMO full cells
using both Gen2 and FEC-FEMC based electrolyte. (C) Cycling performance of dry-
LNMO full cell at 55 °C. Literature summary of LNMO full cell cycling performance using
(D) baseline setups, (E) novel modifications and (F) at high temperature.
With the knowledge gained from the above sections, cells were fabricated to
investigate long-term cycling performance up to 1000 cycles at similar loading. Even with
commercialized carbonate-based electrolyte, our designed dry-LNMO full cell shows a
17
capacity retention of 60% and an average Coulombic efficiency of 99.88% after 1000
cycles at a C/3 rate (Figure 5A and 5B). To move one step further, one type of all-
fluorinated electrolyte 1M LiPF6 in FEC : FEMC = 3 : 7 wt% was applied. The fluorination
of solvent molecules aims to improve the reaction barriers and energies for the
electrolyte/cathode decomposition reactions to further mitigate the parasitic reactions.
Similar types of electrolytes have also been reported in previous research.40-41 It can be
found that both cycling performance and Coulombic efficiency are largely improved to 68%
and 99.96% respectively after 1000 cycles. A comparison of cell cycling performance
between our results and literature reports using both carbonate-based baseline electrolyte
and novel modifications is displayed in Figure 5D and 5E with detailed information listed
in Table S6 and S7. While most of the previous works show less than 300 cycles, several
papers do report more than 300 cycles or higher than 80% capacity retention, albeit with
cathode loadings less than 15 mg/cm2. Elevated temperature cycling is always a challenge
for Mn-base cathode materials. Under such aggressive conditions, the dry electrode with
FEC-FEMC based electrolyte can also deliver a ~70% capacity retention after 100 cycles
whereas most literature work was focused on low loading and half cells (Figure 5C and
5F).42-46 Our work has not only achieved stable long-term cycling after 1000 cycles, but
more importantly, the loading has reached the commercial standard.
18
Figure 6. Schematic of advantages achieved by the dry electrode. Utilization of low-
surface area carbon fiber to reduce side reactions while maintaining robust mechanical
strength and excellent electronic percolation network can be achieved simultaneously in
the dry-coated thick electrode.
The pronounced cycling stability is ascribed to the combined factors of robust
mechanical properties, a highly distributed and interconnected electronic percolation
network, and reduced parasitic reactions (Figure 6). In the dry electrode, strong binding
force enabled by fibrillated PTFE binder helps to maintain the close contact of various
electrode components, which will reduce the cell impedance during cycling. Such binding
force is insufficient in slurry-based electrode especially at large thickness (~150 𝜇m),
which will lead to electrode cracks and failure. In addition, the use of carbon fiber can
effectively link a number of cathode particles to facilitate fast electron transfer and form
an electronic percolation network. Meanwhile, carbon fibers with low surface area reduce
active sites for electrolyte oxidation, which will cause continuous water and HF acid
generation to damage both cathode surface and CEI layer. Furthermore, the partial PTFE
19
coating found on the carbon fiber can further reduce the active sites without affecting
electron transfer.
The advantages achieved in dry electrodes are highly dependent on the appropriate
morphology of both conductive carbon and active materials. With similar degree of binder
fibrillation, active materials with high surface area (e.g. less than 1 𝜇m particle size) may
diminish the effective binder amount per unit area. This will lead to poor cohesion strength
of the electrode and finally result in fabrication failure. As for conductive carbon with low
specific surface area, the ability to well connect active material particles depends on
morphological change during fabrication. As seen from our PFIB-SEM results,
agglomerated VGCF were “straightened” during shear mixing, allowing the fibers to well
connect to a number of LNMO particles and hence building a homogeneous electronic
percolation network. Additionally, structural enhancements attributed to the synergy of
PTFE and VGCF lead to electronic conductivity improvements for ultra-thick dry-coated
electrodes. Finally, it is vital that a carbon-free but surface-porous current collector is used
to remove extra source of electrolyte decomposition at high voltage. The porous surface of
current collector also provides vacancies for particles to adhere into the foil during
calendaring, therefore strengthening the adhesion.
Conclusions
In this work, we developed a PTFE-based dry electrode fabrication method to
prepare high voltage spinel oxide LNMO electrodes. Overcoming major limitations with
slurry-coated electrodes, this process enabled ultra-high loadings (~68 mg/cm2 and ~240
𝜇m) and excellent cycling stability using 3.0 mAh/cm2 level (~21 mg/cm2 and ~90 𝜇m)
electrode at 1000 cycles with both the baseline and a high-performance fluorinated
electrolyte (68% capacity retention after 1000 cycles for the full cell). Based on insights
gleaned from the PFIB-SEM images and the 2D-model, further optimization of this process
will likely bring opportunities to fabricate dry electrodes with an even lower amount of
conductive carbon and binder, therefore improving volumetric energy density. Looking
forward, this methodology can also be applied to other cathode materials operating with
voltages higher than 4.5 V (e.g., Li-rich layer oxide, LiCoMnO4, olivine LiCoPO4, high
20
voltage LCO and NCM) leveraging low surface area conductive carbon with the
appropriate morphology to form effective electronic percolation networks, particularly in
highly loaded electrodes. These considerations confirm the dry electrode method offers a
promising electrode manufacturing solution which is more cost effective, environmentally
benign, and sustainable.
21
Acknowledgements
This work was supported by the Chemours Company. Cathode material in this work was
supported by the U.S. Department of Energy’s Office of Energy Efficiency and Renewable
Energy (EERE) and U.S. Army Tank & Automotive Research Development and
Engineering Command (TARDEC) under the award number: DEEE0008442. The SEM-
EDX and HRTEM in this work were performed in part at the San Diego Nanotechnology
Infrastructure (SDNI) of UCSD, a member of the National Nanotechnology Coordinated
Infrastructure, which is supported by the National Science Foundation (Grant ECCS-
1542148). A.X.C. and D.J.L. acknowledge support from the Air Force Office of Scientific
Research (AFOSR) grant no. FA9550-22-1-0454. A.X.C. acknowledges support from the
UC San Diego President’s Dissertation Year Fellowship. The PFIB in this work was
performed in The Thermo Fisher Scientific Americas NanoPort electron microscopy
facility (located in Hillsboro, Oregon). The XPS in this work were performed at the UC
Irvine Materials Research Institute (IMRI). The ICP-MS, XRD and BET testing in this
work was conducted at Environmental and Complex Analysis Laboratory (ECAL) in the
Chemistry and Biochemistry department in UC San Diego. The authors thank Dr. Marshall
Schroeder for the high performance electrolyte screening and manuscript editing. The
authors thank Neware Instruments for the Neware battery test system. The authors are
grateful for Chemours providing the PTFE binder used in the research. The authors thank
Prof. Zhaoping Liu's group from Ningbo Institute of Materials Technology & Engineering
(NIMTE) for providing the graphite anode.
Authors Contribution
W.Y., M.Z., and Y.S.M. designed the experiments. D.J.K. chose the appropriate PTFE
binder for these experiments and provided PTFE technical content for the manuscript. W.Y.
conducted electrode fabrication, electrochemistry testing, SEM-EDX and all data analysis.
M.C. conducted the modeling and analysis. Z.L. and L.L. conducted the PFIB. W.L. and
W.Y. conducted XPS measurement. A.X.C, D.J.L., W.Y. and B.Sr. conducted the
mechanical testing and analyzed the results. S.B. conducted HRTEM experiments and
analysis. R.S. conducted the ICP measurement. B.S. conducted XRD experiment. Y.-T.C.
22
and D.H.S.T. participated in the scientific and figure discussion. G.R. conducted BET
testing. W.Y. conducted XPS, XRD, PFIB, and ICP data analysis. Y.S.M. and M.Z.
supervised the research. W.Y., M.Z., and M.C. wrote the manuscript. B.G., C.K.W. and
A.S. conducted SEM experiments and did microstructural analysis at Chemours
elucidating the fibrous structure of PTFE/VGCF conducting features. All authors
contributed to the discussion and provided feedback on the manuscript.
Competing interests
Authors declare no competing interests.
Additional information
All data is available in the main text or the supplementary materials. All data needed to
evaluate the conclusions in the paper are present in the paper or the supplementary
materials.
23
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