Concentrated Electrolytes for Enhanced Stability
of Al-Alloy Negative Electrodes in Li-Ion Batteries
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Chan, Averey K. et al. "Concentrated Electrolytes for Enhanced
Stability of Al-Alloy Negative Electrodes in Li-Ion Batteries." Journal
of The Electrochemical Society 166, 10 (June 2019): A1867 © 2019
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http://dx.doi.org/10.1149/2.0581910jes
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Concentrated Electrolytes for Enhanced Stability of Al-Alloy Negative
Electrodes in Li-Ion Batteries
To cite this article: Averey K. Chan et al 2019 J. Electrochem. Soc. 166 A1867
View the article online for updates and enhancements.
Journal of The Electrochemical Society, 166 (10) A1867-A1874 (2019) A1867
Concentrated Electrolytes for Enhanced Stability of Al-Alloy
Negative Electrodes in Li-Ion Batteries
Averey K. Chan, 1,2,∗,zRyoichi Tatara, 3,∗∗,a,zShuting Feng, 4Pinar Karayaylali,5
Jeffrey Lopez, 3Ifan E. L. Stephens, 2,∗∗and Yang Shao-Horn 1,3,5,∗∗∗
1Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, Massachusetts
02139, USA
2Department of Materials, Imperial College London, Royal School of Mines Building, London SW7 2AZ, United
Kingdom
3Research Laboratory of Electronics, Massachusetts Institute of Technology, Cambridge, Massachusetts 02139, USA 4Department of Chemical Engineering, Massachusetts Institute of Technology, Cambridge, Massachusetts 02139, USA 5Department of Mechanical Engineering, Massachusetts Institute of Technology, Cambridge, Massachusetts 02139,
USA
Replacing graphite with alloying Al negative electrodes would allow for the development of high energy density Li-ion batteries. However, large volume changes associated with the alloying/dealloying process often result in pulverization of the electrode and rapid capacity fade during cycling due to the continuous formation of solid electrolyte interphase (SEI) layers and loss of electronic contact. In this study, we report that increasing salt concentration in the electrolyte to> 5 mol dm−3 led to enhanced capacity retention during cycling of Li-Al half-cells, which was accompanied by nearly constant impedance for the Al electrode in lithium bis(fluorosulfonyl)imide (LiFSI)/dimethyl carbonate (DMC) 1:1.1 (mol/mol) superconcentrated electrolyte. X-ray photoelectron spectroscopy (XPS) revealed that a potential hold in the superconcentrated electrolyte formed an SEI layer with a greater LiF concentration than in standard 1 mol dm−3solution. This was supported by Raman spectroscopy of LiFSI solutions in DMC, supplemented with density functional theory calculations, which showed an increased driving force for the reduction of FSI−anions to form LiF from Li+-coordinated DMC complexes with increasing salt concentration. Therefore, the enhanced capacity retention and stability can be attributed to the stability of LiF-rich SEI layers which limit carbonate reduction and charge transfer impedance growth.
© The Author(s) 2019. Published by ECS. 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:10.1149/2.0581910jes]
Manuscript submitted April 15, 2019; revised manuscript received May 24, 2019. Published June 6, 2019.
Current Li-ion battery technology relies on the use of graphite in the negative electrode, which has low average potentials (∼0.1 V vs. Li/Li+) and little irreversible capacity as well as high rate capabil-ity, cycle life and electronic conductivity.1 It is well known that a
solid electrolyte interphase (SEI) forms on the graphite edge plane in ethylene carbonate (EC)-based electrolytes, which allows lithium ions to reversibly intercalate into the graphitic carbon structure. This behavior has been attributed to the reductive decomposition of EC to form an electronically insulating yet ionically conductive SEI layer on the graphite surface, preventing further decomposition of the elec-trolyte and co-intercalation of solvent into the graphite.2–6 Alloying
negative electrode materials, such as Al,7–10Si11and Sn,12–14exhibit
much higher theoretical capacities due to the formation of Li-rich LixM binaries, making them attractive for implementation in advanced high-energy Li-ion batteries.1,8,15However, the large volume changes
required to accommodate the alloying process16,17have been linked to
pulverization of the electrodes, resulting in a loss of electronic contact and destabilization of the SEI.1With these materials, the poor
robust-ness of the SEI repeatedly exposes pristine active material, leading to the continuous formation of electrolyte degradation layers.11
Con-sequently, significant capacity fade has been observed during cycling which has so far hindered the application of alloying negative elec-trode materials in practical Li-ion batteries.7,8,18The electrochemical
properties of several candidate negative electrode materials are com-pared to those of graphite in Table S1. Notably, Si negative electrodes have a very high specific capacity (3579 mAh g−1 for Li15Si4) but undergo the most extreme volume change during lithiation (280%) out of the candidate materials; as such, a large number of battery re-searchers have focused on addressing this problem.1,15,19Despite a
lower specific capacity than Si (993 mAh g−1for LiAl), the density of
∗Electrochemical Society Student Member. ∗∗Electrochemical Society Member. ∗∗∗Electrochemical Society Fellow.
aPresent address: Department of Chemistry and Biotechnology, Yokohama National
University, 79-5 Tokiwadai, Hodogaya-ku, Yokohama 240-8501, Japan.
zE-mail:tatara-ryoichi-nx@ynu.ac.jp;averey.chan15@imperial.ac.uk
Al lends itself to a relatively high volumetric capacity and therefore a high estimated energy increase in a full-cell. Al electrodes also un-dergo a much smaller volume change than Si of 97% upon lithiation. However, the use of Al electrodes presents its own host of challenges, such as the issue of cost, safety and the significant component of resis-tive Al2O3passivation layers in electrodes made using active material nanoparticles, which is a common solution to alleviate fracture in Si electrodes.9Therefore, studying how Al-alloy negative electrodes
can be stabilized, particularly through understanding of the SEI layer which forms on the negative electrode, may assist in the design and development of high energy density Li-ion batteries with long cycle lives.
A potential solution to stabilize the SEI for Al electrodes is the use of concentrated electrolytes, which have attracted recent attention due to the tuneability of physicochemical and electrochemical proper-ties simply by increasing the salt concentration.20Electrolyte solutions
with these high salt concentrations (≥ 3 mol dm−3) have recently been termed “solvent-in-salt” and/or “superconcentrated” electrolytes.20
Generally, Li+ions are coordinated by solvent molecules in the solu-tion to form a [Li(solvent)n]+ structure, where the solvation num-ber n is most often reported to be around four.21–23 Typical
elec-trolyte concentrations used for Li-ion batteries are around 1 mol dm−3 (solvent/Li+ratio∼ 10) and therefore ∼ 60% of solvent molecules are “free” from Li+.24As the salt concentration is increased, the
ac-tivity of free solvent molecules which do not interact with Li+cations decreases.24As a result of this low free solvent activity, interesting
phe-nomena have been widely reported such as high thermal stability,25,26
low volatility,25,27,28suppression of transition metal/reaction
interme-diate/current collector dissolution,29–33expansion of the
electrochem-ical window25,34–36and the formation of compact and stable SEI
lay-ers for Si and Li metal negative electrodes.11These phenomena are
especially prevalent when the solvent/Li+ratio is less than the com-mon solvation number of four (corresponding to a Li salt concentra-tion greater than∼ 3 mol dm−3).11,29Although SEI layer formation
in superconcentrated solutions is not well understood, their stability has been explained by the reduction of salt anions to form a pro-tective LiF film which prevents further electrolyte reductive
decom-position and/or Li dendrite formation.36–39In particular, concentrated
electrolytes with lithium bis(fluorosulfonyl)imide (LiFSI) salt have ex-hibited fast charge transfer and the formation of a protective SEI layer on Si and Li metal due to low Lewis basicity,40small anion size40and
a relatively weak S-F bond in the molecular structure.36,39,41,42The
sta-bility of SEI layers formed from LiFSI salt solutions can be attributed to LiF formation through co-adsorption of dissociated Li+and FSI− and breaking of the S-F bond to leave F(SO2)2N and LiF.43
In this study, galvanostatic cycling and electrochemical impedance spectroscopy (EIS) measurements showed that Li-Al half-cells cycled in LiFSI/DMC 1:1.1 (mol/mol) electrolytes exhibited greater capac-ity retention than those cycled in more dilute solutions as well as nearly constant charge transfer impedance. Through a combination of Raman spectroscopy and density functional theory (DFT) calcula-tions, the solvation structures in superconcentrated and standard elec-trolytes were used to rationalize the observed electrochemical stability by showing that the driving force for the reduction of FSI−anions to form LiF is higher in the superconcentrated electrolyte. Supported by X-ray photoelectron spectroscopy (XPS) and scanning electron microscopy (SEM) analysis, these results suggested that the super-concentrated electrolyte promoted FSI−reduction and therefore the formation of a passivating LiF-rich SEI layer to prevent continuous electrolyte degradation and SEI growth on the negative electrode.
Methods
Experimental methods.—Lithium bis(fluorosulfonyl)imide
(LiFSI,>99%) salt purchased from Oakwood Chemical was used as received. Electrolytes were prepared by mixing LiFSI with dimethyl carbonate (DMC, battery-grade BASF). The residual water contents in the electrolytes were measured by Karl Fischer titration and were below 100 ppm. Viscosities and densities of the electrolytes were measured using a Stabinger viscometer (SVM3001, Anton Paar). Ionic conductivity was determined using a Traceable 23226-505 conductivity meter.
Composite Al electrodes were prepared by mixing 80 wt% Alu-minum powder (Alfa Aesar, 7–15 μm or US Research Nanoma-terials Inc., 100 nm), 10 wt% polyvinylidene fluoride (Kynar) and 10 wt% acetylene black (Chevron) in N-methyl-2-pyrrolidone (Sigma Aldrich). The slurry was then spread onto copper foil (25μm thick-ness) using a gap bar coater and dried at 80°C. Electrodes (1/2 in. diameter, 1.27 cm2area) were punched and dried overnight under dy-namic vacuum (120°C). Active material mass loading of the composite Al electrodes was∼ 1.5 mg cm−2and∼ 0.4 mg cm−2for composite Al nanoparticle electrodes. Aluminum foil electrodes were punched out of aluminum foil (1/2 in. diameter, 16μm thickness) and dried overnight under dynamic vacuum (120°C).
Raman spectra of the electrolytes were measured using a mi-croscope Raman spectrometer system (LabRAM HR, Horiba) with 532 nm laser excitations calibrated with a Si standard. Electrolytes were sealed in a capillary tube and held at 25°C during measurement (T95 system controller, Linkam Scientific). After baseline correction of the spectral range, the integrated intensity of the DMC band (890– 960 cm−1) was normalized. The spectra were then deconvoluted using a Gaussian-Lorentzian (pseudo-Voigt) function.
Coin cells (Hohsen Corp., CR2016) were assembled in an Ar-filled glove box (MBraun,< 0.5 ppm H2O, O2) in a half-cell Li-Al config-uration for galvanostatic measurements of composite electrodes. Alu-minum foil electrodes were charged in two-electrode cells (Tomcell type TJ-AC) in a half-cell Li-Al configuration. Li metal foil (15 mm diameter, 0.16 mm thickness; battery-grade foil, 99.9% purity, Rock-wood Lithium, USA) was used as the counter electrode and was sep-arated from the composite or aluminum foil electrode by a Whatman GF/A glass microfiber separator (19 mm diameter, dried overnight at 150°C under vacuum prior to use) wetted with 100μL of electrolyte. Cells were allowed to rest for 5 hours prior to measurement and then underwent galvanostatic charge/discharge tests from 0.01 VLito 2 VLifor a given number of cycles at a 0.1 C rate (∼ 0.1 mA cm−2geo, 99.3 mA g−1), calculated from the theoretical specific capacity of
LiAl (993 mAh g−1). In this paper, the Li+alloying process into the Al electrode (Li++ Al + e−→ LiAl) is defined as “charging”. Cells were then de-crimped in the glove box and rinsed with 200μL of DMC for further characterization. Aluminum foil electrodes were held at a potential of 0.01 VLiin Tomcell type TJ-AC cells with a cutoff charge of 1 mAh/cm2before being opened in the glove box and rinsed with 200μL of DMC for XPS analysis.
Scanning electron microscopy (SEM) samples were sealed in an aluminum laminate package with a heat sealer inside the glove box for transport, before being opened and quickly placed into the SEM chamber to minimize air exposure. SEM images of the electrode sur-face were taken using a JEOL JSM-5910 scanning electron microscope at a 3 kV operating voltage. Energy dispersive X-ray (EDX) analysis was conducted at a 20 kV operating voltage.
Charged aluminum foil electrodes were transferred from the glove box to the chamber of the X-ray photoelectron spectroscopy (XPS) spectrometer using a sample transfer vessel (ULVAC-PHI, INC.). XPS spectra were collected using a PHI 5000 VersaProbe II (ULVAC-PHI, INC.) with a monochromatized Al Kαsource and a charge neutralizer. Spectra were recorded with a pass energy of 23.5 eV and calibrated with the C1s photoemission peak of adventitious carbon at 285 eV. After subtraction of a Shirley-type background, photoemission lines were fitted using combined Gaussian-Lorentzian functions. All spectra were normalized by fixing the C1s photoemission peak of adventitious carbon (285 eV) to the same value.
For electrochemical impedance spectroscopy (EIS), three-electrode cells were assembled in the glove box with Li metal foil (15 mm diameter), a Whatman GF/A separator (19 mm diameter), a Li4Ti5O12mesh reference electrode (18 mm diameter), a Whatman GF/A separator (19 mm diameter) again, and Al electrode (1/2 in. diameter) from bottom to top, where a mesh Li4Ti5O12reference elec-trode was placed between positive and negative elecelec-trode with two separators and 200μL of electrolyte. A detailed cell configuration can be found in previous work.44A mesh reference electrode was used to
avoid an inhomogeneous electric field during EIS measurement, which is known to cause an artefactual EIS response (ex. “spiral” behavior on Nyquist plots with a Li rod reference electrode).44–48Galvanostatic
and potentiostatic charge and EIS tests were performed using a VMP3 (potentiostat with a frequency response analyser, Biologic) with cells thermally equilibrated by a thermostat chamber (SU-241, Espec) at 25°C. After cell assembly, the Li4Ti5O12mesh reference electrode was electrochemically lithiated (negatively polarized at a constant current of 500μA against the Li metal counter electrode with a cutoff voltage of 1.3 VLi) to a stable reference electrode potential of 1.56 VLi.44,46 The Al working electrode was then charged to an end-of-charge po-tential of 0.1 VLiat 0.1C (99.3 mA g−1) and held at that potential. EIS measurements were carried out at the holding potential (0.1 VLi) every hour with a 10 mV amplitude and a frequency range from∼ 10−1to 106Hz.
Computational methods.—All calculations were performed using
the Gaussian 09 computational package.49Geometries were optimized
at the B3LYP/6-31G(d,p)50,51level of theory; ground states were
ver-ified by the absence of any imaginary frequency. The conductor-like polarizable continuum model (CPCM)52,53was employed to capture
the solvation effects. To mimic the solvation environment in DMC solvent, methyl propanoate was selected as the implicit solvent with the dielectric constant set to 3.1. Single point energy calculations were performed at the B3LYP/6-311++G(d,p) level of theory for vertical reduction energies, where geometry optimization was not allowed for the charged state.
GRed= G
M−− G (M)
Results and Discussion
The concentration-dependent viscosities and ionic conductivities of LiFSI/DMC solutions at 25°C are shown in Figure1a. Concentra-tions, ionic conductivities and viscosities can be found in Table S2.
Journal of The Electrochemical Society, 166 (10) A1867-A1874 (2019) A1869 Slope = 2.40 R2 = 0.997 Coordinated DMC Free DMC If /(If + Ic) Ic /(If + Ic) 1:1.1 1:2 1:4 1:6 1:8 1:10.8 1:24 DMC
Figure 1. (a) Ionic conductivity and viscosity measurements for solutions of LiFSI in DMC at 25°C. Data points are labelled with cDMC/ cLiFSIratios. Viscosities
and densities of the electrolyte were measured using a Stabinger viscometer (SVM3001, Anton Paar). Ionic conductivities were determined using a Traceable 23226–505 conductivity meter. (b) Normalized Raman spectra of LiFSI/DMC solutions measured at 25°C from 700 to 790 cm−1where free FSI−S-N stretching can be assigned to∼ 722 cm−1and (c) from 890 to 960 cm−1, where bands at∼ 935 cm−1can be assigned to coordinated DMC and at∼915.3 cm−1to free DMC. Predicted Raman shifts from DFT calculations for monodentate-DMC (DMC-Mono) and bidentate-DMC (DMC-Bi) were scaled by shifting the DFT-predicted free DMC Raman shift to 915.3 cm−1to match the experimental Raman shift. The isosbestic point at∼ 923 cm−1is marked in blue. (d) Fraction of the integrated intensities of Raman peaks corresponding to free DMC, If/(If+ Ic) and coordinated DMC, Ic/(If+ Ic) from (c). Ifis the integrated intensity of the free DMC peak
centered at∼915.3 cm−1and Icis the integrated intensity of the coordinated DMC peak at∼ 935 cm−1. The dotted line is a linear fit of the concentrations up to
0.25 cLiFSI/ cDMC(LiFSI/DMC 1:4 or more dilute) with R2= 0.997.
The viscosity of solution increases rapidly at concentrations greater than 4 mol dm−3 due to an increasing number of solvated cations with a larger hydrodynamic radius than free solvent molecules in the solution. This effect is combined with the enhanced Coulombic inter-action between solvated Li+cations and FSI−anions with decreasing distance.24,54–57The ionic conductivity first increases with increasing
LiFSI concentration up to∼ 1.5 mol dm−3, which can be attributed to increased mobile Li+concentration as supported by previous re-ports with dimethyl sulfoxide (DMSO),24 sulfolane,58 carbonates29
and ethers.59Beyond 1.5 mol dm−3, the conductivity decreases with increasing Li salt concentration as a result of the decreased mobility of ionic species59associated with increased viscosity.24,54The decrease
in conductivity may also be due to reduced salt dissociation60,61as a
result of direct anion coordination to Li+at high salt concentrations when there are insufficient solvent molecules to fully stabilize the cation. These two factors typically yield a maximum ionic conduc-tivity at concentrations of∼ 1 mol dm−3, as observed extensively for ether- and carbonate-based electrolytes.21,24,54,55
Concentration-dependent solution Raman spectra were measured in order to investigate the Li+solvation structures in LiFSI/DMC solu-tions. The spectrum shown in Figure1cshows that pure DMC exhibits a band centered at∼915 cm−1corresponding to the O-CH3stretching mode of the DMC solvent.29,62Dissolving LiFSI into DMC shifts this
band to 930 – 935 cm−1, with a clear isosbestic point at∼ 923 cm−1, and introduces a band centered at∼ 730 cm−1 corresponding to the FSI− S-N symmetric stretching mode, shown in Figure1b.29,63 The
Raman spectrum of the superconcentrated LiFSI/DMC 1:1.1 solu-tion did not pass through the isosbestic point, which suggests that the coordinated DMC band shifted further to a higher frequency and that the strength of the Li+-DMC interaction increased due to a lack of free solvent to stabilize the Li+cation (a strong Lewis acid). The presence of distinct solvation structures in superconcentrated and standard solutions agrees with results reported for concentrated solu-tions of lithium bis(trifluoromethanesulfonyl)amide (Li[TFSA])/N,N-dimethylformamide56and Li[TFSA]/ DMSO.24
As shown in Figure1b, there is an increase in Raman shift from ∼728 cm−1 to 753 cm−1 as the concentration is increased, which is
often observed in concentrated amide-based solutions.29,60,64,65
Previ-ous work has shown that a Raman shift of∼ 752 cm−1can be assigned to aggregate clusters (FSI−coordinated to≥ 2 Li+) in concentrated LiFSI solutions.29,61In addition, it has been shown that anion bidentate
coordination (two sulfonyl oxygens in FSI−interacting with Li+) is more favorable in concentrated solutions29,54,56due to the stability of
the bidentate structure in the absence of solvent contributions.66It was
reported that standard concentrations of Li[TFSA]/DMSO solutions had high concentrations of free TFSA−anions due to the solvent’s high donor number and the dissociative nature of Li[TFSA] salt.24As
LiFSI is also highly dissociative,67LiFSI/DMSO 1:10 (mol/mol)
so-lution can be used as a reference for the Raman shift for free FSI−. The Raman shift for even the most dilute LiFSI/DMC 1:24 solu-tion still does not reach the value corresponding to free FSI−S-N stretching (∼722 cm−1) found by deconvoluting the Raman spectra for LiFSI/DMSO 1:10 (mol/mol) solution shown in Figure S1. There-fore, it can be inferred that FSI− is still coordinated to Li+ in the solvation structure in standard LiFSI/DMC solutions.
DFT-predicted Raman shifts of bidentate (923 cm−1, two ester oxy-gens in the carbonate molecule interact with Li+, schematic shown in Figure1c), monodentate (959 cm−1, carbonyl oxygen in the car-bonate molecule interacts with Li+, schematic shown in Figure1c) and free DMC (933 cm−1) were compared to the normalized DMC band in Figure1cby scaling the DFT-predicted Raman shift for free DMC to 915.3 cm−1, the value determined experimentally from pure DMC. The experimental Raman shift corresponding to coordinated DMC suggests that predominantly monodentate DMC is found in the solvation structures in both standard and concentrated electrolytes. Considering the molar ratio of LiFSI to DMC, the likely Li+ solva-tion structure in superconcentrated LiFSI/DMC 1:1.1 solusolva-tion would contain one bidentate LiFSI and one monodentate DMC molecule (LiFSI-Bi-DMC-Mo). With this configuration, Li+is coordinated by only 3 oxygens (one oxygen from monodentate DMC and two oxygens from bidentate FSI−), which is lower than the common Li+solvation number of four.24Thus, in LiFSI/DMC 1:1.1 solution, Li+is
energet-ically frustrated and DMC interacts more strongly with Li+than with concentrations cDMC/cLi≥ 4, where Li+is always coordinated by at
1st 15th 30th 1st 15th 30th LiFSI/DMC 1:1.1 LiFSI/DMC 1:4 LiFSI/DMC 1:24 LiFSI/DMC 1:10.8
Figure 2. Galvanostatic charge/discharge curves for Li-Al half-cells using (a) LiFSI/DMC 1:1.1 and (b) LiFSI/DMC 1:10.8 (mol/mol) solutions as the electrolyte
with the 1st, 15thand 30thcycles shown. Galvanostatic cycling was conducted at 25°C and a rate of 0.1C based on the theoretical capacity of 993 mAh g−1in
a cutoff voltage range of 0.01 V to 2 V. (c) Specific discharge (de-lithiation) capacity of Li-Al half-cells vs. cycle number in LiFSI/DMC electrolytes. Molar ratios of LiFSI to DMC are shown. Galvanostatic cycling was conducted at 25°C and a rate of 0.1C based on the theoretical capacity of 993 mAh g−1in a cutoff voltage range of 0.01 V to 2 V. For molar ratios of 1:1.1 and 1:10.8, the data shown is an average of three tests with error bars representing the standard deviation.
least 4 oxygens. As a result, a high frequency shift of the coordinated DMC band is observed, as discussed above.
The fractional integrated intensity contributions of the free (∼915.3 cm−1) and coordinated DMC (∼ 935 cm−1) were used to determine the Li+solvation number, n, where the solvation complex was given by [Li(DMC)n]+. Detailed deconvolution spectra are shown in Figure S2. The fractional integrated intensities of Li+-coordinated DMC increase linearly while those of free DMC decrease with increas-ing molar ratio of LiFSI/DMC for cLiFSI/ cDMC≤ 0.25 (LiFSI/DMC 1:4 or more dilute), as shown in Figure1d. In less concentrated elec-trolytes such as LiFSI/DMC 1:10.8 and 1:24, the fraction of free DMC is much greater because the molar ratio of DMC to LiFSI (cDMC/ cLiFSI) is much higher than the typical solvation number of Li+in aprotic sol-vents (4–5), in agreement with previous work.29On the other hand, the
molar ratio of LiFSI/DMC 1:1.1 in the superconcentrated electrolyte is much greater than the typical solvation number of Li+. Assuming that the Raman scattering coefficients for free and coordinated DMC are identical, one can relate the fractional integrated intensity to n, where Icis the integrated intensity of the coordinated DMC peak, If is the integrated intensity of the free DMC peak and ci is the molar concentration of species i:24 Ic If + Ic = cDMC(c) cDMC = ncLiFSI cDMC
The solvation number n was found by fitting the points where cLiFSI / cDMC ≤ 0.25 in Figure1d, yielding the dotted line shown with a slope of 2.4 for n. However, DFT calculations have shown that the scattering coefficients are not identical for free and coordinated vi-brational modes in DMC; more accurate predictions should divide these solvation numbers by 0.93–0.95.68Thus, we applied this
correc-tion to obtain a solvacorrec-tion number n of 2.5–2.6 (the number of solvent molecules coordinating with Li+), indicating the predominant forma-tion of [Li(DMC)3]+complexes. Therefore, the likely Li+solvation structure in standard solution is one monodentate LiFSI and three monodentate DMC molecules (LiFSI-Mo-3DMC-Mo).
The stability of composite Al electrodes was investigated as a function of concentration via half-cell galvanostatic measurements. The charge/discharge curves with LiFSI/DMC 1:1.1 and 1:10.8 elec-trolytes in Figure2aand Figure2b exhibit the features that are to be expected with Li-Al half-cells, but also show enhanced capacity retention during cycling in the superconcentrated electrolyte. A clear plateau at 0.28 VLicorresponds to the Li alloying reaction (Li++ Al + e−→ LiAl) and a plateau at 0.42 VLicorresponds to the de-alloying reaction (LiAl→ Li++ Al + e−).7,8,69–71The initial charging before
reaching the lithiation plateau at 0.28 VLiis due to reduction of the
oxidized surface film,8while the overpotential observed for the
lithia-tion reaclithia-tion can be associated with the nuclealithia-tion of LiAl phase and the work required to deform the surrounding active material.72The
charge/discharge curves also show that there is no significant polar-ization resistance growth associated with SEI formation in both cases. The difference between the capacities reached with these practical cells (∼ 500 mAh g−1) and the theoretical capacity of 993 mAh g−1is due to a combination of the cutoff voltage, inhomogeneous reaction of the active material72,73and the finite charging rate.74Charge/discharge
curves using LiFSI/DMC 1:4 and 1:24 (mol/mol) are shown in Figure S3. The charge/discharge curves for composite Al nanoparticle elec-trodes in both LiFSI/DMC 1:1.1 and 1:10.8 (mol/mol) electrolytes are shown in Figure S4, which exhibited capacities< 70 mAh g−1and no clear lithiation plateau. Similar pseudocapacitive behavior has been reported in previous work with Al nanoparticle electrodes9,75and has
been linked to an increase in surface area for Al2O3 formation and ability to accommodate mechanical stress during lithiation. As a re-sult, there is a greater cumulative thickness and minimal cracking of Al2O3surface films such that Li+is unable to penetrate through to the active material.9 Active material nanoparticles are therefore not
a viable solution, in this case, to stabilize the electrode material and prevent fracture.
The enhanced capacity in superconcentrated LiFSI/DMC 1:1.1 compared to standard electrolyte concentrations is clear when the spe-cific discharge capacity of the Li-Al half-cells is plotted against cycle number, as shown in Figure2c. Notably in the case of the superconcen-trated electrolyte, there is an initial dip in capacity reaching a minimum at approximately 5 cycles; however, there is an increase in capacity after approximately 15 cycles. Although the ionic conductivities of the standard LiFSI/DMC 1:10.8 and intermediate LiFSI/DMC 1:4 so-lutions are higher than the superconcentrated solution (Figure1a), severe capacity fade was observed due to greater free carbonate ac-tivity and insufficient reduction of FSI−in the electrolyte, leading to the formation of an ineffective SEI which cannot suppress contin-uous electrolyte decomposition. Unfortunately, Al electrodes cycled in the superconcentrated electrolyte dropped to a low capacity of≤ 100 mAhg−1 after 30 cycles, suggesting that the electrolyte is un-able to stabilize the electrolyte-Al interface over extensive cycling (Figure2c).
SEM images were used to discern the state of the electrode after cycling and confirmed that cycled Al electrodes were not stable over extensive cycling in both LiFSI/DMC 1:1.1 and 1:10.8 solutions. The SEM images in Figure3show that after 5 cycles, the Al particles in both superconcentrated and standard electrolytes fractured to a size of ∼ 10 μm from an initial size of ∼ 50 μm after the first cycle. EDX
Journal of The Electrochemical Society, 166 (10) A1867-A1874 (2019) A1871
Figure 3. SEM images of composite Al electrodes after 1, 5 and 15 cycles
in LiFSI/DMC 1:1.1 and LiFSI/DMC 1:10.8 (mol/mol) solutions as the elec-trolyte. Images were taken with an operating voltage of 3 kV and at 500x magnification. The scale bars shown correspond to 50μm.
maps in Figure S5 confirm that the grains in question are indeed Al active material. The first dip in the discharge capacity seen in Figure2c
therefore corresponds to the fracture of the Al particles in the compos-ite electrode. The capacity fade during cycling has been related to the depth of lithiation during charge, with cracking and delamination oc-curring during the volume contraction associated with delithiation.72,75
The lithiated LiAl phase has a brittle Zintl phase structure which if over-lithiated, will experience high stresses making it susceptible to fracture.72SEM images taken after 15 cycles show that particle sizes
are still∼ 10 μm and that fracture has occurred in both electrolytes, despite the large discrepancy in discharge capacity. SEM images of the electrode after 5 and 15 cycles in LiFSI/DMC 1:10.8 solution have much less particle definition than the corresponding samples in the su-perconcentrated electrolyte, which indicates thicker SEI buildup in the standard electrolyte. SEM images of pristine electrodes and after 30 cycles are shown in Figure S6.
EIS measurements were conducted with a three-electrode config-uration to investigate impedance growth related to SEI formation as a function of electrolyte concentration. Two semicircles are typically observed in the Nyquist plot for a Li intercalation or alloying reaction into a composite electrode.9,44It has been proposed that the higher
frequency semicircle (left) corresponds to impedance at the electri-fied interface where Li+migration and adsorption/desorption occurs in the pore structure, coupled with electron migration at the grain boundaries of the composite electrode. The low frequency semicircle (right) has been assigned to the charge transfer resistance through the electrode-electrolyte interface.44Earlier studies have shown that the
low frequency resistance grows with increasing SEI thickness, which is indicative of an increasing energy barrier for charge transfer through the SEI layer.9,44The Nyquist plot in Figure4ashows that the
elec-trode held at 0.1 VLi in the superconcentrated electrolyte exhibited a very stable electrified interface impedance and charge transfer/SEI impedance with negligible increase after 10 hours. In contrast, the Nyquist plot for the electrode held in the standard electrolyte in Fig-ure4bshows that there was a continuous increase in charge transfer resistance linked to continuous growth of the SEI layer. On the other hand, the high frequency component did not increase over time in the standard electrolyte, which is consistent with its assignment to the electrified interface impedance (predominantly associated with Li+
1 hour 10 hours Rs CPEHF CPELF RHF RLF Charge Transfer & EEI Impedance W Electrified Interface Impedance
b) Potential held at 0.1 VLi in LiFSI/DMC 1:10.8
a) Potential held at 0.1 VLi in LiFSI/DMC 1:1.1
1 hour2 hours 3 hours4 hours
5 hours10 hours
Figure 4. Time-dependent Nyquist plots for an Al composite electrode in an
Al | Li4Ti5O12mesh | Li three-electrode cell with (a) LiFSI/DMC 1:1.1 and
(b) LiFSI/DMC 1:10.8 (mol/mol) solutions as the electrolyte. The equivalent circuit for the impedance spectra is shown as an inset. The cell was galvano-statically charged at 93.3 mA g−1(0.1C) and held at a potential of 0.1 VLiat
25°C.
migration in the pore structure) and should remain constant during a potential hold. Therefore, it can be inferred that the SEI formed in the superconcentrated electrolyte acted as a passivating layer which suppressed further electrolyte decomposition, whereas there was con-tinuous SEI growth in the standard electrolyte.
XPS analysis on aluminum foil electrodes held at 0.01 VLiin Li-Al half-cells revealed that a potential hold in the superconcentrated electrolyte formed an SEI layer with a greater LiF component than in the standard electrolyte. An aluminum foil electrode was used in-stead of a composite electrode to avoid any misinterpretation from the presence of conductive carbon and/or PVDF binder.91 The F1s
photoemission lines in Figure5show that LiFSI and LiF, from the degradation of LiFSI, were found in the SEI from both standard and superconcentrated electrolytes. It is clear that after charging in the superconcentrated electrolyte, there was a greater contribution of LiF in the SEI than in the standard electrolyte, indicating that more FSI− anion reduction had occurred. There is also an increase in the LiFSI contribution compared to the standard solution, which is simply due to the higher salt concentration in the LiFSI/DMC 1:1.1 solution.
In the Li1s spectra for the superconcentrated electrolyte, a greater LiF contribution is also observed along with reduced Al2p signals associated with aluminum oxides and aluminum metal. The reduced yet still detectable Al2p signals in the superconcentrated electrolyte indicate that the SEI layer was thicker than in the standard electrolyte after the first charge but remained thinner than the sampling depth of ∼ 10 nm for XPS.78The thick initial SEI layer formed in the
concen-trated electrolyte is likely to remain stable with further cycles, whereas the SEI layer in the standard electrolyte will grow thicker over time due to continuous electrolyte degradation.1,11This conclusion is
sup-ported by the continuous growth of the low frequency charge transfer resistance in standard electrolyte concentrations, in contrast with the effectively constant resistance in superconcentrated solution, as previ-ously shown in Figure4. The contributions of C-O (Eb∼ 286.3 eV),86 O=C-O (Eb ∼ 288.8 eV)86and CO3(Eb ∼ 290.3 eV)81in the C1s spectra can be attributed to deposition of carbonaceous species as-sociated with carbonate reduction and SEI formation in carbonate-based electrolytes.77,92The greater intensity of these contributions in
Binding Energy (eV) Intensity (a.u.) LiFSI/DMC 1:1.1 LiFSI/DMC 1:10.8
Figure 5. XPS spectra of the F1s, O1s, C1s and Li1s/Al2p photoemission lines collected from aluminum foil electrodes held at a potential of 0.01 V with a
cutoff charge of 1 mAh/cm2in Li-Al half-cells with LiFSI/DMC 1:1.1 and 1:10.8 (mol/mol) electrolytes at 25°C. The fitted envelope is shown in black and the measured data points denoted by white circles. F1s spectra (scaling factor= 2.8) were assigned with the following contributions: LiF (Eb∼ 685.1 eV)76
and LiFSI (Eb ∼ 687.8 eV).77,78 O1s spectra (scaling factor = 2.5) were assigned with the following contributions: ROLi (Eb ∼ 531 eV),79,80 surface
O/CO3/O-C=O (Eb ∼ 532 eV)81,82and C-O/O-C=O/OP(OR)3 (Eb ∼ 533.4 eV).76,83,84 C1s spectra (scaling factor= 1) were assigned with the
follow-ing contributions: C-H/C-C (Eb = 285 eV),85 C-O (Eb ∼ 286.3 eV),86 C=O/O-C-O (Eb ∼ 287.6 eV),87,88O=C-O (Eb ∼ 288.8 eV)86 and CO3 (Eb ∼
290.3 eV).81Li1s/Al2p spectra (scaling factor= 0.36) were assigned with the following contributions: LiF (E
b∼ 56.8 eV),77Al (Eb∼ 73.4 eV)89,90and Al2O3
(Eb∼ 76.4 eV).89,90
the standard electrolyte suggests that more carbonate reduction had occurred in the first charge, which is consistent with the higher con-tributions from ROLi (Eb∼ 531 eV)79,80and surface O/CO3/O-C=O (Eb∼ 532 eV)81,82in the O1s spectra.
To connect solvation structure with electrochemical stability, the lowest unoccupied molecular orbitals (LUMOs) of several coordina-tion complexes were calculated. The LUMO can be used to describe the ability to receive electrons and the susceptibility toward reduc-tion; atoms on which the LUMO is localized are the strongest electron acceptors.93–95 The maps of the LUMOs in Figure6show that the
LUMO density for the LiFSI-Bi-DMC-Mo solvation structure (Fig-ure6a) is localized around the FSI−anion, which suggests that the anion will be preferentially reduced. The LUMO density for the LiFSI-Mo-3DMC-Mo structure (Figure6b) is localized on both the FSI− an-ion and DMC molecules such that both FSI−and DMC can be reduced on the negative electrode surface. The calculated reduction energies in Figure6cshow that there is a greater driving force for LiFSI-Bi-DMC-Mo reduction than for LiFSI-LiFSI-Bi-DMC-Mo-3DMC-LiFSI-Bi-DMC-Mo reduction. Therefore, we can infer that superconcentrated LiFSI/DMC solutions are more likely to be reduced than standard solutions. Reduction of the FSI−anion is likely to lead to the formation of a stable, inorganic and compact SEI consisting of LiF,41,42which has been related to inhibited electrolyte
degradation36,38,96,97without an increase in interfacial resistance
dur-ing cycldur-ing.38,40The lower energy of FSI−orbitals may be due to the
shift in electron density from the FSI−anion to the Li+cation, which will lower the LUMO of FSI−. In the case of the LiFSI-Mo-3DMC-Mo solvation structure observed in standard solutions (Figure6b), a small amount of the LUMO density is also centered around the DMC molecules since the FSI−anion does not interact strongly enough to Li+ and consequently retains a high electron density. The electron density around the anion prevents direct reduction of FSI−and makes DMC reduction more energetically favorable. Using natural popula-tion analysis, computed anion charges on the FSI−anion were shown
to be−0.85 a.u. for LiFSI-Bi-DMC-Mo (concentrated) and −0.88 a.u. for LiFSI-Mo-3DMC-Mo (standard). As the computed anion charge is less negative in LiFSI-Bi-DMC-Mo, the results also support the conclusion that more charge transfer occurs between Li+and FSI−in concentrated electrolytes.
Conclusions
Raman spectroscopy was used to investigate the solvation struc-ture of LiFSI in DMC. In standard concentrations, Li+forms solvate [Li(DMC)3]FSI while [Li(DMC)]FSI is formed in superconcentrated solution. A combination of galvanostatic cycling tests and EIS mea-surements showed that Li-Al half-cells cycled in the superconcen-trated electrolyte exhibited higher capacity and capacity retention than more dilute concentrations while maintaining an essentially constant charge transfer impedance. In contrast, the charge transfer impedance increased with time in the standard electrolyte. SEM images showed that the superconcentrated electrolyte was able to prevent buildup of resistive and non-passivating carbonate SEI but was unable to pre-vent fracture of active material particles. XPS analysis was used to detect a greater LiF concentration and a lower concentration of car-bonate degradation products in the SEI layer after a potential hold in the superconcentrated electrolyte, indicating that the enhanced capac-ity retention is the result of a stable LiF-rich interface between the Al active material and the electrolyte. These results were rationalized using DFT calculations, which suggested that the LUMO density is localized around the FSI−anion in the [Li(DMC)]FSI complex, which also had a greater driving force for reduction than the [Li(DMC)3]FSI complex. The conclusions from this work may support progress with Al-alloy negative electrode development and may also be a stepping
Journal of The Electrochemical Society, 166 (10) A1867-A1874 (2019) A1873
a)
c)
b)
Gred = G(M––) – G(M)
Figure 6. Computed reduction energies and maps of the lowest unoccupied
molecular orbitals (LUMOs) for Li+solvation structures in LiFSI/DMC solu-tions: (a) bidentate LiFSI and monodentate DMC (b) monodentate LiFSI and three monodentate DMC. Reduction energies were calculated using B3LYP/6-311++G(d,p)//B3LYP/6-31G(d,p) in implicit methyl propanoate with the di-electric constant set to 3.1. Reduction energies for configurations (a) and (b) are shown in (c). The LUMO maps (isovalue= 0.02) were generated using the optimized geometries obtained at B3LYP/6-31G(d,p). Atom color code: gray – carbon, white – hydrogen, red – oxygen, purple – lithium, yellow – sulfur, blue – nitrogen, cyan – fluorine. The red and green regions represent the pos-itive and negative parts of the orbital wave functions, respectively. Solvation structure (a) is likely in superconcentrated solution, while solvation structure (b) is likely in standard concentrations.
stone in accessing the even higher capacities offered by Si as a negative electrode material.
Acknowledgments
This work was financially supported by BMW Group. This work made use of the MRSEC Shared Experimental Facilities at MIT (XPS and SEM), supported by the National Science Foundation under award number DMR-1419807. A.K.C. gratefully acknowledges the support from the Imperial-MIT Department of Materials Exchange Program. S.F. gratefully acknowledges the Link Foundation for the Energy Fel-lowship. This research used resources of the Extreme Science and Engineering Discovery Environment (XSEDE), which is supported by National Science Foundation grant number ACI-1548562. J.L. gratefully acknowledges support by an appointment to the Intelli-gence Community Postdoctoral Research Fellowship Program at the Massachusetts Institute of Technology, administered by Oak Ridge Institute for Science and Education through an interagency agree-ment between the U.S. Departagree-ment of Energy and the Office of the
Director of National Intelligence. I.E.L.S. acknowledges the ISCF Faraday Challenge project: “Towards a Comprehensive Understand-ing of Degradation Processes in EV Batteries” for fundUnderstand-ing under EP/S003053/1.
ORCID
Averey K. Chan https://orcid.org/0000-0001-6412-7199
Ryoichi Tatara https://orcid.org/0000-0002-8148-5294
Shuting Feng https://orcid.org/0000-0001-5630-7085
Jeffrey Lopez https://orcid.org/0000-0002-6425-5550
Ifan E. L. Stephens https://orcid.org/0000-0003-2157-492X
Yang Shao-Horn https://orcid.org/0000-0001-8714-2121
References
1. M. N. Obrovac and V. L. Chevrier,Chem. Rev., 114, 11444 (2014).
2. R. Fong, U. V. Sacken, and J. R. Dahn,J. Electrochem. Soc., 137, 2009 (1990). 3. M. Gauthier, T. J. Carney, A. Grimaud, L. Giordano, N. Pour, H. H. Chang,
D. P. Fenning, S. F. Lux, O. Paschos, C. Bauer, F. Maglia, S. Lupart, P. Lamp, and Y. Shao-Horn,J. Phys. Chem. Lett., 6, 4653 (2015).
4. D. Aurbach, M. D. Levi, E. Levi, H. Teller, B. Markovsky, G. Salitra, U. Heider, and L. Heider,J. Electrochem. Soc., 145, 3024 (1998).
5. H. Schranzhofer, J. Bugajski, H. J. Santner, C. Korepp, K.-C. Möller, J. O. Besenhard, M. Winter, and W. Sitte,J. Power Sources, 153, 391 (2006).
6. T. Abe, H. Fukuda, Y. Iriyama, and Z. Ogumi,J. Electrochem. Soc., 151, A1120 (2004).
7. B. Ji, F. Zhang, M. Sheng, X. Tong, and Y. Tang,Adv. Mater., 29, 1604219 (2017). 8. Y. Hamon, T. Brousse, F. Jousse, P. Topart, P. Buvat, and D. M. Schleich,J. Power
Sources, 97–98, 185 (2001).
9. X. Lei, C. Wang, Z. Yi, Y. Liang, and J. Sun,J. Alloys Compd., 429, 311 (2007). 10. C. J. Wen, B. A. Boukamp, R. A. Huggins, and W. Weppner,J. Electrochem. Soc.,
126, 2258 (1979).
11. Z. H. Chang, J. T. Wang, Z. H. Wu, M. Gao, S. J. Wu, and S. G. Lu,ChemSusChem,
11, 1787 (2018).
12. R. A. Huggins,J. Power Sources, 81–82, 13 (1999). 13. R. A. Huggins,Solid State Ion., 113–115, 57 (1998).
14. M. Wachtler, M. Winter, and J. O. Besenhard,J. Power Sources, 105, 151 (2002). 15. M. Wang, F. Zhang, C. Lee, and Y. Tang,Adv. Energy Mater., 7, 1700536 (2017). 16. B. A. Boukamp, G. C. Lesh, and R. A. Huggins,J. Electrochem. Soc., 128, 725 (1981). 17. C. K. Chan, H. Peng, G. Liu, K. McIlwrath, X. F. Zhang, R. A. Huggins, and Y. Cui,
Nat. Nanotechnol., 3, 31 (2008).
18. A. N. Dey,J. Electrochem. Soc., 118, 1547 (1971).
19. D. Larcher, S. Beattie, M. Morcrette, K. Edström, J. Jumas, and J. Tarascon,J. Mater. Chem., 17, 3759 (2007).
20. Y. Yamada and A. Yamada,J. Electrochem. Soc., 162, A2406 (2015).
21. K. Ueno, R. Tatara, S. Tsuzuki, S. Saito, H. Doi, K. Yoshida, T. Mandai, M. Matsugami, Y. Umebayashi, K. Dokko, and M. Watanabe,Phys. Chem. Chem. Phys., 17, 8248 (2015).
22. J. M. Alı´a and H. G. M. Edwards,Vib. Spectrosc., 24, 185 (2000). 23. X. Xuan, J. Wang, Y. Zhao, and J. Zhu,J. Raman Spectrosc., 38, 865 (2007). 24. R. Tatara, D. G. Kwabi, T. P. Batcho, M. Tulodziecki, K. Watanabe, H. Kwon,
M. L. Thomas, K. Ueno, C. V. Thompson, K. Dokko, Y. Shao-Horn, and M. Watanabe,
J. Phys. Chem. C, 121, 9162 (2017).
25. K. Yoshida, M. Nakamura, Y. Kazue, N. Tachikawa, S. Tsuzuki, S. Seki, K. Dokko, and M. Watanabe,J. Am. Chem. Soc., 133, 13121 (2011).
26. S. Terada, T. Mandai, R. Nozawa, K. Yoshida, K. Ueno, S. Tsuzuki, K. Dokko, and M. Watanabe,Phys Chem Chem Phys, 16, 11737 (2014).
27. T. Tamura, K. Yoshida, T. Hachida, M. Tsuchiya, M. Nakamura, Y. Kazue, N. Tachikawa, K. Dokko, and M. Watanabe,Chem. Lett., 39, 753 (2010). 28. T. Mandai, K. Yoshida, K. Ueno, K. Dokko, and M. Watanabe,Phys. Chem. Chem.
Phys., 16, 8761 (2014).
29. J. Wang, Y. Yamada, K. Sodeyama, C. Chiang, Y. Tateyama, and A. Yamada, Nat.
Commun., 7 (2016)http://www.nature.com/articles/ncomms12032.
30. N. Tachikawa, K. Yamauchi, E. Takashima, J. W. Park, K. Dokko, and M. Watanabe,
Chem. Commun., 47, 8157 (2011).
31. K. Dokko, N. Tachikawa, K. Yamauchi, M. Tsuchiya, A. Yamazaki, E. Takashima, J. W. Park, K. Ueno, S. Seki, N. Serizawa, and M. Watanabe,J. Electrochem. Soc.,
160, A1304 (2013).
32. Y. Yamada, C. H. Chiang, K. Sodeyama, J. Wang, Y. Tateyama, and A. Yamada,
ChemElectroChem, 2, 1687 (2015).
33. C. Zhang, A. Yamazaki, J. Murai, J. W. Park, T. Mandai, K. Ueno, K. Dokko, and M. Watanabe,J. Phys. Chem. C, 118, 17362 (2014).
34. L. Suo, O. Borodin, T. Gao, M. Olguin, J. Ho, X. Fan, C. Luo, C. Wang, and K. Xu,
Science, 350, 938 (2015).
35. T. Doi, Y. Shimizu, M. Hashinokuchi, and M. Inaba,J. Electrochem. Soc., 163, A2211 (2016).
36. Y. Yamada, K. Furukawa, K. Sodeyama, K. Kikuchi, M. Yaegashi, Y. Tateyama, and A. Yamada,J. Am. Chem. Soc., 136, 5039 (2014).
38. J. Qian, W. A. Henderson, W. Xu, P. Bhattacharya, M. Engelhard, O. Borodin, and J. Zhang,Nat. Commun., 6 (2015).
39. Y. Yamada, M. Yaegashi, T. Abe, and A. Yamada,Chem. Commun., 49, 11194 (2013). 40. S. Terada, K. Ikeda, K. Ueno, K. Dokko, and M. Watanabe,Aust. J. Chem., 72, 70
(2019).
41. R. Miao, J. Yang, X. Feng, H. Jia, J. Wang, and Y. Nuli,J. Power Sources, 271, 291 (2014).
42. G. G. Eshetu, S. Grugeon, G. Gachot, D. Mathiron, M. Armand, and S. Laruelle,
Electrochimica Acta, 102, 133 (2013).
43. M. Wang, L. Huai, G. Hu, S. Yang, F. Ren, S. Wang, Z. Zhang, Z. Chen, Z. Peng, C. Shen, and D. Wang,J. Phys. Chem. C, 122, 9825 (2018).
44. R. Tatara, P. Karayaylali, Y. Yu, Y. Zhang, L. Giordano, F. Maglia, R. Jung, J. P. Schmidt, I. Lund, and Y. Shao-Horn,J. Electrochem. Soc., 166, A5090 (2019). 45. M. Ender, J. Illig, and E. Ivers-Tiffée,J. Electrochem. Soc., 164, A71 (2017). 46. J. Costard, M. Ender, M. Weiss, and E. Ivers-Tiffée,J. Electrochem. Soc., 164, A80
(2017).
47. T. Osaka, D. Mukoyama, and H. Nara,J. Electrochem. Soc., 162, A2529 (2015). 48. Y. Hoshi, Y. Narita, K. Honda, T. Ohtaki, I. Shitanda, and M. Itagaki,J. Power Sources,
288, 168 (2015).
49. M. J. Frisch, G. W. Trucks, H. B. Schlegel, G. E. Scuseria, M. A. Robb, J. R. Cheeseman, G. Scalmani, V. Barone, G. A. Petersson, H. Nakatsuji, X. Li, M. Caricato, A. Marenich, J. Bloino, B. G. Janesko, R. Gomperts, B. Mennucci, H. P. Hratchian, J. V. Ortiz, A. F. Izmaylov, J. L. Sonnenberg, D. Williams-Young, F. Ding, F. Lipparini, F. Egidi, J. Goings, B. Peng, A. Petrone, T. Henderson, D. Ranasinghe, V. G. Zakrzewski, J. Gao, N. Rega, G. Zheng, W. Liang, M. Hada, M. Ehara, K. Toyota, R. Fukuda, J. Hasegawa, M. Ishida, T. Nakajima, Y. Honda, O. Kitao, H. Nakai, T. Vreven, K. Throssell, J. A. Montgomery, J. E. Peralta, F. Ogliaro, M. Bearpark, J. J. Heyd, E. Brothers, K. N. Kudin, V. N. Staroverov, T. Keith, R. Kobayashi, J. Normand, K. Raghavachari, A. Rendell, J. C. Burant, S. S. Iyengar, J. Tomasi, M. Cossi, J. M. Millam, M. Klene, C. Adamo, R. Cammi, J. W. Ochterski, R. L. Martin, K. Morokuma, O. Farkas, J. B. Foresman, and D. J. Fox,
Gaussian 09 Revision A.02, Gaussian Inc., Wallingford CT, (2016).
50. A. D. Becke,J. Chem. Phys., 98, 5648 (1993).
51. C. Lee, W. Yang, and R. G. Parr,Phys. Rev. B, 37, 785 (1988).
52. L. Xing, O. Borodin, G. D. Smith, and W. Li,J. Phys. Chem. A, 115, 13896 (2011). 53. S. Trasatti,Pure Appl. Chem., 58, 955 (1986).
54. R. Tatara, G. M. Leverick, S. Feng, S. Wan, S. Terada, K. Dokko, M. Watanabe, and Y. Shao-Horn,J. Phys. Chem. C, 122, 18316 (2018).
55. K. Ueno, J. Murai, K. Ikeda, S. Tsuzuki, M. Tsuchiya, R. Tatara, T. Mandai, Y. Umebayashi, K. Dokko, and M. Watanabe,J. Phys. Chem. C, 120, 15792 (2016). 56. K. Fujii, H. Wakamatsu, Y. Todorov, N. Yoshimoto, and M. Morita,J. Phys. Chem.
C, 120, 17196 (2016).
57. D. Brouillette, G. Perron, and J. E. Desnoyers,J. Solut. Chem., 27, 32 (1998). 58. K. Dokko, D. Watanabe, Y. Ugata, M. L. Thomas, S. Tsuzuki, W. Shinoda,
K. Hashimoto, K. Ueno, Y. Umebayashi, and M. Watanabe,J. Phys. Chem. B, 122, 10736 (2018).
59. K. Yoshida, M. Tsuchiya, N. Tachikawa, K. Dokko, and M. Watanabe,J. Phys. Chem. C, 115, 18384 (2011).
60. D. M. Seo, O. Borodin, S. Han, P. D. Boyle, and W. A. Henderson,J. Electrochem. Soc., 159, A1489 (2012).
61. S. Han, O. Borodin, D. M. Seo, Z. Zhou, and W. A. Henderson,J. Electrochem. Soc.,
161, A2042 (2014).
62. J. E. Katon and M. D. Cohen,Can. J. Chem., 53, 1378 (1975).
63. K. Kimura, J. Motomatsu, and Y. Tominaga,J. Phys. Chem. C, 120, 12385 (2016). 64. D. W. McOwen, D. M. Seo, O. Borodin, J. Vatamanu, P. D. Boyle, and
W. A. Henderson,Energy Env. Sci, 7, 416 (2014).
65. H. Moon, T. Mandai, R. Tatara, K. Ueno, A. Yamazaki, K. Yoshida, S. Seki, K. Dokko, and M. Watanabe,J. Phys. Chem. C, 119, 3957 (2015).
66. Y. Umebayashi, T. Mitsugi, S. Fukuda, T. Fujimori, K. Fujii, R. Kanzaki, M. Takeuchi, and S. Ishiguro,J. Phys. Chem. B, 111, 13028 (2007).
67. K. Ueno, K. Yoshida, M. Tsuchiya, N. Tachikawa, K. Dokko, and M. Watanabe,J. Phys. Chem. B, 116, 11323 (2012).
68. A. V. Cresce, S. M. Russell, O. Borodin, J. A. Allen, M. A. Schroeder, M. Dai, J. Peng, M. P. Gobet, S. G. Greenbaum, R. E. Rogers, and K. Xu,Phys. Chem. Chem. Phys.,
19, 574 (2017).
69. L. Zhang, L. Chen, H. Luo, X. Zhou, and Z. Liu,Adv. Energy Mater., 7, 1700034 (2017).
70. F. Zhang, B. Ji, X. Tong, M. Sheng, X. Zhang, C. Lee, and Y. Tang,Adv. Mater. Interfaces, 3 (2016)http://doi.wiley.com/10.1002/admi.201670112.
71. X. Tong, F. Zhang, B. Ji, M. Sheng, and Y. Tang,Adv. Mater., 28, 9979 (2016). 72. M. H. Tahmasebi, D. Kramer, R. Mönig, and S. T. Boles,J. Electrochem. Soc., 166,
A5001 (2019).
73. Y. Wen, Y. Zhu, A. Langrock, A. Manivannan, S. H. Ehrman, and C. Wang,Small, 9, 2810 (2013).
74. D. Linden and T. B. Reddy, Eds., Handbook of batteries, 3rd ed., p. 1, McGraw-Hill, New York, (2002).
75. N. S. Hudak and D. L. Huber,J. Electrochem. Soc., 159, A688 (2012).
76. S. Verdier, L. El Ouatani, R. Dedryvère, F. Bonhomme, P. Biensan, and D. Gonbeau,
J. Electrochem. Soc., 154, A1088 (2007).
77. B. Philippe, R. Dedryvère, M. Gorgoi, H. Rensmo, D. Gonbeau, and K. Edström,J. Am. Chem. Soc., 135, 9829 (2013).
78. R. Younesi and F. Bardé,Sci. Rep., 7 (2017).
79. J. Dupin, D. Gonbeau, P. Vinatier, and A. Levasseur,Phys. Chem. Chem. Phys., 2, 1319 (2000).
80. S. Malmgren, K. Ciosek, M. Hahlin, T. Gustafsson, M. Gorgoi, H. Rensmo, and K. Edström,Electrochim. Acta, 97, 23 (2013).
81. A. T. Appapillai, A. N. Mansour, J. Cho, and Y. Shao-Horn,Chem. Mater., 19, 5748 (2007).
82. J. C. Dupin, D. Gonbeau, H. Benqlilou-Moudden, P. Vinatier, and A. Levasseur,Thin Solid Films, 384, 23 (2001).
83. K. Edström, T. Gustafsson, and J. O. Thomas,Electrochim. Acta, 50, 397 (2004). 84. C. Kozlowski and P. M. A. Sherwood,J. Chem. Soc. Faraday Trans. 1 Phys. Chem.
Condens. Phases, 81, 2745 (1985)http://xlink.rsc.org/?DOI=f19858102745. 85. J. Matthew,Surf. Interface Anal., 36, 1647 (2004).
86. H. Ago, T. Kugler, F. Cacialli, W. R. Salaneck, M. S. P. Shaffer, A. H. Windle, and R. H. Friend,J. Phys. Chem. B, 103, 8116 (1999).
87. J. Tomasi, B. Mennucci, and R. Cammi,Chem. Rev., 105, 2999 (2005). 88. G. Kresse and J. Furthmüller,Phys. Rev. B, 54, 11169 (1996).
89. S. Klingner, F. Voigts, W. Viöl, and W. Maus-Friedrichs,Surf. Eng., 29, 396 (2013). 90. S. Oswald, F. Thoss, M. Zier, M. Hoffmann, T. Jaumann, M. Herklotz, K. Nikolowski, F. Scheiba, M. Kohl, L. Giebeler, D. Mikhailova, and H. Ehrenberg,Batteries, 4, 36 (2018).
91. M. Gauthier, P. Karayaylali, L. Giordano, S. Feng, S. F. Lux, F. Maglia, P. Lamp, and Y. Shao-Horn,J. Electrochem. Soc., 165, A1377 (2018).
92. S. Leroy, F. Blanchard, R. Dedryvère, H. Martinez, B. Carré, D. Lemordant, and D. Gonbeau,Surf. Interface Anal., 37, 773 (2005).
93. X. Chen, X. Shen, B. Li, H. Peng, X. Cheng, B. Li, X. Zhang, J. Huang, and Q. Zhang,
Angew. Chem. Int. Ed., 57, 734 (2018).
94. B. Bellina, I. Compagnon, L. MacAleese, F. Chirot, J. Lemoine, P. Maître, M. Broyer, R. Antoine, A. Kulesza, R. Mitri´c, V. Bonaˇci´c-Koutecký, and P. Dugourd,Phys. Chem. Chem. Phys., 14, 11433 (2012).
95. M. Yoosefian and N. Etminan,RSC Adv., 6, 64818 (2016).
96. Y. Yamada, K. Usui, C. Chiang, K. Kikuchi, K. Furukawa, and A. Yamada,ACS Appl. Mater. Interfaces, 6, 10892 (2014).
97. S. Terada, H. Susa, S. Tsuzuki, T. Mandai, K. Ueno, K. Dokko, and M. Watanabe,J. Phys. Chem. C, 122, 16589 (2018).
