Tetrahydrothiophene 1-oxide as highly effective co-solvent for propylene carbonate-based electrolytes

Journal of Power Sources 437 (2019) 226881

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Tetrahydrothiophene 1-oxide as highly effective co-solvent for propylene carbonate-based electrolytes Kristina Oldiges a, b, Julian Michalowsky c, Mariano Grünebaum a, Natascha von Aspern a, b, Isidora Cekic-Laskovic a, Jens Smiatek a, c, Martin Winter a, b, d, Gunther Brunklaus a, b, * a

Helmholtz Institute Münster, IEK-12, Forschungszentrum Jülich GmbH, Corrensstrasse 46, 48149, Münster, Germany Institute of Physical Chemistry, University of Münster, Corrensstrasse 28/30, 48149, Münster, Germany Institute for Computational Physics, University of Stuttgart, Allmandring 3, 70569, Stuttgart, Germany d MEET Battery Research Center, Corrensstrasse 46, 48149, Münster, Germany b c

H I G H L I G H T S

G R A P H I C A L A B S T R A C T

� Ionic conductivity of PC-based electro­ lytes is increased by 13%. � Correlation between Liþ ion speciation and cycling behavior. � Effective SEI and CEI formation on graphite and NCM111 electrodes. � Blends with 10–40 mol-% THT1oxide enabled stable cycling in lithium-ion cells. � Excellent long-term and low tempera­ ture cycling behavior with 15 mol-% THT1oxide.

A R T I C L E I N F O

A B S T R A C T

Keywords: Lithium-ion batteries Non-aqueous electrolytes Co-solvents Propylene carbonate Ion transport Solid electrolyte interphase

Propylene carbonate (PC) together with cyclic sulfur compounds such as tetrahydrothiophene 1-oxide (THT1oxide) as co-solvent and lithium hexafluorophosphate (LiPF6) as conducting salt are introduced as new aprotic liquid electrolytes for lithium-ion batteries. Starting with the single solvent electrolyte LiPF6 in PC, by addition of THT1oxide, the ion transport properties even at temperatures down to 20 � C are improved by the different solvation behavior of Liþ ions due to the high Liþ ion affinity of the sulfinyl (-S¼O) group and by the resulting decrease of the Liþ ion complex size. Electrolytes that contain Liþ ion complexes with both PC and THT1oxide molecules in the solvation shell are able to form protective interphase layers on graphite and NCM111 (LiNi1/3Co1/3Mn1/3O2) electrodes that are both permeable for Liþ ions while ensuring good electronic insulation, thus enabling stable cycling in lithium-ion cells with only minor capacity fading. THT1oxide/PCbased electrolytes afford better long-term as well as low temperature cycling behavior compared to estab­ lished state-of-the-art (SOTA) organic carbonate-based electrolytes. The obtained results allow for the design of new co-solvents for PC and comparable cyclic organic carbonates, and provide a non-toxic and cheap alternative to crown ethers without affecting the Liþ ion transference/transport numbers.

* Corresponding author. Helmholtz Institute Münster, IEK-12, Forschungszentrum Jülich GmbH, Corrensstrasse 46, 48149, Münster, Germany. E-mail address: [email protected] (G. Brunklaus). https://doi.org/10.1016/j.jpowsour.2019.226881 Received 7 May 2019; Received in revised form 24 June 2019; Accepted 11 July 2019 Available online 23 July 2019 0378-7753/© 2019 Elsevier B.V. All rights reserved.

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Journal of Power Sources 437 (2019) 226881

1. Introduction

In contrast, the approach of encasing Liþ ions instead of anions will be an effective strategy, if appropriate co-solvents are applied that are commercially available or readily producible and non-toxic, in order to enable large-scale applications in commercial LIBs. The compound should possess a high Liþ ion affinity, i.e. a distinctly higher donor number (DN) [24,59] than PC (DN ¼ 15.1 kcal mol 1) [60], while maintaining or increasing the actual Liþ ion transference and transport numbers, as well as a lower reduction potential than PC (<0.8 V vs. Li/Liþ) or comparable organic cyclic carbonates. Notably, molecules with sulfinyl (-S¼O) groups possess high donor numbers, e.g., dibutyl sulfoxide (DN ¼ 31 kcal mol 1) and dimethyl sulfoxide (DN ¼ 29.8 kcal mol 1) [60], and prominent Liþ ion affinities. Cyclic organic sulfites, such as ethylene sulfite (ES) [39,61–64], propylene sulfite (PS) [38] and vinyl ethylene sulfite (VES) [61,65], have already been successfully applied as reduction-type electrolyte additives, pos­ sessing a higher reduction potential than PC and therefore suppressing PC co-intercalation due to reduction prior to PC and subsequent SEI for­ mation. Previous studies [66] showed that the cyclic sulfur compound tetrahydrothiophene 1-oxide (THT1oxide) is a particularly promising co-solvent for PC-based electrolytes affording excellent physicochemical properties compared to other commercially available five-membered cyclic sulfur compounds previously utilized in lithium-ion cells, such as ES, tetramethylene sulfone (sulfolane or TMS) [67,68], 1,3,2-dioxathio­ lane 2,2-dioxide (DTD or ethylene sulfate) [69,70] and 1,3-propane sultone (1,3-PS) [27,71,72]. THT1oxide is commercially available and readily synthesizable [73,74], classified as non-hazardous substance [75] and possesses a high flash point (ϑf ¼ 112 � C) [75] and boiling point (ϑb ¼ 235–237 � C) [75], which are comparable to PC (ϑf ¼ 116 � C [76] and ϑb ¼ 242.0 � C [26]). THT1oxide has the potential to be an affordable and non-toxic alternative to crown ethers and to direct future research to cation rather than anion receptors. In this work, the compound was applied as single and co-solvent with PC, considering its ionic mobility, cell performance and interfacial electrochemistry, respectively, also including the formation, dynamics and stability of protective layers on the electrodes. Note that LiPF6 was used as conducting salt to avoid aluminum dissolution, which is a drawback of other conducting salts like lithium bis(trifluoromethylsulfonyl)imide (LiTFSI) [77–81].

The increasing demand for electric vehicles (EVs) and hybrid electric vehicles (HEVs) [1] requires high power and energy density lithium-ion batteries (LIBs) [2,3] that are able to operate in a wide temperature and potential range. Despite the hopes into future positive development of ceramic solid-state [4] and polymeric solid state batteries [5–8] that may enable rechargeable lithium metal battery (LMB) [9,10] anodes [11,12], liquid electrolytes still dominate the market and afford a lot of future opportunities [13–17], and many open research questions are still not answered [18,19]. State-of-the-art (SOTA) non-aqueous aprotic electrolyte solvents generally comprise of ethylene carbonate (EC), which suffers from a high melting point (ϑm ¼ 36.4 � C) [20], and linear organic carbonates, such as dimethyl carbonate (DMC), ethyl methyl carbonate (EMC) or diethyl carbonate (DEC), which are highly flam­ mable and volatile [21–23]. For the purpose of realizing a substantial improvement in terms of safety, ion mobility and cell performance, new electrolyte components have to be characterized and new electrolyte formulations, especially in the presence of propylene carbonate (PC) or other organic cyclic carbonates with comparable characteristics, have to be developed [17]. A trial and error search approach is limited in providing significantly new insight, rendering computational chemistry relevant to establish new strategies for the design of new co-solvents for cyclic organic carbonates such as PC [24]. With regard to their broad applicability, methods like molecular dynamics (MD) simulations and quantum-chemical approaches like density functional theory (DFT) are indispensable for the analysis of ion complexes and the carrier medium, affording better understanding of transport mechanisms and structural properties of the electrolyte solution. PC has been used as solvent for lithium metal batteries from the beginning [25]. Although very desir­ able due to its low cost, wide liquid temperature range (△T ¼ -49 242.0 � C) [26] and good solvation behavior, propylene carbonate cannot be applied as single solvent in LIBs, as it decomposes on graphite anodes at ~0.8 V vs. Li/Liþ under massive graphite expansion due to solvent co-intercalation exfoliation and propylene gas evolution [27–31] due to insufficient formation of the solid electrolyte interphase (SEI) [32–34]. Furthermore, PC has a more than twice as high viscosity (2.5 mPa s at 25 � C [35]) as an EC/DMC (1:1 by wt) blend (1.2 mPa s at 25 � C [36]), impeding ion transport; a low viscosity is important for achieving high battery cell rate and power. Much research on SEI forming additives has been performed in the last years in order to make PC-based electrolytes applicable [37–43]. Functional additives, how­ ever, do not have a crucial impact on the viscosities and overall ionic conductivities of the PC-based electrolytes, as they are merely added in small amounts (� 5 wt-% or vol-%). The only moderate viscosity of PC hence still reduces the ion transport, which has a negative impact on the cycling behavior and battery life time. Crown ethers, such as 12-crown-4 and 15-crown-5, improve the solubility of the conducting lithium salts and therefore the ionic con­ ductivities of the resulting electrolytes [44–48]. Furthermore, they are able to suppress PC reduction at the same time due to their strong sol­ vating ability to Liþ ions and the consequent weakened interactions between Liþ ions and PC molecules, so that co-intercalation of PC into graphite is hindered [49,50]. However, the strong coordination of Liþ ions by crown ethers leads to a decrease of the Liþ ion transference numbers [14]. Apart from this, the main barrier for the application of crown ethers in LIB electrolytes is their costs as well as their high toxicity and the related issues concerning processing and disposal thereof. For that reason, the focus was shifted to anion receptors, such as aza-ethers [51], borates [52–54], boranes [52,53,55] and boronates [56], which are able to improve the achievable ionic conductivity and Liþ ion transference number. Aza-ethers, however, show limited solubility in polar solvents and are unstable with lithium hexafluorophosphate (LiPF6), the typically used electrolyte salt [57,58], whereas the costs and toxicology of boron-based anion receptors may constitute a serious barrier for commercial LIB applications.

2. Experimental 2.1. Materials Tetrahydrothiophene 1-oxide (Alfa Aesar, 97.0%) was pre-dried over molecular sieve (water content < 80 ppm), whereas propylene carbon­ ate (BASF, battery grade) and LiPF6 (BASF, battery grade) were used as received. The electrolytes with 1 M LiPF6, x mol-% THT1oxide and (100-x) mol-% PC (x ¼ 0, 5, 10, 15, 20, 30, 40, 50, 70, 100) were pre­ pared in an argon-filled glove box (H2O and O2 contents < 0.1 ppm, MBRAUN). 1 M LiPF6 in 50 wt-% EC and 50 wt-% DEC (purchased as LP40, BASF, battery grade) served as reference electrolyte. Graphite (High Power, 1.1 mAh cm 2), NCM111 (LiNi1/3Co1/3Mn1/3O2, High Power, 1.0 mAh cm 2) and LMO (LiMn2O4, 1.0 mAh cm 2) electrode sheets were purchased from Custom Cells Itzehoe GmbH, Germany. 2.2. Permittivity measurements Permittivity measurements were carried out with an ALPHA DCM470 permittivity meter by Zadow Electronics. The device was located in a dry room with a water content below 30 ppm and a temperature of approx. 20 � C. The measurement is based on the determination of a phase angle shift from the current response of an applied AC voltage with a frequency of 470 kHz. 2.3. Conductivity measurements Ionic conductivities were measured with a MCS 10 impedance-based 2

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Journal of Power Sources 437 (2019) 226881

conductivity meter (BioLogic) with an integrated frequency response analyzer and a temperature control unit. The device was calibrated before every measurement with a KCl standard solution at 25 � C (cell constant: C ¼ 1.059 cm 1). The measurement was then performed be­ tween 20 � C and 50 � C in 5 � C steps (△ϑ �0.01 � C).

circuit potential to 8 V vs. Li/Liþ with a constant scan rate of 20 μV s 1. The resulting current density between LMO and the lithium metal counter electrode was measured.

2.4. Viscosity measurements

Graphite/NCM1111 cells were cycled in the voltage range 3–4.2 V at 20 � C. For this purpose, coin cells of the type CR2032 containing Sepa­ rion® as separator und 100 μl electrolyte were assembled in a dry room. The graphite (Ø 13 mm) and NCM111 (Ø 12 mm) electrodes were dried under reduced pressure (~10 3 mbar) in a BÜCHI glass oven at 120 � C for 24 h before use. After three formation cycles at 0.2C, the cells were cycled at 1C for charge/discharge 100 cycles. For the electrochemical impedance measurements, two single-layer pouch cells (5 � 5 cm) were assembled with the same materials and cycled in the voltage range from 3 to 4.2 V at 20 � C. 3 cycles at 0.2C were performed with the first cell and 3 cycles at 0.2C and subsequent 100 cycles at 1C with the second cell, considering that both cells had 50% state of charge (SOC) at the end of the procedure.

2.9. Galvanostatic cycling

Viscosities were measured with an Anton Paar MCR 301 rheometer in a dry room (water content below 30 ppm). The device was equipped with a temperature system CTD 450 and a CP50–0.5/TG measuring system with a diameter of 49.947 mm, a cone angle of 0.473� and a distance of 68 μm between cone and lower plate. Viscosity measure­ ments were performed in a temperature range from 20 � C to 50 � C in steps of 10 � C; viscosities were also measured at 25 � C. The shear rates were increased with increasing temperature from 2000 s 1 to 9000 s 1. 2.5. PFG NMR spectroscopy PFG NMR measurements were performed with stimulated echo se­ quences at a Bruker AVANCE III 200 spectrometer using a Bruker Diff50 probe at 25.0 � C (stabilized with �0.1 � C), equipped with a Bruker 7 Li/1H and 19F coil (5 mm). In each case, the gradient strengths were varied from 5 to 1800 G/cm, with gradient pulse length δ of 1 ms and diffusion time Δ of 40 ms, recording 16 scans at relaxation delays of up to 60 s. The self-diffusion coefficients of all electrolyte species were determined by fitting peak intensities as a function of the gradient ac­ cording to equation (1): I ¼ I0 ⋅e

D⋅γ2 ⋅g2 ⋅δ2 ⋅ðΔ

δ 3Þ

2.10. Scanning electron microscopy analysis Scanning electron microscopy (SEM) analysis was performed with a Carl Zeiss AURIGA SEM microscope (Carl Zeiss Microscopy GmbH) to examine the surface morphology of the graphite and NCM111 working electrodes in the pristine state and after cycling. The electrodes were removed from the cell in a dry room with a water content below 30 ppm and were washed three times with 500 μl DMC (BASF, battery grade).

(1)

2.11. Electrochemical impedance spectroscopy

I, I0, D and g are the observed intensity of the NMR signal, its initial intensity, the diffusion coefficient, and the applied gradient strength. The corresponding gyromagnetic ratio γ was set to 1.655 ⋅103 Hz/G and 4.006 ⋅103 Hz/G for 7Li and 19F. Data analysis was performed using the Dynamics Center (Bruker Topspin3 software).

Electrochemical impedance spectroscopy measurements were per­ formed with a VMP3 (BioLogic Science Instruments). Symmetrical and asymmetrical coin cells with graphite and NCM111 electrodes, which were removed from cycled pouch cells with 50% SOC, were assembled and measured in a frequency range between 100 kHz and 10 mHz. The impedances were obtained by fitting the curves in the Nyquist plots using ZView 3.2.

2.6. Raman measurements Raman measurements were performed with a Bruker VERTEX 70 spectrometer equipped with a RAM II Raman module and Nd:YAG laser with a wavelength of 1064 nm and power output of 300 mW. 1000 scans were acquired at a resolution of 2 cm 1, all spectra were stored in the range from 0 cm 1 to 4000 cm 1. Bruker OPUS software was applied for the Raman measurements whereas the Origin 2016 program package was used for data analysis or peak deconvolution.

2.12. X-ray photoelectron spectroscopy analysis The cycled electrodes were transferred into the XPS instrument (Axis Ultra DLD, Kratos, U.K.) without air and water contact and kept under reduced pressure (<10 9 mbar) for 12h. Al Kα radiation with an energy of 1486.3 eV, an angle of 0� (cathode) or 45� (anode) of emission, a pass energy of 120 eV at a 10 mA filament current and a 12 kV filament voltage source energy was used, and the charge neutralizer was switched on to composite the charging of the sample. The sputter depth profiling for the anodes was carried out by using a polyatomic ion gun (coronene as the ion source) with a sputter crater 10 times bigger than the mea­ surement area. The sputtering time was 60, 120 and 600s. For each sample, three (cathode) or two (anode) data points with lateral resolu­ tion of (700 � 300) mm were taken and arithmetically averaged. The XPS spectra were fitted by using the CasaXPS software (Version 2.3.16 PR 1.6, Casa Software Ltd., U.K.), where the C 1s C-H/C-C peak (284.5 eV) was chosen as internal standard for the binding energy (BE) calibration.

2.7. Cyclic voltammetry Cyclic voltammetry experiments were performed with a VMP3 (BioLogic Science Instruments). Three-electrode Swagelok® cells were used with graphite and NCM111 as working electrodes (Ø 12 mm) and lithium metal as counter (Ø 12 mm) and reference (Ø 5 mm) electrodes. The stability of graphite was measured in the potential range from opencircuit potential (OCP) to 0.005 V vs. Li/Liþ and of NCM111 in the range from OCP to 4.2 V vs. Li/Liþ applying a scan rate of 20 μV s 1 at 20 � C. The Whatman® glass microfiber separator (Grade GF/D) was wetted with 200 μl electrolyte.

2.13. Differential scanning calorimetry

2.8. Linear sweep voltammetry

Differential scanning calorimetry (DSC) measurements were per­ formed with a DSC Q2000 measuring device. The samples were pre­ pared in hermetic aluminum pans. After an equilibration at 25 � C and an isothermal step for two minutes, heat flow was measured three times in the temperature range from 150 � C to 120 � C. Helium was used as ambient gas.

Linear sweep voltammetry (LSV) measurements were performed with a VMP3 (BioLogic Science Instruments) at 20 � C to determine oxidative stabilities. Three-electrode Swagelok® cells were used, where the potential difference between the LMO working electrode (Ø 12 mm) and lithium metal reference electrode was increased from the open 3

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Journal of Power Sources 437 (2019) 226881

3. Results and discussion

that is the motion of complexes or aggregates with net charges, the contributions of Liþ and PF6 ions seem to determine the observed vis­ cosity and ionic conductivity behavior. The minimum viscosity and maximum ionic conductivity are mainly determined by PF6 anions and PC molecules, as they also show maximum self-diffusion coefficients of (15.3 � 0.2) 10 11 m2 s 1 and (19.5 � 0.2) 10 11 m2 s 1 (Fig. S2) at a molar fraction of 30 mol-% THT1oxide, respectively. The self-diffusion coefficients of Liþ ions, which are the most relevant species for ion transport in LIB electrolytes, are approximately constant up to a THT1oxide content of 40 mol-%, suggesting that the higher viscosity of the THT1oxide solvent has a negative impact on ion mobility only with higher THT1oxide contents. The addition of 10–40 mol-% THT1oxide leads to increased PF6 self-diffusion coefficients, rendering these nonaqueous aprotic electrolytes most suitable for application in LIBs and LMBs. In contrast to electrochemical impedance spectroscopy, PFG NMR spectroscopy considers the diffusion of all involved species, i.e. charged and neutral, including ion pairs and higher aggregates, from which “ideal” conductivities without the presence of ion correlations may be estimated via the Nernst-Einstein equation: P N ⋅e2 ⋅ ððzj Þ2 cj Dj Þ σ NMR ¼ A (3) kB ⋅T

3.1. Ionic mobility and interactions in PC/THT1oxide blends Propylene carbonate possesses a high relative permittivity (εr ¼ 66.2 at 20 � C, εr ¼ 64.9 at 25 � C [82]), enabling sufficient ion solvation and salt dissociation. However, PC suffers from moderate viscosity (2.5 mPa s at 25 � C [35]) which essentially determines the ionic conductivity and suppresses ion transport, in particular at low temper­ atures [15]. In general, facilitating high ionic mobility is desired in order to reduce polarization effects and therefore to enable reversible long-time cycling of battery cells. It is known that the addition of low viscosity solvents like DMC with a low relative permittivity (εr ¼ 3.20 at 25 � C [82]) fosters more effective ion transport in mixtures with high permittivity, high viscosity solvents [83,84]. This synergistic effect leads to a maximum of the ionic conductivity, where the optimal composition has to be identified dependent on the individual molecular species [24]. Linear organic carbonates are however flammable and may reduce the safety of the cell [21], going along with the lacking property of forming an effective SEI and enabling stable cycling with graphite, which re­ quires additional functional SEI additives. In view of this background, the cyclic compound THT1oxide seems not to be the ideal co-solvent for PC in order to expect a conductivity maximum, as the relative permittivity at 20 � C is 44 and the viscosity even higher than that of PC. However, a maximum ionic conductivity of (6.96 � 0.02) mS cm 1 and minimum viscosity of (6.4 � 0.5) mPa s with a fraction of 30 mol-% THT1oxide were observed at 25 � C (Fig. 1), 0.81 mS cm 1 higher and 0.85 mPa s lower compared to 1 M LiPF6 in PC electrolyte. The only moderate ionic conductivity of electrolytes based on PC as single solvent is therefore increased by 13%, rendering THT1oxide/PC-based electrolytes more attractive for application in LIBs and LMBs. In the absence of an external electric field, the solvated Liþ ions constitute the slowest diffusing species in the investigated electrolytes, leading to transport numbers tNMR Liþ between 0.32 � 0.02 and 0.35 � 0.02 for every electrolyte blend (Table 1), which were calculated from selfdiffusion coefficients measured by pulsed field gradient (PFG) NMR (Fig. S2 and Table 1) applying Equation (2): tNMR ¼P i

ci Di ððzj Þ2 cj Dj Þ

where NA, e and kB are the Avogadro constant, the elementary charge, and the Boltzmann constant, respectively. The obtained values are depicted in Table 1. Deviations from the “ideal” behavior can be quan­ tified by the degree of ion dissociation α [15], which is denoted as the ratio between the “real” impedance-based conductivity σImpedance and the “ideal” conductivity σNMR :

α¼

σImpedance σ NMR

(4)

Molar conductivities of the considered electrolytes are presented as a function of the amount of THT1oxide (see Fig. S2). The dissociation degrees for all considered electrolytes with different amounts of THT1oxide are ~0.8, indicating well-dissociated conducting salts in the single solvents PC and THT1oxide and in their blends, respectively. As it was recently discussed for multicomponent solutions [24,85], ion dissociation-association equilibria are crucially affected by the local composition of the solution around the ions, often referred to as “ion speciation”. With regard to the constant results for α, it can be assumed that PC and THT1oxide have a comparable distribution around the ions. Furthermore, it is known that higher relative permittivities lead to an increased ion dissociation [86]. Here, both solvents show relatively high relative permittivities (PC: εr ¼ 66.2, THT1oxide: εr ¼ 44 at 20 � C). The difference of 22.2 seems to have no impact on the dissociation degree. The amount of free charge carriers therefore does not change signifi­ cantly and hence cannot account for the increased ionic conductivities. The Liþ ion speciation and underlying molecular interactions of all involved species were determined to explain why self-diffusion co­ efficients and ionic conductivities attain maximum values and viscos­ ities minimum values upon addition of 30 mol-% THT1oxide. Experimental determination of the average coordination numbers from Raman spectra was not possible due to several peak overlaps. Also, MD simulations indicated that a large portion of THT1oxide molecules co­ ordinates to two or more Liþ ions (see Fig. S3), yielding variable sol­ vation shells of single cations, though the actual first shell coordination numbers for the THT1oxide/PC-based as well as single solvent-based electrolytes could be estimated from the obtained MD data. The extracted coordination numbers are depicted in Fig. 2a and Table S1. MD simulations showed that the Liþ ions are surrounded by 5–6 PC molecules when PC is used as a single solvent. Due to the small size of Liþ ions and the resulting high surface charge density, as well as the absence of polarization effects in classical MD force fields, the Liþ ion is difficult to model with a classical molecular dynamics approach.

(2)

where zj, ci;j and Di;j are the charge, the concentration and the selfdiffusion coefficient of the ionic species i and j, respectively. Since impedance measurements explicitly measure correlated charge transfer,

Fig. 1. Ionic conductivities and viscosities of electrolytes containing 1 M LiPF6 in x mol-% THT1oxide and (100-x) mol-% PC with x ¼ 0, 5, 10, 15, 20, 30, 40, 50, 70, 100 at 25 � C. The conductivity and viscosity values are accurate within �0.02 mS cm 1 and �0.5 mPa s, respectively. 4

K. Oldiges et al.

Journal of Power Sources 437 (2019) 226881

Table 1 Summary of self-diffusion coefficients and transport numbers of the ionic species, molar ionic conductivities obtained from impedance and PFG NMR measurements and the degree of ion dissociation for electrolytes containing 1 M LiPF6 in x mol-% THT1oxide and (100-x) mol-% PC with x ¼ 0, 5, 10, 15, 20, 30, 40, 50, 70, 100. The values are accurate within �0.2 10 11 m2 s 1 (D(Liþ) and D(PF6 )), �0.02 (t NMR ), �0.02 S cm2 mol 1 (σImpedance ), �0.04 S cm2 mol 1 (σNMR ) and �0.04 (∝), respectively. i Amount of THT1oxide [mol-%]

D(Liþ) [10

0 5 10 15 20 30 40 50 70 100

7.4 7.4 7.4 7.5 7.5 7.5 7.5 7.0 5.6 4.4

11

m2 s

1

]

D(PF6 ) [10

11

m2 s

1

]

13.9 13.9 14.6 15.0 15.3 15.3 14.0 13.4 10.8 9.2

tNMR Liþ

tNMR PF

σImpedance [S cm2 mol 1]

σNMR [S cm2 mol 1]

α

0.35 0.35 0.34 0.33 0.33 0.33 0.35 0.34 0.34 0.32

0.65 0.65 0.66 0.67 0.67 0.67 0.65 0.66 0.66 0.68

6.15 6.21 6.42 6.51 6.77 6.96 6.49 5.65 4.86 4.05

8.00 8.00 8.26 8.45 8.56 8.56 8.08 7.66 6.16 5.11

0.77 0.78 0.78 0.77 0.79 0.81 0.80 0.74 0.79 0.79

6

Fig. 2. a) Coordination numbers of PC and THT1oxide with Liþ and PF6 ions, as well as b) the sum of coordination numbers of Liþ and PF6 ions in electrolytes containing 1 M LiPF6 in x mol-% THT1oxide and (100-x) mol-% PC with x ¼ 0, 5, 10, 15, 20, 30, 40, 50, 60, 70, 80, 85, 90, 95, 100, derived from MD simulations. The inserted solid lines are guides to the eye.

Therefore, although the simulation results do indeed bear quantitative meaning, they should nonetheless be interpreted in a more qualitative and complementary manner [87,88]. Yet, the utilized OPLS/AA force field adequately depicts the qualitative behavior and relative trends, further verified by comparison to the present experimental results, revealing similar features with the same molar ratios of THT1oxide and PC. By adding even small amounts of THT1oxide to the solution, THT1oxide replaces PC in the almost six-membered Liþ ion solvation complex. The high affinity of the sulfinyl (-S¼O) group leads to the observation that THT1oxide stepwise replaces the other PC molecules at increasing amounts of THT1oxide in the electrolyte. With ~15 mol-% THT1oxide, the same amount of PC and THT1oxide molecules are located around the Liþ ion, whereas PC is almost completely replaced with 30 mol-% THT1oxide, yielding a complex of Li(PC)0.7(THT1ox­ ide)3.7 (see Table S1). Hence, the first solvation shell of Liþ ions with a molar fraction of 30 mol-% THT1oxide is composed of 84% THT1oxide and only 16% PC molecules (see Table S1). The DFT calculations showed that average Liþ-O distances in the geometry-optimized Li(PC)5 com­ plexes (0.208 nm) are higher than in Li(THT1oxide)4 complexes (0.195 nm), reflecting the high Liþ ion affinity of THT1oxide. Moreover, the corresponding free energies of geometry-optimized Li(PC)5 and Li (THT1oxide)4 clusters as shown in Fig. S1 (△G ¼ 52.094 keV and △G ¼ 68.799 keV) highlight a stronger binding free energy between the Liþ ions and the corresponding THT1oxide molecules when compared to coordinating PC molecules. Therefore, a transition from a five- or six- to a four-membered complex takes place, which is better observed when plotting the sum of Liþ and PF6 ion coordination

numbers, respectively (Fig. 2b). From a fraction of 30 mol-% THT1oxide on, lithium is coordinated to four THT1oxide molecules, indicating a smaller hydrodynamic radius of the complex, which is not changed with higher amounts of THT1oxide, whereas the PF6 ion complex constantly contains ~11 ligands around the anion (see Fig. 2b). This leads to the assumption that the decrease of ionic conductivities and self-diffusion coefficients after reaching the maximum and increase of the electro­ lyte viscosity may be attributed to a viscosity effect in the surrounding solvent instead of changes within the complexes. The coordination number is not the only factor that determines the diffusion behavior of the species. Rather, charge-transfer processes, local relaxations and polarization effects may contribute to the diffusion properties of the ions [89], which is also known for the diffusion of Naþ and Cl ions in aqueous solutions [90,91] and in case of ionic liquids [92]. Considering the experimental self-diffusion coefficients (see Fig. S2), MD data suggest that the first coordination shell of Liþ ions is largely saturated with THT1oxide at the same molar ratio at which Liþ ion diffusion is showing its decrease, i.e. at 30 mol-% THT1oxide. Therefore, higher amounts than 30 mol-% THT1oxide are not recommendable for applications in lithium-ion cells. Notably, MD data revealed that the portion of THT1oxide molecules that are not bound within a Liþ or PF6 ion-solvent complex and only part of the carrier medium rises quickly for THT1oxide molar fractions higher than 30 mol-%. This is in accor­ dance with the increasing self-diffusion coefficients of THT1oxide from that point on. PC has a maximum self-diffusion coefficient of 1.95 10 10 m2 s 1 with 30 mol-% THT1oxide, as almost every PC molecule is removed from the Liþ ion complex. The addition of THT1oxide to the 5

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Journal of Power Sources 437 (2019) 226881

PC-based electrolytes induced a larger number of unbound PC mole­ cules, which increase the average self-diffusion of the part of the solu­ tion that is not bound in a solvation complex, assuming the role of a carrier medium. Both the viscosity minimum and conductivity maximum can be ascribed to the different solvation behavior of Liþ ions and the resulting composition of the carrier medium, whereas the addition of THT1oxide has no influence on the solvation behavior of the PF6 ions. At molar fractions lower than 30 mol-% THT1oxide, the sulfur compound is almost completely coordinated within Liþ ion complexes and hardly influences the viscosity of the electrolyte, but only the size of the dissolved complexes. Therefore, a decrease of the Liþ ion complex size as well as an increase of solvent viscosity are competitive effects, determining the minimum electrolyte viscosity and thus achievable maximum ionic conductivity [15]. For reaching the conductivity maximum, the Liþ-THT1oxide complex has to be sufficiently small, but the amount of THT1oxide and its higher viscosity should not have a negative impact on the ionic conductivity of the resulting electrolyte.

the oxidation peak is shifted toward higher potentials with increasing cycle number. The observation that lithium metal cells with THT1oxide/ PC blends could be cycled despite the presence of THT1oxide clearly suggests that the electrolyte is able to form a functional cathode elec­ trolyte interphase (CEI [93–95]) on NCM111 electrodes. See Refs. [96, 97] with regard to the practical meaning of Coulombic efficiencies with nickel-based cathode materials. LSV data of 1 M LiPF6 in THT1oxide/PC electrolytes in LMO/Li cells (Fig. S6) revealed that already small amounts of the sulfur compound in a PC-based electrolyte led to a decrease of the oxidation potential from ~5.3 V vs. Li/Liþ for PC as single solvent to ~4.5 V vs. Li/Liþ. All up­ coming constant current cycling experiments in graphite/NCM111 cells were hence conducted with a cut-off voltage of 4.2 V. THT1oxide/PC blends led to reversible cycling in lithium metal cells, but also in graphite/NCM111 lithium-ion cells. The specific discharge capacities as a function of the cycle number are depicted in Fig. 4a. Due to the irreversible process of the graphite electrode in the electrolyte with THT1oxide as single solvent, the corresponding cycling data is not provided. THT1oxide as functional electrolyte additive (in gen­ eral < 5 wt-% or vol-%) was not sufficient to enable stable cycling in considered LIB cells. Also, the amount of 70 mol-% was too high. Lithium-ion cells with 50 mol-% THT1oxide, however, showed sub­ stantial capacity fading, while blends with 10–40 mol-% THT1oxide enabled effective SEI and CEI formation, affording specific discharge capacities of 95–110 mAh g 1 (with respect to the NMC111 cathode in balanced LIB full cells) and almost no capacity fading within charge/ discharge 100 cycles. Considering the results obtained from the ionic mobility investigations, the maximum molar fraction for attaining improved ion transport properties should be 30 mol-%. The 1 M LiPF6 in 15 mol-% THT1oxide and 85 mol-% PC electrolyte is highlighted as the most promising one, yielding the highest specific discharge capacity of ~110 mAh g 1 and almost no capacity fading. The Coulombic effi­ ciency was ~98.9% after three formation cycles and higher than 99.8% for the cycles 7–100. Relating these observations to possible complex formation in the electrolyte blends (Fig. 2 and Table S1), it is noticeable that the LIB cells apparently only run if the Liþ ion coordination com­ plexes are composed of two solvents. Indeed, PC has to be available in the complex for reduction and polymerization reactions, whereas THT1oxide is required for hindering PC to co-intercalate into graphite. Thus, electrolytes containing less than 10 mol-% and more than 40 mol% THT1oxide were not suitable for LIB cells. The best cell performance was attained with Li(PC)2.8(THT1oxide)2.5 complexes that comprise of almost the same amount of PC and THT1oxide. Established organic carbonate-based electrolytes, such as 1 M LiPF6 in 50 wt-% EC and 50 wt-% DEC (denoted as LP40), allowed for up to 10 mAh g 1 higher

3.2. Electrochemical behavior and cell performance of PC/THT1oxide blends PC-based electrolytes are often incompatible with graphite anodes, constituting the most important aspect that is required to be improved. As blends with THT1oxide offered excellent ionic mobility, the effect on the cycling performance and on the behavior with electrodes is of current interest. Cyclic voltammetry measurements confirmed that the 1 M LiPF6 in PC electrolyte gets reduced at ~0.8 V vs. Li/Liþ, and 1 M LiPF6 in THT1oxide also exhibited decomposition reactions starting at ~0.8 V vs. Li/Liþ (Fig. 3a). The irreversible process is clearly recog­ nizable in the following cycles (Fig. S4, left). Sharp reversible lithium de-/intercalation peaks were only observed in case of THT1oxide/PC blends (Fig. 3b). Here, an electrolyte containing 15 mol-% THT1oxide and 85 mol-% PC was used, since this blend showed improved ion transport behavior and the Liþ ion complexes contained almost the same amount of PC and THT1oxide molecules. In the first cycle, a small peak at ~0.8 V vs. Li/Liþ could be observed (see Fig. S4, right) indicating irreversible decomposition reactions and formation of a protective layer on the graphite electrode. THT1oxide/PC blends exhibited reversible cycling with graphite anodes, but also with transition metal electrodes such as NCM111 (see Fig. S5). PC as a single solvent and the blend with 15 mol-% THT1oxide afforded stable constant current cycling with NCM111 in the range from OCP to 4.2 V vs. Li/Liþ, whereas THT1oxide as a single solvent was not appropriate and led to reduced oxidation and reduction peaks, where

Fig. 3. Cyclic voltammograms of graphite/Li cells containing a) 1 M LiPF6 in PC, THT1oxide or 15 mol-% THT1oxide and 85 mol-% PC (1st cycle), and b) 1 M LiPF6 in 15 mol-% THT1oxide and 85 mol-% PC (1st, 3rd and 5th cycle), measured with a scan rate of 20 μV s 1 at 20 � C. 6

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Journal of Power Sources 437 (2019) 226881

Fig. 4. Constant current cycling profiles of graphite/NCM111 cells containing a) 1 M LiPF6 in x mol-% THT1oxide and (100-x) mol-% PC with x ¼ 5, 10, 15, 20, 30, 40, 50, 70 as well as b) 1 M LiPF6 in 15 mol-% THT1oxide and 85 mol-% PC and in 50 wt-% EC and 50 wt-% DEC, respectively, at 20 � C. After three formation cycles at 0.2C, the cells were cycled at a rate of 1C.

specific discharge capacities, possibly due to higher ionic mobilities and more effective film formation at 20 � C. Though cells containing the organic carbonate-based electrolyte showed higher specific discharge capacities in the first 750 charge/discharge cycles, cells with the elec­ trolyte containing 15 mol-% THT1oxide had a better long-term cycling behavior, yielding 80% capacity retention after 1111 cycles, whereas cells based on LP40 reached 80% after 871 cycles (Fig. 4b). 3.3. SEI and CEI formation and dynamics An effective SEI on the graphite and CEI on the NCM111 electrode was formed in case of electrolyte blends with 10–40 mol-% THT1oxide, so that LIB cells could be cycled without capacity fading within 100 charge/discharge cycles (see Fig. 4a). Both co-solvents, PC and THT1oxide, decomposed at potentials ~0.8 V vs. Li/Liþ (see Fig. 3a). DFT calculations yielded the energies of the lowest unoccupied molec­ ular orbitals (LUMO), amounting to 0.61 eV and 0.52 eV for PC and THT1oxide, respectively, indicating that THT1oxide may be reduced at marginally lower potentials than PC, so that PC in principle could be decomposed prior to THT1oxide, in this way forming a SEI. For the purpose of gaining insight into SEI and CEI formation, composition and stability, SEM and XPS analysis as well as impedance measurements were performed before constant current cycling, after three formation cycles at 0.2C and after formation and subsequent 100 cycles at 1C. Data obtained from impedance measurements (see the Nyquist plots in Fig. S9) of symmetric graphite/graphite cells corroborated that a SEI was formed. Fitting the first semicircle in the Nyquist plot with two resistances connected in series and one constant phase element (CPE) [98] parallel to the second resistance (see the equivalent circuit in Fig. S8), the interphase resistances were derived (Fig. 5). Due to overlaps, the other semicircles could not be unambiguously identified, thus an analysis of individual charge-transfer resistances would lead to an over-interpretation and was not performed. The data revealed that the interphase resistance on the graphite electrodes increases from (5.3 � 0.1) Ω after 3 charge/discharge cycles to (9.5 � 0.2) Ω after 103 cycles. The SEI is therefore slowly growing or modifying with increasing cycle number. During SEM analysis, a morphology change of the formed SEI was observed, which should have a significant impact on the interphase resistances. As seen in Fig. 6, the SEI was almost completely formed during the first three formation cycles, covering the surface of the electrode with a dense layer. Upon further cycling, the SEI encased the particles of the active material. In agreement with the results from the cyclic voltammetry experiments, SEI formation was

Fig. 5. Interphase impedances of symmetrical graphite/graphite (SEI, left) and NCM111/NCM111 (CEI, middle) cells as well as graphite/NCM111 cells (SEI þ CEI, right) with cycled electrodes and 1 M LiPF6 in 15 mol-% THT1oxide and 85 mol-% PC as electrolyte. Error bars are shown in the figure.

initiated by constant current cycling and not chemically, as after OCP no interphase resistance was visible in the Nyquist plots. Also, in the first cycle of the cyclic voltammograms of graphite/Li cells, a small peak at ~0.8 V vs. Li/Liþ was observed indicating electrochemically induced formation of the SEI. Symmetric NCM111/NCM111 cells showed that a distinctly thinner CEI could be formed electrochemically, which has a comparable interphase resistance and thus a comparable thickness and morphology after 3 cycles ((1.3 � 0.1) Ω) and 103 cycles ((1.5 � 0.3) Ω). These observations are in agreement with SEM micrographs of NCM111 electrodes, which do not reflect morphology changes (see Fig. S7). Similarly, the cyclic voltammograms of NCM111/Li cells in Fig. S5 indicated no further electrolyte decomposition in the considered potential range after formation. The impedance data from graphite/NCM111 cells highlighted the impact of both layers on the entire interphase resistance, changing from (8.3 � 0.6) Ω to (19.7 � 7.7)Ω, which clearly is attributable to the SEI formation. XPS analysis of both graphite and NCM111 electrodes was performed to determine the composition and thickness of the corresponding surface 7

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Journal of Power Sources 437 (2019) 226881

Fig. 6. SEM micrographs of graphite electrodes: a)-b) pristine, c)-d) after 3 formation cycles at 0.2C, e)-f) after 3 formation cycles at 0.2C þ 100 cycles at 1C. Constant current cycling was performed with graphite/NCM111 cells containing 1 M LiPF6 in 15 mol-% THT1oxide and 85 mol-% PC.

layers. These parameters were calculated according to the model developed by Niehoff et al. [99,100]. The obtained results are depicted in Fig. S10. The SEI featured a thickness of ~5 nm ((5.5 � 0.6) nm after 3 cycles, (4.9 � 0.1) nm after 103 cycles), whereas the CEI was notably thinner exhibiting a thickness of ~1 nm ((1.0 � 0.3) nm after 3 cycles, (1.4 � 0.1) nm after 103 cycles). For comparison: electrolytes based on EC and VC only led to solid electrolyte interphase thicknesses of up to 3.3 nm [101]. The little higher SEI thickness as well as the different morphology and relative compositions with THT1oxide/PC electrolytes could have a slightly negative impact on the specific discharge capacities compared to EC-based electrolytes [102], but obviously implicated higher stability and therefore improved long-term cycling (see Fig. 4). The surface layers are almost completely formed during the formation cycles, as no considerable change of surface thickness between 3 and 103 cycles could be observed. The slight increase of interphase re­ sistances therefore is ascribed to morphology and/or compositional changes [102]. The higher error margins of SEI and CEI thicknesses after 3 cycles, as shown in Fig. S10, suggested that the layer was more inhomogeneous than after 103 cycles. This conclusion could be drawn, as XPS measurements were performed on three different positions on the cathode and two positions on the anode. Organic components within all surface layers amounted to ~66 at-%, indicating good permeability for solvated Liþ ions [103]. Obtained XPS data revealed that the organic parts are composed of a polymer comprising ether groups, possibly PEO, indicating that a polymerization reaction took place on both electrodes. The SEI did not contain sulfur compounds, whereas the inorganic part of the CEI comprised metal sulfites and sulfates (see other components in Fig. S10). These salts have a positive impact on the lithium-ion cell, as they are known as electronic insulators that are able to effectively suppress the continued oxidation of solvent molecules. Thus, on both electrodes, THT1oxide/PC blends afforded surface layers that are permeable for Liþ ions and electronically insulating at the same time. Since sulfur containing compounds were not part of the SEI, it appears that THT1oxide constitutes an initiator for a polymerization reaction of PC. Tetrahydrothiophene as potential reaction product is highly volatile and could therefore not be detected on the electrodes. The SEI formation mechanism is different compared to electrolytes that contain the related additive sulfolane, which decomposes during the first cycle, thereby yielding a sulfide-containing SEI on graphite [67].

crystallization and melting peak was observed down to 150 � C and the glass transition temperature was even shifted toward lower tempera­ tures upon addition of 15 mol-% THT1oxide (from 96 � C to 100 � C), hence indicating better low temperature ionic conductivities. Indeed, higher ionic conductivities could be attained with THT1oxide/PC elec­ trolytes at lower temperatures (Fig. 7b). Upon inspection of the achievable ionic conductivities in the tem­ perature range from 20 to 50 � C, the THT1oxide/PC blend with 30 mol-% THT1oxide showed the highest conductivity values in the entire temperature range, e.g., an ionic conductivity of (11.54 � 0.02) mS cm 1 at 50 � C and of (1.48 � 0.02) mS cm 1 at 20 � C. With decreasing amount of THT1oxide, the ionic conductivities of the blends at high temperatures are comparable to each other. At low temperatures, the ionic conductivities are decreasing, so that a value of only (0.99 � 0.02) mS cm 1 at 20 � C for electrolytes based on PC as single solvent was observed. The curve in the Arrhenius plot is more distinctly decreasing, while the addition of THT1oxide leads to a more linear behavior. With 70 mol-% THT1oxide, the ionic conductivities at 20 � C are again comparable with the values for electrolytes with PC as single solvent. The viscosities show an opposite trend to the ionic conductivities in the temperature range from 20 to 50 � C (see Fig. S11). The 1 M LiPF6 in 50 wt-% EC and 50 wt-% DEC electrolyte shows ionic conductivities, which are comparable to THT1oxide/PC blends at high temperatures. A slightly higher ionic conductivity of the LP40 electrolyte at 20 � C could be the reason for the higher specific discharge capacities within the first 750 charge/discharge cycles. At 0 � C, the ionic conductivities amount to 3.71 mS cm 1, which is 8% higher than the values for 30 mol-% THT1oxide blends. However, long-term as well as low temperature cycling of THT1oxide/PC blends were better compared to common organic carbonate-based electrolytes, as it can be seen in Fig. 4b and Fig. S12. The higher ionic conductivities of LP40 at 0 � C, therefore, are not indicative of better cycling behavior. Instead, the 1 M LiPF6 in 15 mol-% THT1oxide and 85 mol-% PC electrolyte convinced by suffi­ ciently high specific discharge capacities at 0 � C, possibly due to more effective SEI formation at lower temperatures, rendering the investigated electrolytes highly attractive for low temperature applications. 4. Conclusions The applicability of propylene carbonate for LIBs containing carbo­ naceous electrodes was enabled by increasing the ionic conductivity of 1 M LiPF6 in PC by 13% on the one hand and by enabling effective SEI and CEI formation on graphite and NCM111, respectively, on the other hand. This is attributed to the different solvation behavior of Liþ ions and the resulting composition of the carrier medium, caused by the application of the cyclic sulfur compound tetrahydrothiophene 1-oxide as co-solvent, opening new possibilities for the design of PC-based electrolytes.

3.4. Low temperature behavior of PC/THT1oxide blends Propylene carbonate is considered as preferable non-aqueous aprotic electrolyte solvent due to its excellent stability at low temperatures. The compound is known to be liquid above 49 � C [26], still possessing reasonable ionic conductivities and viscosities. DSC measurements revealed that PC-based electrolytes with 15 mol-% THT1oxide have comparable stabilities in the low temperature range (Fig. 7a). No 8

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Journal of Power Sources 437 (2019) 226881

Fig. 7. a) Heat flows of 1 M LiPF6 in x mol-% THT1oxide and (100-x) mol-% PC with x ¼ 0, 15, 100 in the temperature range from 150 � C to 120 � C obtained from DSC measurements. b) Arrhenius plots of ionic conductivities of THT1oxide/PC blends (0, 15, 30, 70, 100 mol-% THT1oxide) with 1 M LiPF6 as conducting salt measured between 20 and 50 � C. The conductivities of 1 M LiPF6 in 50 wt-% EC and 50 wt-% DEC were added for comparison. The values are accurate within �0.02 mS cm 1. The dashed lines are guides to the eye.

THT1oxide showed a comparable behavior to crown ethers, since both compounds possess high Liþ ion affinity, replace PC in the Liþ ion solvation complexes and improve solubility of the lithium salt and ionic conductivity of the electrolyte. THT1oxide, however, is classified as nonhazardous and does not significantly affect the adequate Liþ ion trans­ ¼ 0.35 � 0.02). SEI port number of the 1 M LiPF6 in PC electrolyte (t NMR Liþ and CEI revealed better low temperature and long-term cycling compared to common EC-based electrolytes, where a fraction of 10–30 mol-% THT1oxide appeared most promising. The obtained results indicate that PC is decomposed in the first cycle and that the SEI for­ mation reaction is just initiated by THT1oxide, as sulfur containing compounds were not detected on the anode. PC is thus required for reduction and polymerization and THT1oxide might suppress PC decomposition via initiation of a polymerization reaction. The ratio of PC and the co-solvent THT1oxide has to be selected with the intention that Liþ ion complexes contain both THT1oxide and PC molecules in the solvation shell, in the best case the same amount of both, so that in case of solvent blends with 15 mol-% THT1oxide and 85 mol-% PC rather high specific discharge capacities of ~110 mAh g 1 as well as superior cycling behavior could be achieved. The highest ionic conductivities were observed instead with a molar fraction of 30 mol-% THT1oxide ((6.96 � 0.02) mS cm 1), where PC is almost completely replaced in the Liþ ion solvation complex. This work on THT1oxide introduces a possible strategy to design new co-solvents for PC and comparable cyclic organic carbonates such as 1,2-butylene carbonate. The molecules should exhibit distinctly higher donor numbers than PC (DN ≫ 15.1 kcal mol 1) for ensuring high Liþ ion affinity that hinders PC to co-intercalate with the solvated Liþ ions into graphite. Also, they should have a lower reduction poten­ tial than PC (<0.8 V vs. Li/Liþ) and be non-toxic as well as readily producible. The presence of sulfur atoms in the molecular structure can also have a positive impact on the electron insulation properties of the CEI. The higher viscosity of the co-solvent is of minor impact, since the overall viscosity of the electrolyte blend is only affected with higher amounts of the co-solvent. The approach of encasing Liþ ions instead of anions by design of proper electrolyte solvent composition is therefore highly promising and encouraging. THT1oxide/PC-based electrolytes have a potential to replace SOTA non-aqueous electrolytes based on highly volatile and flammable linear organic carbonates. Based on these findings,

computational studies could generate related molecules with compara­ ble and even better characteristics, and are currently in progress. Cooperative work between simulations and experiments is therefore highly recommended in the field of electrolyte research. Acknowledgements Special thanks are offered to Verena Naber and Debbie Berghus, who performed the DSC measurements. Johannes Thienenkamp is acknowl­ edged for assistance with NMR analysis. This work was supported by the state of NRW [Ministerium für Kultur und Wissenschaft des Landes Nord-rhein-West-fa-len, grant number 433]. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.226881. References [1] M.A. Hannan, M.M. Hoque, A. Hussain, Y. Yusof, P.J. Ker, State-of-the-Art and energy management system of lithium-ion batteries in electric vehicle applications: issues and recommendations, IEEE Access 6 (2018) 19362–19378. [2] J. Betz, G. Bieker, P. Meister, T. Placke, M. Winter, R. Schmuch, Theoretical versus practical energy: a plea for more transparency in the energy calculation of different rechargeable battery systems, Adv. Energy Mater. (2018) 1803170. [3] R. Wagner, N. Preschitschek, S. Passerini, J. Leker, M. Winter, Current research trends and prospects among the various materials and designs used in lithiumbased batteries, J. Appl. Electrochem. 43 (2013) 481–496. [4] F. Zheng, M. Kotobuki, S. Song, M.O. Lai, L. Lu, Review on solid electrolytes for all-solid-state lithium-ion batteries, J. Power Sources 389 (2018) 198–213. [5] K.S. Ngai, S. Ramesh, K. Ramesh, J.C. Juan, A review of polymer electrolytes: fundamental, approaches and applications, Ionics 22 (2016) 1259–1279. [6] S.B. Aziz, T.J. Woo, M.F.Z. Kadir, H.M. Ahmed, A conceptual review on polymer electrolytes and ion transport models, J. Sci.: Adv. Mater. Devices 3 (2018) 1–17. [7] L. Imholt, T. D€ orr, P. Zhang, L. Ibing, I. Cekic-Laskovic, M. Winter, G. Brunklaus, Grafted polyrotaxanes as highly conductive electrolytes for lithium metal batteries, J. Power Sources 409 (2019) 148–158. [8] K. Borzutzki, J. Thienenkamp, M. Diehl, M. Winter, G. Brunklaus, Fluorinated polysulfonamide based single ion conducting room temperature applicable geltype polymer electrolytes for lithium ion batteries, J. Mater. Chem. A 7 (2019) 188–201. [9] L. Imholt, D. Dong, D. Bedrov, I. Cekic-Laskovic, M. Winter, G. Brunklaus, Supramolecular self-assembly of methylated rotaxanes for solid polymer electrolyte application, ACS Macro Lett. 7 (2018) 881–885.

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