ACTAUNIVERSITATIS

UPSALIENSISUPPSALA

2017

Digital Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology 1525

Non-aqueous Electrolytes andInterfacial Chemistry in Lithium-ion Batteries

CHAO XU

ISSN 1651-6214ISBN 978-91-554-9931-0urn:nbn:se:uu:diva-319425

Dissertation presented at Uppsala University to be publicly examined in Room 2005,Ångström Laboratory, Lägerhyddsvägen 1, Uppsala, Wednesday, 14 June 2017 at 09:15 forthe degree of Doctor of Philosophy. The examination will be conducted in English. Facultyexaminer: Professor Brett Lucht (Department of Chemistry, University of Rhode Island).

AbstractXu, C. 2017. Non-aqueous Electrolytes and Interfacial Chemistry in Lithium-ion Batteries.Digital Comprehensive Summaries of Uppsala Dissertations from the Faculty of Science andTechnology 1525. 72 pp. Uppsala: Acta Universitatis Upsaliensis. ISBN 978-91-554-9931-0.

Lithium-ion battery (LIB) technology is currently the most promising candidate for powersources in applications such as portable electronics and electric vehicles. In today's state-of-the-art LIBs, non-aqueous electrolytes are the most widely used family of electrolytes. In thepresent thesis work, efforts are devoted to improve the conventional LiPF6-based electrolyteswith additives, as well as to develop alternative lithium 2-trifluoromethyl-4,5-dicyanoimidazole(LiTDI)-based electrolytes for silicon anodes. In addition, electrode/electrolyte interfacialchemistries in such battery systems are extensively investigated.

Two additives, LiTDI and fluoroethylene carbonate (FEC), are evaluated individually forconventional LiPF6-based electrolytes combined with various electrode materials. Introductionof each of the two additives leads to improved battery performance, although the underlyingmechanisms are rather different. The LiTDI additive is able to scavenge moisture in theelectrolyte, and as a result, enhance the chemical stability of LiPF6-based electrolytes even atextreme conditions such as storage under high moisture content and at elevated temperatures.In addition, it is demonstrated that LiTDI significantly influences the electrode/electrolyteinterfaces in NMC/Li and NMC/graphite cells. On the other hand, FEC promotes electrode/electrolyte interfacial stability via formation of a stable solid electrolyte interphase (SEI) layer,which consists of FEC-derivatives such as LiF and polycarbonates in particular.

Moreover, LiTDI-based electrolytes are developed as an alternative to LiPF6 electrolytesfor silicon anodes. Due to severe salt and solvent degradation, silicon anodes with the LiTDI-baseline electrolyte showed rather poor electrochemical performance. However, with the SEI-forming additives of FEC and VC, the cycling performance of such battery system is greatlyimproved, owing to a stabilized electrode/electrolyte interface.

This thesis work highlights that cooperation of appropriate electrolyte additives is an effectiveyet simple approach to enhance battery performance, and in addition, that the interfacialchemistries are of particular importance to deeply understand battery behavior.

Keywords: Lithium-ion batteries, electrolyte, electrolyte additives, electrochemistry,interfacial chemistry

Chao Xu, Department of Chemistry - Ångström, Structural Chemistry, Box 538, UppsalaUniversity, SE-751 21 Uppsala, Sweden.

© Chao Xu 2017

ISSN 1651-6214ISBN 978-91-554-9931-0urn:nbn:se:uu:diva-319425 (http://urn.kb.se/resolve?urn=urn:nbn:se:uu:diva-319425)

List of Papers

This thesis is based on the following papers, which are referred to in the text by their Roman numerals.

I LiTDI: A Highly Efficient Additive for Electrolyte Stabiliza-

tion in Lithium-Ion Batteries C. Xu, S. Renault, M. Ebadi, Z. Wang, E. Björklund, D. Guyo-mard, D. Brandell, K. Edström and T. Gustafsson Chemistry of Materials 29 (2017) 2254–2263

II The Role of LiTDI Additive in LiNi1/3Mn1/3Co1/3O2/graphite Lithium-ion Batteries at Elevated Temperatures C. Xu, F. Jeschull, W. R. Brant, D. Brandell, K. Edström and T. Gustafsson In Manuscript

III Improved Performance of Silicon Anode for Li-Ion Batter-ies: Understanding the Surface Modification Mechanism of Fluoroethylene Carbonate as an Effective Electrolyte Addi-tive C. Xu, F. Lindgren, B. Philippe, M. Gorgoi, F. Björefors, K. Edström, T. Gustafsson Chemistry of Materials 27 (2015) 2591–2599

IV A hard X-ray photoelectron spectroscopy study on the solid

electrolyte interphase of a lithium 4,5-dicyano-2-(trifluoromethyl)imidazolide based electrolyte for Si-electrodes F. Lindgren, C. Xu, A. Andersson, K. Edström, T. Gustafsson, F. Björefors Journal of Power Sources 301 (2016) 105-112

V SEI Formation and Interfacial Stability of a Si Electrode in a LiTDI-Salt Based Electrolyte with FEC and VC Additives for Li-Ion Batteries F. Lindgren, C. Xu, L. Niedzicki, M. Marcinek, T. Gustafsson, F. Björefors, K. Edström, and R. Younesi ACS Applied Materials & Interfaces 8, 15758–15766, 2016

Disclaimer: Parts of this thesis are based on my licentiate thesis entitled All silicon lithium-ion batteries (Uppsala University, 2015)

Reprints were made with permission from the respective publishers.

Contributions to the papers

Paper I. Conceived the idea, planned and performed most of the experi-ments and data analysis. Involved in all discussions and wrote the manuscript.

Paper II. Conceived the idea, planned and performed most of the experi-

ments and data analysis. Involved in all discussions and wrote the manuscript.

Paper III. Conceived the idea, planned and executed most of the experi-ments and data analysis. Involved in all discussions and wrote the manuscript.

Paper IV. Conceived the idea, planned and performed PES measurements. Involved in all discussions and manuscript writing.

Paper V. Involved in experiments and all discussions. Contributed to manuscript writing.

Other publications to which the author has contributed that are not included in the present thesis:

Interface layer formation in solid polymer electrolyte lithium batteries: an XPS study C. Xu, B. Sun, T. Gustafsson, K. Edström, D. Brandell, M. Hahlin Journal of Materials Chemistry A 2 (2014) 7256-7264 At the polymer electrolyte interfaces: the role of the polymer host in interphase layer formation in Li-batteries B. Sun, C. Xu, J. Mindemark, T. Gustafsson, K. Edström, D. Brandell Journal of Materials Chemistry A 3 (2015) 13994-14000 Flexible freestanding Cladophora nanocellulose paper based Si anodes for lithium-ion batteries Z. Wang, C. Xu, P. Tammela, J. Huo, M. Strømme, K. Edström, T. Gustafs-son, L. Nyholm Journal of Materials Chemistry A 3 (2015) 14109-14115 Conducting Polymer Paper-Based Cathodes for High-Areal-Capacity Lithium-Organic Batteries Z. Wang, C. Xu, P. Tammela, P. Zhang, K. Edström, T. Gustafsson, M. Strømme, L. Nyholm Energy Technology 3 (2015) 563-569 A high pressure x-ray photoelectron spectroscopy experimental method for characterization of solid-liquid interfaces demonstrated with a Li-ion battery system. J. Maibach, C. Xu, S. K. Eriksson, J. Åhlund, T. Gustafsson, H. Siegbahn, H. Rensmo, K. Edström, M. Hahlin Review of Scientific Instruments 86 (2015) 044101

Contents

1. Lithium-ion batteries ........................................................................... 11

2. Scope of this thesis .............................................................................. 17

3. Non-aqueous liquid electrolytes .......................................................... 183.1 Lithium salts ............................................................................... 183.2 Solvents ...................................................................................... 203.3 Additives .................................................................................... 22

4. The solid electrolyte interphase ........................................................... 25

5. Methodology ........................................................................................ 285.1 Battery preparation ..................................................................... 285.2 Galvanostatic battery cycling ..................................................... 295.3 Photoelectron spectroscopy ........................................................ 315.4 Other characterizations ............................................................... 33

6. Improving conventional LiPF6-based electrolyte with additives ......... 366.1 LiTDI as a highly effective electrolyte additive ......................... 36

6.1.1 Moisture scavenging effects .................................................. 366.1.2 Cycling performance ............................................................. 396.1.3 The graphite/electrolyte interface .......................................... 426.1.4 The NMC/electrolyte interface .............................................. 46

6.2 FEC as SEI-forming additive for silicon anodes ........................ 476.1.1 Electrochemistry and surface morphology ............................ 476.1.2 SEI composition ..................................................................... 49

7. Developing LiTDI-based electrolytes for silicon electrodes ............... 527.1 Electrochemical performance ..................................................... 527.2 Silicon/electrolyte interfacial chemistries .................................. 54

7.2.1 Initial SEI formation on silicon with pure LiTDI-based electrolyte ............................................................................................ 547.2.2 Surface chemistry on silicon with additives .......................... 55

Conclusions and outlook ............................................................................... 59

Sammanfattning på svenska .......................................................................... 62

Acknowledgements ....................................................................................... 65

Reference .................................................................................................. 67

Abbreviations

AFM atomic force microscopy BEV battery electric vehicles C.E. coulombic efficiency CB carbon black CE counter electrodes CMC-Na carboxymethyl cellulose sodium salt CV cyclic voltammetry DEC diethyl carbonate DFT density functional theory DMC dimethyl carbonate EC ethylene carbonate EIS electrochemical impedance spectroscopy EMC ethyl methyl carbonate ESW electrochemical stability window FEC fluoroethylene carbonate FT-IR Fourier transform infrared spectroscopy HOMO highest occupied molecular orbital IEA International Energy Agency IMFP inelastic electron mean free path LCO lithium cobalt oxide, LiCoO2 LFP lithium iron phosphate, LiFePO4 LIB lithium-ion batteries LiBOB lithium bis(oxalato)borate LiTDI lithium 2-trifluoromethyl-4,5-dicyanoimidazole LiTf lithium trifluoromethanesulfonate LiTFSI lithium bis(trifluoromethylsulphonyl)imide LMO lithium manganese oxide, LiMn2O4 LNMO lithium nickel manganese oxide, LiNi0.5Mn1.5O4 LTO lithium titanate, Li4Ti5O12 LUMO lowest unoccupied molecular orbital MCMB mesocarbon microbeads NCA lithium nickel-cobalt-aluminum oxide LiNi0.8Co0.15Al0.05O2 NMC lithium nickel-manganese-cobalt oxide, LiNixMnyCo(1-x-y)O2 NMR nuclear magnetic resonance spectroscopy OCV open circuit voltage

PC propylene carbonate PES photoelectron spectroscopy RE reference electrode SBR styrene-butadiene rubber SEI solid electrolyte interphase SEM scanning electron microscopy TEM transmission electron microscopy TGA thermogravimetric analysis TOF-SIMS time of flight-secondary ion mass spectroscopy VC vinylene carbonate WE working electrode XPS X-ray photoelectron spectroscopy XRD X-ray diffraction

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1. Lithium-ion batteries

Continuous development of modern society leads to massive energy con-sumption, and currently, worldwide energy supply relies heavily on fossil fuels, namely oil, natural gas and coal.1–3 Figure 1a illustrates the total pri-mary energy supply by different fuels in 2014, during which oil, coal and natural gas dominantly contributed to 31.3%, 28.6% and 21.2%, respective-ly.2 However, fossil fuels are not renewable, and the reserves are finite, which means that it is only a matter of time until they will be completely consumed. Moreover, combustion of fossil fuels inevitably generates the primary greenhouse of CO2, which is already causing observable global cli-mate change.4, 5 Therefore, seeking alternatives in terms of renewable energy resources has already been widely acknowledged as an urgent issue for hu-man society.

Figure 1. (a) The total primary energy supply by different fuels in the year 2014. (b)The total oil consumption by sector in the year 2014. © IEA (International Ener-gy Agency) 2016 Key World Energy Statistics, IEA publishing.

Among the three main fossil fuels, oil contributed the most to the total primary energy supply, and consumption of oil is dominated by transport (shown in Figure 1b).1, 2 Replacing oil with renewable resources in transpor-tation is therefore of great importance to reducing oil consumption. Re-chargeable batteries, which can be discharged to power a load and recharged many times, are therefore particularly interesting candidates for such appli-cation. In practice, the installed batteries should provide as much energy as

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possible, so that the vehicle has sufficient range to reach destination. Conse-quently, energy density, which is a measure of a certain amount of stored energy in a given unit of volume or weight, is a critical parameter for the battery technologies in automobile applications. Among the existing electro-chemical energy storage technologies, lithium-ion batteries (LIBs) possess the highest energy density compared to lead-acid or nickel-metal hydrides, as summarized in Figure 2. It should be noted that the values for energy densities were obtained in 2009, and today, after continuous development in battery chemistry and engineering, state-of-the-art lithium-ion batteries can already achieve a volumetric energy density of more than 650 Wh/L at the cell level (Panasonic NCR18650B, nominal voltage of 3.6 V, nominal capac-ity of 3350 mAh).6 Multiple lithium-ion battery chemistries have been suc-cessfully developed and applied to battery electric vehicles (BEV); a selec-tive summary of commercially available BEVs in the US market is presented in Table 1.

Figure 2. Comparison of the volumetric and gravimetric energy densities between different rechargeable battery technologies. Reproduced from Ref 7 with permission of The Royal Society of Chemistry.

Although realization of electro-mobility has become one of the most im-portant driving forces in developing advanced lithium-ion batteries, the LIB technology was initially applied to portable electronics such as mobile phones, laptops and cameras, after its commercialization by Sony in 1991.8 Since then, performance of lithium-ion batteries has been significantly im-proved in almost every aspect – energy density, cycle life, safety, etc., and LIBs are dominating the market of power supplies for such portable elec-tronics. In addition, the LIB technology has been adopted to emerging re-newable energy systems such as wind and photovoltaics. For such applica-tions, LIBs are commonly used to balance consumption and generation of electricity.

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Table 1. A selective summary of battery electric vehicles on the US market. Repro-duced from Ref 9 with permission of The Electrochemical Society.

Manufacturer Model Battery Size (kWh)

Battery Chem-istry* Battery Supplier Vehicle

Range (km)

Tesla S 60-100 C/NCA Panasonic/Tesla 334-508

BMW i3 22, 33 C/NMC Samsung/Bosch 129, 183

Nissan Leaf 24, 30 C/LMO (C/NMC)

AESC and LG Chem 135, 172

Chevrolet Spark 19 C/LFP A123 132

Mitsubishi I 16 LTO/LMO Toshiba 100

* NCA = LiNi0.8Co0.15Al0.05O2, C = graphite, NMC = LiNixMnyCo(1-x-y)O2, LMO = LiMn2O4, LFP = LiFePO4, LTO = Li4Ti5O12

A conventional lithium-ion battery is composed of one positive (also re-ferred as cathode) and one negative (anode) electrode, and these two elec-trodes are separated by an electrically insulating membrane which is soaked with non-aqueous liquid electrolyte. A schematic drawing of a lithium-ion battery is shown in Figure 3. During charging, the cathode material is elec-trochemically oxidized, and Li-ions are extracted from the material and move from the cathode to the anode through the electrolyte. Electrons move from the cathode to the anode in the outer circuit for charge compensation. Simultaneously, Li-ions and electrons are both transported to the anode, and Li-ions are inserted into the anode material. These processes are reversed during discharging, where the electrons in the outer circuit give rise to a useful electrical current.

Figure 3. Working principle and basic components of a conventional lithium-ion battery.

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Cathode materials LiCoO2 (LCO) was used as the cathode material in first generation of lithi-um-ion batteries commercialized by Sony.8 The layer-structured LCO has a reasonably good energy density, which is attributed to a theoretical capacity of 136 mAh/g (corresponding to extracting and inserting about 0.5 Li per LiCoO2) and a nominal potential of 3.9 V (vs. Li+/Li. All potentials stated in this thesis are, unless specifically stated, measured versus the Li+/Li redox couple).10, 11 However, this material suffers from high cost, toxicity and safe-ty hazard from O2 evolution after over-charge.12, 13 Metal substitutions (Mn, Ni and/or Al) are effective approaches to mitigate these issues, and a few compounds have already been widely adopted in commercial lithium-ion batteries, such as LiNi0.33Mn0.33Co0.33O2 (NMC 111) and LiNi0.8Co0.15Al0.05O2 (NCA).14–17 The redox reactions of the NMC 111 material are summarized as: Delithiation:

LiNi0.33Mn0.33Co0.33O2→ Li(1-x)Ni0.33Mn0.33Co0.33O2 + xLi+ + xe"

Lithiation:

Li(1-x)Ni0.33Mn0.33Co0.33O2 + xLi+ + xe"→ LiNi0.33Mn0.33Co0.33O2

More recently, lithium-rich layered oxides, commonly denoted as (1-x)Li2MnO3·xLiMO2 (M = Mn, Co, Ni), have attracted tremendous attention due to their high energy density (gravimetric capacity of ~250 mAh/g , aver-age voltage of ~ 3.7 V).17–20 Poly-anionic type compounds are also attractive alternative cathode materials to the layered LiMO2 (M = Mn, Co, Ni and/or Al). In particular, the olivine LiFePO4 (LFP) has already been increasingly utilized in commercial applications, primarily due to its outstanding thermal stability.16, 17, 21–23 However, LFP is limited by its poor electronic and ionic conductivity, as well as its relatively low working potential (3.5 V).21, 22 Anode materials Graphite is an excellent anode material and is dominatingly used in most of today’s commercial lithium-ion batteries. Graphite is a Li-intercalation mate-rial and has a theoretical capacity of 372 mAh/g, which corresponds to full lithiation to the LiC6 phase.24 Its low irreversible capacity, high coulombic efficiency, good cycle life, low cost, and low average working potential (0.125 V), are significant advantages for its use as an anode material. Li4Ti5O12 (LTO) is another intercalation anode material which is commer-cially applied. However, LTO has a theoretical capacity of 175 mAh/g and 1.5 V operating potential, resulting in a much lower energy density in practi-cal cells.25 Nevertheless, the LTO anode material generally exhibits excellent

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lifetime. This is mainly attributed to its zero-strain nature, which means that no volume change occurs upon lithiation/delithiation, and the high working potential largely avoids undesired electrolyte degradation.25

Alloying anode materials, on the other hand, generally have significantly higher gravimetric and volumetric capacity compared to the widely utilized graphite, and therefore are promising candidates to further improve the ener-gy density of today's LIBs.26 In particular, silicon has attracted the most at-tention. The electrochemical reactions of a silicon anode are described as:

Lithiation: xLi++xe"+Si→LixSi

Delithiation: LixSi→xLi++xe"+Si

Upon full lithiation to the Li15Si4 phase, silicon has a theoretical capacity of 3579 mAh/g gravimetrically and 2194 Ah/L volumetrically at room tempera-ture.27 This corresponds to a 34% increase in volumetric energy density in the full cell stack, theoretically, if graphite is replaced with silicon, as esti-mated by Obrovac and Chevrier using LCO as the cathode.26 However, the practical use of silicon anode materials is severely limited by the large irre-versible capacity, low coulombic efficiency and poor cycle life.26, 28, 29 Such poor electrochemical performances are primarily attributed to the large vol-ume change in particle size during cycling. More specifically, a full lithiation to Li15Si4 leads to a volume expansion of 280%.27 Consequently, this leads to multiple serious issues, such as bulk particle pulverization and break-down of electrical contacts between silicon particles, current collectors as well as conductive additive.30–33 Moreover, the solid electrolyte interphase (SEI) formed on the silicon particles is repeatedly damaged and repaired, which is accompanied by continuous electrolyte degradation.33–35 This issue is further discussed in Chapter 4: The solid electrolyte interphase. Various approach-es have been investigated to counter these problems, such as capacity-limited cycling, nano-engineering, surface modifications and coatings, inter-metallic compounds, functional binders and SEI-forming additives.26 For instance, reducing silicon particles to nano-size is a highly effective approach to miti-gate the issue of particle pulverization.30, 36, 37 Functional binders, such as carboxymethyl cellulose (CMC), poly(acrylic acid) and alginates, which all contain carboxyl groups, have been shown to be able to enhance cycling performance of silicon materials.38–42 It is believed that the carboxyl group in these binders can be chemically bonded to the native surface of silicon parti-cles, which consists of silicon oxide and silanol, and such bonding can sig-nificantly improve the mechanical and electrochemical stability of silicon electrodes.26, 41, 43, 44

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Electrolytes The electrolyte in LIBs governs Li+ transport between the cathode and the anode during electrochemical cycling. Generally, the electrolyte consists of lithium salts and solvents or a Li+ host material. LIB electrolytes can be cat-egorized into the following classes: non-aqueous liquid electrolyte, gel elec-trolyte, solid polymer electrolyte, ionic liquid and ceramic electrolyte. This thesis will be focused on the conventional non-aqueous liquid electrolytes, which are the most widely used electrolytes in today's commercial lithium-ion batteries. Such electrolytes are in general composed of one or a few lith-ium salt(s), low molecular weight carbonate solvents, as well as various ad-ditives in small quantities.45, 46 Since electrolytes are the main focus of this thesis, more detailed introductions on the salt, solvent and additive are indi-vidually given in Chapter 3: Non-aqueous liquid electrolytes.

As an example, the commercially available Selectilyte™ LP40 electrolyte from BASF consists of 1 M LiPF6 in the solvent mixture of ethylene car-bonate (EC) and diethyl carbonate (DEC) (1/1, weight ratio). Compared to the other types, such non-aqueous electrolytes possess significant ad-vantages, in particular a wide electrochemical stability window (ESW) and high ionic conductivity over a wide temperature range.45, 46 The ionic con-ductivity of state-of-the-art non-aqueous electrolytes can reach the level of 10 mS/cm at room temperature, and support battery operation in the range of -30 to +60 °C.46 Ideally, the ionic conductivity should be solely contributed by Li+ to prevent depletion of Li+ concentration at the electrode surface, which may cause cell failure at high cycling rate. In practice, however, both Li+ cations, anions as well as charged multiplets contribute to charge transport in the electrolyte. The lithium ion transference number (tLi) is a measure of the fraction of the current facilitated by Li+ in the electrolyte. Conventional electrolytes with LiPF6 in a moderate concentration of around 1 M typically have a Li+ transference number of 0.2 to 0.5, depending on the solvents and measurement methods.46–48 In such electrolytes, Li+ is predomi-nately coordinated to the carbonyl oxygen from the carbonate solvents due to the strong interaction between the oxygen lone-pair and Li+, as confirmed by multiple studies with various techniques.46, 49–52 Moreover, the coordination number of Li+ in conventional carbonate electrolytes is typically found to be 4-6, depending on the salt concentration.46

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2. Scope of this thesis

The primary scope of this thesis work is developing advanced electrolytes for lithium-ion batteries and understanding electrode/electrolyte interfacial chemistries. Studies have been carried out on electrolyte additives as well as alternative lithium salts. The electrode/electrolyte interfaces in such novel electrolytes have been investigated with multiple advanced characterization techniques.

Key findings from the five publications listed in List of Papers are sum-marized in Chapter 6 and 7. The results are categorized into two groups according to the different baseline electrolytes:

1. Conventional LiPF6-based electrolytes: Paper I and II studied the

electrolyte stabilization mechanism of the LiTDI additive, as well as its effects on electrochemistry and interfacial chemistries in NMC/Li and NMC/graphite cells. Paper III investigated the use of FEC as the SEI-forming electrolyte additive for silicon electrodes and the underlying interfacial chemistry.

2. LiTDI-based electrolytes: Paper IV analyzed the chemical composi-tion of the SEI layer on silicon anodes in the batteries which used a LiTDI-based electrolyte. Paper V further investigated the influence of SEI-forming additives FEC and VC on the electrochemical per-formance and interfacial chemistry on silicon anodes.

The molecular structures of four key chemicals used in this thesis are

summarized in Figure 4.

Figure 4. Molecular structures of LiPF6, LiTDI, FEC and VC.

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3. Non-aqueous liquid electrolytes

3.1 Lithium salts Several lithium salts were explored at the early stage of developing function-ing lithium-ion batteries, such as LiAsF6, LiBF4, LiPF6, lithium bis(oxalato)borate (LiBOB), lithium trifluoromethanesulfonate (LiTf), LiClO4, etc.45 Among all these candidates, LiPF6 showed the most well bal-anced properties, which lead to its dominating position as the primary salt in today's state-of-the-art non-aqueous electrolytes. These essential properties are: (1) weak Li+-PF6

- interactions which ensures sufficient Li+-PF6- disasso-

ciation in solution, resulting in a high Li+ transport ability; (2) wide electro-chemical stability window; (3) ability of passivating the aluminum current collector; and (4) formation of stable solid electrolyte interphases.45, 46, 53 However, LiPF6 is still far from the perfect salt. In particular, LiPF6 exhibits rather poor thermal stability, and decomposes to LiF and PF5 at elevated temperatures. Moreover, the P-F bond in LiPF6 as well as the thermal de-composition product PF5 are rather labile towards hydrolysis by inevitable trace amounts of moisture in batteries. These hydrolysis reactions lead to the formation of highly corrosive HF, which is a detrimental species for both active and inactive materials in batteries. These reactions are summarized as:

LiPF6 ⇌ LiF + PF5

PF5 + H2O → PF3O + 2HF

In addition, such parasitic reactions can not only be initiated by H2O, but also the aprotic solvents in the electrolyte, especially linear alkyl carbonates such as DEC.54 The poor thermal stability of LiPF6-based electrolytes results in continuous electrolyte degradation, which consequentially leads to per-formance fading due to loss of active lithium and cell impedance increase in practical lithium-ion batteries.55 Therefore, the poor thermal stability of to-day's non-aqueous electrolytes is considered as one of the primary reasons for accelerated battery aging at elevated temperatures.55–57

Numerous alternative lithium salts with enhanced thermal stability have been developed with the aim of replacing LiPF6 as the primary salt in elec-trolytes. However, LiPF6 is still holding the dominating position in today's commercial lithium-ion batteries, and the situation is unlikely to change in the near future.46 This is mainly due to the fact that such improved stability

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is usually achieved at the cost of losing some other essential properties. For instance, LiBOB is an extensively investigated alternative to LiPF6, and was originally proposed as a promising salt for elevated temperature applications. This is attributed to its high thermal stability (up to 275 °C), and it is also a fluorine-free anion, which eliminates the possibility of HF formation in the electrolyte.58 Indeed, improved performance of batteries which use LiBOB-based electrolytes compared to LiPF6 has been reported in multiple stud-ies.58–60 However, its intrinsic limitations, namely limited solubility in car-bonate solvents and anodic stability only up to 4.2 V have severely restricted its wide application.45 Lithium bis(trifluoromethylsulphonyl)imide (LiTFSI) is another alternative which also exhibits significantly better thermal stabil-ity, showing no decomposition below 360 °C.45 However, electrolytes that consist of LiTFSI at moderate concentrations, e.g. 1 M, in carbonate solvents are not able to passivate aluminum foil, which is used as the current collector for the cathode, at potentials above 4.2 V.45 Krämer et al. concluded that aluminum undergoes anodic dissolution reactions, which are accompanied by various electrolyte decomposition reactions.61 Therefore, the application of LiTFSI is severely limited, especially for high voltage cathode materials which operate above 4.2 V. Besides LiBOB and LiTFSI, which are the most representative examples of borate and imide type salts, respectively, there are more salts from other families such as phosphates, heterocyclic anions and aluminates which have been proposed and successfully developed.46 However, these newly developed alternatives all have their own limitations, and the examples of LiBOB and LiTFSI illustrate the importance of ful-filling all the crucial criteria of a lithium salt in order to achieve successful application.

More recently, lithium 2-trifluoromethyl-4,5-dicyanoimidazole (LiTDI) has been developed as another alternative and shows superior thermal stabil-ity as well as chemical stability against water.62–65 Niedzicki et al. reported that the LiTDI electrolyte at a lower salt concentration (0.31 mol/kg in EC/DMC, 1/2 wt/wt) shows similar Li+ conductivity as compared to the conventional LiPF6 in the same solvent mixture.66 Such performance pre-sents a significant advantage of LiTDI over other salts for large-scale com-mercial application, as much lower amounts of salt are required for electro-lyte production. Moreover, a LiTDI electrolyte (0.5 M, PC solvent) has been demonstrated to be electrochemically stable up to 4.6 V vs. Li+/Li against aluminum.62 Therefore, LiTDI-based electrolytes can be particularly attrac-tive for cathode materials which operate at high voltages, and have already showed promising performance in various lithium-ion battery chemistries.66–

68 In particular, Paillet et al. demonstrated that LFP/LTO and NMC/LTO full cells which used 1 M LiTDI in EC/DEC (3/7, v/v) without any additional electrolyte additives, already exhibited as good performance as the ones using standard LiPF6 electrolytes.69

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However, LiTDI based electrolytes, especially in the absence of SEI-forming additives, exhibit rather poor compatibility with low working poten-tial anode materials, for instance graphite and silicon.69–71 Severe electrolyte decomposition has been observed in such battery chemistries. This leads to high irreversible capacities and low coulombic efficiencies during the initial cycles, as well as poor capacity retention during extended cycling.69–71 Shkrob et al. proposed that, on the lithiated graphite surface, LiTDI under-goes sequential de-fluorination reactions, which results in a thick and ionic-resistive surface layer on the electrode.70 This proposal was based on DFT calculations and XPS experimental evidence which shows large quantities of LiF on the graphite electrodes which had been cycled with LiTDI-based electrolyte.70–72 With the presence of SEI supporting additives, such as FEC and VC, LiTDI-based electrolytes are, however, able to deliver comparable performance to the conventional LiPF6 electrolyte in graphite/Li cells.69

3.2 Solvents The solvent is the other essential component of the electrolyte in lithium-ion batteries besides the salt. In today's state-of-the-art non-aqueous electrolytes, the solvent is primarily a mixture of multiple low molecular weight car-bonates, for instance, EC, propylene carbonate (PC), dimethyl carbonate (DMC), DEC and ethyl methyl carbonate (EMC).45 For a well-functioning electrolyte in lithium-ion batteries, the solvent should meet the following criteria in terms of physical properties: (1) it should have a high dielectric constant (e) in order be able to dissolve salts to a sufficient concentration; (2) it should have a low viscosity (h) to ensure facile ion transport; (3) it should have a wide electrochemical stability window and good compatibility with electrode materials; (4) it should have a low melting temperature (Tm) and a high boiling temperature (Tb) so that it remains liquid over a wide tempera-ture window; (5) it should have a high flash point for good safety; (6) it should be nontoxic and economically acceptable. Some properties of the commonly used carbonate solvents are summarized in Table 2.

The strategy of mixing cyclic carbonates with linear carbonates aims at achieving balanced properties, since they both have their advantages and limitations. For instance, the cyclic carbonates, i.e. EC and PC, have much higher dielectric constant compared to the linear carbonates, as listed in Ta-ble 2. On the other hand, without involving any linear carbonates, which have lower viscosity, the high viscosity of EC and PC severely limits their practical use, especially at high salt concentrations. Introducing linear car-bonates, however, significantly lower the flash point (FP) of the electrolyte, which is an essential parameter for battery safety. As listed in Table 2, EC and PC have a FP of 145.5 ± 4 and 135.5 ± 4 °C, respectively, whereas, for

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example, the FP of the standard electrolyte of 1 M LiPF6 EC/DMC (1/1, wt/wt) is as low as 25.5 ± 1 °C.73 Table 2. Physical properties of carbonate solvents used in lithium-ion battery elec-trolytes.

Solvent M.W. e at 25 °C45 h (cP) at 25 °C45 Tm (°C)45 Tb

(°C)45 Tf (°C)73

EC 88 89.78 1.90 (40 °C) 36.4 248 145.5 ± 4

PC 102 64.92 2.53 -48.8 242 135.5 ± 4

DMC 90 3.107 0.59 (20 °C) 4.6 90.5 16 ± 1

DEC 118 2.805 0.75 -74.3 126 33 ± 1

EMC 104 2.958 0.65 -53 107.5 23.5 ± 1

Moreover, another important reason for mixing cyclic and linear carbonates is the electrochemical performance. In particular, EC has, for a long time, been considered as an indispensable component in the batteries which use carbonaceous anode materials such as mesocarbon microbeads (MCMB), graphite, etc.45, 74 This is mainly attributed to the fact that EC is able to be reduced and simultaneously form a stable solid electrolyte interphase (SEI) layer on the surface of such electrodes.24, 74 Detailed insights into the SEI are introduced later in Chapter 4: The solid electrolyte interphase. However, EC-based electrolytes suffer from a few critical issues, particularly that high amounts of EC in the electrolyte (usually 30% to 50%) significantly limit the operation temperature range due to the high melting point of EC. More spe-cifically, the Li+ conductivity of EC-containing electrolytes drops signifi-cantly when the temperature decreases.75 This issue is especially of concern for EV applications which require a wide operating temperature range.

Gmitter et al. investigated the possibility of EC-free electrolytes by using EMC as the predominant solvent.76 To achieve a stable SEI on the graphite electrode without EC, a small amount of vinylene carbonate or fluoroeth-ylene carbonate, which are SEI-forming additives and referred to as "ena-blers", were introduced into the electrolyte. This strategy has been further extensively explored by J. Dahn's group.77–79 For the high voltage NMC442/graphite cells, which operate to 4.4 V, Ma et al. concluded that a solvent blend of 95% EMC with 5% FEC enabler showed the best combina-tion of properties in terms of cell impedance, gas evolution, voltage drop, capacity fading and the enabler consumption.79

It is widely acknowledged that the oxidation stability of the conventional carbonate based electrolytes is limited to 4.5 V, above which the electrolyte oxidation at the surface of the cathode becomes more severe. Fluorination of

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carbonates is a common approach to increase the oxidation stability as the energy level of the highest occupied molecular orbital (HOMO) is generally lowered when F is introduced into carbonate molecules. Zhang et al. pre-sented several electrolytes specifically designed for 5 V LiNi0.5Mn1.5O4 (LNMO) based battery chemistries. These electrolytes contained partially fluorinated carbonates and a highly fluorinated ether.80 LNMO/LTO cells with such electrolytes exhibited dramatic performance improvements at an elevated temperature of 55 °C, as compared to those which used the standard EC/EMC-based Gen2 electrolyte.80 Xia et al. also demonstrated that the fluorinated electrolytes out-performed the conventional ones in high voltage NMC442/graphite cells tested to 4.5 V.81, 82

Apart from using alternative carbonates, several other classes of solvents have been explored with the aim to, completely or partially, replace car-bonates to achieve certain performance improvements. For instance, ethers generally have low viscosities, and are therefore considered as good candi-dates for promoting Li+ ionic conductivity.45 However, their practical appli-cations are limited by the poor oxidation stability (up to 4 V).83, 84 Sulfones and sulfoxides, on the other hand, generally have excellent oxidation stabil-ity as well as high dielectric constants. However, these solvents are limited by the inability to form a stable SEI on graphite and their high viscosities, which result in low ionic conductivity.45, 46 Such examples illustrate the dif-ficulty of meeting all the essential criteria and the complexity in developing well-functioning electrolytes.

3.3 Additives Due to the demanding requirements for the electrolyte in lithium-ion batter-ies, it is nearly impossible to achieve satisfactory performance with only the basic components of a single lithium salt and standard solvents. Incorporat-ing small amounts of other components (usually considered as lower than 10%), which are referred as additives, without altering the skeletal composi-tion, is therefore an economical and convenient way to improve perfor-mance.46 Various additives have been reported over the past several years, and can be generally categorized into the following families based on the functionality: additives for anode, additives for cathode, flame retardants and overcharge protection additives. Among all these additives, additives for anode and cathode which aim to stabilize the electrode/electrolyte interfaces are the most extensively explored for lithium-ion batteries. In particular, vinylene carbonate (VC) and fluoroethylene carbonate (FEC) have been investigated in this thesis work.

VC is by far the most successful electrolyte additive for improving the sta-bility of the electrode/electrolyte interface and was originally proposed by the company SAFT85. Later on, Aurbach et al. investigated the effects of VC on

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synthetic graphite anodes as well as spinel LiMn2O4 and LiNiO2 cathodes.86 The authors concluded that VC is particularly beneficial for the graphite an-ode, although it also showed effectiveness on the two mentioned cathodes. Moreover, numerous studies show that VC is a highly effective additive also for improving the performance of alloy-type anode materials such as Si and Sn.87–91 On the other hand, Dahn and co-workers showed more pronounced performance improvements in NMC/Li cells as compared to graphite/Li cells, suggesting that VC is more effective for the NMC cathode rather than the graphite anode.92 Nevertheless, it has been widely accepted that such perfor-mance improvements for various electrode materials is primarily attributed to the formation of polymeric species from the degradation products of VC, alt-hough the exact mechanism still under debate.86, 93, 94

Fluoroethylene carbonate (FEC) is another electrolyte additive which has been extensively studied. In addition to graphite, FEC has been proven to be highly effective for various types of silicon-based materials, such as silicon thin film, nano-particles, nano-wires, etc.95–102 With the presence of FEC in the electrolyte, silicon-based batteries were found to be superior in terms of capacity retention as well as coulombic efficiency, compared to the ones which used FEC-free electrolytes.

Unlike VC, which decomposes and forms an organic-dominated surface layer, the degradation products of FEC generally contain significant amounts of inorganic LiF. Besides LiF, formation of polymeric species is also usually observed. Naikai et al. proposed that the FEC decomposition is a two-electron reduction mechanism, which eventually leads to the formation of Li2CO3, LiF and polyacetylene.96 Etacheri et al., however, proposed that FEC decomposition is rather a one-electron reaction, meaning that FEC was firstly derived into LiF and VC, which subsequently undergoes an electro-chemical reduction induced polymerization and forms poly(VC).97 More recently, Jung et al. introduced a four-electron mechanism for reduction of the FEC molecule where the reactions lead to formation of CO2, H2, LiF, Li2O, Li2CO3 and a partially cross-linked polymer.101 Although the exact reaction mechanism is still unclear, it has nonetheless been widely accepted that the surface layer derived from FEC on silicon electrodes is mainly com-posed of polymeric compounds with embedded LiF domains.

A few groups also reported that the use of FEC as electrolyte additive or co-solvent is able to improve the performance of cathode materials, especial-ly high voltage cathodes such as LiCoPO4, LiNi0.5Mn1.5O4 (LNMO) and LiNi0.33Mn0.33Co0.33O2 (charged up to 4.6 V).103–105 On the other hand, Akte-kin et al. showed that FEC had rather limited influence on electrochemical performance of the LNMO/LTO battery chemistry.106

Besides the effects of FEC on battery performance, there are also a few other important facts worth noting for the practical application of FEC. Jung et al. reported that FEC was continuously consumed during electrochemical cycling.101 It should be mentioned that the examined batteries used silicon-

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carbon composited anodes. Once FEC was totally consumed, a significant increase in cell polarization and a very accelerated capacity fade were ob-served. Such findings highlight the importance of the amount of FEC in the electrolyte for lithium-ion batteries, especially the ones containing silicon-based anode materials. Moreover, Kim and co-workers convincingly showed that excess FEC is likely to undergo defluorination in the presence of Lewis acids (e.g. PF5 in LiPF6-based electrolytes), and therefore promotes the for-mation of highly corrosive HF and various other acids, especially at elevated temperatures.107 As shown in Figure 5, the color of two electrolytes which contained both FEC and LiPF6 changed dramatically after storage, and these electrolytes also became highly acidic. As a result, FEC-based LCO-NMC/Si-C full cells exhibited worse cycling performance, displaying signif-icantly more severe metal dissolution as compared to the EC-based cells.107 Online electrochemical mass spectrometry revealed that the gas generation from batteries which use EC-based and FEC-based electrolytes are different, and excessive amounts of FEC can lead to more gas evolution.108, 109

Figure 5. (a) FEC, (c) FEC/DEC (3/7, wt/wt), and (e) DEC without salt, and (b) FEC, (d) FEC/DEC (3/7, wt/wt), and (f) DEC with 1.5 M LiPF6 salt before and after storage at 60 °C for 3 days. (g) Indicator chart for the pH scale. Reprinted from Ref 107 with permission from Elsevier.

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4. The solid electrolyte interphase

As mentioned previously, one essential criterion for an ideal electrolyte in lithium-ion batteries is a wide ESW, which is defined by the energy differ-ence between the HOMO and LUMO of the electrolyte, as illustrated in Fig-ure 6.

Figure 6. A schematic illustration of the energy diagram of a conventional lithium-ion battery.

For conventional LiPF6-based liquid electrolytes, the ESW is predominately attributed to the stability window of the carbonate solvents.45, 46, 110 These are thermodynamically unstable at especially low potentials when the carbonate molecules are coordinated to Li+ in the electrolyte.45, 46, 110 For example, computational studies suggest that the presence of Li+ in bulk electrolytes significantly promotes the reduction reaction of EC.110, 111 As a result, car-bonate electrolytes will be continuously reduced at the anodes which operate at low average potentials (e.g. 0.125 V for graphite). This issue was one of the major obstacles in the early development stage of rechargeable lithium-ion batteries, until Fong, von Sacken and Dahn introduced the key finding of EC, which decomposes at a potential of ~0.8 V but simultaneously form a stable passivating layer on carbonaceous anodes.24, 112, 113 This beneficial layer is referred as the solid electrolyte interphase (SEI), and is able to great-ly suppress continuous electrolyte degradation and therefore improves the cyclability of batteries. Ideally, the SEI should be electronically insulating to prevent electron transfer, and therefore mitigate further electrolyte reduction. At the same time, it should be Li+ conductive so that Li+ transport can be

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facilitated during cycling. Moreover, this layer should be uniform, stable and flexible to protect the electrode and accommodate volume changes of the electrode during cycling.113

The SEI concept was originally introduced by Peled, and it was proposed to describe the surface layer formed when alkali and alkaline earth metals were in contact with electrolytes in non-aqueous battery systems.114 This layer serves as an interphase between the electrode and electrolyte and has similar properties as a solid electrolyte, i.e. ionically conductive but electron-ically insulating.114 This concept has, later on, been adapted to rechargeable lithium and lithium-ion batteries, and tremendous amounts of effort have been devoted to understanding the SEI in such battery systems.46, 110 Specifi-cally, the formation mechanism, morphology, chemical composition and influence on battery electrochemistry are of general interest. Even so, the nature of the SEI has still not been fully revealed due to its complexity as well as the difficulty in its characterization, since the thickness of the SEI layer is usually only a few Å to a couple of nanometers.110, 112, 113 Moreover, certain species in the SEI are sensitive to air and therefore require an air-free environment for sample preparation, transfer as well as measurement in ana-lytical studies.115 Surface sensitive characterization techniques, for instance X-ray photoelectron spectroscopy (XPS), atomic force microscopy (AFM), time of flight-secondary ion mass spectroscopy (TOF-SIMS), are frequently used to investigate SEI layers.110, 116–120 In addition, more conventional meth-ods such as scanning electron microscopy (SEM), transmission electron mi-croscopy (TEM), Fourier transform infrared spectroscopy (FT-IR), Raman spectroscopy, etc., are also widely utilized.46, 110, 113, 121

Nevertheless, the fact that the SEI formed on graphite electrodes general-ly contains both organic and inorganic species, which are degradation prod-ucts of the salt and solvents in the electrolyte, has been widely accepted.46, 110 For instance, the SEI layer formed on graphite (using LiPF6 in EC-based electrolyte) is usually composed of polycarbonates and polyethylene oxide, as well as inorganic compounds such as Li2O, Li2CO3, LiF, lithium alkox-ides and lithium alkyl carbonates.116, 122, 123

The SEI layer has significant influence on the electrochemical perfor-mance of lithium-ion batteries. Despite the fact that it is beneficial for the battery in terms of preventing continuous electrolyte degradation, the for-mation of the SEI consumes a given amount of Li-inventory in the battery since most of the SEI components contain Li+. The formation processes also involve, to a large extent, electrolyte reductions which consume charge (electrons) during the reactions. Therefore, SEI formation normally results in irreversible capacity loss, especially during the first electrochemical cycle. Moreover, this passive layer strongly affects the battery performance in terms of capacity retention, rate capability, etc.110, 112, 113

The emergence of high capacity alloying anode materials such as Si and Sn raises further challenges for the stability of the SEI layer formed on these

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materials.26 This is primarily due to the nature of these materials, which ex-perience extreme volume changes during electrochemical cycling, thereby resulting in continuous electrolyte degradation and spontaneous SEI for-mation. Such detrimental effects result in fast fading of cycling performance, which is generally observed for silicon-based anode materials. The SEI layer derived from conventional electrolytes (LiPF6 as salt, EC-based) on silicon anodes is found to be similar to that on graphite, and is mainly composed of LiF, lithium alkoxides, polyethylene oxides and carbonates such as Li2CO3, lithium alkyl carbonates and polycarbonates.34, 121, 124–126 On the other hand, such an SEI is not stable enough to accommodate large volume changes during cycling of silicon materials, and continuous electrolyte degradation is generally observed as a result.26, 121, 125 Among all strategies to improve the performance of silicon anodes (introduced in Chapter 1), SEI-forming addi-tives such as FEC and VC are found to be particularly effective by directly enhancing the stability of the SEI layer.88, 96–98, 100

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5. Methodology

5.1 Battery preparation All cells investigated in this thesis work were assembled in two-electrode pouch cells, unless otherwise stated. A picture of such a pouch cell and a schematic drawing of the key battery components in the cell are shown in Figure 7a and b, respectively. Each cell was composed of two composite electrodes with a porous separator in between, which was pre-soaked with liquid electrolyte. All these components were placed into a laminated alumi-num foil bag with two additional metal tabs attached to each electrode, and then vacuum-sealed into pouch cells. More specifically, both tabs were nick-el foil for silicon/Li cells (Paper III, IV and V) whereas for NMC/Li and NMC/graphite cells (Paper I and II), one tab was aluminum (connected to NMC) and the other was nickel (connected to Li or graphite). An external current was applied/withdrawn from the two metal tabs.

Figure 7. (a) A picture of a pouch cell used in this thesis work. (b) Schematic repre-sentation of the components in a pouch cell.

Composite electrodes were prepared by casting a slurry onto a metal current collector using the doctor-blading technique. The electrode slurry was a mix-ture of active material, binder, conductive additive and liquid solvent. These different components and their compositions in silicon, NMC and graphite electrodes, respectively, are listed in Table 3. Copper foil was used as cur-rent collector for silicon and graphite electrodes, and carbon coated alumi-num foil was used for NMC electrodes. Circular electrodes were punched out after 12 h pre-drying at 60 °C, and were then further dried at 120 °C for 12 h inside an argon-filled glovebox (O2 < 2 ppm, H2O < 1 ppm).

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Table 3. A summary of the compositions of the electrodes used in this thesis work.

Electrodes Active material Binder Conductive additive

Silicon Silicon nanoparticle CMC-Na Carbon black SuperP

NMC LiNi0.33Mn0.33Co0.33O2 Kynar 2801 Carbon black SuperP

Graphite Graphite CMC-Na, SBR* Carbon black SuperP

* Weight ratio of CMC-Na:SBR=1:1 The electrolytes used in this thesis work can be categorized into two clas-

ses, one using LiPF6 as the primary lithium salt (Paper I, II and III) and the other instead using LiTDI (Paper IV and V). More specifically, the baseline LiPF6-based electrolyte was LP40 (1 M LiPF6, EC/DEC, 1/1 wt/wt, BASF) and the LiTDI-based electrolyte was 0.6 M LiTDI, EC/DMC,1/2 (v/v). Based on these, electrolytes consisting of additional additives were also in-vestigated, and it should be noted that all additives (LiTDI in Paper I and II, and FEC in Paper III, FEC and VC in Paper V) were introduced in weight percentage.

5.2 Galvanostatic battery cycling Galvanostatic cycling was the electrochemical characterization technique predominately used for the evaluation of battery performance in this thesis work. During a galvanostatic test, a constant current is applied/withdrawn to/from the outer circuit which is connected to a cell, and voltage response is monitored as a function of time. Figure 8 shows such voltage profiles of a NMC/Li cell at the 1st and 100th cycles. The current is terminated once the test reaches certain pre-set conditions, such as voltage or time. The details of different cut-off settings which have been used can be found in each of the corresponding publications. For instance, the NMC/Li cell shown in Figure 8 was stopped at a constant voltage of 4.3 V for the charging processes, whereas discharging was stopped at 3.0 V.

The magnitude of the applied current (I, (mA)) is calculated based on the mass (m, (g)) and theoretical capacity (Qtheo, (mAh/g)) of the active material, as well as the desired charging/discharging rate, as described in Equation 1. The cycling rate is commonly referred as the C-rate (C/n, in which, n (h) corresponds to the time required for a full charge or discharge). Alternative-ly, current density (j, (mA/g)) has also been widely used instead of current, and can be calculated from Equation 2.

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I= Qtheo×m n (1)

j= I m = Qtheo n (2)

Figure 8. Voltage profiles of a NMC/Li cell at the first and 100th cycle. The cycling was performed at 1C rate and 55 °C.

Critical characteristics of a lithium-ion battery can be obtained from a gal-vanostatic experiment. Practical capacity, which is a measure of the actual amount of charge that can be stored in a certain amount of the material, can be calculated from the current and charging/discharging time (Equation 3). The NMC/Li cell shown in Figure 8 was cycled at a current density of 150 mA/g (corresponding to 1C rate, based on a theoretical capacity of 150 mAh/g).

Q=j × t (3)

Coulombic efficiency (C.E.) is another important parameter for evaluating battery performance. It essentially describes the reversibility of battery cy-cling. C.E. is the ratio of the output capacity (Qout) by a battery to the input capacity (Qin), as expressed in Equation 4.

C.E.=Qout Qin ×100% (4)

The input process is referred to the initial electrochemical process of a cy-cling test, and the output is therefore the subsequent one. For instance, the first cycle of the NMC/Li cell (Figure 8) shows a low C.E. of 89.7%, which means that 10.3% of the initial charge was lost. Seidlmayer et al. demon-strated that this is primarily due to kinetic hindrance in the NMC 111 materi-

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al when the fully intercalated state is approached, as the fully lithiated state can only be achieved with an extended voltage hold at low potentials.127

5.3 Photoelectron spectroscopy The photoelectron spectroscopy (PES) technique is based on the photoelec-tric effect, which describes the process of photoelectron emission when a material is irradiated by light. More specifically, when the used light is in the X-ray region, this technique is widely known as X-ray photoelectron spec-troscopy (XPS) or electron spectroscopy for chemical analysis (ESCA). Dur-ing a XPS experiment, kinetic energies of emitted photoelectrons and the amount of electrons as a function of the kinetic energy are examined by an electron analyzer. A schematic representation of this working principle of XPS is illustrated in Figure 9a.

Figure 9. (a) A schematic presentation of the working principle of the PES and (b) an overview spectrum of a pristine silicon composite electrode.

The binding energies (EB), which can be calculated from the measured kinet-ic energies with Equation 5, are rather defined for the photoelectrons from specific atomic core levels, meaning that PES is able to identify the ele-mental composition at the surface of the sample:

EB = hν − EK − φ (5)

where EK is the kinetic energy of the emitted photoelectron and j is the work function, which is dependent on both the spectrometer and the sample. As an example, a survey spectrum of a pristine silicon composite electrode is demonstrated in Figure 9b, which shows the presence of C, O and Si on the sample. Furthermore, binding energies of the same core level electrons may vary depending on the specific chemical states, known as the ''chemical shift''. The inset in Figure 9b shows the Si 2p spectrum, which consists of contributions from the elemental Si (Si-Si) and the native oxide species (Si-Ox).

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The inelastic electron mean free path (IMFP, λ) describes how far an electron can travel in a solid before losing energy. In this work, the probing depth of PES is referred to 3λ, as this includes 95% of the total elastically emitted photoelectrons. Due to the very limited mean free path for electrons in solid materials, the probing depth of XPS is generally rather limited. Ow-ing to these particular properties, the highly surface sensitive PES is one of the most widely used techniques for investigating the chemical composition of the thin SEI layer in lithium-ion batteries.

Figure 10. Universal curve of the electron IMFP as a function of the kinetic energy of the electron.

In addition, the IMFP, and therefore the probing depth of a certain core lev-el, is dependent on the kinetic energy of electrons, or in other words, the excitation energy as the binding energy and work function are rather defined for one core level measurement. It is well described by the universal curve of the IMFP (shown in Figure 10) that it increases when the electron energy increases. In conventional XPS experiments, the kinetic energies of the pho-toelectrons are generally higher than 100 eV. Conventional, also known as ‘in-house’ XPS, is limited to certain excitation sources, such as AlK𝛼 and MgK𝛼, which means that the probing depth is fixed for each core level. Syn-chrotron facilities, however, provide the possibilities of varying the probing depth in PES experiments by altering the excitation energy. In this thesis work, PES measurements have been carried out using three set-ups (listed in Table 4). The most surface-sensitive measurements were carried out at the I411 beamline at MAX-lab with low photon energy (soft) X-ray radiation. At the KMC-1 beamline, BESSY, bulk properties were examined as excita-tion energies up to 12 keV are available.

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Table 4. A summary of the three PES setups used in this thesis work.

Setups Excitation energy

Beamline I411 at MAX-lab, Lund, Sweden 50 – 1500 eV

In-house ESCA AlKα (1487 eV)

Beamline KMC-1 at BESSY, Berlin, Germany 2000 – 12000 eV

For all PES measurements, samples were washed with DMC three times prior to measurements in order to remove electrolyte residues. Air-tight transfer chambers were used for all sample transfers to avoid air-exposure. The obtained spectra were curve fitted with Gaussian-dominated (< 10% Lorentzian) peaks using the Igor Pro software. All spectra were energy cali-brated with respect to the hydrocarbon (C-C/C-H) C 1s peak at 285 eV.

5.4 Other characterizations

Thermogravimetric analysis Thermogravimetric analysis (TGA) measures the mass variation of a sample as a function of temperature during, e.g. a linear temperature ramp in a cer-tain atmosphere. Therefore, TGA is widely utilized to examine the thermal stability of a material. More specifically, in Paper I, TGA was used to quan-tify the water content in hydrated LiTDI as well as its thermal stability. The experiments were performed under a N2 gas flow.

Fourier transform infrared spectroscopy (FTIR) Fourier transform infrared spectroscopy (FTIR) is a widely used method for identifying vibrational modes of functional groups in molecules. In a FTIR experiment, the sample is irradiated with infrared light, and the absorption spectrum is recorded as a function of wavenumber. In Paper I, FTIR with an attenuated total reflectance attachment (ART-FTIR) was used to identify LiTDI and also trace the evolution of its signature vibrational modes during air exposure.

Nuclear magnetic resonance spectroscopy Nuclear magnetic resonance (NMR) spectroscopy is an indispensable tech-nique for characterizing molecule structures. NMR spectroscopy relies on the NMR phenomenon, in which nuclei absorb and re-emit electromagnetic radiation (usually radio frequency pulses) in a magnetic field. Such nuclei, referred to as NMR-active nuclei, should have an overall spin quantum

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number that is non-zero, such as 1/2, 1, 3/2, etc. The most common NMR active nuclei are, for instance 1H, 13C, 19F, 31P, 29Si. In this thesis work, Pa-per I in particular, 19F and 31P NMR spectroscopy was applied to evaluate the decomposition of LiPF6-based electrolytes.

Scanning electron microscopy In a scanning electron microscopy (SEM) experiment, the sample is irradiat-ed by a high energy and sharply focused electron-beam in high vacuum, while simultaneously emitted signals such as secondary electrons and backscattered electrons, are recorded. This technique is extensively used to examine the surface morphology of a specimen. In this thesis work, SEM was used to study the morphological characteristics of the electrodes after electrochemical cycling with different electrolytes.

Electrochemical impedance spectroscopy In an electrochemical impedance spectroscopy (EIS) experiment, an electro-chemical cell is subjected to an alternating voltage (usually of a small mag-nitude, e.g. 10 mV) at various frequencies, and the current response is rec-orded and analyzed. At a given frequency ω, the obtained results can be rep-resented as in Equation 6:

Z(ω) = Zre-jZim (6)

Where Zre and Zim are the real and imaginary parts of the impedance Z(w), and j=√-1. More specifically, Zre and Zim correspond to the ohmic resistance and the capacitive contribution, respectively. In this thesis work, EIS results were presented in Nyquist plots, which display Zim vs. Zre for different values of w.

EIS results of conventional lithium-ion batteries are, in general, compli-cated and difficult to interpret. Nevertheless, it is still rather straightforward to compare the total impedance between different cells according to the magnitude of Z. Furthermore, the Randles equivalent circuit model is a widely used method to distinguish different contributions from the battery components. This approach was adopted in Paper I to differentiate and es-timate the charge transfer resistances at NMC/electrolyte and Li/electrolyte interfaces.

X-ray diffraction X-ray diffraction (XRD) is an extremely powerful tool for studying the crys-tal structure of a material. The condition under which the X-ray diffraction phenomenon occurs can be described using Bragg’s law (Equation 7):

nλ = 2d sinθ (7)

where l is the wavelength of the incident X-ray, d is the inter-atomic dis-tance between a set of lattice planes, and 2q is the angle between the incident

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and scattered X-ray beam. In a XRD experiment, the sample is irradiated with an X-ray beam and the scattered radiation is recorded as a function of q. The inter-planar distance can be calculated using Bragg’s law, and the crys-tal structure can be determined. XRD was applied in Paper II to trace the evolution of the NMC unit cell during electrochemical cycling. In addition, the obtained unit cell parameters were used to estimate the Li+ content in the post-cycled NMC material.

Cyclic voltammetry In a cyclic voltammetry experiment, the potential of the working electrode is linearly scanned at a certain rate, and the current response is recorded as a function of the potential. The peak current and peak potential in the obtained cyclic voltammogram can be used to evaluate the reaction thermodynamics, kinetics, etc. In Paper III, CV experiments were conducted to compare the difference in reduction potential of the electrolyte with and without the FEC additive.

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6. Improving conventional LiPF6-based electrolyte with additives

6.1 LiTDI as a highly effective electrolyte additive

6.1.1 Moisture scavenging effects LiTDI has previously been studied as an alternative lithium salt to LiPF6, where LiTDI has displayed higher stability against water.65 Here, the effects of exposing dehydrated LiTDI to air have been investigated. Figure 11 shows the characteristic vibrational mode evolution of LiTDI and H2O after different air exposure times (relative humidity of ∼75%). The three regions are attributed to the nitrile stretching mode [n(C≡N)], the H2O bending mode [n2(H-O-H)], and the N-C-N bending mode [in the imidazole ring, nu(N-C-N)], respectively, as annotated in Figure 11.

Figure 11. FTIR spectra of dehydrated LiTDI and air-exposed LiTDI at various exposure times. The LiTDI･H2O was obtained by drying hydrated LiTDI at 80 °C in vacuum for 2 hours. Reproduced with permission from Paper I. Copyright 2017 American Chemical Society.

The spectrum of the dehydrated LiTDI shows no noticeable signal in the water bending region, indicating that it was sufficiently dried. After the an-

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hydrous LiTDI was exposed to air for 5 minutes, one new peak appears at 1635 cm−1, which corresponds to a H2O bending signal. This suggests that LiTDI is able to absorb water once it is in contact with air. In the meantime, this hydration process causes peak shifts in both the nitrile and N-C-N vibra-tional modes. Furthermore, as the exposure time progresses to 40 min, a different water uptake process starts to occur, as indicated by emerging sig-nals in the C N stretching and H2O bending regions. Interestingly, no no-ticeable change is observed from the nitrogen ring-related bending signal. These results suggest that there are at least two distinct water absorption (WA) processes for LiTDI. The rst process (WA-1) occurs during the first 20 min of air exposure, involving both the nitrile and the imidazole ring groups. The second process (WA-2) starts later and is mainly involving the nitrile groups.

Figure 12. TGA results of the hydrated LiTDI, the LiTDI solid after air-exposure for 5 days, and that dried at 80 °C afterwards. Reproduced with permission from Paper I. Copyright 2017 American Chemical Society.

This mechanism is supported by TGA results, which reveal that there are also two distinct water removal (WR) steps from the hydrated LiTDI solid (shown in Figure 12). The first occurs below 75 °C (WR-1, 14.5 wt% loss) and the other between 75 and 140 °C (WR-2, 7.2 wt% loss). In addition, the molar ratio between these two types of absorbed water and the LiTDI is close to 2:1:1, indicating that the two LiTDI hydrates are monohydrate LiTDI·H2O and trihydrate LiTDI·3H2O. Based on these findings, LiTDI was studied as a moisture scavenging additive, with the aim to improve the sta-bility of LiPF6-based electrolytes.

The poor stability of LiPF6-based electrolytes against moisture is illus-trated by the severe degradation of the water contaminated LP40 electrolyte,

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as shown in Figure 13. In this experiment, 2000 ppm (in weight) of H2O was added to the conventional LP40 electrolyte, which is transparent and color-less in its initial stage (Figure 13a). However, after being stored for 100 days, the color of the water-enriched LP40 electrolyte turned to dark brown, indicating that the electrolyte underwent rather severe degradation (Figure 13b). 31P NMR characterizations were performed on the fresh LP40 and the H2O enriched equivalents after storage for 35 days; the corresponding spec-tra are shown in Figure 13d. The 31P spectrum of the fresh LP40 contains only a strong septuplet signal centered at -143.5 ppm (sept, 1JP−F = 711 Hz), which is attributed to the PF6

− anion. After H2O enrichment and storage for 35 days, two new signals appeared. More specifically, the new triplet cen-tered at -15.3 ppm (1JP−F = 949 Hz), and the doublet signal centered at -7.2 ppm (1JP−F = 900 Hz) are corresponding to OPF2OR and OPF(OR)2 type compounds, respectively. Such compounds are commonly observed degrada-tion products from LiPF6-based electrolytes. According to quantification analysis of the result, these degradation products account for 25% of the total phosphorus species.

Figure 13. (a) Images of the fresh LP40 and LP40 with 2% LiTDI. (b) Images of LP40 with 2000 ppm of H2O and LP40 with 2% LiTDI and 2000 ppm of H2O after storage for 100 days. (c) Images of LP40 and LP40 with 2% LiTDI after storage at 55 °C for 35 days. (d) 31P NMR spectra of the fresh LP40 and LP40 stored for 35 days with 2000 ppm of H2O. (e) 31P NMR spectra of the fresh LP40 with 2% LiTDI and LP40 stored for 35 days with 2% LiTDI and 2000 ppm of H2O. Reproduced with permission from Paper I. Copyright 2017 American Chemical Society.

Moreover, a similar phenomenon of color change was also observed for the electrolyte stored at an elevated temperature of 55 °C. The color turned to dark brown for the LP40 electrolyte which was stored at 55 °C for 30 days, as shown in Figure 13c.

With the presence of 2 wt% LiTDI in the LP40 electrolyte, however, both the water enriched and 55 °C stored electrolytes showed negligible color change, and remained light yellow. The 31P NMR spectra in Figure 13d

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further confirmed that no LiPF6 hydrolysis species were produced after stor-age for 35 days, even though the electrolyte was contaminated by 2000 ppm of H2O. It should be noted that 2000 ppm of H2O and 2 wt% of LiTDI are approximately 1:1 in molar ratio, indicating that the moisture scavenging effect of LiTDI is highly effective. In addition, the coordination between such absorbed water molecule to LiTDI is presumably similar to the struc-tures of the LiTDI·H2O monohydrate.

6.1.2 Cycling performance The previous section demonstrated that LiTDI can be used as a highly effec-tive electrolyte additive for improving the chemical stability of conventional LiPF6-based electrolytes. Here, the effects of the LiTDI additive on the elec-trochemical performance of NMC/Li and NMC/graphite cells at an elevated temperature of 55 °C are discussed.

Electrochemical performance of NMC/Li cells

Figure 14. (a) Discharge capacity retention of the NMC/Li cells with LP40, 1% LiTDI and 2% LiTDI electrolytes. (b and c) Voltage profiles of the NMC/Li cells using LP40 and 1% LiTDI electrolytes, respectively. Reproduced with permission from Paper I. Copyright 2017 American Chemical Society.

Three electrolytes, as-received LP40, LP40 with 1% LiTDI (denoted 1% LiTDI), and LP40 with 2% LiTDI (denoted 2% LiTDI), were studied in NMC/Li cells at 1C rate and 55 °C, and the obtained electrochemical per-

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formances are presented in Figure 14. Figure 14a shows that the LP40 cell exhibited a continuous and significantly faster capacity decay as compared to the 1% LiTDI and 2% LiTDI cells. 1% LiTDI additive seems to be suffi-cient to truly boost the cycling performance of NMC/Li cells at 55 °C. Fur-thermore, without the LiTDI additive, charge and discharge voltage pro les changed drastically during long-term battery cycling, as shown in Figure 14b. More specifically, the voltage profile of the charging process is gradu-ally shifting to higher voltages, whereas the discharging slope is shifting to lower values. Such evolutions indicate a continuous increase in voltage hys-teresis and in cell impedance. When 1% LiTDI was introduced into the elec-trolyte (Figure 14c) much less change in the charge and discharge voltage profiles could be seen. Table 5. Equivalent circuit fitting results for the cells after 250 electrochemical cy-cles using LP40, 1% LiTDI and 2% LiTDI electrolytes, respectively. Reprinted with permission from Paper I. Copyright 2017 American Chemical Society.

LP40 1% LiTDI 2% LiTDI

RLE (Ω) 4.4 3.6 4.1

Rct,Li (Ω) 103.2 96.1 72.6

Rct,NMC (Ω) 322.6 163.2 195.8

Sum (Ω) 430.2 262.9 272.6

Equivalent circuit fitting was carried out from the EIS results, and the ob-tained results are summarized in Table 5. It should be mentioned that the assignment of the charge transfer resistances to the Li/electrolyte (Rct,Li) and NMC/electrolyte (Rct,NMC) interfaces is based on EIS experiments in three-electrodes cell setups, which used NMC as the working electrode (WE), and Li as both the reference (RE) and counter electrodes (CE). RLE is primarily attributed to the ohmic resistance of the liquid electrolyte, but also electrical contact resistance. The fitting results convincingly show that the total im-pedance of the LP40 cell is much higher than those of the two LiTDI-containing cells. These large differences can predominantly be attributed to the charge transfer resistance at the NMC/electrolyte interface (Rct,NMC), which is 322.6 for the LP40 cell, but 163.2 and 195.8 for the 1% LiTDI and 2% LiTDI cells, respectively. These findings strongly indicate that the LiTDI plays a significant role in the electrode/electrolyte interfacial chemis-tries, especially at the NMC/electrolyte interface in NMC/Li cells.

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Electrochemical performance of NMC/graphite cells The effects of the LiTDI additives are investigated in NMC/graphite cells in Paper II. NMC/graphite cells, which used LP40 and 1% LiTDI electrolytes, respectively, were evaluated at 55 °C and C/2 rate, and electrochemical re-sults are presented in Figure 15. The NMC/graphite cell using the as-received LP40 electrolyte exhibited a rather poor cycle life with gradually decreasing capacity. After 30 cycles, a catastrophic fading occurred. An improved cycling performance was achieved by the 1% LiTDI cell, showing more than 80 cycles with rather stable capacity retention. The gradual capac-ity decrease is believed to be primarily attributed to the loss of active lithium caused by parasitic reactions such as electrolyte decomposition and SEI for-mation. The following sudden capacity loss, however, is predominantly caused by cell impedance growth.128, 129

Figure 15. (a) Discharge capacity retention and coulombic efficiency of NMC/graphite cells using LP40 and 1% LiTDI electrolytes at 55 °C, respectively. The capacity is calculated on the NMC material. (b) Voltage profiles of the NMC/graphite cells using LP40 and 1% LiTDI electrolytes at 30th cycle.

Although the capacity retention was improved with the LiTDI additive, the coulombic efficiencies during the first 20 cycles were found to be lower than those of the LP40 cells (shown in Figure 15a). This implies that electro-chemically irreversible reactions occurred to a larger extent when LiTDI was present in the electrolyte. Such an observation is consistent with previous reports which showed that batteries using LiTDI-based electrolytes without any additional SEI-forming additives generally have rather large irreversible capacity loss within the initial cycles.68–72 Shkrob et al. proposed that LiTDI undergoes de-fluorination reactions at reductive potentials, resulting in a SEI layer which shows adverse effects on battery cycling.70 The coulombic effi-ciency of the 1% LiTDI cell increased gradually, and reached a maximum of ~99.6% by the 50th cycle. Figure 15b presents voltage profiles of the two NMC/graphite cells for the 30th cycle and surprisingly these two profiles are dramatically different, although the capacities of the two cells (shown in Figure 15a) are rather similar. Specifically, the LiTDI cell showed a consid-erably lower voltage hysteresis and a shorter constant voltage step as com-

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pared to the LP40 cell, indicating a lower cell impedance. Such behavior is highly similar to that of the NMC/Li cells, which also showed a prolonged cycle life and suppressed voltage hysteresis increase when the LiTDI addi-tive was applied.

EIS characterization in three-electrode cells were also performed to eval-uate the difference between the cells with and without the LiTDI additive, as well as to identify the contributions from the two electrode/electrolyte inter-faces, i.e. NMC/electrolyte and graphite/electrolyte. The obtained EIS spec-tra of the LP40 and 1% LiTDI cells after 20 cycles are shown in Figure 16a and b, respectively.

Figure 16. Nyquist representation EIS spectra of NMC/Li/graphite cells with LP40 (a) and 1% LiTDI (b) electrolytes after 20 cycles, respectively. Experiments were performed using three-electrode cells, which were composed of NMC as the work-ing electrode, Li as the reference electrode and graphite as the counter electrode.

Overall, dramatically higher cell impedance was found for the LP40 cell compared to that for the 1% LiTDI cell. Moreover, the spectra of the full cell impedance are highly similar to that obtained only from the NMC cathode, whereas the contribution from the graphite anode is negligible (details of impedance contributions from graphite are presented in Figure S3 in Paper II). This indicates that the impedances of both full cells are dominated by the cathodes, which is in good agreement with the results from NMC/Li cells. Liu et al. also reported previously that the cathode dominated the cell im-pedance growth in NMC/graphite cells after cycling at 55 °C.130

6.1.3 The graphite/electrolyte interface The morphology of graphite electrodes from NMC/graphite cells with LP40 and 1% LiTDI electrolytes after 100 cycles was examined by SEM, and the obtained images are shown in Figure 17. The surface of the graphite elec-trode from the LP40 cell (Figure 17a) is found to be rather inhomogeneous.

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In particular, two distinct areas, one brighter than the other, can be noticed and are highlighted in the image. Figure 17b shows that in the brighter area, graphite particles are unevenly covered by spherical and toroidal shaped grains (µm sized), as well as larger fragments which are likely agglomerates of these grains. Such newly formed substances can be attributed to decom-position products of the electrolyte, and the rather large amount is presuma-bly due to the elevated cycling temperature which stimulates electrolyte deg-radation. On the other hand, the darker area shows rather distinct graphite particles without grains or agglomerates on top, as displayed in Figure 17c.

Figure 17. SEM images of graphite electrodes after cycling using LP40 (a-c) and 1% LiTDI (d, e) electrolytes, respectively. Figures a and d are in low magnification, and figures b, c and e are in high magnification.

The morphology of the graphite electrode after cycling is drastically affected by the incorporation of the LiTDI additive. Figure 17d shows that the graph-ite electrode of the LiTDI cell is rather homogenous over a large area. In addition, the formed SEI layer is also rather uniformly covering the graphite

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as well as carbon black particles, which are all clearly distinguishable as shown in Figure 17d.

If comparing the two graphite anodes with and without the LiTDI addi-tive, dramatic differences are not only observed in the surface morphology, but also in the chemical compositions of the formed SEI layers. Atomic con-centrations of various elements on the two graphite electrodes are calculated from in-house XPS measurements, and summarized in Figure 18. It worth mentioning that no metal signals of Ni, Mn and Co are detected on any of the two graphite electrodes.

Figure 18. Atomic concentrations of various elements on the graphite electrodes from NMC/graphite cells with LP40 and 1% LiTDI electrolytes after 100 cycles.

With 1% LiTDI additive in the electrolyte, the amount of fluorine is approx-imately twice as much as that on the LP40 graphite electrode, whereas the oxygen concentration is considerably lowered. Since the oxygen species in the SEI layer are commonly attributed to the reduction products of carbonate solvents, the smaller amount of oxygen suggests that solvent degradation is mitigated when LiTDI is present in the electrolyte. This is most likely due to the improved chemical stability of the electrolyte as discussed in Section 6.1, or formation of a more protective SEI layer with the LiTDI additive. Interestingly, for the graphite electrode in the 1% LiTDI cell, the molar ratio of nitrogen to phosphorus is found to be 1:1, whereas this ratio in the bulk 1% LiTDI electrolyte is approximately 1:5. This means that the graphite surface is rather rich in the amount of nitrogen as compared to that of phos-phorus, suggesting that LiTDI, and/or its derivatives, are significantly more concentrated on the graphite electrode surface as compared to LiPF6. These findings are therefore strong indications of LiTDI participating in the SEI formation on graphite electrodes.

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Figure 19. (a) F 1s, (b) N 1s and (c) C 1s XPS spectra of graphite electrodes from NMC/graphite cells with LP40 and 1% LiTDI electrolytes after 100 cycles, respec-tively.

High resolution F 1s, N 1s and C 1s spectra of the graphite electrodes from the NMC/graphite cells with LP40 and 1% LiTDI electrolytes, respectively, are presented in Figure 19. F 1s spectra show that the relative amounts of different fluorine containing species are rather similar, and both graphite electrodes are dominated by LiF. However, as the atomic concentration of fluorine on the graphite electrode from the 1% LiTDI cell is almost twice of that in the LP40 sample (Figure 18), the absolute amount of LiF is therefore much higher on the 1% LiTDI sample as compared to the LP40 sample. As discussed above, Shkrob et al. previously proposed that LiTDI undergoes a one- or two-step de-fluorination reactions at reductive potentials based on density functional theory (DFT) calculations and XPS experiments.68, 70, 71 The findings in this study are therefore in good agreement with this proposed mechanism. Figure 19b shows that the N 1s spectrum of the 1% LiTDI graphite surface is composed of two peaks, which are located at ~401 and ~399 eV. The latter one is attributed to the nitrogen in the intact imidazolate ring or nitrile.71 The other N 1s signal at ~401 eV is therefore attributed to LiTDI-derived compounds which contain, possibly, ammonium cations (de-noted as C-N+ in Figure 19b) according to the binding energy.131–133 C 1s spectra in Figure 19c further supports the formation of such C-N species, which have a signal located at ~286.6 eV.

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6.1.4 The NMC/electrolyte interface In contrast to large differences between the graphite electrodes, surface mor-phology and surface atomic concentration of the two NMC electrodes dis-play high similarities (shown in Figure S8 and Figure 5 in Paper II, respec-tively).

Figure 20. (a) F 1s, (b) N 1s and (c) C 1s spectra of the NMC electrodes from the NMC/graphite with LP40 and 1% LiTDI electrolytes after 100 cycles, respectively.

Nevertheless, selected core level XPS spectra (F 1s, N 1s and C 1s) from the two NMC electrodes are presented in Figure 20. Figure 20a shows that the F 1s spectra of the two NMC electrodes are rather similar. Since the atomic concentrations of fluorine are almost identical, the total amounts of LiF on the two NMC electrodes should therefore also be similar. This is therefore in contrast to the finding at the graphite anode, which shows a substantially higher amount of LiF with the presence of LiTDI. These results suggest that the de-fluorination reaction of LiTDI mainly occurs on the graphite surface at low potentials. Interestingly, similar nitrogen-containing species are also found on the NMC electrode surface, as indicated by the N 1s and C 1s spec-tra shown in Figure 20b and c. This is presumably because LiTDI undergoes different decomposition routes on the NMC cathode, most likely oxidation reactions, although the nitrogen products are similar in composition as those on the graphite anode. It is also possible that LiTDI is first reduced and de-fluorinated on graphite, and the produced radicals or their derivatives diffuse to the NMC electrode and simultaneously form a surface layer. The decom-position mechanisms of LiTDI at different electrode surfaces should be thor-oughly studied further in the future.

To briefly summarize, it has, in this section, been described that the LiTDI additive is able to prevent bulk electrolyte degradation. In particular, it strongly influences the chemical composition and morphology of the SEI layer on the graphite anode. Moreover, decomposition products of LiTDI are

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also detected on the surface of NMC electrode, indicating that LiTDI has effects on the NMC/electrolyte interfacial chemistry as well.

6.2 FEC as SEI-forming additive for silicon anodes This section describes the use of FEC as a SEI-forming additive in conven-tional LiPF6-based electrolytes for silicon anodes. The effects on electro-chemical performance, morphology and SEI composition are evaluated and discussed. The presented results in this section are summarized from Paper III.

6.1.1 Electrochemistry and surface morphology

Figure 21. (a) Gravimetric capacity and coulombic efficiency of silicon/Li cells with and without FEC. (b) Nyquist plot representations of the EIS spectra of the sili-con/Li cells with and without FEC after 85 cycles. Reproduced with permission from Paper III. Copyright 2016 American Chemical Society.

Silicon/Li cells, which contained LP40 and the LP40 with 10 wt% FEC elec-trolytes (denoted as FEC/LP40), respectively, were investigated via gal-vanostatic cycling. The capacity retention and coulombic efficiencies of these two cells are presented in Figure 21a. The silicon/Li cell with conven-tional LP40 electrolyte showed continuous and fast capacity decay. A rather low specific capacity of ~780 mAh/g was obtained by the end of the electro-chemical testing, i.e. after 85 cycles. Moreover, the coulombic efficiency was at a rather low level of ~97% after 50 cycles. Such poor performance is essentially attributed to the large volume change of silicon during cycling. In addition, the SEI layer derived from this baseline electrolyte LP40 is not sufficient to accommodate such large volume variation, so that electrolyte degradation continuously occurs as the cycling progresses.26, 121, 125 However, with the presence of 10 wt% of FEC in the electrolyte, both capacity reten-

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tion and coulombic efficiency of the silicon/Li cell were much improved as compared to those of the LP40 cell. More specifically, no significant capaci-ty loss was observed, and the cell achieved a coulombic efficiency of ~99% for all 85 cycles. Furthermore, the EIS spectra of two cells after 85 cycles, presented in Figure 21b, show that the FEC/LP40 cell had lower cell imped-ance as compared to the LP40 cell. Such findings are in good agreement with the less pronounced capacity decay when FEC was incorporated into the LP40 electrolyte.

Figure 22. SEM images of the silicon electrode with LP40 (a, b) and FEC/LP40 (c, d) electrolytes after 85 cycles, respectively. Reproduced with permission from Paper III. Copyright 2016 American Chemical Society.

Surface morphologies of the silicon electrodes after 85 cycles with LP40 (Figure 22a, b) and FEC/LP40 electrolytes (Figure 22c, d) were character-ized using SEM technique. Numerous cracks of various sizes were formed on the silicon electrode after cycling with the conventional LP40 electrolyte, whereas the FEC/LP40 cycled electrode retained a smooth surface, as shown in the low magnification images (Figure 22a, c). Such large cracks may lead to the breakdown of the electronic conducting network within the electrode, resulting in increased cell impedance. Significant differences between these two electrodes are also observed at high magnification. In particular, surface voids on the LP40 sample were severely blocked by the newly formed SEI layer, which unevenly covered the electrode surface as well as individual silicon particles. This may cause difficulties for the silicon particles to access electrolyte, and the amount of electrochemically active silicon material is

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therefore decreased. By adding 10% FEC into the LP40 electrolyte, this is-sue was successfully mitigated. The surface morphology was more well pre-served after extended cycling, showing a conformal SEI layer coating, which is most likely derived from the FEC additive.

6.1.2 SEI composition In order to further understand the effects of FEC additive, the chemical com-positions of the SEI layers after 85 cycles using either LP40 or FEC/LP40 electrolytes were thoroughly examined with PES. The atomic concentrations of in the two silicon electrode surfaces were found to be significantly differ-ent. The overview spectrum of the FEC/LP40 silicon electrode shows a pre-dominating F 1s signal (located at ~690 eV), while the spectrum of the LP40 sample contains a strong O 1s peak (at ~530 eV), as presented in Figure 23a. More specifically, the F 1s spectrum of the FEC/LP40 sample, shown in Figure 23b, reveals a dominating contribution from LiF. F 1s signals of the LP40 cycled silicon electrode consist of LiPF6 and also LiF, but in com-parable intensities. Since the only difference between these two systems was the 10 wt% FEC additive in the electrolyte, the additional amount of LiF should be generated from the decomposition of FEC. Indeed, such observa-tions have also been reported previously, and de-fluorination and formation of LiF are recognized as the key signatures of the FEC additive.96, 97

Figure 23. Survey (a) and F 1s spectra (b) of the silicon electrodes with FEC/LP40 and LP40 electrolytes after 85 cycles, respectively. The excitation energy was 2005 eV. Reproduced with permission from Paper III. Copyright 2016 American Chemical Society.

Besides fluorine, carbon is another essential element for understanding the SEI layer as the majority of SEI components contain carbon. C 1s spectra of the FEC/LP40 and LP40 cycled silicon electrodes, obtained from measure-ments at photon energies of 430, 2005, and 6015 eV, are presented in Figure 24. First of all, the characteristic signal of carbon black (CB, ~282.8 eV in Figure 24), which is commonly used as a reference for estimating the SEI

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layer thickness, is only clearly visible in the 6015 eV experiments for the FEC/LP40 silicon electrode. A similar result is also found for the LP40 sam-ple. This indicates that both SEI layers produced were already quite thick after 85 cycles. Various other carbon-containing species were also formed. The SEI layers on both silicon electrodes were found to be rather homoge-nous in terms of the carbon-containing species, across the three depths, since no signals emerge or disappear apart from the carbon black signal. More specifically, such carbon species are hydrocarbon (C-C/C-H at 285 eV), C-O (~286.8 eV), O-C-O (~288.1 eV), OCO2 (~289.6 eV) and in particular, poly-carbonate compounds derived from FEC (290.8 eV, highlighted in red). In Paper I, this FEC derivative is attributed to -CHF-OCO2- type species, but it also has been proposed that FEC forms polyacetylene or poly-VC type of compounds on silicon anodes.96, 97

Figure 24. C 1s spectra of the silicon electrodes with FEC/LP40 and LP40 electro-lytes after 85 cycles, measured with photon energies of 430 (a), 2005 (b) and 6015 eV (c), respectively. Reproduced with permission from Paper III. Copyright 2016 American Chemical Society.

In contrast, for the LP40 samples without additive, the peak with the highest binding energy is located at 290 eV, which is commonly attributed to con-ventional carbonate species (Li2CO3 and lithium alkyl carbonates). Moreo-ver, the other C 1s peaks from the FEC/LP40 samples are slightly different in binding energy positions as compared to those in the LP40 case. This in-dicates that these compounds are of different chemical compositions; most likely due to that degraded FEC segments take part in the formation process.

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To shortly summarize, the chemical composition and morphology of the SEI layer on silicon electrodes, with and without the presence of the FEC addi-tive, are found to be rather different as illustrated in Figure 25. In particular, incorporation of FEC increases the amount of LiF, and also contributes to the formation of different carbon-containing compounds on silicon. This SEI layer homogeneously covers the silicon particles, and is strong enough to prevent formation of large cracks. With the standard LP40 electrolyte, the SEI layer on the silicon electrode is instead dominated by organic-type spe-cies, and unevenly distributed on the electrode.

Figure 25. Schematic representations of the SEI formation on silicon anodes with FEC/LP40 (a) and LP40 (b) electrolytes, respectively. Reprinted with permission from Paper III. Copyright 2016 American Chemical Society

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7. Developing LiTDI-based electrolytes for silicon electrodes

In Chapter 6, the electrolyte stabilization additive LiTDI and the SEI-forming additive FEC were discussed. Both studies have shown that incorpo-rating additives is an effective while still simple approach to improve today's conventional LiPF6-based electrolytes. Alternatively, replacing LiPF6 with other, chemically more stable, lithium salts is also a highly attractive path to achieve better electrolyte performance. In Chapter 7, LiTDI is studied as the primary lithium salt in non-aqueous electrolytes for silicon-based batteries based on results in Paper IV and V.

7.1 Electrochemical performance The electrochemical performance of LiTDI-based electrolytes, with and without SEI-forming additives, were examined in silicon/Li cells, and the electrochemical results are shown in Figure 26. Although LiTDI has been demonstrated to be stable towards moisture and heat, the pure LiTDI-based electrolyte exhibited rather poor compatibility with the silicon anode. The delithiation capacity during the first cycle was ~600 mAh/g, which corre-sponds to a very low coulombic efficiency of 50%. This indicates that 50% of the charge was not recovered, which is most likely due to significant elec-trolyte decomposition on the silicon anodes. Furthermore, Figure 26a shows that such a silicon/Li cell can only be cycled for 62 cycles, during which the coulombic efficiencies were also rather low, at a level of 96-98%.

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The well-known SEI-forming additives FEC and VC were incorporated with the purpose of improving the interfacial stability of the LiTDI-based electro-lyte for silicon anodes. The corresponding cycling performance is summa-rized in Figure 26c and d. The delithiation capacity of the first cycle was then increased to more than 800 mAh/g, and the coulombic efficiency was also improved to 68%. More importantly, long term cycling performance of the silicon anode was dramatically enhanced. The silicon/Li battery with FEC and VC additives maintained high delithiation capacities, as well as a high coulombic efficiency of 99% during extended cycling. This demon-strated that LiTDI-based electrolytes can be comparable to the state-of-the-art LiPF6-based electrolytes with additives.

Figure 26. Cycling performance of silicon/Li cells with LiTDI baseline electrolyte (a, b), and the electrolyte contained additives (c, d), respectively. The baseline electrolyte is 0.6 M LiTDI EC/DMC (2/1, v/v). Reproduced with permission from Paper V. Copyright 2016 American Chemical Society.

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7.2 Silicon/electrolyte interfacial chemistries

7.2.1 Initial SEI formation on silicon with pure LiTDI-based electrolyte

Figure 27a shows the potential profile during the first cycle of a silicon/Li cell with the baseline LiTDI electrolyte. Similar to the results presented in Section 7.1, this cell chemistry showed a poor reversibility in the electro-chemical performance. The specific capacities of the lithiation and delithia-tion processes are 2700 and 1500 mAh/g, respectively, which corresponds to a low coulombic efficiency of 55.6%.

Figure 27. (a) Potential profile of a silicon/Li cell with the baseline LiTDI electro-lyte at the first cycle. (b) F 1s and (c) C 1s spectra of the silicon anodes at various state of charge in the first cycle. Reproduced from paper IV, Copyright 2016, with permission from Elsevier.

Five representative silicon samples at various states of charge were studied further with PES. The potentials of these samples are highlighted in Figure 27a, and their corresponding F 1s and C 1s spectra are shown in Figure 27b and c, respectively. The OCV sample, originating from a silicon electrode soaked in a silicon/Li cell but without any electrochemical testing, mainly shows the characteristics of a LiTDI baseline electrolyte and the as-prepared silicon electrode. After being reduced to 0.7 V, a drastic increase of the con-tribution from LiF is found in the F 1s spectrum, compared to that of the OCV sample. This finding is in good agreement with the observation from the NMC/graphite cell which contains LiTDI as an additive in the electrolyte (Section 6.1.3). Both studies suggest that LiTDI undergoes de-fluorination at

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low potentials, thereby producing LiF. Distinct evolutions in the C 1s spectra are observed when the silicon electrode is further reduced to 0.13 V. In par-ticular, the contributions from the C-O/C-N and OCO2 species are much increased, whereas the carbon black signal becomes almost invisible. This indicates that the electrode surface is, to a large extent, covered by electro-lyte decomposition products such as carbonates and C-O/C-N species. As the cycling continues to 0.01 V and 0.9 V, little significant changes are observed in the F 1s and C 1s spectra. This shows that the LiTDI baseline electrolyte undergoes rather severe degradation on silicon electrodes at reductive poten-tials, and that decomposition products are found to be originated from both the LiTDI salt and carbonate solvents.

7.2.2 Surface chemistry on silicon with additives As described in Section 7.1, the cycling performance of silicon/Li can be greatly improved by introducing the SEI forming additives FEC and VC to the LiTDI baseline electrolyte. Figure 28 summarizes the atomic concentra-tions of two silicon electrodes which were cycled using the LiTDI baseline electrolyte with and without these additives. The sample with additives shows higher amounts of fluorine and lithium, whereas the oxygen, nitrogen and carbon contents are lower, compared to the additive-free one. This is highly similar to the observations on silicon electrodes using LP40 electro-lyte with and without FEC additive (Paper III), which also revealed an in-creased amount of fluorine at the electrode surface when the FEC additive was used. The decreased amount of oxygen compounds suggests a sup-pressed degradation of carbonate solvents as compared to the baseline elec-trolyte.

Figure 28. Atomic concentrations of various elements of silicon electrodes with and without FEC and VC additives after 50 cycles. The results are calculated from the in-house XPS measurements. Reproduced with permission from Paper V. Copyright 2016 American Chemical Society.

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Figure 29. (a) F 1s and (b) C 1s spectra of silicon electrodes used LiTDI baseline electrolyte with FEC and VC additives after 1, 25 and 50 cycles, respectively. (c) F 1s and (d) C 1s spectra of silicon electrode with LiTDI baseline electrolyte after 50 cycles. Reproduced with permission from Paper V. Copyright 2016 American Chem-ical Society.

Furthermore, evolution of more detailed chemical compositions of the SEI on silicon are illustrated by the F 1s and C 1s spectra, which are shown in Figure 29a and b, respectively. No major differences are observed in F 1s or C 1s spectra when comparing the silicon electrodes after 1, 25 and 50 cycles. In contrast to the F 1s spectra of the silicon anode after 50 cycles with the electrolyte containing additives, the sample from the pure LiTDI baseline electrolyte system consisted of considerably less LiF. The additive free sam-ple showed rather comparable amounts of LiF and C-F species, which can be primarily attributed to LiTDI and its derivatives. The additional LiF can most likely be attributed to the presence of FEC, which is also known for producing LiF-rich SEI layers on silicon anodes, as also concluded from Paper III. Moreover, the carbon-containing species produced in these two battery chemistries were also found to be different. More specifically, char-acteristic carbonate species derived from the FEC and VC additives were only found on the silicon electrodes which originated from the cells contain-ing such additives. For the electrode with baseline electrolyte, the carbonate species are presumably attributed to conventional Li2CO3 or lithium alkyl carbonates.

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The major limitation of the most widely used LiPF6 salt is its poor chemical stability, and the subsequent formation of the highly corrosive HF. Philippe et al. reported that after long term storage or cycling in conventional LiPF6-based electrolyte, the silicon oxide surface transformed to SiOxFy, which is a strong indication of surface corrosion by HF.128 LiTDI has been proven to be highly stable towards water, and is therefore unlikely to form HF during battery storage or cycling. Here, Si 1s spectra are analyzed to examine the stability of the oxide layer after extended cycling using the LiTDI baseline electrolyte with FEC and VC additives.

Figure 30 presents the Si 1s spectra of the pristine silicon electrodes and the electrodes after 1 and 100 cycles, respectively. The pristine electrode shows a dominating signal at 1839.6 eV, which is attributed to bulk silicon. In addition, the peak at 1844.2 eV (highlighted in red) corresponds to the SiO2 surface layer. After the first electrochemical cycle, a shift to lower binding energy is observed for both these two signals. This is presumably indicating certain degrees of lithiation of these compounds, and the for-mation of lithium silicate and LixSi, as illustrated in Figure 30. More im-portantly, no SiOxFy was found on the electrode surface even after extended cycling to 100 cycles.

Figure 30. Si 1s spectra of the pristine silicon electrode, and the silicon electrodes after 1 and 50 cycles using LiTDI baseline electrolyte with FEC and VC additives. Reproduced with permission from Paper V. Copyright 2016 American Chemical Society.

These results show that LiTDI-based electrolytes are much less likely to generate HF during cycling compared to the conventional LiPF6-based elec-trolyte. This could be of particular importance for certain cathode materials,

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such as LiMn2O4 and NMC materials, which are known for problems with metal dissolution.

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Conclusions and outlook

This thesis focused on developing non-aqueous electrolytes for lithium and lithium-ion batteries, as well as understanding the effects on electrochemis-try and interfacial chemistries in these cells. The first part of the work, Chapter 6, extensively discussed LiTDI and FEC additives for the widely used LiPF6-based electrolytes. Chapter 7 summarized the efforts in develop-ing LiTDI-based electrolytes for silicon anodes. The present thesis also showed that synchrotron radiation PES spectroscopy techniques can provide valuable information on the chemical composition of electrode surfaces.

Improving conventional LiPF6-based electrolytes Two electrolyte additives, i.e. LiTDI and FEC, were investigated individual-ly in this part of the thesis work. LiTDI was found to be beneficial for the chemical stability of LiPF6-based electrolytes. FEC, on the other hand, was utilized as an additive for enhancing the interfacial stability of silicon anodes with conventional LiPF6 based electrolyte.

The poor chemical stability of conventional LiPF6-based electrolyte at el-evated temperatures or in the presence of moisture still remains an unsolved issue for today’s lithium-ion batteries. LiTDI is found to be able to strongly absorb the moisture in the electrolyte, and therefore greatly suppress the hydrolysis of LiPF6. With the much improved chemical stability of the elec-trolyte, NMC/Li and NMC/graphite cells exhibited an enhanced cycling performance at elevated temperatures. In addition, LiTDI also influences the interfacial chemistries at both the graphite/electrolyte and NMC/electrolyte interfaces. The increase in impedance at the NMC/electrolyte interface was, in particular, mitigated with the presence of LiTDI additive.

Another issue of the LiPF6-based electrolyte is the poor interfacial com-patibility with silicon anode materials. As a SEI-forming additive, FEC de-composes at silicon/electrolyte interfaces and simultaneously forms a stable SEI layer on the surface of the silicon anode. The SEI layer consisted of a high content of LiF, FEC-derived poly-carbonates as well as other species from electrolyte decomposition. This layer homogeneously covers the silicon electrode surface, and successfully mitigates electrolyte decomposition as compared to FEC-free electrolytes. Consequently, significantly improved cycling performance of silicon/Li cells was achieved with the addition of FEC, accompanied with a slow increase in cell impedance.

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Developing LiTDI-based electrolytes for silicon anodes Since the conventional LiPF6-based electrolyte, especially the salt, is still far from the perfect electrolyte, developing alternative electrolytes with alterna-tive salts is another approach to improve the performance of lithium-ion batteries. LiTDI-based electrolytes were therefore studied in silicon/Li cells. However, such a baseline electrolyte, i.e. without any additives, showed rather poor compatibility with silicon anodes. The low coulombic efficiency in the first electrochemical cycle is a strong indication of severe electrolyte decomposition. Photoelectron spectroscopy revealed that the formed SEI layer was composed of decomposition products of both LiTDI and carbonate solvents. In particular, PES studies of the electrode with the presence of LiTDI, either utilized as additive or primary salt in the electrolyte, showed a LiF-rich surface layer on anodes.

SEI forming additives of FEC and VC were introduced to the LiTDI baseline electrolyte, and much improved cycling stability was observed in the silicon/Li cell. This is primarily attributed to the more stable SEI layer derived from the additives as compared to that formed from the decomposi-tion of the baseline electrolyte. This finding suggests that such additives are also applicable to electrolytes using salts other than LiPF6. In addition, no SiOxFy species were found on the surface of silicon electrodes after extended cycling, indicating that LiTDI-based electrolytes are much less likely to pro-duce the highly corrosive compound HF during battery operation.

Outlook This thesis illustrated that introducing functional electrolyte additives is a highly effective and simple approach to enhance battery performance. In addition, better insight into the interfacial chemistry in lithium-ion batteries is of great importance to further understand battery behavior. The FEC addi-tive is particularly important for electrode materials that require highly stable SEI layers, such as silicon. The moisture scavenging LiTDI additive, on the other hand, can be used for batteries in elevated temperature applications, or conventional cells so that the efforts required to dehydrate battery compo-nents can be reduced. LiTDI-based electrolytes can be attractive for cell chemistries that require high stability of the lithium salt, and use high poten-tial anode materials such as LTO.

Moreover, photoelectron spectroscopy is a powerful tool for studying in-terfacial chemistries in lithium-ion batteries. With synchrotron-based radia-tion, non-destructive depth profiling of the chemical composition of surface layer on electrodes can be realized.

Along with all the findings, there are also many unaddressed issues and potential research topics worthy of further investigations. Two interesting examples in particular are highlighted here:

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1) Stability issues of the electrolytes containing FEC Kim et al. reported that the color of electrolytes which contained both FEC and LiPF6 changed dramatically to black after storage at elevated tempera-tures, and such degraded electrolytes contained various acids including the highly corrosive HF.110 In experiments carried out in this PhD thesis work, a rather fast change in color was also observed from the LP40 electrolyte with 10% FEC after storage at room temperature in glovebox. A better under-standing of the degradation process as well as finding effective approaches to prevent such degradation are particularly important for battery chemistries using FEC.

2) Poor coulombic efficiencies during initial cycles of NMC/graphite cells with the LiTDI additive

The present thesis work (Chapter 7) as well as a previous study by Paillet et al. have revealed that LiTDI-based electrolytes generally have rather poor compatibility with low working potential anode materials such as graphite and silicon.72 Indeed, Figure 15a in Section 6.2.2 showed that the cou-lombic efficiency during the initial cycles of NMC/graphite cells with the LiTDI additive are also rather low. This would lead to a larger consumption of the active lithium inventory in practical lithium-ion batteries. Additional cooperating SEI-forming additives along with small amounts of LiTDI in conventional LiPF6-based electrolyte could be a viable approach to address this issue.

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Sammanfattning på svenska

Dagens samhälle är i stor utsträckning beroende av icke förnybara fossila bränslen i form av olja, naturgas och kol. Det är bara en tidsfråga innan dessa bränslen är fullständigt uttömda på vår planet. Dessutom bidrar den massiva användningen av fossila bränslen till en produktion av enorma mängder CO2, vilken är en betydande växthusgas som redan bidragit till globala klimatförändringar. Att hitta alternativa och förnybara energikällor är därför en extremt angelägen och brådskande uppgift. På senare år har mängden energi producerad från exempelvis solceller och vindkraft ökat markant. Till detta kommer också det faktum att transportsektorn, som do-minerar konsumtionen av olja, också behöver laddningsbar teknik för energi-lagring för att till fullo ersätta den stora användningen av fossila bränslen.

Figur I. Schematisk bild av ett litium-jonbatteri.

Återuppladdningsbara litium-jonbatterier, som under urladdning kan använ-das för att driva ett fordon och som också kan laddas många gånger med bibehållen hög energitäthet, är perfekta för elfordon. Livslängden hos ett litium-jonbatteri är ett väsentligt kriterium för denna sorts tillämpningar, och livslängden är synnerligen beroende av stabiliteten hos elektrolyten och yt-kemin i batteriet. Elektrolyten är en nyckelkomponent i batteriet, där dess huvudsakliga funktion är transporten av Li+-joner mellan de två aktiva elek-troderna. De vanligast förekommande elektrolyterna i dagens högpresterande litium-jonbatterier är organiska elektrolyter, som framför allt består av ett

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eller ett par litiumsalter och organiska lösningsmedel. Gränsskiktet mellan elektrolyten och elektroden är det område där de elektrokemiska reaktioner-na sker, och därför är dess egenskaper avgörande för batteriets prestanda. Dessa två materialkemiska områden, det vill säga elektrolyter och deras mot-svarande ytkemi i litium-jonbatterier, är huvudsakligt fokus i den här dok-torsavhandlingen, och visas översiktligt i Figur I. Mer specifikt har två fa-miljer av organiska elektrolyter studerats, där den avgörande skillnaden varit vilket litiumsalt som använts: antingen litiumhexafluorofosfat (LiPF6) eller litium-2-trifuorometyl-4,5-dicyanoimidazol (LiTDI). Att förbättra dagens konventionella elektrolyter LiPF6-baserade elektrolyter är det vanligaste valet för dagens litiumbatterier. Dessa elektrolyter lider emellertid av dålig stabilitet gentemot fukt eller för-höjda temperaturer. För att förbättra stabiliteten tillsattes små mängder med LiTDI, och det upptäcktes att LiTDI kunde fånga in vatten och därmed för-hindra nedbrytningen av LiPF6. Figur II visar att den konventionella elektro-lyten brutits ned rejält efter förvaring vid 55 °C, vilket syns från att färgen förändrats från färglös till brun. Elektrolyten med tillsats av LiTDI förblir å andra sidan oförändrad. Battericyklingsegenskaperna för elektrolyten för-bättrades drastiskt med tillsatsen av LiTDI. Dessutom visade forskningen att även mycket små mängder LiTDI hade en betydande påverkan på de före-ningar som bildades vid både anodens och katodens gränsskikt.

Figure II. (a) Foto på elektrolyter bestående av LP40 (vänster), och LP40 med 2wt% LiTDI (höger). (b) Foto på LP40, och LP40 med 2wt% LiTDI efter förvaring vid 55 °C i 30 dagar. LP40 innehåller saltet LiPF6.

Ytterligare ett problem med konventionella flytande elektrolyter är deras dåliga stabilitet vid gränsskikten gentemot anoder av kisel. Elektrolyterna genomgår en kontinuerlig nedbrytning på kiselytan, vilket leder till en snabb försämring av batteriets prestanda. För att motverka detta studerades kemi-kalien FEC som en tillsatts för att skapa bättre och stabilare gränsskikt. Med FEC kunde de kiselbaserade batterierna uppvisa mycket goda cyklingsegen-skaper där kapaciteten förblev hög, liksom den coulombiska effektiviteten. Utforskning av alternativ LiTDI studerades också som det huvudsakliga litiumsaltet i elektrolyten, för att fullständigt kunna ersätta LiPF6, som har förhållandevis dålig kemisk stabilitet. LiTDI-baserade elektrolyter undersöktes därför – för första gången

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– gentemot kiselanoder. Kiselbaserade celler med denna sorts elektrolyter uppvisade mindre god prestanda beroende på omfattande nedbrytning av elektrolyten. Därför introducerades gränsskiktsbildande additiv som FEC och VC i den LiTDI-baserade elektrolyten, och batteriprestandan förbättra-des påtagligt. Detta är främst beroende på att gränsskiktet mellan elektrod och elektrolyt stabiliseras i närvaro av sådana tillsatser. I den här avhand-lingen visas betydelsen av elektrolytens sammansättning för stabiliteten hos gränsskikten i litium-jon batterier och hur detta påverkar batteriets funktion.

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Acknowledgements

First and foremost, I would like to thank my main supervisor, Torbjörn Gus-tafsson, for giving me the opportunity to work on this project, and all the support throughout the Ph.D study. Thank you for always trusting me as an independent researcher and giving me large freedom in academic research. I really enjoyed all these years working with you.

I would like to thank my co-supervisors, Kristina Edström and Daniel Brandell, for all the fruitful discussions and helpful suggestions. I would also like to thank Dominique Guyomard and Bernard Lestriez for all your help during my stay in Nantes. It was truly a fantastic experience. Special thanks to Maria Hahlin for introducing me to the fascinating world of XPS. I would like to acknowledge Leif Nyholm for sharing your knowledge of electro-chemistry. Thanks to Fredrik Björefors for the great teaching experience I had in your course. I would like to thank Zhaohui Wang for always enlight-ening me with new ideas and sharing your knowledge.

I would like to also thank my talented collaborators: Fredrik Lindgren, Zhaohui, Stéven Renault, Bertrand Philippe, Bing Sun, and Julia Maibach. Your creative ideas and meticulous experimental work carried me and all these work would never be possible without your contribution. Special thanks to Fredrik, I learned a lot from your creative way of thinking. I would also like to acknowledge Henrik Eriksson for your support in our lab.

I would like to acknowledge all the ‘‘synchrotron team’’ members for bringing fun as well as being supportive during the synchrotron beam times: Maria, Julia, Reza Younesi, Bertrand, Fredrik L., Tim Nordh, Katarzyna Ciosek Högström, Sara Malmgren, Burak Aktekin and Erik Björklund.

It is my great pleasure to share our office with you, Andreas Blidberg and Tim. I would like to thank you both for all the fun discussions on scientific issues, life outside of work as well as helping me out with Swedish. Thanks a lot to Jia Liu for always sharing positive attitudes and happiness.

Many thanks to my colleagues and friends at the Ångström Laboratory: Adam, Alina, Anti, Bing Cai, Cesar, Charlotte, Chenjuan, David, Fabian, Fang, Gabi, Girma, Habtom, Jeff, Jonas Mindemark, Linus, Mario, Matt, Mahsa, Mohammed, Richard, Ruijun, Sara Munktell, Solveig, Shruti, Shuai-nan and Peng, Stéven, Viktor Nilsson, Viktor Renman, Wei Wei, Yiming, Yuxia, Yu Zhang and Zhen Qiu, for all the enjoyable discussions, coffee breaks and after-works.

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Special thanks to my friends for all the wonderful ‘‘Sanguosha’’ time we had together: Bing C. & Kenneth, Bing S., Chenjuan, Jia, Ruijun & Qiuhong, Yang, Yiming, Yuan and Zhaohui. Thanks to my friends back in China, espe-cially Zheng Qian, for always sharing happiness and cheering me up. Con-gratulations with the arrival of your healthy newborn!

I would like to thank my parents for always trusting and supporting me in every aspect. I would also like to thank my sister and brother in law for bringing laughter and looking after our parents. Last but not least, Bao, thank you for all the happiness and support you brought. My life has never been more enjoyable and I will always be with you. I love you all.

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