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Rechargeable Lithium-Ion Batteries
Rechargeable lithium-ion batteries (LIBs) have become widely used power sources not only for portable devices but also for electrical vehicles. Differently from primary batteries, the electrochemical reaction is reversible, therefore the cells can be recharged several times and the delivered energy can be stored (6).
Traditional LIBs contain four main components: positive electrode, negative electrode, non-aqueous electrolyte and separator. Transition metal oxides or phosphate based materials are generally used as positive electrodes (LiCoO2, LiNi1/3Co1/3Mn1/3O2, LiMn2O4 and LiFePO4), graphite is commonly used as negative electrode, non-aqueous electrolytes contain lithium salts (LiPF6, LiClO4, LiTFSI, etc.) and carbonate based solvents (ethylene carbonate, propylene carbonate, etc.) and separators are made of polymers (polyethylene, polypropylene). The electrolyte is a medium that can only allow ion transport, whereas it is an insulator for electron transport (6).
Li-ion Battery Working Principle and Characteristics
A typical LIB works following the principle that upon charging of the battery, Li ions are extracted from the positive electrode, move inside the electrolyte and insert into the negative electrode (Figure I – 1). During lithium extraction, oxidation processes take place at the positive electrode and electrons flow through the external circuit from positive to negative electrode, where reduction processes happen. Thus, electricity is produced by the electrochemical reaction inside the cell. The inverse process happens during the discharge of the battery.
When the electrochemical cell is charging, oxidation process at positive electrode (i.e., LiMO2) can be represented by LiMO2 → Li1-xMO2 + xLi+ + xe- eq. I – 1
While, reduction process at negative electrode (i.e., graphite) can be represented by C6 + xLi+ + xe- → LixC6
The overall reaction is the sum of two half-cell reactions, eq. I – 2 LiMO2 + C6 ⇄ Li1-xMO2 + LixC6 eq. I – 3
The free energy of the system, which is the driving force of the redox processes, can be expressed based on the assumption that cell is at equilibrium, ∆ = . . ℎ eq. I – 4 where n is number of electrons exchanged, F is Faraday constant (96485 C.mol-1) and ℎ is the theoretical voltage between electrodes. At equilibrium state, there is no current flow, therefore open circuit voltage ( ) is approximately equal to theoretical voltage. During charge and discharge, current evolves; therefore, working voltage, which is the actual operating voltage of the battery, differs from the open circuit voltage.
Among rechargeable storage systems, Li-ion battery has the highest specific energy, which is the product of the capacity and average discharge voltage. To maximize the specific energy of the cell, active materials should provide large discharge capacity and/or can be cycled at high discharge potentials. The theoretical capacity can be defined as the charge (or lithium ions) storage ability of active materials, which is calculated by ℎ = . eq. I – 5 3600.
where Qth is theoretical capacity (mAh.g-1), n is the number of electrons (or lithium ions) exchanged during charge-discharge, F is Faraday constant (96485 C.mol-1) and Mw is the molecular mass of active material (g.mol-1). The capacity of the active material is strongly affected by the C-rate (1/h), which can be defined as a measure of the current rate at which a battery is discharging and can be represented by = ℎ ( ℎ ) = eq. I – 6 where i is current in A, theoretical capacity is expressed in Ah. For example, a cell charging at a C/10 rate needs 10h to reach its full capacity.
For industrial applications, Li-ion cells should have long cycle life, which is usually defined by the number of charge-discharge cycles that the cell can offer before its capacity falls below 80% of its nominal capacity.
To develop new high performance and affordable Li-ion batteries, some important parameters must be considered, such as
1. Selection of widely available and low-cost precursors for the active materials
2. Development of new high-energy materials with cost effective and scalable production methods
3. Improvement on battery cycle (shelf) life
4. Understanding of battery safety and failure mechanisms
5. Discovery of easy recycling methods
Positive Materials for Li-ion Batteries
In 1980s, TiS2 was reported as the first intercalating compound by Whittingham (7) et al. after it demonstrated the ability to intercalate and de-intercalate lithium ions reversibly. Though TiS2 shows excellent cyclability, its average potential is rather low (2.1 V vs Li+/Li), consequently the energy density is not sufficient enough for high energy applications (8). Thereafter, study of oxide type compounds gains interest because they could operate at relatively high potentials (> 4 V). These efforts results in the discovery of LiCoO2, which was reported by Goodenough (9) et al. and later commercialized by Sony in lithium- ion cells in 1991. However, a maximum of 0.5 Li ions could be deinserted from the layered LiCoO2 without destabilizing the structure; therefore the practical capacity is only half of its theoretical capacity (274 mAh.g-1) (9). Increasing cobalt prices in the market is also another disadvantage. Having high theoretical capacity and cheaper price of nickel, LiNiO2 is seen as an alternative material (10,11). The studies show that the reversibility of LiNiO2 is lower than the one of LiCoO2 (12), thus manganese substitution is suggested. Thereafter, the best composition, which is found to be LiNi1/3Co1/3Mn1/3O2 (and called NMC111), is discovered as a solid solution of LiCoO2, LiNiO2 and LiMnO2 compositions (13–16). LiNi1/3Co1/3Mn1/3O2 delivers a capacity up to 160 mAh.g-1 at 0.1 C rate in a potential range of 2.5 V – 4.4 V vs Li+/Li (17). Here, the nickel redox provides high capacity, cobalt improves cycling stability and manganese enhances thermal stability. In parallel to LiNi1/3Co1/3Mn1/3O2, layered LiNi0.80Co0.15Al0.05O2 is also studied (18–20). This material delivers 200 mAh.g-1 at 0.1 C rate in a potential range of 3 V – 4.5 V (21) and Al substitution increases the thermal stability (22).
Moreover, polyanionic LiFePO4 is also studied in the literature (23–25). Though LiFePO4 shows excellent cycling stability, its practical energy density is quite limited.
Having lower cost (<10 $.kg-1) and relatively stable structure compared to layer oxides, spinel LiMn2O4 attracts research interests and was also commercialized by Moli Energy. However, its practical capacity is rather small (~140 mAh.g-1). To utilize the advantages of different chemistries, blend of LiMn2O4 with LiNi1/3Co1/3Mn1/3O2 or LiNi0.80Co0.15Al0.05O2 have also been suggested (26,27). Addition of small amount of layer oxides suppresses the Mn dissolution as well as improves the capacity retention of LiMn2O4 cells (28).
High Energy Positive Materials
To extend the driving ranges of the electrical vehicles, specific energy of current lithium ion batteries should be further increased. The specific energy is related to both the capacity and the operating potential. Common approaches are focused on the use of high capacity materials, for instance Ni-rich NMC (LiNixCoyMnzO2, with x≥0.6) and Li-rich (or Mn rich) NMC [xLi2MnO3. (1-x)LiMO2] or high potential positive materials, such as LNMO (LiNi0.5Mn1.5O4).
• Ni-rich NMC (LiNi0.80Co0.10Mn0.10O2)
The capacity of the typical NMC111 (LiNi1/3Co1/3Mn1/3O2) is reported as 160 mAh.g-1, which is smaller than the theoretical value (275 mAh.g-1) upon cycling at 0.1 C rate in a potential range between 2.5 V – 4.4 V vs Li+/Li (3). Whereas, its Ni-rich version, LiNi0.80Co0.10Mn0.10O2 (NMC811) delivers much higher specific capacity (200 mAh.g-1 ) upon cycling at 0.1 C rate in a potential range between 3 V- 4.3 V (29). Despite the capacity and cost advantages, NMC811 shows less thermal stability, which is related to the structure stability. Upon cycling, oxygen release causes structural variation and thermodynamically stable rock salt LixNi(1-x)O2 structure forms on electrode surface (30). That increases cell resistance and lowers the cycling stability. Surface coating techniques (31), doping (32) and core-shell structural design (33) have been proposed to solve these problems.
• Li-rich (or Mn-rich) NMC [xLi2MnO3. (1-x)LiMO2]
To improve structural stability and specific capacity of NMC, Thackeray et al. (28) suggest to prepare the solid solution of LiMO2 with electrochemically inactive Li2MnO3 composition. The new lithium rich composition delivers unusual charge capacity exceeding 250 mAh.g-1 between 2 V and 4.8 V vs Li+/Li. (34,35). Such high capacity is explained with the presence of cationic (Ni and Co) redox processes up to 4.4 V and anionic (O2) redox processes between 4.4 V and 4.8 V. Unfortunately, the first charge capacity could not be recovered at the end of the discharge and low initial columbic efficiency is observed. That is explained by irreversible oxygen loss (Li2O) from the Li2MnO3 structure upon delithiation. Besides, such oxygen loss also causes structural modifications, layered Li2MnO3 changes into spinel LiMnO2 and voltage decay is observed due to Mn4+ to Mn3+ transitions (36).
• High Voltage Spinel LNMO (LiNi0.5Mn1.5O4)
Spinel type materials are also considered as high energy materials (37–39) due to their high operating potentials (4.7 V) even though their practical capacity is rather small (140 mAh.g-1). Moreover, these materials could be easily produced, they show high rate capability and superior safety behavior. Those properties have already attracted many research interests. Despite such advantages, cycling at high potential causes electrolyte decomposition, that later impedes cycling stability (40). Therefore, new compatible electrolyte must be discovered for further application of these materials. In addition to these materials, recently, Li-rich disordered rocksalt oxides and sulfides have been considered as promising high energy positive materials for lithium-ion batteries (41,42). Their cation disordered rocksalt structure is so robust that less structural problems are expected. Moreover, rocksalt oxides offer capacities exceeding 250 mAh.g-1 at high average potentials (3.6 V), while rocksalt sulfides offer excellent capacities (>400 mAh.g-1) at relatively low average potentials (2.3 V). Their specific energy is considerably higher than current materials. Therefore, many researches have recently been started to develop new Li-rich cation disordered rocksalt compositions.
Cation Disordered Rocksalts as Positive Materials
The first cation disordered rocksalts: stoichiometric compositions
In 1998, Obrovac et al. discovered cation disordered rocksalts during the preparation of layer oxides compositions, LiMO2 (where M= Ni, Co, Mn and Fe) with high energy ball milling (43). Milling of LiCoO2 longer than 16 hours, results in a α-LiFeO2 type (cation-disordered rocksalt) structure rather than a α-NaFeO2 type (cation-ordered rocksalt) structure (Figure I – 2). The same trend is observed for all other compositions (LiNiO2, LiMnO2, LiFeO2 and LiTiO2); longer milling always results in formation of disordered rocksalt phases. They also observe that disordered structures leads to smaller discharge capacities (< 70 mAh.g-1) and large irreversibility at the first cycle (Figure I – 2). The reason of such low capacities is attributed to poor lithium ion conduction in disordered lattice.
Thereafter, Li2MTiO4 (M: Ni, Co, Fe, Mn, V) type disordered rocksalts are reported at the beginning of 2004 (44-49). To eliminate diffusion limitations in disordered rocksalts, new synthesis approaches concentrate on low temperature synthesis methods, such as citric acid assisted sol-gel method (44–46), acid base reactions (47), hydrothermal synthesis (48) and molten salt synthesis (48,49). With these soft chemistry methods coupled with carbon coating strategies, average particle size becomes less than 100 nm and electronic conductivity increases. Thus, much higher practical capacities are reached. For instance, among Li2MTiO4 compositions, Li2NiTiO4 delivers a charge capacity of 181 mAh.g-1, upon cycling between 2.5 V- 4.8 V (Figure I – 3). However, 50% of charge capacity is lost at the end of the discharge and only 25% of discharge capacity is obtained at the end of the 3rd cycle (50). That is explained by irreversible structural modifications upon cycling.
Besides, Li2CoTiO4 shows high reversible capacity of 219 mAh.g-1 at relatively low cycling rate (0.05C), but the large polarization at the first cycle could not be reduced despite the carbon coating strategies (51).
In brief, the practical capacities of stoichiometric rocksalts are still much lower than the theoretical capacities. The reason is attributed to poor lithium diffusion through cation disordered structure. To overcome lithium conduction challenges in cation disordered rocksalts and to understand lithium ion conduction mechanism in these materials, theoretical calculations are required.
Lithium diffusion mechanism through cation disordered rocksalts
Cation disordered rocksalt crystallises in a face centered cubic (fcc) structure with a space group of Fm3̅ . In this fcc structure, cations (lithium or transition metals) statistically occupy 4a positions (½ ½ ½) of octahedral interstitials, whereas anions occupy 4b positions (0 0 0) as shown in the Figure I – 4.
Previous studies describes the lithium conduction mechanism into fcc crystal lattice with a theory of Oh-Th-Oh transition in which lithium ions migrate from one octahedral site (Oh) to another octahedral site with a relaxation at tetrahedral site (Th) (52,53). Moreover, height of the tetrahedron is reported as an important parameter. Tetrahedron height should be tall enough that the energy barrier of transition between Oh-Th sites becomes smaller, thus lithium ions can migrate easily between the sites. This energy barrier changes based on cation distribution around tetrahedral sites (54).
In 2014, A. Urban et al. perform density functional theory calculations to describe lithium conduction mechanism in disordered rocksalts (54,55). They describe that both 1-TM and 2-TM channels are inactive for lithium conduction due to strong repulsion between the cations, whereas 0-TM channels are active for lithium conduction due to low repulsion between lithium ions (or lithium-vacancy). Thus, energy barrier of transition between Oh-Th sites is smaller. Besides, they report that increasing lithium to metal (Li/Me) ratio might create a percolation network through 0-TM channels.
Table of contents :
General Introduction
Chapter I. State of the art
I.1. Rechargeable Lithium-Ion Batteries
I.1.a. Li-ion Battery Working Principle and Characteristics
I.1.b. Positive Materials for Li-ion Batteries
I.1.c. High Energy Positive Materials
I.2. Cation Disordered Rocksalts as Positive Materials
I.2.a. The first cation disordered rocksalts: stoichiometric compositions
I.2.b. Lithium diffusion mechanism through cation disordered rocksalts
I.3. Li-rich Cation Disordered Rocksalt Oxides
I.3.a. Li3NbO4 based cation disordered rocksalt oxides
I.3.b. Li2TiO3 based cation disordered rocksalt oxides
I.3.c. Li4MoO5 based cation disordered rocksalt oxides
I.4. Li-rich Cation Disordered Rocksalt Sulfides
I.4.a. Li2TiS3 based cation disordered rocksalt sulfides
I.4.b. Li3NbS4 based cation disordered rock salt sulfides
I.5. Redox Activity in Disordered Rocksalts
I.6. Conclusion
Chapter II: Experimental Procedures
II.1. Synthesis Methods
II.2. Structural Characterizations
II.3. Electrochemical Characterizations
Chapter III: Development of New Li2.2NiTi0.2Nb0.6O4 Positive Material in Disordered Rocksalt Structure
III.1. Synthesis Improvement
III.1.a. Simple tempering (long firing) and quenching strategies
III.1.b. Effect of washing (type of solvent)
III.1.c. Effect of molten salt to precursor ratio
III.1.d. Molten salt synthesis with a different eutectic salt (LiCl-KCl)
III.1.e. Effect of calcination temperature and duration
III.2. Understanding Studies
III.2.a. Ex Situ X-ray Diffraction Analysis
III.2.b. In Situ X-ray Diffraction Analysis
III.3. Conclusion
Chapter IV: Developing Li-rich Rocksalt Metal Sulfides-Selenides: Li2TiSexS3-x
IV.1. Disordered Li2TiS3 as a Positive Electrode
IV.1.a. Mechanochemical Synthesis of Li2TiS3 Powders
IV.1.b. Electrochemical Performance of Li2TiS3
IV.1.c. Air Sensitive Property of Li2TiS3
IV.2. Selenium Substituted Li2TiS3
IV.2.a. Synthesis Improvements
IV.2.b. Electrochemical performances of Li2TiSexS3-x
IV.3. Conclusion
Chapter V: Characterization of Li2TiSexS3-x Compositions
V.1. Operando Neutron Diffraction of Li2TiS3
V.2. Understanding Degradation: Study of the Structural Property by ex situ XRD
V.2.a. Ex situ XRD and SEM Analyses of Li2TiS3 electrodes
V.2.b. Ex situ XRD and SEM Analyses of Li2TiSexS3-x electrodes
V.3. XPS Analysis of Li2TiSexS3-x Compositions
V.3.a. Methodology: Analysis of XPS Spectra
V.3.b. Survey analysis of Li2TiSexS3-x samples
V.3.c. High Resolution XPS Analysis of Li2TiSexS3-x Powders
V.3.d. XPS Analysis of Li2TiSexS3-x Pristine Electrodes
V.3.e. XPS Analysis on Li2TiSexS3-x Electrodes at Different States of Charge
V.4. Discussion on the XPS Results
V.5. Conclusion
General Conclusion and Perspectives
Bibliography
Appendix
