synthesis and characterization of new positive electrode
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Synthesis and characterization of new positive electrodematerials for Li-ion high energy applications
Yagmur Celasun
To cite this version:Yagmur Celasun. Synthesis and characterization of new positive electrode materials for Li-ion highenergy applications. Chemical and Process Engineering. Université Grenoble Alpes [2020-..], 2020.English. �NNT : 2020GRALI047�. �tel-03080208�
THÈSE
Pour obtenir le grade de
DOCTEUR DE L’UNIVERSITE GRENOBLE ALPES
Spécialité : 2MGE : Matériaux, Mécanique, Génie Civil, Electrochimie
Arrêté ministériel : 25 mai 2016
Présentée par
Yagmur CELASUN
Thèse dirigée par Sébastien MARTINET, Ingénieur de Recherche
Co-encadrée par Jean-François COLIN, Ingénieur de Recherche
et David PERALTA, Ingénieur de Recherche
préparée au sein du Laboratoire LM (CEA/LITEN) Grenoble dans l'École Doctorale I-MEP2 - Ingénierie Matériaux, Mécanique, Environnement, Energétique, Procédés, Production
Synthèse et caractérisation de nouveaux matériaux d'électrode positive pour des applications Li-ion à haute énergie Thèse soutenue publiquement le « 30 Septembre 2020 » devant le jury composé de :
Madame Fannie, ALLOIN Directeur de Recherche, LEPMI, Université Grenoble Alpes, Présidente de jury
Monsieur Sylvain FRANGER Professeur des Universités, ICMMO, Université Paris-Sud, Rapporteur
Monsieur Philippe MOREAU Professeur des Universités, IMN, Université de Nantes, Rapporteur
Monsieur Sébastien MARTINET Ingénieur de recherche, CEA Grenoble LITEN, Directeur de thèse Et des membres invités : Monsieur Jean-François COLIN Ingénieur de recherche, CEA Grenoble LITEN, Invité
Monsieur David PERALTA Ingénieur de recherche, CEA Grenoble LITEN, Invité
Acknowledgments
This thesis study was often exciting, sometimes difficult and rarely stressful for me. I have
always felt that I have been so lucky because I have had a great thesis director and talented
supervisors. I would like to thank my thesis director, Sebastien Martinet, for his valuable
insights and fruitful discussions during my project meetings. I always found him as a very
professional and gentleman person. His vision and attitude also guided me how to become a
charismatic leader. I would also like to thank to my project supervisors, Jean François Colin
and David Peralta for their valuable supports during my thesis. Jean François’s deep
knowledge in crystalline materials and David’s practical synthesis ideas shaped this project as
it is today and produced valuable outcomes. With their supervision, I learned how to continue,
where to stop and when I should push forward. I always admired their respect to individuals,
thinking style and way of criticism. Though I have not seen them before my interview, the
harmony between us was always great. I thank again for their leadership in my project and I
will never forget the time I spent with them.
As like other PhD theses, we found interesting materials, but then struggled in the sea of
unknowns. At that time, Prof. Claude Delmas, who is a precious visitor of CEA LITEN, shared
his rich knowledge with us and showed different aspects of the problems. I would like to thank
him for our fruitful discussions during our team meetings.
When I needed a practical information or opinion about any method or knowledge, I have got
some hints from my lovely office mate, Carole Bourbon. I would like to thank her for her joyful
attitude and the great time that we spent together. I wish to her happy retirement in 2020!
I would also like to thank Anass Benayad, for his significant help for XPS measurements and
discussions on the results, Laurent Gonon for his help for Raman Spectroscopy, Celine
Barchasz and Alice Robba for the discussions about lithium sulfur system. I would also like to
extend my thanks to all my colleagues in LITEN/STB, Sebastien Launois, Vasily Tarnopolskiy,
Ognjen Hajndl, Camille Guerin, Laurent Bernard, Guillaume Lefevre and Guillaume Tonin,
Gaelle Besnard, Damien Guegan, Benjamin Mercier and Oumaima Chaouachi for their
friendship and personal support during my thesis.
2
I would like to thank to Parviz Hajıyev for his valuable support in my personal and
professional life. I never forget our discussions during lunches and coffee breaks in CEA. We
always found crazy ideas to create new materials and also shared the problems that we faced
during the thesis.
I also would like to thank to my friends, Mauro, Alessandro, Taki, Krishna, Nono, Youcef,
Abdel, Tom, Shu Chen and Nora from “Lets Speak French” activity. I could improve my
French speaking skills during our weekly meetings, and I spent wonderful time when we were
having a barbecue, playing football in the grass, skiing Rhone Alpes, watching French movies
and sightseeing small French cities.
At stage of the submitting the thesis manuscript, the stress level increased. Sarp Mamıkoglu
helped me to correct the language of the manuscript at final stage and encouraged me to get
rid of the stress. I would like to thank to him for his support.
I would also like to thank CEA LITEN, especially to our department manager, Sebastien
Patoux, and lab manager, Didier Bloch, for their financial support for this thesis study. With
their support, I could have participated to various conferences, presented my thesis work,
obtained necessary scientific formations for my project and improved my French speaking
skills.
My special thanks go to my family for their encouragements, advices and prays during my
thesis study. With their precious support, I have always believed in myself to achieve the
things that I started, and I could have strived against negativity, depression and
procrastination.
3
Table of Contents
General Introduction .......................................................................................................................... 6
Chapter I. State of the art ................................................................................................................ 10
I.1. Rechargeable Lithium-Ion Batteries ........................................................................................ 10
I.1.a. Li-ion Battery Working Principle and Characteristics ..................................................... 10
I.1.b. Positive Materials for Li-ion Batteries ................................................................................ 13
I.1.c. High Energy Positive Materials ........................................................................................... 14
I.2. Cation Disordered Rocksalts as Positive Materials .............................................................. 16
I.2.a. The first cation disordered rocksalts: stoichiometric compositions ............................... 16
I.2.b. Lithium diffusion mechanism through cation disordered rocksalts ............................. 18
I.3. Li-rich Cation Disordered Rocksalt Oxides ........................................................................... 21
I.3.a. Li3NbO4 based cation disordered rocksalt oxides............................................................. 21
I.3.b. Li2TiO3 based cation disordered rocksalt oxides .............................................................. 24
I.3.c. Li4MoO5 based cation disordered rocksalt oxides ............................................................ 25
I.4. Li-rich Cation Disordered Rocksalt Sulfides ......................................................................... 27
I.4.a. Li2TiS3 based cation disordered rocksalt sulfides ............................................................. 27
I.4.b. Li3NbS4 based cation disordered rock salt sulfides .......................................................... 29
I.5. Redox Activity in Disordered Rocksalts ................................................................................. 30
I.6. Conclusion .................................................................................................................................... 35
Chapter II: Experimental Procedures ............................................................................................. 36
II.1. Synthesis Methods .................................................................................................................. 36
II.2. Structural Characterizations .................................................................................................. 39
II.3. Electrochemical Characterizations ....................................................................................... 51
Chapter III: Development of New Li2.2NiTi0.2Nb0.6O4 Positive Material in Disordered
Rocksalt Structure ............................................................................................................................. 57
III.1. Synthesis Improvement .......................................................................................................... 61
III.1.a. Simple tempering (long firing) and quenching strategies ............................................ 63
III.1.b. Effect of washing (type of solvent) .................................................................................. 66
III.1.c. Effect of molten salt to precursor ratio ............................................................................ 68
III.1.d. Molten salt synthesis with a different eutectic salt (LiCl-KCl) .................................... 70
III.1.e. Effect of calcination temperature and duration ............................................................. 73
4
III.2. Understanding Studies ........................................................................................................... 76
III.2.a. Ex Situ X-ray Diffraction Analysis ................................................................................... 78
III.2.b. In Situ X-ray Diffraction Analysis .................................................................................... 79
III.3. Conclusion................................................................................................................................. 84
Chapter IV: Developing Li-rich Rocksalt Metal Sulfides-Selenides: Li2TiSexS3-x ................ 86
IV.1. Disordered Li2TiS3 as a Positive Electrode .......................................................................... 87
IV.1.a. Mechanochemical Synthesis of Li2TiS3 Powders ............................................................ 87
IV.1.b. Electrochemical Performance of Li2TiS3 .......................................................................... 92
IV.1.c. Air Sensitive Property of Li2TiS3 ....................................................................................... 95
IV.2. Selenium Substituted Li2TiS3 .............................................................................................. 100
IV.2.a. Synthesis Improvements ................................................................................................. 101
IV.2.b. Electrochemical performances of Li2TiSexS3-x ............................................................... 109
IV.3. Conclusion............................................................................................................................... 120
Chapter V: Characterization of Li2TiSexS3-x Compositions ...................................................... 122
V.1. Operando Neutron Diffraction of Li2TiS3 ........................................................................... 122
V.2. Understanding Degradation: Study of the Structural Property by ex situ XRD ......... 126
V.2.a. Ex situ XRD and SEM Analyses of Li2TiS3 electrodes ................................................... 126
V.2.b. Ex situ XRD and SEM Analyses of Li2TiSexS3-x electrodes ............................................ 128
V.3. XPS Analysis of Li2TiSexS3-x Compositions ........................................................................ 132
V.3.a. Methodology: Analysis of XPS Spectra .......................................................................... 132
V.3.b. Survey analysis of Li2TiSexS3-x samples ........................................................................... 134
V.3.c. High Resolution XPS Analysis of Li2TiSexS3-x Powders ................................................ 137
V.3.d. XPS Analysis of Li2TiSexS3-x Pristine Electrodes ............................................................ 141
V.3.e. XPS Analysis on Li2TiSexS3-x Electrodes at Different States of Charge ....................... 144
V.4. Discussion on the XPS Results ............................................................................................. 148
V.5. Conclusion ................................................................................................................................ 150
General Conclusion and Perspectives ......................................................................................... 151
Bibliography ..................................................................................................................................... 156
Appendix ........................................................................................................................................... 170
5
General Introduction
6
General Introduction
In recent years, energy requirements have increased in-line with population growth in
the world. To preserve our resources and to create a sustainable world that is less
dependent on fossil fuels, application of renewable energy technologies has been
expanded. However, sustainability and efficiency of these technologies highly depend
on the success of energy storage systems (1,2).
Moreover, we are now faced with global warming and its result, climate change.
Changing our lifestyle, such as building smart homes and clean transportation are seen
as possible solutions to reduce our carbon footprint. Many automakers have already
started new initiatives, some of them even became actual battery producers (3). Such
attempts are linked to vast investments into battery research and manufacturing,
therefore batteries are becoming a hot topic again, after their first implementation in
cell phones in 1990s.
Figure 0 - 1: Cost breakdown of Li-ion cell (4)
However, electric vehicles are still expensive, since the battery pack is the single major
cost component in an electric vehicle. To reduce electric vehicle costs, US Department
of Energy has set a mission for EV producers that the cost of battery packs should be
reduced to 125 $.kWh-1 until 2022 (3). In order to realize this mission, new battery
chemistries should be discovered, volume of production should be augmented, cell
7
design should be optimized and manufacturing costs should be decreased. In fact, the
most significant contributor to battery costs are raw material costs (50-70%). Among
them, positive electrode is reported to have the highest percentage of costs (27%) (4)
(Figure 0 - 1). NMC [Lithium Nickel Manganese Cobalt Oxide (LiNixMnyCozO2)], NCA
[Lithium Nickel Cobalt Aluminium Oxide (LiNixCoyAlzO2)] are common positive
materials for electrical vehicles, while LCO [Lithium Cobalt Oxide (LiCoO2)] is
common positive electrode for portable applications. Such materials contain
considerable amount of cobalt and the price of cobalt is now an important concern
because of its scarcity, cost and mining conditions in Democratic Republic of Congo.
Thus, research on new cathode material focus on Co-poor or even Co-less materials.
However, it is not so easy since Co is usually needed to stabilize the layered structure,
therefore new chemistries should be discovered.
Figure 0 - 2: Specific energy of materials from material to pack level (5)
As electro-mobility applications have volume, weight and size constraints, a high
energy density battery cell is mandatory (Figure 0 - 2). The positive material (cathode)
is the limiting factor of the battery, so its specific energy has to be increased (5).
Currently there are only a few materials that allow battery systems reaching high
specific energies (250 Wh.kg-1). Among them, we can cite Li-rich NMC, Ni-rich NMC
and high-voltage spinel materials. Despite their high capacity, low cycling stability and
8
safety problems associated to structural deformations are also reported for these
materials. This highlights the importance of discovering new high energy materials.
Recently, disordered rocksalt type positive electrodes appear to be promising because
disordered rocksalt oxides offer capacities exceeding 250 mAh.g-1 at high average
potential (3.6 V vs Li+/Li), while disordered rocksalt sulfides offer excellent capacities
(>400 mAh.g-1) at relatively low average potential (2.3 V vs Li+/Li). Consequently, their
energy density is much higher than today’s commercial materials. This attracts
research interest; thus, many groups are now working on this subject to bring new
materials into the market.
In this work, we study two different disordered rocksalt families: disordered rocksalt
oxides and sulfides materials. We apply different methods to synthesize each family
and optimize our processes to have better electrochemical performances. Then, we
perform various characterizations to understand their electrochemical and structural
properties.
This thesis work is divided in five chapters:
The first chapter introduces Li-ion technology and current high-energy positive
materials briefly. Thereafter, it describes disordered rocksalt positive materials and
their properties in detail. It also explains why stoichiometric compositions failed before
and what Li-rich compositions offer today. Finally, it shows promising electrochemical
properties of Li-rich disordered rocksalts.
The second chapter contains the detail of the synthesis processes, sample preparations
and procedure of electrochemical, morphological and structural characterizations.
The third chapter contains synthesis optimizations to improve electrochemical
performance of disordered rocksalt Li2.2NiTi0.2Nb0.6O4. As the model material is
prepared, ex situ and in situ X-ray diffraction (XRD) studies are performed to
understand the structural evolution upon cycling. Finally, this chapter explains why
electrochemical performances of such materials are limited at the end of the first cycle.
9
The fourth chapter describes preparation of disordered rocksalt Li2TiS3 powders with
mechanochemical synthesis. To solve cycling stability problem of Li2TiS3, selenium
substituted materials are prepared with new synthesis process and their
electrochemical performances are compared with bare Li2TiS3.
The fifth chapter contains structural characterizations, such as operando neutron
diffraction and ex situ XRD to investigate lithium insertion/extraction mechanism in
Li2TiS3 during cycling. To elucidate redox processes in both Li2TiS3 and selenium
substituted materials, powders and electrodes are analyzed with ex situ X-ray
photoelectron spectroscopy (XPS).
In the last part of the thesis, we summarize our findings on disordered rocksalt oxide
and sulfide materials. By referring our experience in those materials, we present a
perspective and share our suggestions for future studies.
10
Chapter I. State of the art
I.1. 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).
I.1.a. 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.
11
Figure I - 1: Rechargeable Li ion battery: transfer of lithium ions from positive
electrode (LixMO2) to negative electrode (LixC6) during charge process
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 eq. I - 2
The overall reaction is the sum of two half-cell reactions,
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,
∆𝐺 = 𝑛. 𝐹. 𝑈𝑡ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 eq. I - 4
where n is number of electrons exchanged, F is Faraday constant (96485 C.mol-1) and
𝑈𝑡ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 is the theoretical voltage between electrodes. At equilibrium state, there is
no current flow, therefore open circuit voltage (𝑈𝑜𝑐𝑣) is approximately equal to
12
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
𝑄𝑡ℎ =
𝑛.𝐹
3600.𝑀𝑤
eq. I - 5
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
𝐶 = 𝑐ℎ𝑎𝑟𝑔𝑒 (𝑜𝑟 𝑑𝑖𝑠𝑐ℎ𝑎𝑟𝑔𝑒) 𝑐𝑢𝑟𝑟𝑒𝑛𝑡
𝑡ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 𝑐𝑎𝑝𝑎𝑐𝑖𝑡𝑦 =
𝑖
𝑄 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.
13
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
I.1.b. 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
14
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).
I.1.c. 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
15
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
16
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.
I.2. Cation Disordered Rocksalts as Positive Materials
I.2.a. 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.
Figure I - 2: X-Ray diffractogram of LiCoO2 and electrochemical performance of LiCoO2 that
is prepared by ball milling (43) at different milling durations
17
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.
Figure I - 3: Electrochemical performance of Li2NiTiO4 electrodes that are cycled at a rate of
C/50 between 4.8 V and 2.5 V vs Li+/Li (50) at the first three cycles
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
18
disordered rocksalts and to understand lithium ion conduction mechanism in these
materials, theoretical calculations are required.
I.2.b. Lithium diffusion mechanism through cation disordered rocksalts
Cation disordered rocksalt crystallises in a face centered cubic (fcc) structure with a
space group of Fm3̅𝑚. 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.
Figure I - 4: Cation disordered rocksalt structure: face centered cubic structure that contains
lithium (Li), transition metal (TM) and oxygen (O) ions
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).
19
Figure I - 5: Cation arrangements around tetrahedral sites: 0-TM, 1-TM and 2-TM channels
Figure I - 5 shows different cation arrangements around tetrahedral sites, for instance
0-TM represents the tetrahedral sites that are coordinated by only lithium ions without
any transition metals (TM), whereas 1-TM and 2-TM represent the tetrahedral sites
that are coordinated by one and two transition metals respectively.
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.
20
Figure I - 6: 0-TM percolation probability (a) and accessible Li content based on Li/Me ratio
(b) in Li-rich disordered rocksalt Li1.211Mo0.467Cr0.3O2 (56)
In another work, they describe the Li/Me threshold as 1.10 to form a percolating
network along 0-TM channels (56). As the material becomes richer in lithium, the
probability for Li+ ions finding an available space in the lattice increases due to the
percolating network formed through the lithium diffusion channels (Figure I - 6). They
also indicate that a new Li-rich rocksalt Li1.211Mo0.467Cr0.3O2 can deliver a reversible
capacity of 300 mAh.g-1 associated to its Li-rich content.
This result perfectly illustrates that:
1. Only reduced capacities are obtained in stoichiometric rocksalts (Li/M=1.0)
because the amount of lithium is not sufficient enough to create percolating
network through lithium diffusion channels.
2. To achieve higher capacities, cation disordered rocksalts should contain excess
lithium amount, hence sufficient lithium conduction paths can form.
21
I.3. Li-rich Cation Disordered Rocksalt Oxides
I.3.a. Li3NbO4 based cation disordered rocksalt oxides
After theoretical calculations reveal the critical Li/Me ratio as 1.10, recent research
interests focus on new Li-rich cation disordered rocksalts that are prepared from
electrochemical inactive ordered rocksalts such as Li2TiO3, Li3NbO4, Li4MoO5 with
transition metal substitution. These materials are not electrochemically active as
cathode materials because they do not have redox active centers, all transition metals
are stabilized at their highest valence state, thus they cannot be oxidized to upper
valence states. For instance, in the case of cation ordered Li3NbO4, which is white
powder and electrochemically inactive, almost no capacity is observed upon cycling
(Figure I - 7).
Figure I - 7: Electrochemical performance of cation ordered Li3NbO4 (57)
Yabuuchi et al. report different overlithiated rocksalts, Li1.3NbxMO2 (where, M: Ni, Fe,
Mn, Co; x=0.3 or 0.43) prepared in disordered rocksalt structure by transition metal
substitution (57). Such materials are synthesized in micronic size and they deliver
smaller capacities (~140 mAh.g-1) based on single electron redox reactions (Me2+/Me3+
or Me3+/Me4+). To increase electronic conductivity, samples are coated with carbon by
mechanical milling, particle size is reduced to submicronic size. Thanks to these
22
improvements, capacities have been increased to 250-300 mAh.g-1, two times higher
than before (57). The compositions are shown in Table I - 1. Among all, Mn doped
Li3NbO4 (Li1.3Nb0.3Mn0.4O2) shows the lowest polarization with a reversible capacity of
260 mAh.g-1 at room temperature (Figure I - 8). Similar composition
(Li1.25Nb0.25Mn0.50O2) is studied by another group as well (58).
Table I - 1: Summary of cation disordered rocksalt oxides reported in the literature
Type Composition T (°C) rate
Initial
charge
capacity
Reversible
Capacity
(mAh.g-1)
Potential
window
(V)
Reference
Li1.30Nb0.30Mn0.40O2 260 260 [57]
Li1.30Nb0.30Fe0.40O2 300 250 [57]
Li1.30Nb0.43Ni0.27O2 280 280 [57]
Li1.30Nb0.43Co0.27O2 310 310 [57]
Li1.30Nb0.30V0.40O2 180 180 4.8 - 1.5 [59]
Li1.30Nb0.30Mn0.40O2 60 10 mA.g-1 380 270 4.8 - 1.5 [57]
Li1.25Nb0.25Mn0.50O2 55 10 mA.g-1 280 290 4.8 - 1.5 [58]
Li2TiO3 - Li2MnO3 Li1.20Ti0.40Mn0.40O2 5 mA.g-1 360 310 4.8 - 1.5 [41]
Li1.20Ti0.40Fe0.40O2 5 mA.g-1 315 210 4.8 - 1.5 [41]
Li1.20Ti0.40Fe0.40O2 10 mA.g-1 350 250 4.8 - 2.0 [60]
Li6/5Ti2/5Mo2/5O2 10 mA.g-1 130 290 4.0 - 1.0 [61]
Li1.20Ni0.333Ti0.333Mo0.133O2 20 mA.g-1 250 230 4.8 - 1.0 [55]
Li3NbO4 - Li3MoO4 Li9/7Nb2/7Mo3/7O2 25 10 mA.g-1 150 280 4.0 - 1.0 [61]
Li3MoO4 Li4/3Mo6+2/9Mo3+
4/9O2 25 10 mA.g-1 180 330 4.0 - 1.0 [61]
Li3MoO4 - LiCrO2 Li1.211Mo0.467Cr0.3O2 25 C/20 280 280 4.3 - 1.5 [56]
Cation Disordered Rocksalt Oxides
50
4.8 - 1.010 mA.g-1
25
Li3NbO4 based
Substitution with
Me3+ and Me2+
Li2TiO3 - Li3MoO4
25
23
Figure I - 8: Electrochemical performances of cation disordered Li1.3Nb0.3Mn0.4O2 cells cycled at
a rate of 10 mA.g-1 between 4.8 V and 1.5 V vs Li+/Li, at 60 °C (57) with SEM image of
Li1.3Nb0.3Mn0.4O2 powder
Despite its high energy density (950 Wh.kg-1 at positive active material level), poor
cyclability is observed at room temperature. When the cells are cycled at 60 °C, inital
charge capacity reaches 380 mAh.g-1 (Figure I - 8). Origin of such a high capacity in
Li1.3Nb0.3Mn0.4O2 electrodes is investigated by X-ray absorption spectroscopy (XAS) (57).
Results indicate that both transition metal and oxide ions play an active role during
cycling, therefore large reversible capacities result from both cationic and anionic
redox processes.
Figure I - 9: Electrochemical performances of cation disordered Li1.3Nb0.3V0.4O2 cells cycled at a
rate of 10 mA.g-1 between 4.8 V and 1.5 V vs Li+/Li (59)
24
The same authors report another disordered rocksalt Li1.3Nb0.3V0.4O2, which is prepared
by substitution of Li3NbO4 with vanadium metal (59). Differently from other
compositions (Table I - 1), this material delivers smaller reversible capacity (170
mAh.g-1) because only cationic redox (V3+ / V4+) is active during cycling (Figure I - 9).
Such result suggests that anionic redox is not always active for all the materials, thus
capacities can vary based on compositions.
I.3.b. Li2TiO3 based cation disordered rocksalt oxides
B. Li et al. studied disordered rocksalt Li2TiO3, which is electrochemically inactive.
Therefore, almost no capacity is obtained at slow cycling rates though it is synthesized
in nanometric size (60). The reason of poor electrochemical activity is attributed to
thermodynamic limitations. To overcome such limitations, substitution of Li2TiO3 with
Fe and Mn is suggested. Thereafter, Fe-substituted Li2TiO3 (Li1.2Ti0.4Fe0.4O2) and Mn-
substituted Li2TiO3 (Li1.2Ti0.4Mn0.4O2) cation disordered rocksalt compositions are
prepared, as illustrated in Figure I - 10 (41,60).
Figure I - 10: Electrochemical performances of Li1.2Ti0.4Fe0.4O2 and Li1.2Ti0.4Mn0.4O2 cells that
are cycled at a rate of 5mA.g-1 between 4.8 V- 1.5 V vs Li+/Li at 50 °C (41)
25
In Li1.2Ti0.4Fe0.4O2 electrodes, initial charge capacity reaches 350 mAh.g-1 and a potential
plateau is observed at 4 V. During discharge, polarization increases, thus discharge
capacity reduces to 250 mAh.g-1. Similar observation is made on Li1.2Ti0.4Mn0.4O2
electrodes: discharge capacity (320 mAh.g-1) is slightly smaller than the charge capacity
(360 mAh.g-1) at the first cycle though subsequent cycles are fully reversible. The
difference between first and subsequent cycles in Li1.2Ti0.4Mn0.4O2 is much smaller
compared to the one in Li1.2Ti0.4Fe0.4O2. The author explains the reason with the
reversibility of anionic redox process. Anionic redox process is almost reversible for
Li1.2Ti0.4Mn0.4O2, whereas it is irreversible in Li1.2Ti0.4Fe0.4O2 electrodes.
In summary, Li1.2Ti0.4X0.4O2 (X: Fe3+, Mn3+) compositions can deliver large reversible
capacities at high operating potentials. However, it should be noted that those results
are obtained at high temperatures (50 °C), thus electrode kinetics is still slow.
I.3.c. Li4MoO5 based cation disordered rocksalt oxides
Another group prepares cation disordered in the Li4MoO5-LiMoO2 system that has the
potential to deliver high capacity based on multi electron (Mo3+ / Mo6+) redox reactions.
As a comparison, Li3NbO4-LiMoO2 and Li2TiO3-LiMoO2 solid solutions are also
prepared (61).
Figure I - 11: Electrochemical performances and capacity retention of Li6/5Ti2/5Mo2/5O2,
Li9/7Nb2/7Mo3/7O2 and Li4/3Mo6+2/9Mo3+4/9O2 cells cycled at a rate of 10mA.g-1 between 4.8 V-
1.0 V at 55 °C (61)
26
Figure I - 11 shows the electrochemical performance of three samples. Among them,
Li4/3Mo6+2/9Mo3+4/9O2 delivers the highest capacity as its lithium content is greater than
Nb and Ti based compositions (61). This result supports the previous observation that
Li-rich content forms percolating network along diffusion channels, thus lithium
conduction increases as well as the capacities (56). Redox activity of the three
compositions are examined with XAS and they confirm that redox activity is based on
pure cationic redox process (Mo3+/Mo6+). Nevertheless, it should be noted that despite
their large reversible capacities (280 - 340 mAh.g-1), operating potential is quite low (~3
V). In addition, the fact that you need to run a first discharge before getting the full
capacity could seriously impede their application in full cells. However, these
materials have good cycling stability, which is attributed to structural stability (61). Ex
situ XRD results reveal that cation disordered structure is rigid, therefore no change in
the crystalline structure is detected and only 1.4% volumetric change is observed at the
end of the second cycle. In brief, Li4MoO5 based cation disordered rock salts can offer
large reversible capacities based on pure cationic redox reactions at relatively low
operating potentials. Their capacity retention is excellent thanks to rigid rocksalt
structure.
27
I.4. Li-rich Cation Disordered Rocksalt Sulfides
I.4.a. Li2TiS3 based cation disordered rocksalt sulfides
In the literature, cation disordered rocksalt sulfides are also studied. For instance, A.
Sakuda et al. mechanochemically synthesize Li2TiS3 powders at different milling
durations (Figure I - 12). Cation disordered rocksalt phase is obtained after 40 h milling
without any remaining precursors (62).
Figure I - 12: (a) X-Ray diffractograms of Li2TiS3 powders milled for different durations, b)
comparison of experimental and simulated diffraction pattern of Li2TiS3 (62)
Figure I - 13: Electrochemical performance of Li2TiS3 electrode at a rate of C/10 between 3.0 V
and 1.5 V vs Li+/Li (62)
Coin cell tests show that Li2TiS3 cells can provide a charge capacity of 274 mAh.g-1 and
discharge capacity of 425 mAh.g-1 based on 1.6 and 2.5 electrons exchanged upon
28
cycling at a rate of C/10 between 3 V and 1.5 V (Figure I - 13). At the second cycle,
reversible 400 mAh.g-1 charge-discharge capacities are obtained at 2.3 V vs Li+/Li.
Capacity retention of Li2TiS3 is compared in both liquid and solid-state electrolyte
systems. When the cells are cycled in carbonate-based electrolytes, they observe that
cycling stability gradually decreases during cycling and 80% capacity is reached at the
end of the 20th cycle (Figure I - 14). The reason of capacity fading in liquid cells is still
unknown. In contrast, superior cycle life is obtained in solid state cells upon cycling at
a rate of 0.5 mA.cm-2. Thus, 98.9 % capacity remains at the end of the 50th cycle (based
on the difference between cycle 5 and 50). Therefore, such materials are suggested to
be used in solid state batteries (63).
Figure I - 14: Capacity retention of Li2TiS3 electrodes (a) in carbonate-based electrolyte at a
rate of C/10 between 3.0 V and 1.5 V vs Li+/Li and (b) in solid state electrolyte at 50°C (62) at
different rates
To discover structural evolution upon cycling, ex situ X-ray diffraction analysis of the
electrodes cycled at different states of charge is performed. Peak intensities of Li2TiS3
reduced and structure became amorphous at the end of the charge, but it returned to
cubic rocksalt structure at the end of discharge (64). Similar observation is reported for
cation disordered Li8FeS5 electrodes as well (65). Cation disordered structure allows
reversible structural change from disordered rocksalt to amorphous phase.
29
I.4.b. Li3NbS4 based cation disordered rock salt sulfides
Same author also reports a disordered rocksalt Li3NbS4 prepared with the same
synthesis recipe (62,66). Disordered rocksalt Li3NbS4 has slightly bigger lattice
parameters (5.13 Å) than Li2TiS3 (5.06 Å).
When the cells are cycled at a rate of C/10 between 3 V – 1 V, a reversible capacity of
486 mAh.g-1 is achieved corresponding to 4.5 exchanged electrons (Figure I - 15).
Cycling curve is reversible in subsequent cycles and its energy density is reported as
high as 923 Wh.kg-1. Li3NbS4 cells show better cycling stability than Li2TiS3 cells, 97%
capacity is reached at the end of the 50th cycle (66). However, average potential of
Li3NbS4 is smaller (2.20 V) than average potential of Li2TiS3 (2.30 V). At high current
rates (1000 mA.g-1), superior cycling stability is observed.
Figure I - 15: Specific capacity vs cell voltage of Li3NbS4 electrodes and capacity retention vs
number of lithium ions exchanged during cycling (66)
Different sulfides, Li2TiS3-Li3NbS4 and Li2SnS3-Li3NbS4 are also reported by the same
author and all of them are synthesized in cation-disordered rocksalt structure (67,68).
Li(8+x)/3Ti(4-4x)/3NbxS4 present reversible capacities higher than 350 mAh.g-1 corresponding
to three electrons process between 3.0 V and 1.5 V (67). On the contrary, tin-substituted
composition, Li3-x/3Nb1-xSn4x/3S delivers smaller capacities with increasing amount of tin
in the composition (68). The reason is explained with low electronic conductivity of tin
30
(7.4x10-7 S.cm-1). The author concludes that sufficient electronic conductivity is
necessary to activate anionic (sulfur) redox process.
Overall, disordered rocksalt sulfides can provide large reversible capacities, however
sufficient electronic conductivity is necessary to activate anionic (sulfur) redox process.
Moreover, capacity fading in the cells containing liquid electrolytes is a major problem
for all disordered rocksalt sulfides. The reason of such a degradation in liquid cells is
still unknown and needs to be elucidated.
I.5. Redox Activity in Disordered Rocksalts
Previously, large capacities delivered by disordered rocksalts were attributed to
activity of both cationic and anionic redox processes. This implies that such processes
should be understood well to design new high-energy chemistries (41,60,69).
Redox activity in a composition can be outlined with the electronic states of the
materials. This can be explained with the metal-ligand band structure (Figure I - 16).
Density of states of a material contains M (metal) at conduction band and L (ligand) at
valence band below energy of fermi level (EF). When purely cationic redox is active,
metal-ligand bands are well separated from each other, therefore there is no
overlapping with these two bands (69). High ionicity between these two bands allows
the metal acting as a redox center and providing necessary electrons during Li
insertion.
Figure I - 16: Cationic redox (Metal (Mn+ d) ligand (Ln- p) band structure) and anionic redox
(Metal (M(n-1)+ d) ligand (L(n+1)- p) band structure)
31
When anionic redox is active, high covalency of metal-ligand band is present. P orbital
of ligand is overlapping with d orbital of metal atom, so ligand transfers electrons, fills
d orbital of metal atom and leaves behind oxidized species (or holes). That means part
of transition metal is reduced to lower valence state, this is called Reductive Coupling
Effect in the literature (70). Such effect is reported for both Li-rich layered (69–74) and
cation-disordered rocksalts (62,73) when anionic redox is activated.
• Anionic Redox in Cation Disordered Rocksalt Oxides
As previously mentioned, Li1.2Ti0.4Mn0.4O2 electrodes delivered large reversible
capacities. To examine such large reversible capacities in Li1.2Ti0.4Mn0.4O2 electrodes,
XAS study is conducted by Li et al. (60). XAS spectra of Ti shows that valence state of
titanium does not change, therefore Ti does not contribute to charge transfer. Whereas,
XAS spectra of Mn and oxygen progressively changed upon charging (Figure I - 17).
This result suggests that charge capacity is produced by both cationic (Mn3+ / Mn4+) and
anionic (O2- / On-) redox activities.
Figure I - 17: XAS spectra of O K-edge and Mn L-edge of Li1.2Ti0.4Mn0.4O2 electrodes during
charge-discharge process (60)
Same study is conducted to examine redox activities in Li1.2Nb0.3Fe0.4O2 electrodes
(Figure I - 18). At the beginning of charge, oxygen species are partially reduced,
whereas iron species are partially oxidized (Fe3+ / Fe4+). However, such reduced oxygen
species disappear at the end of the charge, O K-edge spectra shows similar results for
32
pristine and charge states. This result indicates that anionic redox is activated at the
beginning of the charge, and responsible for the charge transfer. However, the reverse
process is not observed during discharge, indicating anionic redox is not fully
reversible in Li1.2Nb0.3Fe0.4O2 electrodes. Besides, electrode morphology changes at the
end of the discharge (Figure I - 18). This also indicates that such irreversible anionic
redox leads to oxygen release that causes side reactions on the electrode surface as well
as reorganization of particle morphology.
Figure I - 18: Li1.2Nb0.3Fe0.4O2 electrodes at the end of the charge and discharge: Fe L-edge, O
K-edge spectra (41) and TEM images of Li1.2Nb0.3Fe0.4O2 before and after cycling at 50 oC
33
Overall, it seems that anionic redox activity in cation-disordered rocksalt oxides is not
always reversible for all the compositions and irreversible anionic redox can result in
smaller reversible capacities. The reversibility of anionic redox is still unknown,
therefore, further studies are required to elucidate such phenomenon.
• Anionic Redox in Cation Disordered Rocksalt Sulfides
To utilize anionic redox, it is necessary to increase covalency between metal and
ligand, this requires using high electronegative (or less electropositive) metal atom and
less electronegative ligand. For instance, production of persulfide (S22-) is much easier
compared to oxide ions (O22-) because sulfur p orbitals are relatively higher than oxide
p orbitals (69). This creates high covalency between transition metal-sulfur due to
overlapping metal d band and ligand p band, thus anionic redox process is easily
triggered in rocksalt sulfides (69).
Figure I - 19: Coin cell test (a), ex situ X-ray diffraction analysis (b), pair distribution
function analysis of Li2TiS3 electrodes (c) and Ti-S and S-S bonds obtained from ab initio
molecular dynamics calculations (d) (64)
34
Sakuda et al. conducted ex situ X-ray diffraction analysis, pair distribution function
analysis and ab initio molecular dynamics calculations to explain redox activity in
Li2TiS3 electrodes (64). Pair distribution function analysis of electrodes at different states
of charge indicates that at medium range order (r > 5 Å), peak intensities reduces
during lithium extraction and increases during lithium insertion (Figure I - 19c). This
is aligned with ex situ XRD result in which a reversible structural transition is observed
(Figure I - 19b). At short ranges (r < 5 Å) in Figure I - 19c, they observe that covalent S-
S bonds reversibly form and dissociate during charge and discharge correspondingly.
Formation of S-S bonds is attributed to reversible sulfur redox process, sulfur oxidizes
during charge and reduces during discharge. Formation of S-S bonds during charge
are also observed in ab initio molecular dynamics simulations, S atoms lost more than
one lithium neighbor and make covalent S-S bond. Such simulations are done based
on “random” lithium extraction and insertion mechanisms upon cycling. They observe
that atomic distance of Ti-S does not change, S-S distance of S coordinating around Ti
atoms shortens and lengthens during charge and discharge respectively. Such
arrangement allows flexible structural transition from amorphous to rocksalt
structure.
As a summary, anionic redox process is highly reversible in disordered rocksalt
sulfides, flexible structural transition allows such systems to reversibly exchange the
lithium ions.
35
I.6. Conclusion This chapter describes lithium ion technology and current positive electrode materials.
To further increase the specific energy of lithium ion batteries, new high energy Li-rich
and Ni-rich NMC, high-potential spinel LNMO compositions are developed and
extensively studied. Alternatively, Li-rich cation disordered rocksalts have recently
attracted research interests after they have shown promising electrochemical
performances. In this thesis project, we study two different cation disordered families:
rocksalt oxides and sulfides.
As we previously mentioned, Li-rich cation disordered rocksalt oxides can offer large
capacities (≥ 250 mAh.g-1), however there is always irreversibility after first charge. The
origin of such irreversibility is still unknown. To answer this question, we study a
selected composition, Li-rich disordered rocksalt Li2.2NiTi0.2Nb0.6O4. At first, we apply
various synthesis optimizations to improve its electrochemical performance and the
later, we investigate its structural properties with in situ X-ray diffraction technique.
In the second part of the thesis project, we study Li-rich cation disordered rocksalt
Li2TiS3 that have shown unusual electrochemical performances (425 mAh.g-1 at 2.3 V).
Despite its large reversible capacities, strong capacity fading is observed during
cycling. The reason of poor capacity retention is still unknown, therefore our first aim
is to examine degradation problem of Li2TiS3 electrodes with structural (ex situ and
operando neutron diffraction) and morphological (SEM and EDX) studies. To improve
its capacity retention, we propose substitution of Li2TiS3 with selenium, thus we
prepare new compositions (Li2TiSexS3-x), which have shown better capacity retention
than Li2TiS3.
Furthermore, we still lack crucial information about the nature of redox processes in
Li2TiS3 electrodes. Thus, our second aim is to perform electrochemical
characterizations on both Li2TiS3 and Li2TiSexS3-x electrodes. Meanwhile, we perform
surface characterizations on electrodes through X-ray photoelectron spectroscopy to
detect redox active species and to elucidate active redox processes in those materials.
36
Chapter II: Experimental Procedures
This chapter describes the synthesis methods, structural and electrochemical
characterization techniques that are used throughout the thesis. The principles of each
technique, sample preparations and experimental parameters are described.
II.1. Synthesis Methods
Molten salt Synthesis
In the previous chapter, electrochemical performances of disordered rocksalts showed
size dependent properties, thus low temperature synthesis routes were suggested to
prepare these materials. To prepare rocksalt oxides, we select molten salt synthesis
because it allows us preparing ceramic powders at low temperatures. According to the
synthesis procedure, precursors are mixed with an eutectic salt mixture and the
resulting powder is calcined at a temperature above the melting point (Figure II - 1).
During the calcination, the salt melts and acts as the reaction medium, thus reaction
takes place at much lower firing temperatures due to high solubility and diffusivity of
the active species inside the salt (75). After the synthesis, the salt is removed by washing
of the powder with a solvent, such as water.
Figure II - 1: Molten salt synthesis of Li2.2NiTi0.2Nb0.6O4 powder
37
Here are the important parameters for low temperature molten salt synthesis (75):
• Molten salt should be stable, shouldn’t react with precursors during the
synthesis
• Solubility of species in molten salt should be high enough to obtain
targeted phases.
• For low temperature synthesis, molten salt should be selected as it has
low melting point
Mechanochemical Synthesis
Mechanochemical synthesis allows preparing powders in crystalline, meta-stable and
amorphous phases. Among different millers, SPEX and planetary millers are
commonly used in laboratory scale production. In this thesis study, planetary ball
milling is used to prepare disordered rocksalt sulfur-selenium compositions
(Li2TiSexS3-x) that are prepared with a 62:1 (wt%) ball to powder ratio and milled at 510
rpm for 15h, 20h, 25h and 40 h (Figure II - 2). All the samples are prepared in a glovebox
under inert argon atmosphere to avoid air contamination. Appropriate amount of
precursors is weighed and placed into the jar that is then transferred from glovebox to
planetary miller (PM 100) operating under room temperature (Figure II - 2). After the
synthesis, powder is recovered in the glovebox.
Figure II - 2: Mechanochemical synthesis of Li2TiSexS3-x compositions
Rotating disc of the planetary miller induces planetary-earth like movement of the jar
(Figure II - 3). In such a configuration, the jar rotates at its own axes with an inverse
direction of the one of the rotating disc. This creates a centrifugal force that allows the
38
balls to travel freely inside the jar, hit the jar wall as well as each other. As a result,
collision energy is produced. Meanwhile powder particles are trapped between the
balls, consequently they are grinded and rewelded due to mechanically activated
solid-state reactions (76).
Figure II - 3: Mechanochemical synthesis experimental set-up: stainless steel jar, zirconia
balls and planetary miller (PM 100), scheme of grinding medium in PM100
The parameters that has influence on phase and the morphology of the final product
are described below (76):
Grinding medium: Type and size of the balls affect the energy inside the jar, thus phase
of the powder can change. In general, bigger balls create stronger shocks and the
impact energy on the particles becomes greater. Besides, hard materials, such as
stainless steel, hardened steels, tempered steel are frequently used to increase weight
of the grinding medium as they create enough impact energy. In our experiments, we
use zirconium jar with zirconium balls to avoid any reaction between sulfur and steel.
Ball to powder mass ratio: Depending on the ball to powder ratio, energy input can
change. High ball to powder ratio induces high energy inside the jar.
Milling speed: High milling speeds creates high energy input into the powders,
consequently, temperature inside the cell increases. High temperature can be beneficial
to allow the ingredients diffusing enough to trigger chemical reactions. Whereas it
might also cause undesired reactions as well, such as the formation of impurities in the
39
final product if milling takes longer durations. Therefore, powders should be milled at
optimum milling speed.
Milling duration: Milling duration often depends on the milling speed and powder
load (ball to powder ratio). In general, reactions are completed at shorter durations
when the powders are milled at high speeds or at high loads (high ball to powder
ratios) as the input energy is sufficient. It is also noted that milling powders for longer
durations increase the contamination in the final product as the balls often wear out.
II.2. Structural Characterizations
Powder X-ray Diffraction (XRD)
The crystal structure of the synthesized powders is determined by X-ray diffraction
using a BRUKER AXS D8 diffractometer using Cu Kα (λ=1.540 Å) anticathode, the
sample is scanned in the range of 8-80° using a step size of 0.02°. As the sulfides are air
sensitive, all the samples are prepared in the glovebox and the sample surface is
protected with Kapton®, which is a polyimide film, shown in Figure II - 4. Kapton® has
high transmittance to X-rays, however its signal still can be detected at small 2𝜃 angles
(0-28°) in the diffraction patterns of the powders. Besides, its tightness to air, and more
specifically moisture, is limited to couple of hours. Therefore, duration of the X-ray
diffraction is limited to a maximum of three hours.
Figure II - 4: X-ray diffraction analysis: sample holder for powder and electrode
40
X-ray diffraction is a useful technique to determine the crystal structure of the
crystalline powders. Each powder has a unique crystal structure that contains parallel
crystalline planes designated by Miller indexes (hkl) that are separated by “dhkl”
distance called interreticular distance.
Figure II - 5: Scheme of X-ray diffraction technique and Bragg’s law
As the incident X-ray wave interacts with the powder, each plane reflects X-rays
differently, hence this creates a constructive interference. For instance, two different X-
ray waves with a wavelength of λ hit a crystal structure at incident angle θ, first wave
will reflect from top atomic plane, the second wave will reflect from bottom plane
(Figure II - 5). Since the second wave must cross deeper into the bottom plane, it travels
further than the second wave, its distance is much bigger than the first wave. If this
distance is equal to nλ then these two waves are in phase, therefore they constructively
interfere. This verifies Bragg’s law (Eq. II - 1).
Bragg’s law: nλ = 2 dhkl sinθ (Eq. II - 1)
where n: number of X-ray waves, λ is X-ray wavelength, dhkl is distance between the
crystalline planes and θ is the reflected angle.
The most important use of X-ray diffraction method is to identify the materials. Each
material has a characteristic XRD pattern that can be expressed as a fingerprint for
phase identification. There are two important parameters for the phase identification:
peak positions (d-spacing) and peak intensity. Peak positions allow determining the
41
symmetry and cell parameters, whereas peak integrated intensity provides
information about the atomic positions and the occupancies.
X-ray diffractogram provides various information about powders, such as peak
positions, peak profile, peak height and peak width, as illustrated in Figure II - 6.
Figure II - 6: X-ray diffractogram that provides various information about the crystalline
materials
Peak intensities (Ihkl) are proportional to structural factor (Fhkl), it is expressed below
Ihkl = S. jhkl.LPhkl.|Fhkl|2 (Eq. II - 2)
where S: scale factor, jhkl: multiplicity of the reflection, LPhkl: Lorenzian factor and
polarization corresponding diffractometer configuration, Fhkl: structural factor
Structural factor can be defined as
(Eq. II - 3)
Structural factor provides information about the type of the atoms (fj) and the atomic
positions (u, v and w) within the unit cell. Therefore, the target of the structural
analysis is to compare observed and calculated structural factors.
42
The size of the crystalline domains (or crystallite size) can be calculated by using
Scherrer equation (Eq. II - 4). In the equation, β(θ) corresponds to peak widths obtained
from the full-width at half maximum (FWHM) corrected from the instrumental
broadening at corresponding angle (2θ), K is Scherrer constant that varies between
0.62 to 2.08 (0.94 for cubic systems) and λ is X-ray wavelength. The crystallite sizes (L)
can be calculated at each angle, but its mean value is frequently used for a convenience.
𝛽(𝜃) =𝐾𝜆
𝐿𝑐𝑜𝑠𝜃 (Eq. II - 4)
In this thesis project, Fullprof program has been mainly used (77). The peak positions
and peak profiles depend on the crystal structure and the space group
correspondingly. At first, XRD patterns are fitted using Le Bail method in which peak
shape and positions are refined iteratively until minimum Chi-square value (< 5) is
reached. In this way, peak positions are fixed and the unit cell parameters are
calculated with high accuracy. To perform Rietveld refinement and simulated XRD
pattern, structure models (crystal information file - cif) are taken from the databases
(Crystallography Open Database). Rietveld refinement is performed to minimize the
difference between the observed and the calculated data by refining the positions and
occupancies of the atoms in the crystal.
In the case of the in situ analysis, the data are refined using Jana2006 to enable the
correction of the effect of the variable divergence slit (78). The refinement has been
carried using the Fundamental Parameter Approach, in which the instrumental
resolution is calculated from the description of the main parameter of the optical set-
up. This allows a direct correction of the peak profile.
• Ex situ XRD analysis
Ex situ XRD analysis is another way to detect structural evolution in disordered
rocksalts. Coin cells are cycled in the standard procedure that is described in Section
2.3 and stopped at different states of charge. After cycling, cells are dismounted in the
glovebox, electrodes are recuperated from coin cells, they are washed with dimethyl
43
carbonate (DMC) and dried inside the glovebox. After drying, electrodes are placed on
the surface of a glass plate and covered with Kapton® to protect from air atmosphere
(Figure II - 4).
• In situ XRD analysis
To observe structural evolution in disordered Li2.2NiTi0.2Nb0.6O4 during cycling, in situ
XRD analysis is performed using Biologic Instrument and BRUKER diffractometer.
The cell is cycled in GITT (Galvanostatic Intermittent Titration Technique) mode.
Figure II - 7: Scheme of in situ XRD analysis with GITT mode of cycling
The cell is charged at a rate of C/50 for 30 min, then the current is stopped and it follows
a relaxation for 210 min without current pulses. During cell relaxation, XRD analysis
is performed (Figure II - 7). Thereafter, same procedure is applied for discharging the
cell. The loop is arranged until the cell fully charge-discharged.
Figure II - 8: In situ cell assembly
The in situ cell contains 30-40 mg of active material (70wt% Li2.2NiTi0.2Nb0.6O4 and 30%
Carbon black) as positive electrode, glass fiber (Whatman®) disk as separator, LP100 (1
M LiPF6 in EC: PC: DMC= 1:1:3) as non-aqueous electrolyte, Li foil as negative
44
electrode (Figure II - 8). Beryllium, which is transparent to X-rays, is used as current
collector and a piece of aluminum foil (2 µm) is used to protect beryllium from
oxidation at the potentials higher than 4 V. The sample powder is prepared by mixing
70% active materials with 30% carbon, and then dried in a Buchi® oven at 60 °C under
vacuum for 24 hours.
Operando Neutron diffraction
To observe the structural evolutions in the composition of Li2TiS3 and to detect where
extra Li+ ions go into the structure, we perform operando neutron diffraction at Institute
Laue Langevin (ILL) in Grenoble. The measurements are performed on the D20
diffractometer (high resolution configuration) at a wavelength of 1.594 Å between 1°
and 170° with a step of 0.05°. The cell is cycled at a rate of C/10 in a range from 3.5 V
to 1.0 V.
Neutron diffraction is a complementary technique to X-ray diffraction. Differently
from X-rays, neutrons interact with the nucleus of the atoms, instead of electrons. Each
atom has a specific coherent elastic scattering length that can be positive or negative
value. This property allows determining light elements (Li, B, H, C, O and F),
differentiating between the elements that have close atomic numbers as well as the
isotopes in the crystalline structure. Moreover, it is a useful technique to characterize
the materials in operando diffraction mode, both electrochemical and structural
characterizations are simultaneously performed. However, large amount of active
material should be subjected to the neutron beams to get sufficient intense diffraction
peaks. In general, working with the cells containing large amount of powders is often
challenging, it often results in large cell polarization and smaller capacities. Therefore,
electrochemical performance of neutron cell must be optimized beforehand. Besides,
separator and electrolyte in the neutron cell contain hydrogen that presents high
inelastic coherent scattering cross section and thus create a large diffraction
background. To avoid such large background, electrolytes are usually prepared from
deuterated solvents that does not contain any 1H.
45
Figure II - 9: Operando neutron diffraction cell assembly79 and experiment in ILL
The cell used for neutron diffraction is designed by M. Bianchini79 (Figure II - 9). The
bottom part of the cell is made of Ti-Zr alloy. The neutron scattering lengths of titanium
(bTi= -3.37x10-12) and zirconium (bZr= 7.16x10-12) are opposite in sign and almost
related to 2:1. Such alloy is prepared as Ti2.08Zr that becomes transparent to neutrons,
hence it does not contribute to the background of diffraction pattern.
• Cell preparation for operando neutron diffraction
The cell contains 250 mg of sample powder (Li2TiS3), positive electrode, glass fiber
(Whatman®) disk as separator, deuterated electrolyte (1 M LiPF6, d-PC: d-DMC= 1:3,
Eurisotop), which is dried on molecular sieves, and lithium metal as negative electrode
(Figure II - 9). The positive electrode is prepared containing 70% active materials and
30% carbon and it is dried in Buchi® at 60 °C vacuum for 24 hours.
Before the experiment in ILL, we studied and optimized the electrochemistry of Li2TiS3.
At the beginning, we observed a large polarization and poor reversibility when the
cells were cycled in a standard schedule (3 V – 1.5 V). Thereafter, we optimized the
amount of the powder and found an adapted cycling schedule (3.5 V – 1 V).
46
Scanning Electron Microscopy (SEM)
Morphology of synthesized powders are investigated by Scanning Electron
Microscopy equipped with a field electron gun (MEB LEO 1530 Gemini). The surface
of the sample is scanned through focused beam of electrons. As electrons interact with
the atoms, electrons are removed from the inner shell and empty place is filled by
electrons at higher shell energy level, resulting in an energy release. Electrons emitted
by the atoms are collected by detectors and the detected signal is transferred into
information about surface topography. Two different types of electrons are emitted
from the surface: secondary and backscattered electrons. Secondary electrons are
collected close to the specimen surface; hence they create high resolution images.
Backscattered electrons emerge from deeper locations within the specimen; therefore,
they produce low resolution images compared to secondary electrons. As the intensity
of backscattered electron signal depends on the atomic number (Z), they provide
information about the element distribution in the sample.
Energy-Dispersive X-ray Spectroscopy (EDS)
To perform the chemical mapping and determine elements in the compositions, we
used Energy-Dispersive X-ray Spectroscopy using 15 kV in MEB LEO 1530 Gemini.
The surface of the sample is bombarded with focused beam of electrons and emitted
X-rays are collected by the detectors. This provides the information about the element
distribution in a specific region of the sample. Although this technique is easy and
practical to determine the elements in a composition, most of the instruments cannot
detect light elements (Z<10) precisely, thus quantitative analysis of powders is often
difficult.
• Sample preparation for SEM and EDS Analyses
A specific sealed sample holder is used to transfer air sensitive samples from glovebox
to the SEM vacuum chamber. The samples are placed on the surface of adhesive carbon
tape that has been already sticked onto SEM stub (Figure II - 10). Data is recorded for
47
various magnifications (2000, 5000, 10000 and 30000) by using 5 kV accelerating
voltage.
Figure II - 10: Sealed sample holder for SEM of sulfide electrodes
Raman Spectroscopy
To observe chemical bonds in different disordered rocksalt sulfur-selenium
compositions, Raman measurements are carried out on a LabRam HR instrument
(Horiba Jobin-Yvon) in backscattering geometry. 632 nm laser excitation wavelength
on a laser power of 10% filter (≈ 0.17 mW) is used. To prevent heating of the materials
by the laser beam and to obtain a sufficient signal to noise ratio, short acquisition times
(10 s) are used with multiple accumulations (30).
Working mechanism of Raman Spectroscopy is based on inelastic scattering, so called
“Raman Effect”. As the sample is irradiated by photons with a single frequency, the
atoms vibrate and rotate as a response of electromagnetic field that is formed upon
photon-sample interaction. As a result, photons are scattered at different frequencies.
In Raman spectroscopy, one can know the characteristics of each material from
• Raman frequencies, which allows distinguishing the species from each other.
• Width of Raman peaks, which gives information about the quality of crystals
and amount of plastic deformations
• Polarization of Raman peak, which is more related to crystal symmetry and
orientation.
48
Different laser can be used for Raman light source. As the wavelength of the laser gets
smaller, signal quality increases, however it also causes heating of the sample, thus the
sample is more easily damaged.
• Sample preparation for Raman Spectroscopy
Powders are prepared in the glovebox and a sealed sample holder is used for Raman
measurements. The powder is put in the center of the sample holder that is closed by
a cover containing a transparent glass that is invisible to the laser. Sample holder is
transferred from glovebox to the LabRam HR instrument coupled with Atomic Force
Microscopy (AFM) machine. AFM probe is used to select the particle region, and then
laser beam is focused on selected particles.
X-ray Photoelectron Spectroscopy (XPS)
To determine redox active species and to elucidate the redox mechanism in disordered
rocksalt sulfur-selenium compositions (Li2TiSexS3-x), X-ray Photoelectron Spectroscopy
(XPS) is performed using Versaprobe II PHI 5000 (ULVAC-PHI) spectrometer
equipped with a 100 μm focused monochromatic Al Kα X-ray source (1486.6 eV). At
the beginning, powders are analyzed and fitting models are developed. The same
fitting models are applied to the electrodes at different states of the charge.
Principle of XPS is based on the photoelectronic effect in which sample is irradiated by
photons with an energy of hυ and electrons are ejected from the core levels with a
kinetic energy of Ek. The kinetic energies of the ejected electrons are measured, binding
energies are calculated by the spectrometer (Figure II - 11).
Binding energy of core electron (Eb) is expressed by Einstein relationship (Eq. II - 5):
Eb = hѵ –E
k - Ф
sp Eq. II - 5
where Ek: kinetic energy of ejected electron, which is measured by spectrometer, hυ:
energy of photons, Фsp: work function induced by the analyzer.
49
Figure II - 11: Principle of X-ray photoelectron spectroscopy
XPS provides information about
• Chemical composition: Each element has unique core electrons energies. By
measuring the binding energies of core electrons (at K, L or M core levels),
elements present on the sample can be identified. Such elemental analysis is
performed by the “Survey” mode of XPS technique and binding energies can be
found in XPS database.
• Chemical state: The binding energy of a core electron is sensitive to the chemical
environments of the elements. If the same element is bonded to a different
chemical species, the binding energy of its core electron can be different. This
arises from different electronegativities of the atoms in a bond. As the electron
density of valence is shifted to more electronegative (electron-withdrawing)
atom, less energy is required to extract an electron from such electron-enriched
environment. Thus, the kinetic energy of the ejected electrons from the more
electronegative atom increases, but the binding energy of core electron
decreases. Such change in the binding energy results in a shift of core peak at
XPS spectrum. From this observation, the change in the valence states of the
species can be detected by High Resolution XPS technique.
50
• Quantification: As the number of photoelectrons depends on the atomic
concentration of the elements present in the sample, XPS also allows
quantifying the chemical composition. Atomic concentration (Ci) can be
expressed as below, where Ii: peak intensity (peak area after background
removal) and Si: sensivity factor
(Eq. II - 6)
• Sample preparation for XPS analysis
Electrodes after cycling are dismounted in the glovebox, washed with DMC (dimethyl
carbonate) and let to dry overnight. Then electrodes are cut into small pieces and
placed into sample holder. Later, it is transferred into vacuum chamber using a transfer
vessel to avoid any air contamination (Figure II - 12).
Figure II - 12: Sample holder for XPS analysis and the XPS machine
The first aim is to identify the elements present in the samples, therefore survey spectra
is recorded over a spectral range of 0−1200 eV using a pass energy of 117 eV,
corresponding to an energy resolution of 1.6 eV and an acquisition time of 0.5 s.eV-1.
The high-resolution spectral analysis is carried out using a pass energy of 23 eV
allowing an energy resolution of 0.5 eV.
51
II.3. Electrochemical Characterizations
Electrochemical Cell Descriptions
• Preparation of oxide electrodes
The electrodes are prepared according to 80:10:10 formulation: 80% active mass of
material (Li2.2NiTi0.2Nb0.6O4) is mixed with 10% Carbon SP (5321) in a mortar. To have
a bulk electrode and to adjust the viscosity of the slurry, 10% of PVDF (Polyvinylidene
fluoride) binder is added with few drops of NMP (N-methyl-2-pyrrolidone) solvent.
Then, slurry is coated on an Al current collector (20 μm) by using 100 μm doctor blade
and left to dry in the oven at 55 °C for 12 hours. Later, electrodes are cut (14 mm disk
diameter), pressed (10 tons), weighed (loading > 4 mg) and dried in the Buchi® oven at
80 °C for 48 hours under vacuum.
• Preparation of sulfide-selenide based electrodes
Unlike oxide-based electrodes, sulfide and selenide-based electrodes are prepared in a
glovebox under inert argon atmosphere to avoid air contamination. Electrodes are
prepared based on the same formulation (80:10:10) as described previously for
preparation of oxide electrodes. Slurry is coated manually on the Al foil by using 100
µm doctor blade and left to dry in the glovebox for three days. Later, electrodes are cut
(14 mm), pressed (10 tons), weighed (loaded > 2 mg) in the glovebox and dried in the
Buchi® oven at 60 °C for 12h under vacuum.
• Coin cell preparation
The electrochemical characterizations of synthesized powders are performed in coin
cell, which is shown in Figure II - 13. The cells are prepared by using 14 mm diameter
positive electrodes, 16 mm Li foil as a negative electrode and 150 μl electrolyte of LP100
containing 1 M LiPF6 in EC:PC:DMC, where EC: ethylene carbonate; PC: propylene
carbonate; DMC: dimethyl carbonate, in volumetric ratio of 1:1:3 respectively. Two
52
types of separators (Φ = 16.5 mm) are used; polypropylene separator (Celgard® 2400)
and polyolefin separator (Viledon®).
Figure II - 13: Coin cell assembly
Electrochemical Techniques
• Galvanostatic Cycling (GC)
To evaluate the specific capacity, average potential and cyclability of rocksalt oxides
and sulfides, galvanostatic cycling technique is used. This technique also gives
information about the potential-composition relations, such as number of lithium
intercalated, system reversibility and diffusion properties (Figure II - 14). To observe
such properties, common approach is to cycle the cells into fully charged and
discharged states with a constant current rate of C/n, which is related to the
intercalation of lithium ions through the electrodes in n hours. The working electrode
undergoes an oxidation process by applying a positive current to the electrochemical
cell until it reaches the high cut off potential, then it goes through a reduction process
by applying a negative current until it reaches the low cut off potential. Difference
between charged and discharged states gives information about the (ir)reversibility of
the system.
53
Figure II - 14: Galvanostatic cycling
• Galvanostatic cycling schedule
Disordered Li2TiS3 is cycled at a rate of C/10 in a potential range from 3 V to 1.5 V vs
Li+/Li, whereas disordered Li2.2NiTi0.2Nb0.6O4 is cycled at a rate of C/50 in a potential
range of 4.8 V and 2 V. ARBIN cycling instrumentation that operates at 22 °C is used
to cycle the cells. To detect if the materials have diffusion limitations, high temperature
cycling schedule is carried out. In that case, coin cells are placed into an oven that
operates at 55 °C and is connected to the ARBIN instrumentation.
• Galvanostatic Intermittent Titration Technique (GITT)
To study the thermodynamics and kinetics of disordered Li2.2NiTi0.2Nb0.6O4,
Galvanostatic Intermittent Titration Technique (GITT) is used. This technique allows
obtaining galvanostatic profile at equilibrium potential.
In GITT technique, the cell is charged (or discharge) with a small constant current for
a certain time, then the current is interrupted and the cell is given a rest until the
potential is stabilized (Figure II - 15). This procedure is repeated until the cell reached
fully charged or discharged states. With imposed current, lithium is deinserted (or
54
inserted) into the active material, thus potential changes based on concentration
gradient. Without no imposed current, potential reaches its thermodynamically
equilibrium condition with ion transport that occurs from the bulk to the electrodes.
Figure II - 15: GITT mode of cycling
The equilibrium potential is key for various investigations in electrodes. Most of the
active materials undergoes structural modifications upon cycling that adversely affect
the electrochemical performances. With galvanostatic mode of cycling, it is quite
difficult to detect the actual start of such modifications, one can only detect the
modifications in the electrodes at the end of the charge and discharge states. On the
contrary, GITT allows performing X-ray diffraction analysis at the equilibrium
potentials, thus it makes possible to detect step by step the structural modifications
based on lithium (de)insertion process. It also allows determining diffusion and kinetic
properties of electrodes at the equilibrium potentials.
55
• Cycling Voltammetry (CV)
To determine redox potentials, electrode kinetics and reversibility of sulfur-selenium
compositions, cycling voltammetry technique is applied. According to this technique,
cell is submitted to a potential evolving linearly at a specific scan rate (V.s-1) and the
current evolves as a response of potential change. This gives information about the
different redox processes that take place at the electrodes.
Figure II - 16: Cyclic Voltammetry test, parameters obtained in a voltammogram
When an electrochemical cell is scanned with a scan rate (µV.s-1), positive current (iox)
is measured for an oxidation reaction, whereas a negative current (ired) is measured for
reduction reaction in current-potential curve (cyclic voltammogram). Besides,
oxidation (Eox) and reduction (Ered) redox potentials can be identified in cyclic
voltammogram (Figure II - 16). The difference between these two potentials (∆E) gives
indications on the reversibility of the redox reactions.
For a Nernstian system, redox reactions are considered to be reversible, thus the
potential difference is expressed as below
∆𝐸 = 𝐸𝑟𝑒𝑑 − 𝐸𝑜𝑥 +56.5
𝑛 𝑚𝑉 (Eq. II - 7)
where n is the number of electrons exchanged.
56
Figure II - 17: Voltammogram of the electrodes scanned from slow scan rates (5 µV.s-1) to
higher scan rates (100 µV.s-1) with gradual increments, calculation of potential difference for
each scan rates to determine electrode kinetics at slower and higher scan rates
In our experiments, two different cells are cycled with a slow scan rate (5µV.s-1) to
observe lithium (de)insertion process in the sulfur-selenium compositions. To
investigate the electrode kinetics at faster scan rates and reversibility of electrode
kinetics, the cells are scanned from slow scan rates (5 µV.s-1) to higher scan rates (100
µV.s-1) with gradual increments (Figure II - 17). The amplitude of the current increases
with scan rates, indicating the relation between electrode kinetics of the redox
reactions with applied scan rates that is expressed in Randles Sevcik equation for a
reversible process:
Ip= (2.687x105)n3/2v1/2D1/2AC (Eq. II - 8)
Ip is peak current, n is number of transferred electrons, v is scan rate (V.s-1), D is
diffusion coefficient (cm2.s-1), A is electrode area (cm2) and C is concentration (mol.L-1)
By calculating potential difference (∆E) for each scan rate, diffusion coefficient of the
species can be calculated using Randles-Sevcik equation (Eq. II - 8). Thus, the electrode
kinetics of different species can be distinguished from each other.
57
Chapter III: Development of New
Li2.2NiTi0.2Nb0.6O4 Positive Material in
Disordered Rocksalt Structure
This chapter describes molten salt synthesis of Li2.2NiTi0.2Nb0.6O4 as well as structural
and electrochemical characterizations. To improve the capacity and develop a model
material, various synthesis optimizations are carried out. Thereafter, understanding
studies on model material are performed.
Introduction
In the first chapter, we mention that Li-rich disordered rocksalt oxides can provide
large reversible capacities at relatively high operating potentials. After those results,
the study of disordered rocksalts has been started in CEA LITEN. Several Li-rich
disordered rocksalt Li2+1/3*zNiTi1-4/3*zNbzO4 compositions have been prepared with
molten salt synthesis. Such compositions are a derivative of Li2MTiO4 composition,
where M: Ni, Co, Fe, Mn (80–84). Among them, Li2.2NiTi0.2Nb0.6O4 is selected in this thesis
project. The first target of the project is to enhance the electrochemical properties of
Li2.2NiTi0.2Nb0.6O4 with optimizing synthesis parameters to prepare a model material.
The next target is to perform in situ X-ray diffraction analysis on model material to
examine a major problem in cation disordered rocksalt oxides, “irreversibility at the
first cycle”.
Rocksalt Li2.2NiTi0.2Nb0.6O4 is synthesized based on the same recipe as in the literature
(48). Li(CH3COO).4H2O (99.0% sigma), Ni(CH3COO)2.4H2O (99% Acros), TiO2 (99%
Sigma), Nb2O5 (99.0% Acros) are weighed in molar ratio of 2.2:1:0.2:0.3 and put into 50
ml stainless steel jar that contains zirconia balls and mixed at 350 rpm for two hours.
After two hours, eutectic NaCl-KCl salt in a salt ratio to precursor ratio of 4:1 is added
into the mixture and then mixing is repeated for two hours. As the mixing is
58
completed, mixture is heated to 300 °C for 2 hours to remove acetate from the
precursors and later it is calcined at 670 °C for 3 hours. Between two heating steps, the
powder is taken from the oven and grinded a few minutes in mortar to have a
homogeneous mixture. The calcination temperature is slightly above the melting point
of eutectic salt (667 °C). After the synthesis, powder is washed with hot water (60 °C)
to remove the remaining salt, filtrated and dried in the oven.
Figure III - 1: Powder XRD patterns of disordered rocksalt Li2.2NiTi0.2Nb0.6O4 powder: before
(1) and after (2) washing with water
Figure III - 1 shows XRD patterns of Li2.2NiTi0.2Nb0.6O4 powder before and after
washing. In unwashed powder, diffraction peaks of disordered rocksalt
Li2.2NiTi0.2Nb0.6O4, NaCl and KCl are detected. By washing the powder with water,
remaining salt is totally removed and only rocksalt cubic phase is observed. Besides,
sharp diffraction peaks indicate good crystallinity of Li2.2NiTi0.2Nb0.6O4 powder.
Figure III - 2 shows XRD pattern of Li2.2NiTi0.2Nb0.6O4 powder in Fullprof program
profile matching. Li2.2NiTi0.2Nb0.6O4 powder is indexed in a rocksalt (cubic) phase
(space group, Fm3̅m) with lattice parameter, a= 4.1842 Å.
59
Figure III - 2: XRD pattern of disordered rocksalt Li2.2NiTi0.2Nb0.6O4 powder in Fullprof
program profile matching
To investigate the electrochemical properties of rocksalt Li2.2NiTi0.2Nb0.6O4, electrodes
are prepared with the procedure explained in Chapter II.3. Three coin cells are cycled
at a rate of C/50 between 4.8 V and 2.0 V vs Li+/Li at room temperature (22 °C).
The electrochemical redox process of Li2.2NiTi0.2Nb0.6O4 can be described as below:
Li2.2NiTi0.2Nb0.6O4 ⇄ Li2.2-xNiTi0.2Nb0.6O4 + xLi+ + xe- (Eq. III - 1)
Considering only Ni2+/Ni4+ redox is active during cycling, the theoretical capacity based
on 2 moles of Li+ exchanged per formula is 264 mAh.g-1.
In Figure III - 3, disordered rocksalt Li2.2NiTi0.2Nb0.6O4 delivers 120 mAh.g-1 charge and
74 mAh.g-1 discharge capacities, corresponding respectively to 0.90 Li+ and 0.56 Li+
exchanged upon cycling. The irreversible capacity at the first cycle is 38%. On the
contrary, second cycle is fully reversible and less polarized, but a much smaller
capacity of 69 mAh.g-1 is delivered. Moreover, the cycle life at room temperature is
rather poor (Figure III - 4). Those results indicate that synthesis improvement is
required to achieve better specific capacities.
60
Figure III - 3: Cell performances of disordered rocksalt Li2.2NiTi0.2Nb0.6O4 electrodes cycled at a
rate of C/50 between 4.8 V-2.0 V vs Li+/Li at 22 °C
Figure III - 4: Cycling life of disordered rocksalt Li2.2NiTi0.2Nb0.6O4 electrodes cycled at a rate
of C/50 between 4.8 V-2.0 V vs Li+/Li at 22 °C
61
III.1. Synthesis Improvement
To improve electrochemical performance of disordered rocksalt Li2.2NiTi0.2Nb0.6O4
material, different synthesis parameters are considered (Figure III - 5).
Figure III - 5: Summary of the parameters considered for synthesis improvement
• Simple tempering (long firing): Rocksalt Li2.2NiTi0.2Nb0.6O4 is prepared using
four different metal oxides reactants. To prepare homogeneous and well-
crystalline ceramic powder, simple tempering, in another words long firing, can
be applied. By letting the ingredients well diffuse during the synthesis, pure
rocksalt phase can be obtained.
• Quenching: Disordered rocksalts have shown electrochemical performances
dependent on particle size; especially diffusion is more problematic in powders
that contain big particles. With fast cooling approach, particle size growth can
be controlled, thus lithium diffusion length becomes shorter, ionic limitations
might be reduced. Consequently, much greater capacities could be achieved.
• Effect of washing (type of solvent): Rocksalt Li2.2NiTi0.2Nb0.6O4 is prepared by
molten salt synthesis, thus washing step is mandatory to remove remaining
salts in the composition. Water is a stable solvent and solubility of various salts
62
in water is quite high, thus it can dissolve many different salts. However,
washing can cause loss of lithium in the composition, for instance lithium ions
in the powders can exchange with protons from the water. This might reduce
the specific capacity; therefore, new solvents should be searched to eliminate
that problem.
• Effect of molten salt to precursor ratio: Molten salts have strong solving
capability; thus, they act as a medium in which diffusion rate of the precursors
is much higher. Ratio of salt to precursor can affect the precursors solubility;
therefore, the effective ratio should be determined.
• Type of molten salt: Each eutectic salt has a different melting point and
solubility. Use of eutectic salt that has a low melting point can reduce the
synthesis temperature and allows preparing powders in smaller particles size.
However, salt should be stable against reactants, should not react with the
precursors throughout the synthesis and it should have enough solubility
power to dissolve the reactants in order to prepare the desired phases.
• Effect of calcination temperature and duration: Low temperature synthesis
methods are useful to synthesize powders in smaller particles size. Finding
optimum calcination temperature and duration is important to optimize the
capacity of Rocksalt Li2.2NiTi0.2Nb0.6O4.
63
III.1.a. Simple tempering (long firing) and quenching strategies
To improve our synthesis recipe, we first consider quench strategy. For a comparison,
we also prepare a reference material prepared by classical approach. Three samples
are prepared with the same recipe, however different firing and cooling steps are
applied.
• The first sample (classical/reference) is heated at 670 °C for 3 hours and let to
cool down in the oven
• The second sample (quench) is heated at 670 °C for 3 hours and it is taken from
the oven and let to cool down outside the oven (no cooling agent is used)
• The third sample (long calcination and quench) is heated at 670 °C for 6 hours
and it is taken from the oven and let to cool down outside the oven.
Figure III - 6: XRD patterns of Li2.2NiTi0.2Nb0.6O4 powders with different strategies: classical
calcination (3h), calcination (3h) with quench and long calcination (6 h) with quench
After calcination, the three samples are washed with water and X-ray diffraction
patterns are compared in Figure III - 6. All samples present the targeted phase with
small amount of impurities. NaNbO3 is present in every sample and is not affected by
the different thermal treatments. Small amount of Li2CO3 is only found in the slowly
64
cooled samples indicating that the carbonation may occur during cooling. Surprisingly
a peak coming from an ordered rocksalt, isostructural to Li3Ni2NbO6 is found for the
material quenched after a long calcination, although it was thought that this would
favor the disordered phase.
Figure III - 7: Cell performances of Li2.2NiTi0.2Nb0.6O4 electrodes cycled at a rate of C/50
between 4.8 V-2.0 V vs Li+/Li at 22 °C
Figure III - 7 shows cell performances of the samples prepared by different strategies
at the first cycle. The “classical” samples show the worst behavior with a very high
polarization that may rise from the formation of Li2CO3 during the cooling. On the
contrary, quenched samples shows lower polarization. Between the two quenched
samples, the samples that was submitted to the shorter thermal treatment exhibits
much higher capacities with a charge capacity of 144 mAh.g-1 and a discharge capacity
of 105 mAh.g-1 upon cycling at a rate of C/50 between 4.8 V-2.0 V vs Li+/Li. Charge
capacity is almost 15% higher than the charge capacity of classical material, whereas
discharge capacity is 30% higher than discharge capacity of classical material.
Moreover, we observe irreversibility at the first cycle for quenched sample is much
smaller (27%) (Figure III - 7). The lower performances of the sample quenched after 6h
may be attributed to the presence of ordered phase that is known to present much
poorer ionic conductivity.
65
Afterwards, we investigate the electrochemical performances of the best quenched
sample at different temperatures (Figure III - 8). At 55°C, both charge and discharge
capacities improve, a charge capacity of 193 mAh.g-1 and a discharge capacity of 144
mAh.g-1 are delivered. This result indicates that ionic and electronic conductivities of
disordered rocksalt Li2.2NiTi0.2Nb0.6O4 might improve at high temperature,
subsequently higher capacities can be reached. Besides, the irreversibility at first cycle
is 27% at 22 °C and 25% at 55 °C. Cycling at high temperature does not eliminate the
irreversibility at the first cycle.
Figure III - 8: Comparison of cell performances of Li2.2NiTi0.2Nb0.6O4 cells (powder fired 3h and
quenched) at 22 °C and 55°C, electrodes cycled at a rate of C/50 between 4.8 V - 2.0 V vs
Li+/Li
Figure III - 9 shows cycling stability at room (22 °C) and high (55 °C) temperatures.
Cycling stability is rather good at room temperature, but discharge capacity is quite
small (96 mAh.g-1). Whereas, cycling stability at 55 °C is reduced and capacity fades
quickly. Poor cycling stability at high temperatures has been already reported in
different positive materials in the literature and it was attributed to the degradation of
active materials (81,85).
66
Figure III - 9: Evolution of discharge capacities of Li2.2NiTi0.2Nb0.6O4 (quenched) at 22 °C,
electrodes cycled at a rate of C/50 between 4.8 V-2.0 V vs Li+/Li
III.1.b. Effect of washing (type of solvent)
In our experiments, we always observe asymmetric X-ray diffraction peaks after
washing our powders. Therefore, we have another assumption that during washing,
some of the lithium ions might have exchanged with protons. Such protons might
interfere with lithium upon insertion/disinsertion and can produce HF that causes
accelerated degradation of the cell. Therefore, our next target is to change our washing
strategy, finding another solvent. To determine effect of washing to the synthesis, two
powders are prepared separately with the same recipe and they are washed either with
ethylene glycol or with water.
67
Figure III - 10: XRD patterns of Li2.2NiTi0.2Nb0.6O4 powders prepared with the same recipe
separately and washed with different solvents: water and ethylene glycol. Kα2 contributions
have been stripped with EVA software for clarity reason. Arrows indicates possible
contributions from H+ exchanged phase.
Figure III - 10 shows the XRD patterns of two different batches after washing. No
remaining salt is detected in diffraction patterns of washed powders, both ethylene
glycol and water effectively remove eutectic salt in powders. Reflections from the EG-
washed samples are broader and may cover small impurities, but they present a
symmetrical profile.
In Figure III - 11, we observe that charge capacity of powder washed with ethylene
glycol is slightly greater, but its discharge capacity is smaller than the powder washed
with water. Although discharge capacity reduces with ethylene glycol washing, we are
now sure that there is no proton-lithium exchange in our system. Therefore, we decide
to wash the powders with ethylene glycol in further synthesis.
68
Figure III - 11: Half-cell performances of Li2.2NiTi0.2Nb0.6O4 washed with different solvents:
water (1) vs ethylene glycol (2), electrodes cycled at a rate of C/50 between 4.8 V and 2.0 V vs
Li+/Li at 22 °C, the first two cycles
III.1.c. Effect of molten salt to precursor ratio
Besides finding a new washing solvent, we examine the effect of salt to precursor ratio
to the synthesis at the same time. According to our classical approach, salt to precursor
ratio is 4:1 (molar basis), whereas new samples are prepared in lower (2:1) and higher
(10:1) salt to precursor ratios. The prepared mixtures are heated at 300 °C for 2 hours,
calcinated at 670 °C for 3 hours and quenched. The powders are washed with water,
filtered and dried.
Figure III - 12 shows the XRD patterns of the samples that are prepared in different
salt to precursor ratios. In each case, disordered rocksalt Li2.2NiTi0.2Nb0.6O4 phases are
successfully obtained. Such result indicates that even smaller salt to precursor ratio
(2:1) is enough to prepare those phases. However, an impurity phase indexed as
NaNbO3 is observed in diffraction patterns of salt to powder ratios (4:1) and (10:1). In
contrast, we do not observe NaNbO3 impurity in powder synthesized in low salt to
69
precursor ratio (2:1). This might be because smaller amount of salt is used for the
synthesis, such low amount might not be seen in the diffraction patterns.
Figure III - 12: XRD patterns of Li2.2NiTi0.2Nb0.6O4 powders prepared by different salt
(NaCl:KCl) to precursor ratios: 2:1 (1), 4:1 (2) and 10:1 (3)
Figure III - 13: Half-cell performances of Li2.2NiTi0.2Nb0.6O4 prepared by different precursor to
salt ratios: 2:1, 4:1 and 10:1, electrodes cycled at a rate of C/50 between 4.8 V and 2.0 V vs
Li+/Li at 22 °C, at the first cycle
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As illustrated in Figure III - 13, capacities of Li2.2NiTi0.2Nb0.6O4 powders prepared with
salt to precursor ratios of 4:1 and 2:1 are almost the same, whereas much smaller
charge and discharge capacities are obtained from the powder prepared with the
higher salt to precursor ratio (10:1). Those results support our assumption that eutectic
NaCl-KCl salt reacts with the precursors and produce an impurity phase that blocks
the lithium diffusion and leads to smaller capacities. Therefore, as a next step for the
synthesis improvement, we decide to change the eutectic NaCl-KCl with another
eutectic salt that is more stable and does not react with the reactants during the
synthesis.
III.1.d. Molten salt synthesis with a different eutectic salt (LiCl-KCl)
In the literature, different eutectic salts can be found (Table III - 1). Among them, we
select 0.59LiCl-0.41KCl eutectic salt that has a low melting point (354 °C.) We apply
the same synthesis procedure that was described before, but we just change the type
of molten salt. Instead of using eutectic NaCl-KCl salt, 0.59LiCl-0.41KCl eutectic salt in
a salt to precursor ratio of 2:1 is used to synthesize Li2.2NiTi0.2Nb0.6O4 powders.
Table III - 1: Melting point of salts and eutectic salts
At the beginning, firing at low temperature is considered because the melting point of
the new salt is 354 °C. Two different powders are prepared based on the same recipe.
One sample is heated at 450 °C for 12 hours, thereafter it is quenched. Another is
SaltEutectic
molar ratio
melting
point/°C
LiNO3 255
LiOH 477
LiCl 610
LiBr 552
LiOH-LiCl 0.65 : 0.35 277
LiCl-KCl 0.59 : 0.41 354
LiOH-LiNO3 0.40 : 0.60 243
NaCl 801
KCl 771
NaCl-KCl 1.0 : 1.0 667
71
synthesized with the classical method; it is heated at 300 °C for 2 hours and calcined at
670 °C for 3 hours and quenched. The powders are washed with ethylene glycol,
filtered and dried.
Figure III - 14: XRD patterns of Li2.2NiTi0.2Nb0.6O4 powders prepared with different salts:
eutectic NaCl-KCl and 0.59LiCl-0.41KCl at 450 °C and 670 °C
Figure III - 14 shows XRD patterns of three different powders prepared by eutectic
NaCl-KCl and 0.59LiCl-0.41KCl salts. When the powder is prepared by 0.59LiCl-
0.41KCl salt and heated at 450 °C for 12 hours, no disordered rocksalt phase is formed.
The small peaks detected in diffraction patterns belongs to TiO2 and Nb2O5 precursors.
Such a result indicates that diffusion of titanium and niobium oxide precursors are still
low in the 0.59LiCl-0.41KCl salt, even though melting point of 0.59LiCl-0.41KCl is
lower (354 °C). On the contrary, disordered rocksalt phase is observed in the powder
prepared at 670 °C. That result concludes that much higher temperatures are required
to achieve disordered rocksalt phase.
In Figure III - 15, new sample delivers similar charge capacity (145 mAh.g-1) as classical
material, however the discharge capacity is slightly smaller (95 mAh.g-1).
72
Figure III - 15: Half-cell performance of Li2.2NiTi0.2Nb0.6O4 prepared by different eutectic salts:
NaCl-KCl and 0.59LiCl-0.41KCl, electrodes cycled at a rate of C/50 between 4.8 V and 2.0 V
vs Li+/Li at 22 °C, at the first cycle
Figure III - 16: Morphology of Li2.2NiTi0.2Nb0.6O4 powders synthesized with: NaCl-KCl (1:1)
and 0.59LiCl-0.41KCl salts
73
In Figure III - 16, both powders contain agglomerates of spherical particles. For a
synthesis at the same temperature (670 °C), particles become slightly bigger using
eutectic LiCl-KCl salt for the synthesis. This could be related to the diffusion of species;
ion diffusion rate of reactants could be higher in LiCl-KCl salt melt.
In overall, we successfully synthesize Li2.2NiTi0.2Nb0.6O4 in disordered rocksalt phase
using a different eutectic salt (0.59LiCl-0.41KCl) and the electrochemical performances
are almost similar. Therefore, we decide to continue synthesizing Li2.2NiTi0.2Nb0.6O4
powders with 0.59LiCl-0.41KCl salt.
III.1.e. Effect of calcination temperature and duration
Another approach to improve the capacity of disordered rocksalt Li2.2NiTi0.2Nb0.6O4 is
to synthesize the powders at low temperatures. In classical method, the mixtures are
heated at 300 °C to remove the acetate and decomposed at 670 °C. To investigate the
effect of temperature to the synthesis, two samples are prepared and decomposed at
570 °C and 670 °C separately. To discover the effect of firing duration, three samples
are heated at 670 °C for different durations: standard 3 hours (standard), 1 hour and
0.5 h (30 mins).
Figure III - 17 shows samples calcined at different temperatures and durations. In XRD
pattern of the sample calcined at 570 °C, dissymmetrical diffraction peaks are
observed, indicating that new sample contains another phase. The reason might be
slow diffusion of the reactants at 570 °C. For three powders calcined at 670 °C at
different durations, disordered rocksalt phases are successfully obtained. Such result
indicates that diffusion of reactants is high enough at 670 °C, disordered rocksalt phase
can be achieved even at short firing durations (< 3 hours).
74
Figure III - 17: XRD patterns of Li2.2NiTi0.2Nb0.6O4 powders synthesized different synthesis
temperatures and durations: powder synthesized at 670 °C for 3 hours and at 670 °C for 1
hour, 570 °C 3 hours and 670 °C 30 mins (0.5 hour)
Figure III - 18: Cell performances of Li2.2NiTi0.2Nb0.6O4 synthesized at 670 °C for different
durations: 1 h (a), 3 h (b) and 0.5 h (c)
Figure III - 18 shows cell performances of Li2.2NiTi0.2Nb0.6O4 synthesized at 670 °C for
different durations. For Li2.2NiTi0.2Nb0.6O4 powders synthesized for 1 h, reversible
capacities improve, a charge capacity of 166 mAh.g-1 and a discharge capacity of 108
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mAh.g-1 are obtained. This might be related to the size of the particles; smaller particles
should have been obtained at shorter firing duration. As the lithium ion diffusion
length is reduced, thus specific capacities are enhanced. However, this is not valid for
Li2.2NiTi0.2Nb0.6O4 powders synthesized for 0.5 h, as it is explained with non-
homogeneous phase formation previously.
Figure III - 19: Cycling life of Li2.2NiTi0.2Nb0.6O4 synthesized at 670 °C for different durations:
3 h, 1 h and 0.5 h
In Figure III - 19, electrodes show similar cycling life, except sample synthesized for
0.5 h. Here, we conclude that cell performance of Li2.2NiTi0.2Nb0.6O4 is improved by
synthesizing the powders at shorter firing duration (1 h).
76
III.2. Understanding Studies
After optimizing the electrochemical performance of Li2.2NiTi0.2Nb0.6O4, structural
studies are conducted to understand why the capacity of this material is limited.
Li2.2NiTi0.2Nb0.6O4 shows high irreversibility at the first cycle, but a good reversibility
for the following cycles. The reason of the irreversibility at the first cycle is unknown.
As previously shown, lithium diffusion in Li2.2NiTi0.2Nb0.6O4 is better at high
temperatures (55 °C) compared to room temperatures (22 °C), we draw the hypothesis
that Li2.2NiTi0.2Nb0.6O4 might have lithium diffusion limitations. At the end of the
charge, Li ions are extracted from the shell of the particles and material’s core remains
richer in lithium (Figure III - 20). During discharge, lithium ions insert into the
structure and fill empty spaces in the material surface. As the Li diffusion is limited,
Li accumulates faster at the surface than in the bulk and may create a diffusion barrier.
Figure III - 20: Hypothesis on lithium diffusion mechanism through Li2.2NiTi0.2Nb0.6O4
To test this hypothesis, the cell is cycled at variable charge potentials beginning from
a low potential of 3.9 V up to 4.8 V with a fixed discharge potential (2 V), as illustrated
in Figure III - 21.
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Figure III - 21: Cell performance of Li2.2NiTi0.2Nb0.6O4 that was cycled at variable charge
potentials beginning from a low potential of 3.9 V up to 4.8 V with a fixed discharge potential
(2 V) at a rate of C/50 vs Li+/Li at 22 °C
Interestingly, we observe that the irreversibility arises at low state of charge (SOC) and
almost no charge and discharge capacities are obtained upon cycling between 3.9 V -
2 V. At potentials greater than 4.2 V, lithium exchange process becomes more
reversible upon cycling. However, following our hypothesis, we would expect to see
a good reversible behavior at low state of charge and an increase in the irreversible
capacities with increasing SOC. After these findings, we conclude that the result does
not align with our hypothesis. Therefore, the irreversibility might arise from the
structural modifications instead of lithium diffusion limitations. Now, we have
another hypothesis in which there might be a structural change at the end of the
charge. Thus, new structure allows only few lithium ion exchange. This can be further
investigated with ex situ and in situ X-ray diffraction techniques.
78
III.2.a. Ex Situ X-ray Diffraction Analysis
To perform ex situ X-ray diffraction, three electrodes are prepared and one electrode
is charged to 4.8 V while another is discharged to 2 V. Detailed procedure can be
found in Chapter II.2.
Figure III - 22: Ex situ XRD analysis of electrodes at different states of charge: pristine
electrode, electrodes at end of charge (4.8 V) and discharge (2 V)
Figure III - 22 shows ex situ XRD analysis of electrodes at different states of charge.
After aligning aluminum peaks, we observe that both peak intensities and lattice
parameters reduce at the end of charge (4.147 Å). Besides, shape of the diffraction peak
at 63.5° becomes much broader, indicating that charged electrode might contain
another phase. On the contrary, lattice parameters of discharged electrode are slightly
bigger than the pristine electrode (4.187 Å). Additionally, the diffraction pattern of
discharged electrode is quite different from the pristine electrode. This result indicates
that structural change occurring during cycling is not completely reversible. It seems
that final composition might not have the same structure as the original composition.
Thereafter we decide to investigate the change of lattice parameters step by step with
in situ X-ray diffraction.
79
III.2.b. In Situ X-ray Diffraction Analysis
Figure III - 23 shows cycling curve of in situ cell containing Li2.2NiTi0.2Nb0.6O4 powder
that is cycled with GITT mode of cycling (for cell preparation, see Chapter II.2). The
cell is cycled for 0.5 hour, and then current is interrupted. Afterwards, the cell is rested
at 3.5 hours and X-ray diffraction is conducted during relaxation.
Figure III - 23: Cycling curve of in situ cell containing Li2.2NiTi0.2Nb0.6O4 powder, GITT mode
of cycling
Figure III - 24 shows global view in situ XRD diffraction pattern of Li2.2NiTi0.2Nb0.6O4
powder. In diffraction patterns, aluminum and beryllium peaks are identified as well
as the disordered rocksalt phase. For clarity sake, some diffraction patterns (pristine,
end of charge and discharge) are selected and shown in Figure III - 25.
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Figure III - 24: Global view of in situ XRD diffraction pattern of Li2.2NiTi0.2Nb0.6O4 powder
Focusing on the high 2θ region (62°- 64°) in Figure III - 25, we have similar observations
previously described in ex situ X-ray diffraction section. The peak intensities reduce
and the diffraction pattern of charged powder seems to contain another disordered
rocksalt besides disordered rocksalt Li2.2NiTi0.2Nb0.6O4. To be sure about the second
phase, we decide to refine lattice parameters using Fullprof program profile matching.
81
Figure III - 25: In situ X-ray diffraction pattern of Li2.2NiTi0.2Nb0.6O4 powder at different
states of charge: pristine, end of charge and discharge
At first, we start to do fitting the pattern according to our previous assumption that
second disordered rocksalt phase is progressively forming during charging. However,
we detected a discrepancy in the fitting results, thus we were not able to find a good
fit after few steps of lithiation. Afterwards, we assume that the second phase should
have been at the beginning, in other words initial powder (pristine state) might have
already contained two phases before cycling. The refinement of the initial powder is
then carried with both hypotheses, showing a real improvement when introducing a
second phase (Figure III - 26).
82
Figure III - 26: Results of the refinements of the initial powder. Top: with 1 phase, bottom:
with 2 phases
We find that the two phases have very close lattice parameters (Phase 1: 4.1824 Å,
Phase 2: 4.1848 Å), explaining why this phase has not been detected before. During
charge, the lattice parameters are decreasing in a different way and we observe
different slopes (Figure III - 27a).
83
Figure III - 27: Change of (a) lattice parameters and (b) size of the crystalline domains of
Li2.2NiTi0.2Nb0.6O4 powder during in situ XRD analysis
Lattice parameters increase during discharge, but they are slightly bigger than those
in the pristine states (Figure III - 27a). This result indicates that the final phases are
different from the pristine states. The irreversible structural change may thus be one
of the causes of the irreversible capacity observed during cycling.
Meanwhile, we calculate the size of the crystalline domains in Li2.2NiTi0.2Nb0.6O4
powder using Scherrer equation described in Chapter II.2. In Figure III - 27b,
crystalline sizes are decreasing, indicating that powder pulverizes upon charging. It
seems that disordering, such as dislocations or defects happen in the material. During
discharge, we observe abnormal change, a parabolic distribution of crystalline sizes.
The change in crystalline sizes is irreversible as the final crystalline sizes are much
smaller than those in pristine states. This apparition of defects (either dislocation or
long-range disordering) may also be a cause of the capacity fading of this material.
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III.3. Conclusion
The objective of this chapter was to develop overlithiated disordered rocksalt
Li2.2NiTi0.2Nb0.6O4 with high capacity. Our initial synthesis has resulted in smaller
charge (120 mAh.g-1) and discharge (74 mAh.g-1) capacities than the theoretical
capacity (264 mAh.g-1). To reach desirable capacities, we optimize our synthesis
procedure considering various synthesis parameters.
Table III - 2: Comparison of initial and final synthesis parameters
After optimization of the synthesis parameters (Table III - 2), we finally succeed to
obtain charge and discharge capacities of 166 mAh.g-1 and 108 mAh.g-1 at room
temperature respectively. However, we observe that there is always a high
irreversibility at the end of the first cycle. At first, we suppose that irreversibility arises
from lithium diffusion problem. After testing our hypothesis, we obtain a result that
does not align with our first hypothesis. Thereafter, we test our second hypothesis that
we consider a possible structural change during charge-discharge.
We observe from ex situ XRD result that there should be a second disordered phase
forming during charge process. To validate this, we also perform in situ XRD analysis
of electrodes at different SOC. Refinement of the initial powder indicates that second
disordered rocksalt phase has already been present in the pristine state. In addition,
we find out that the final phases at the end of the discharge are different from those in
the pristine states. From the study of size in the crystalline domains, we conclude that
disordering happening during cycling causes irreversible structural modifications.
85
In general, Li-rich disordered rocksalt oxides discussed in the literature offer large
reversible capacities due to their high amount of lithium content. Our findings show
that this could be possible if the material has pure disordered structure that can allow
lithium ions exchange reversibly during cycling. It seems that diffusion limitations in
disordered rocksalt oxides have been partly solved with overlithiation, nevertheless
preparing those materials in pure disordered structures might gain importance for
future studies.
86
Chapter IV: Developing Li-rich
Rocksalt Metal Sulfides-Selenides:
Li2TiSexS3-x
The first part of this chapter describes mechanochemical synthesis of Li2TiS3
compound as well as structural, morphological and electrochemical characterizations
of this material. To improve the cycle life of Li2TiS3, we propose selenium substitution
with mechanochemical synthesis. Therefore, second part of the chapter describes
synthesis optimization of sulfides-selenides solid solutions, structural and
morphological characterizations and the electrochemical performance comparison of
new materials with Li2TiS3 material.
Introduction
Metal sulfide family is extensively studied in the literature, especially after layered TiS2
showed unique lithium ion intercalating property (86–88). In parallel to TiS2, layered TiS3
was also studied. Upon discharging the cells, three lithium ions are intercalated into
TiS3 (Eq. IV – 1 and Eq. IV - 2), the higher capacity of 80mAh was reached without
precision on the mass of active material (89–91). Unfortunately, this capacity was not
recovered upon charging the cells because two Li ions were not reversible, thus rapid
capacity fading was observed. This was attributed to the structural instability of
layered TiS3 during cycling (62).
TiS3 + 2Li+ +2e- → Li2TiS3 (Eq. IV - 1)
Li2TiS3 + Li+ + e- →Li3TiS3 (Eq. IV - 2)
In 2014, A. Sakuda et al. prepared Li2TiS3 as a disordered rocksalt phase by
mechanochemical synthesis and they observed that this material could provide 425
mAh.g-1 capacity based on 2.5 Li+ ion reversible exchange process. Beyond its
interesting electrochemical performances, authors reported that the cycling life of this
87
material was poor. Rapid capacity fading was observed in the liquid electrolytes,
whereas 92% capacity was maintained after 80 cycles in solid-state electrolytes.
Therefore, this material was suggested for solid-state batteries rather than for
conventional batteries. In this chapter, our first aim is to develop disordered rocksalt
Li2TiS3 with mechanochemical synthesis. Various structural, morphological and
electrochemical characterizations are performed to understand the properties of
Li2TiS3 compound. To solve cycle life problem, we propose substitution of Li2TiS3 with
selenium. Therefore, this chapter also describes the development of a new synthesis
process and new selenium substituted compositions (Li2TiSexS3-x).
IV.1. Disordered Li2TiS3 as a Positive Electrode
IV.1.a. Mechanochemical Synthesis of Li2TiS3 Powders
Li2TiS3 samples are prepared by using the same synthesis recipe as in the literature (62).
Sulfides are known as air sensitive materials; therefore, all samples are prepared in
argon filled glovebox.
In the literature, Li2TiS3 powders are synthesized with planetary miller at milling speed
of 510 rpm for a range of milling durations (from 20 h to 100 h) (62). Pure disordered
rocksalt phase is obtained for milling durations longer than 40 hours though
considerable amount of zirconium contamination is seen in the X-ray diffractogram.
At the beginning of the thesis, our target is to synthesize Li2TiS3 with the same
synthesis recipe, but we would also like to improve the purity of the powders. As
already mentioned in Chapter II.1, powder milling at longer durations causes
undesired reactions and contaminations in the final product as the balls often wear
out. Thus, we consider synthesizing the powders at shorter durations. This can be done
using slightly bigger zirconium balls (5 mm) and jar (50 ml) than those in the literature
(4 mm and 40 ml) with a fix ball-to jar ratio (62:1, in weight). Our synthesis recipe to
prepare Li2TiS3 powders is the following one: 0.4363 g lithium sulfide (Li2S, Sigma
Aldrich, 98%) and 1.0635 g titanium sulfide (TiS2, Sigma Aldrich, 99.8%) are weighed
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and placed into 50 ml stainless steel jar that contains 285 zirconia balls (5 mm in
diameter). Then, the jar is closed tightly and transferred from an argon-filled glovebox
to the planetary ball-milling machine. To find optimum milling conditions, Li2TiS3
powders are prepared using various milling durations: 15h, 20h, 25h and 40h at 510
rpm. After the milling, the jar is transferred from planetary ball-milling machine to the
glovebox to avoid the air contamination during powder recovery. Thereafter, samples
are analyzed with XRD to identify the phase of the powders. Sample preparation for
XRD is done in the glovebox and Kapton film is used to protect the sample from air
contaminations. The procedure is described in Chapter II.2.
Figure IV - 1: X-ray diffraction patterns of disordered rocksalt Li2TiS3 powders prepared with
planetary miller at 510 rpm for 15h, 20h, 25h and 40h durations and SEM picture of powder
milled for 20 hours
In Figure IV - 1, XRD patterns of the four different Li2TiS3 powders are quite similar,
almost the same peak positions and intensities are observed. SEM images of Li2TiS3
powder (milled for 20h) are also displayed in Figure IV - 1. As seen in the images, the
powder contains agglomerates that consist of nanometer sized spherical particles.
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Figure IV - 2: X-ray diffraction pattern of disordered rocksalt Li2TiS3 powder (milled for 20 h)
profile matching in Fullprof program
To find the lattice parameters, X-ray diffraction data of Li2TiS3 powder milled at 20
hours is selected to be used in Fullprof profile matching program (Figure IV - 2). At
first, XRD pattern is fitted to optimize the peak positions and intensities iteratively
until the high accuracy fitting (Chi2<5) is reached. The lattice parameters of Li2TiS3 is
found to be 5.0855 Å, which is very close to the one reported in the literature (62). To
optimize the global peak shape and intensities as well as to refine the lattice
parameters, Rietveld refinement is considered. After searching the structure model
(cif-crystal information file) of Li2TiS3 in databases, we could not find corresponding
cif of Li2TiS3. Afterwards, we decide to use the same Wyckoff positions (WP) of
disordered rocksalt (Fm3̅m) Li2TiO3 for Rietveld refinement. The first refinement is
performed with sof (site occupancy factor) modified using stoichiometry of the
composition (66.6% Li and 33.4% Ti metals and 100% S ligand in Li2TiS3 composition).
Accurate occupancies are obtained after high accuracy fitting (Chi2= 1.479) is reached
(Figure IV - 2). The summary data of Rietveld refinement is illustrated in Table IV - 1.
90
Table IV - 1: Rietveld refinement parameters of Li2TiS3 (WP: Wyckoff positions, sof: site
occupation factor, Biso: atomic displacement parameter), global Chi2 = 1.902
Afterwards, electrochemical performances of the four different samples are tested. The
detailed procedure of slurry coating and coin cell preparation can be found in Chapter
2.3. To have reproducible results, three identical cells are cycled at a rate of C/10
between 3 V and 1.5 V vs Li+/Li.
Figure IV - 3: Half-cell performances of Li2TiS3 cells (powders synthesized at 510 rpm for
different durations) at the first cycle, coin cells cycled at a rate of C/10 between 3.0 V – 1.5 V
Li+/Li at 22 °C
Figure IV - 3 shows the electrochemical performances of the different Li2TiS3 electrodes
at the first cycle and the results are summarized in Table IV - 2. For each material,
discharge capacities are found greater than charge capacities, thus there is always an
extra capacity in Li2TiS3 samples. Discharge average potentials are quite similar.
91
Overall, we select the 20 hours milling condition as optimum milling duration for our
future synthesis though 25 hours milling in Table IV - 2 seems more effective. This
discrepancy arises from the fact that the first synthesis is made with fresh balls at 20
hours milling duration. We later find out that there is not a big difference in terms of
cell performance and capacity retention between milling 20h and 25h when the milling
is operated with used balls.
Table IV - 2: Coin cell results of Li2TiS3 compositions synthesized by ball milling for 15h, 20h,
25h and 40h , where CQ1: first charge capacity, DQ1:first discharge capacity, Vav: average
discharge potential at the first cycle
Table IV - 3: ICP-AES (Inductively coupled plasma atomic emission spectroscopy) of Li2TiS3
powders mechanochemically synthesized at different milling durations. Data is normalized
based on fixed Ti amount (Ti=1.00 mol) in the structure.
To investigate the real stoichiometry and the purity of the final products, Li2TiS3
powders are analyzed with ICP-AES technique and results are shown in Table IV - 3.
After data is normalized based on fixed Ti amount (Ti=1.00 mol), quite diverse atomic
ratios are observed in the powders. We find that powders contain a small trace of
zirconium contamination (0.4 - 4%). The atomic ratios of lithium and sulfur in the
powders increase with milling durations (from 15 hours to 25 hours). However, we
observe that atomic ratios obtained from the sample milled at 40 hours are rather out
15 h 20h 25h 40h
CQ1 (mAh.g-1) 271 300 308 99
DQ1 (mAh.g-1) 334 340 354 267
Vav_discharge (V) 2.23 2.22 2.20 2.10
Li2TiS3
92
the error bar. This may arise from the analysis. ICP samples are dissolved in acidic
aqueous solution, so a part of sulfur may have evolve as H2S and leave the sample
before the analysis.
In summary, four different Li2TiS3 powders are mechanochemically synthesized as
disordered rocksalt phase without any remaining precursors. With our optimized
recipe, we can achieve desired phases at shorter milling durations (< 40 hours) than
those in the literature (62). This can be explained by different size of balls, as slightly
bigger balls are used for the mechanochemical synthesis. As bigger balls create
stronger shocks, the impact energy on the particles increases. This allows reaching
desired phases at shorter milling durations. When the electrochemical performances
of powders are investigated, no direct relation between milling durations and
zirconium contamination as well as the performances is observed. Among all milling
durations, we decide to further synthesize our powders milling for 20 hours.
IV.1.b. Electrochemical Performance of Li2TiS3
Now, we investigate the electrochemical performance of Li2TiS3 electrodes in detail
(Figure IV - 4). At the end of the first charge, 1.50 Li+ ions are extracted from Li2TiS3
and the capacity reaches 250 mAh.g-1.
Figure IV - 4: Half-cell performances and cycling life of Li2TiS3 (unpressed) electrodes cycled
at a rate of C/10 between 3 V - 1.5 V vs Li+/Li: the first (black color), second (red color) and
third (blue color) cycles
93
During the discharge, 2 Li+ ions are inserted into the structure, the composition
changes into Li2.5TiS3 and 338 mAh.g-1 of discharge capacity are delivered. The second
and third cycling curves are fully reversible and extra capacity is observed at the end
of the discharge. That is also observed in different rocksalt sulfide compositions
(Li3NbS4 and Li2SnS3) that show larger discharge capacity than the charge capacity
(92,93). However, the origin of this extra capacity is still unknown. Moreover, Li2TiS3
electrodes show poor cycling stability, rapid capacity fading is observed after three
cycles (Figure IV - 4). Poor cycle life of Li2TiS3 cells is already mentioned in the
literature (62). One approach to improve cycle life of Li2TiS3 electrodes is to study the
effect of electrode pressing. To do so, the electrodes are pressed under 10 tons in the
glovebox and they are tested in coin cells.
Figure IV - 5: Effect of pressing to the cycling life: Half-cell performances of Li2TiS3 electrodes
pressed electrodes and cycling life of pressed and unpressed Li2TiS3 electrodes cycled at a rate
of C/10 between 3 V - 1.5 V vs Li+/Li: the first (black color), second (red color) and third (blue
color) cycles
Figure IV - 5 shows electrochemical performances of pressed Li2TiS3 electrodes under
10 tons in a cold-press machine in the glovebox. Here, pressed electrodes deliver much
higher charge capacity (300 mAh.g-1) compared to the unpressed electrodes (250
mAh.g-1). In addition, reversible 340 mAh.g-1 charge and discharge capacities are
obtained at the end of the second and third cycles. Reduction of electrode porosity by
pressing should increase electronic conductivity, thus reversible capacities improve.
94
Besides, we observe that pressing the electrodes also have a positive impact on cycle
life, which is improved considerably. As the pressed electrodes have improved
electronic conductivity, we expected that their rate capability should improve as well.
To investigate this, two different cells are charged to 3 V at C/10 rate. (Figure IV - 6).
To determine cell behavior at high rates, a signature test is used: the cells are
discharged directly at a rate of 10D, when the discharge cut-off voltage is reached, the
current is then reduced to 5D, and this protocol is then continued with 2D, D, D/2, D/3,
D/5 and finally D/10 consecutive rates to reach a full discharge of the material. Between
two discharge phases, electrodes are rested for two hours. In Figure IV - 6, we observe
that Li2TiS3 electrodes retain 74% of discharge capacity at D rate, 64% capacity at 5D
rate and 51% capacity at 10D rate.
Figure IV - 6: Cycling curve of signature test: Current (mA) vs Test time (h) and rate
capability of pressed Li2TiS3 electrodes that are charged at a rate of C/10 to 3 V and discharged
by different discharge (D) rates to 1.5 V (duration of rest time is two hours between two
discharge phases)
The reason of poor cycle life Li2TiS3 cells may be due to sulfur reaction with the
electrolyte. In the literature, it is reported that lithium sulfur batteries also suffer from
poor cycle life (94–96). That is attributed to reaction of sulfur with carbonate-based
electrolytes, therefore tetraglyme-dioxolane-based electrolytes are suggested (97). In
our experiments, we decide to test Li2TiS3 electrodes in tetraglyme-dioxolane based
electrolytes as well.
95
Figure IV - 7: Capacity and average potential of Li2TiS3 electrodes cycled at a rate of C/10
between 3 V - 1.5 V in two different electrolytes: carbonate-based electrolyte (1 M LiPF6 in
EC:PC:DMC 1:1:3 in volume) and tetraglyme-dioxolane based electrolyte (1M LiTFSI +
0.1M LiNO3 in TEGDME:DIOX 1:1 in volume)
Figure IV - 7 shows Li2TiS3 electrodes cycled in two different electrolytes, such as
carbonate-based electrolyte (1 M LiPF6 in EC:PC:DMC 1:1:3 in volume) and
tetraglyme-dioxolane based electrolyte (1M LiTFSI + 0.1M LiNO3 in TEGDME:DIOX
1:1 in volume). Here we observe that cycle life of Li2TiS3 electrodes in carbonate-based
electrolyte is better than the one in tetraglyme-dioxolane-based electrolyte. Besides,
same charge average (2.46 V) and discharge average (2.23 V) potentials are observed
in the two systems. Thus, we decided to keep testing the electrodes with carbonate-
based electrolyte.
IV.1.c. Air Sensitive Property of Li2TiS3
Li2TiS3 is an air-sensitive material, thus this property slows down the rate of process
from material synthesis to coin-cell preparations because it requires always working
under inert atmosphere, often in an Ar-filled glovebox. For industrial perspective,
material stability is important, especially for upscaling process.
In this part, the air-sensitivity of Li2TiS3 is examined. Five different electrodes are
prepared for XRD analysis in the glovebox, later they are left in a fume hood and a dry
room for one day and one week durations. Here, working under fume hood represents
working under air condition. As the dry room has less humidity (dew point: -40oC)
96
compared to air atmosphere, working in dry room conditions will show oxygen
sensitivity rather than humidity.
Figure IV - 8: X-ray diffraction patterns of Li2TiS3 electrodes at different environments:
pristine electrode, after 1 day in dry room, after 1 week in dry room, after 1 day in fume hood,
after 1 week in fume hood
Figure IV - 8 shows XRD pattern of pristine Li2TiS3 electrodes exposed under different
durations and atmospheres. When the electrodes remain for 1 day in the dry room,
peak intensity at 2θ = 36° decreases, but the crystallinity is conserved. The same is
observed at the electrodes remained for 1 week in the dry room. In contrast, almost no
peaks are detected in electrodes remained under fume hood. We detect various TixOy
compounds in electrodes that remained for 1 week under fume hood, indicating that
sulfur and lithium atoms react with air and they may produce H2S gas and amorphous
Li2CO3 compounds that are not observed in the X-ray diffraction patterns.
Figure IV - 9 and Figure IV - 10 show coin cell test results of Li2TiS3 electrodes exposed
under different durations and atmospheres. Electrodes remaining in dry room deliver
almost similar charge and discharge capacities for the first 8 cycles. Thereafter, small
capacity fading is observed in electrodes remained in the dry room. In contrary,
97
smaller capacities at the first cycle are observed in electrodes remaining in fume hood.
Such results indicate that Li2TiS3 material is more sensitive to humidity and it is rather
stable under oxygen environment. This also indicates the possibility of production and
slurry coating of this material in dry room.
Figure IV - 9: (a) Coin cells tests results of standard Li2TiS3 electrodes (black) in a comparison
of electrodes after 1 day (green) and 1 week (red) in dry room
Figure IV - 10: Coin cells tests results of Li2TiS3 electrodes after 1 day (blue) and 1 week
(black) in fume hood
98
• Lithiation of Li2TiS3 with Butyl-lithium Reaction
As previously mentioned regarding Li2TiS3 electrodes, there is always extra capacity
at the first cycle, thus the final composition is different from the initial composition.
Such compositional deviation disappears at the end of the second cycle, fully reversible
lithium intercalation is observed. To eliminate the irreversibility at the first cycle and
increase charge-discharge capacities, we aim at preparing Li2TiS3 in its more lithiated
version, (Li3TiS3). In early studies on TiS2 and TiS3 systems, Dines et al. reported that
butyl-lithium reaction was an effective way for lithiation of the positive materials (6).
Therefore, lithiation of Li2TiS3 is made by butyl-lithium experiment using the same
procedure as in the literature.
• Experimental procedure of Butyl-lithium reaction
Based on the reaction in Eq. IV - 3, 0.025 mol Li2TiS3 powder, 14.60 ml cyclohexane and
magnetic stirrer are put into a small reactor (20 ml). 0.0375 mol butyl lithium (50%
excess) is added slowly into the reactor, the reactor is closed with its cover and left to
stir at 450 rpm in the glovebox (Figure IV - 11a).
Butyl-lithium reaction:
nC4H9Li + Li2TiS3 → Li3TiS3 + ½ C8H18 (Eq. IV - 3)
Mechanochemical reaction:
3/2Li2S + Ti + 3/2S → Li3TiS3 (Eq. IV - 4)
After stirring for a week, the mixture is transferred into small flasks (2 ml) and put into
a centrifuge (Figure IV - 11b and Figure IV - 11c). They are stirred under 14000 rpm for
20 min, then the powders settle in the bottom, while excess cyclohexane solution is
removed from the flask. To remove the organic impurities, the powder is washed with
anhydrous cyclohexane three times (Figure IV - 11d). Thereafter, it is left to dry for two
days in the glovebox. As a comparison, we also prepare Li3TiS3 samples with
mechanochemical synthesis.
99
Based on reaction in Eq. IV - 4, 1.5 mol of Li2S, 1 mol of Ti and 1.5 mol of S are weighed
and placed into the jar. The jar is transferred into planetary ball milling machine and
milling conditions are set to 510 rpm for 20 hours.
Figure IV - 11: Butyl lithium experimental set-up: (a) stirring for a week, (b) taking the
mixture into the small flask, (c) centrifuge of the samples, (d) washing the powder with
cyclohexane
Figure IV - 12 shows the XRD patterns of Li2+xTiS3 powders prepared by butyl lithium
experiment and ball milling. We observe Li2S precursor in X-ray diffraction patterns of
powder prepared by mechanochemistry. It seems that powder prepared by butyl-
lithium also contain small amount of Li2S, which is a side product. The presence of Li2S
impurity indicates that it is more difficult to stabilize Ti3+ in the disordered rocksalt
structure.
Figure IV - 12: XRD pattern of Li2+xTiS3 comparison of butyl lithium reaction vs ball milling
100
Figure IV - 13: Half-cell performances of LixTiS3 that is lithiated by butyl lithium experiment,
electrodes cycled at a rate of C/10 between 3 V - 1.5 V vs Li+/Li; the first cycle (black color),
second cycle (red color) and third cycle (blue color)
In Figure IV - 13, we observe that Li2+xTiS3 cells show almost fully reversible charge-
discharge curve though capacities are rather smaller (~110 mAh.g-1) than expected.
This may be caused by the unsuccessful coating, because slurry gelified during
preparation. It seems that unreacted Li2S stays on the particle surface and reacts with
PVdF binder.
IV.2. Selenium Substituted Li2TiS3
As already mentioned, the cycling stability of Li2TiS3 is rather poor and therefore it
should be improved. We consider that substitution can be a promising solution. For
instance, sulfur anions can be replaced with an alternative anion. Looking at the
periodic table, selenium comes into prominence as a novel prospected material since
it belongs to the same periodic group (6A), hence it is expected to show similar
elemental activity as sulfur. In the literature, it is also well-known that selenium
substitution improves cycling stability of Li2S cells (98,99). However, before the
synthesis, one must consider the theoretical capacity of new compositions. Theoretical
capacities based on two lithium ions exchanged during cycling are reduced by
101
selenium substitution since atomic mass of selenium (78.97 g.mol-1) is much bigger
than atomic mass of the sulfur (32.06 g.mol-1) (Table IV - 4). Therefore, we first consider
the following compositions Li2TiSeS2 and Li2TiSe2S.
Table IV - 4: Theoretical capacities of Se substituted compositions based on two Li+ ions
exchange process
To synthesize selenium substituted Li2TiS3 samples, necessary precursors are searched
from the producers, however neither Li2Se nor TiSe2 are commercially available.
Therefore, alternative synthesis routes are searched. In the literature, the same
precursor problem is mentioned in the development of Li2Se batteries and powders
are synthesized by liquid synthesis routes that are quite challenging and require more
than one step (98). To make one pot and much faster synthesis, we aim at preparing
new compositions with mechanochemical synthesis using metallic precursors, such as
metallic Ti and Se precursors rather than molecular TiSe2 precursor.
IV.2.a. Synthesis Improvements
Before preparing new materials, it is essential to adapt our previous synthesis of Li2TiS3
to the use of metallic Ti and S precursors instead of TiS2. In this way, we can compare
this new synthesis route with the standard route where there are already enough
structural and electrochemical characterizations. In fact, TiS2 precursor is quite
expensive (22 €.g-1) compared to metallic Ti and S precursors (Ti: 6 €.g-1, S: 2 €.g-1),
therefore using metallic precursors would also bring benefits in terms of
manufacturing costs. However, it should be noted that these prices are obtained from
102
the Sigma Aldrich for lab-scale powders, therefore they may not be representative for
industrial powders.
• Wet Mechanochemical Synthesis of Li2TiS3
At the beginning, we aimed at synthesizing Li2TiS3 powder with Li2S powder, metallic
Ti and S powders. Our four different synthesis trials show that metallic precursors do
not react with Li2S precursor, they are rather deposited on the jar surface. To eliminate
this, we consider milling of the precursors with anhydrous solvent. In our first trial,
we try to synthesize Li2TiS3 through a wet mechanochemical synthesis for 20 h and 25
h milling durations. However, we always detect traces of remaining Li2S precursor
with rocksalt phase in XRD patterns. Pure rocksalt Li2TiS3 phase is finally obtained
when the powder is wet milled for 32 hours with already used balls (100).
Synthesis Procedure in detail: 0.436 g Li2S (98%, Sigma Aldrich), 0.435 g Ti (99.98%,
Alfa Aesar) and 0.609 g S (99.99%, Alfa Aesar) precursors are placed into 50 ml stainless
steel jar that contains 285 zirconia balls (5 mm in diameter). Thereafter, anhydrous
hexane (99.9%, Sigma Aldrich) is poured into the jar until covering the entire ball
surface. Milling is operated at 510 rpm for 32 hours. The milling run time is set as 1
min, but pause time is set as 4 min. After milling, jar is transferred from planetary ball
milling machine to glovebox. To evaporate hexane, jar remains open for 3 days in the
glovebox. Then, powder is recovered in the glovebox.
Figure IV - 14 shows the XRD pattern of Li2TiS3 prepared by wet and dry
mechanochemical syntheses. In the two different synthesis processes, Li2TiS3 powders
are successfully achieved in disordered rocksalt phase. Both powders have similar cell
parameters (awet: 5.0853 Å and adry: 5.0855 Å). Moreover, coin cell test results in Figure
IV - 15 show similar charge and discharge capacities of Li2TiS3 prepared by wet and
dry mechanochemical syntheses. Those results confirm that disordered rocksalt Li2TiS3
can be prepared by a new synthesis process that is more flexible and less expensive.
With this new process, different metal sulfides, Li2MS3 or selenides, Li2MSe3 (M: Fe, Ni,
103
Co, Nb, Mo, etc) and different chalcogenides, Li2TiX3, (X: S, Se, Te, etc.) can be prepared
(100,101).
Figure IV - 14: Comparison of wet and dry mechanochemical syntheses: XRD pattern of
Li2TiS3 powders prepared by wet and dry mechanochemical syntheses
Figure IV - 15: Comparison of half-cell performances of Li2TiS3 cells (powders prepared by wet
and dry mechanochemical syntheses)
104
• Wet Mechanochemical Synthesis of Li2TiSeS2
0.336 g Li2S (98%, Sigma Aldrich), 0.410 g TiS2, 0.175 g Ti (99.98%, Alfa Aesar) and 0.578
g Se (99.9%, Alfa Aesar) precursors are placed into 50 ml stainless steel jar that contains
285 zirconia balls (5 mm in diameter). Then, appropriate amount of anhydrous hexane
(99.9 %, Sigma Aldrich) is added until covering the entire ball surface. Milling is
operated at 510 rpm for 20 hours.
• Wet Mechanochemical Synthesis of Li2TiSe2S
0.274 g Li2S, 0.285 g Ti (99.8%, Alfa Aesar) and 0.941 g Se (99.9%, Alfa Aesar) precursors
are placed into 50 ml stainless steel jar that contains 285 zirconia balls (5 mm in
diameter). Then, appropriate amount of anhydrous hexane (99.9 %, Sigma Aldrich) is
added until covering the entire ball surface. Milling is operated at 510 rpm for 20 hours.
After milling, jar is transferred from planetary ball milling machine to glovebox. To
evaporate hexane, jar remains open for 3 days in the glovebox. Then, powder is
recovered in the glovebox.
Figure IV - 16: XRD patterns of Li2TiSexS3-x powders prepared with wet mechanochemical
synthesis
105
XRD patterns of sulfur selenium solid solutions (Li2TiSexS3-x) in Figure IV - 16 show
that selenium substituted samples are successfully obtained in disordered rocksalt
phase by wet mechanochemical synthesis process. XRD patterns of Li2TiSeS2 and
Li2TiSe2S compositions shift to small 2Ѳ angles, indicating that lattice parameters
expand with selenium substitution since the ionic radius of selenium (Se2-, 198 pm) is
larger than the ionic radius of sulfur (S2-, 184 pm). To confirm if selenium is successfully
substituted into Li2TiS3 material, diffraction patterns of Li2TiSeS2 and Li2TiSe2S
powders should be simulated. Before simulating the diffraction patterns, lattice
parameters are found with Profile Matching in Fullprof program. Patterns are fitted
using the same .pcr file generated for Li2TiS3 powder and the lattice parameters are
found after high accuracy fitting (Chi2<5) is reached (Figure IV - 17 and Figure IV - 18).
Afterwards, Rietveld refinement is done using the modified occupancies based on the
stoichiometry of each compositions. The obtained lattice parameters and occupancies
are summarized in Table IV - 5 and Table IV - 6.
Figure IV - 17: XRD pattern of Li2TiSeS2 profile matching in Fullprof program
106
Table IV - 5: Rietveld refinement parameters of Li2TiSeS2 (WP: Wyckoff positions, sof: site
occupation factor, Biso: atomic displacement parameter)
Figure IV - 18: XRD pattern of Li2TiSe2S profile matching in Fullprof program
Table IV - 6: Rietveld refinement parameters of Li2TiSe2S (WP: Wyckoff positions, sof: site
occupation factor, Biso: atomic displacement parameter)
107
To confirm if selenium is incorporated into the crystal lattice, diffraction patterns are
further simulated. Previously found lattice parameters and occupancies are used to
generate simulated diffraction patterns of Li2TiSeS2 and Li2TiSe2S (Figure IV - 19).
Figure IV - 19: Simulated XRD patterns of Li2TiSexS3-x powders
Figure IV - 20: Experimental and simulated peak intensity ratio of (311) and (222) planes,
I(311) / I(222) for Li2TiS3, Li2TiSeS2 and Li2TiSe2S powders
108
Simulated diffraction patterns in Figure IV - 19 look like the experimental diffraction
patterns (Figure IV - 16). We observe again the shift of diffraction patterns into small
2Ѳ angles as well as the increase in peak intensities with selenium substitution. This is
even more characteristic for doublet at 61°, where peak intensity of (311) plane
increases. To confirm selenium incorporation into crystal lattice, the ratio of peak
intensities at (311) and (222) planes, I(311)/I(222), seen in experimental and simulated
diffraction patterns should be calculated for three samples as the (111) plane is biased
by Kapton®. In Figure IV - 20, we observe that the ratio of I(311)/I(222) is almost the same
for Li2TiSeS2 and Li2TiSe2S powders though it is slightly different for Li2TiS3 powder.
This result supports successful selenium substitution into Li2TiS3 material.
SEM images of Li2TiSe2S and Li2TiSeS2 powders are represented in Figure IV - 21. Like
Li2TiS3 powder, selenium substituted samples contain agglomerates that consist of
spherical particles. It seems that the agglomerates tend to have a plate-like morphology
with selenium substitution.
Figure IV - 21: SEM images of Li2TiSeS2 (left) and Li2TiSe2S (right) powders
109
IV.2.b. Electrochemical performances of Li2TiSexS3-x
• Galvanostatic Tests of Li2TiSexS3-x
The electrochemical performances of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes are
investigated considering the possible redox processes (Figure IV - 22).
Figure IV - 22: Half-cell performances and cycle life of Li2TiS3, Li2TiSeS2 and Li2TiSe2S
electrodes cycled at a rate of C/10 between 3 V – 1.5 V vs Li+/Li at 22 oC
110
Figure IV - 22 shows coin cell tests results at the first three cycles of Li2TiS3 upon cycling
at a rate of C/10 between 3 V – 1.5 V vs Li+/Li. As previously described, 1.77 Li+ ions
are extracted from Li2TiS3 at end of initial charge and 300 mAh.g-1 of capacity is
delivered. Here, we explain the possible redox mechanism based on the assumptions
that all the atoms are stabilized at their valence states (Ti4+ and S2-) in pristine electrode
and there is no loss of sulfur atoms during charge-discharge. During charge, sulfur
redox should be active since Ti is already at his maximum valence of Ti4+. Therefore,
charge capacity should have been produced by anionic (sulfur) redox, some part of the
S2- atoms oxidizes into S22- (Eq. IV - 5). During discharge, 2.0 Li+ ions are inserted into
the structure and the composition changed into Li2.23TiS3. Discharge capacity should
have been either provided by anionic redox or both anionic (S2-/S22-) and cationic
(Ti4+/Ti3+) redox processes (Eq. IV - 6). After the first cycle, cycling curve becomes
reversible and 338 mAh.g-1 of capacity is delivered. As previously mentioned, 80% of
capacity is reached at the end of the 7th cycle.
Charge: Li2Ti4+S2-3 → Li0.23Ti4+ (S2-)1.23(S22-)0.885 + 1.77 Li+ + 1.77 e- (Eq. IV - 5)
Discharge: Li0.23Ti4+ (S2-)1.23(S22-)0.885 + 2 Li+ + 2 e-→ Li2.23 (Ti4+ )x(Ti3+)1-x (S2-)3-y(S22-)y
with 2.(0.885-y)+x =2 (Eq. IV - 6)
Figure IV - 22 also shows coin cells tests results of Li2TiSeS2 upon cycling at a rate of
C/10 between 3 V – 1.5 V vs Li+/Li. Li2TiSeS2 electrodes show reversible charge and
discharge capacities similar to Li2TiS3 electrodes. During charge, 1.39 Li+ ions are
extracted from the structure and capacity is recorded as 179 mAh.g-1. If one holds the
same hypothesis to explain redox process of Li2TiS3, the charge capacity should result
from anionic redox that is either sulfur (S2- / S22-) or selenium (Se2-/ Se22-) redox process.
During discharge, more than 2.0 Li+ ions are inserted into the structure, the
composition changed into Li3.08TiSeS2 and 310 mAh.g-1 is delivered. Discharge capacity
should be provided by either anionic (S2-/S22- or Se2-/Se22-) or both anionic and cationic
(Ti4+/Ti3+) redox process. The second and third cycles are fully reversible similar to
Li2TiS3 and cycling life is relatively better than Li2TiS3, 80% of capacity is reached at the
end of 10th cycle.
111
Coin cells tests results at the first three cycles of Li2TiSe2S upon cycling at a rate of C/10
between 3 V – 1.5 V vs Li+/Li show the same trend that is observed for Li2TiSeS2
electrodes. During charge, 1.32 Li+ ions are extracted from the structure and capacity
is recorded as 149 mAh.g-1. In contrary to Li2TiSeS2, much more lithium ions, 3.44 Li+
ions, are inserted into the cubic rocksalt structure and a much higher discharge
capacity (379 mAh.g-1) is obtained. The final composition becomes as Li(4.17)TiSe2S,
which can be regarded as the average (theoretical) composition. More than 3 Li+ uptake
cannot be explained with the same hypothesis, Ti3+/Ti4+ redox process. Therefore, this
should be investigated with further characterizations. Besides, we observe rapid
capacity fading in Li2TiSe2S cells, 80% of capacity is reached at the end of 2nd cycle.
Table IV - 7: Summary of structural and coin cell test results of Li2TiS3, Li2TiSeS2 and
Li2TiSe2S compositions: lattice parameters and cells cycled in galvanostatic mode at a rate of
C/10 between 3 V – 1.5 V vs Li+/ Li
Table IV - 7 shows a summary of structural and coin cell test results of for Li2TiS3,
Li2TiSeS2 and Li2TiSe2S compositions. For Li2TiS3, charge capacity is smaller, whereas
discharge capacity is equal to theoretical capacity (339 mAh.g-1) since rocksalt structure
allows two Li+ reversible exchange. For Se substituted samples, smaller charge
capacities are obtained at the first cycle. On the contrary, discharge capacities are
greater than the theoretical capacities (261 mAh.g-1 for Li2TiSeS2, 213 mAh.g-1 for
Li2TiSe2S). This indicates the presence of extra capacity at the end of the discharge.
Besides, there is a correlation between degree of Se substitution and the formation of
extra capacity at the end of the discharge. Li2TiSe2S shows the highest discharge
capacity. Given average potentials represent an average of charge and discharge
112
potentials during 10 cycles. Average potentials decrease with selenium substitution,
for instance Li2TiSe2S showed the lowest charge and discharge average potentials
(2.24V and 1.98 V respectively).
• Cyclic Voltammetry Tests of Li2TiSexS3-x
To describe electrochemical processes and identify the electron transfer kinetics, cyclic
voltammetry tests of Li2TiS3, Li2TiSeS2 and Li2TiSe2S cells are conducted at various scan
rates (5 µV.s-1, 10 µV.s-1, 20 µV.s-1, 50 µV.s-1 and 100 µV.s-1) between 3 V – 1.5 V vs Li+/
Li.
Figure IV - 23: Redox potentials comparison of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes at
the first (5µV.s-1) and the second (10 µV.s-1) cycles between 3 V -1.5 V vs Li+/Li
Figure IV - 23 shows the cyclic voltammograms of Li2TiS3, Li2TiSeS2 and Li2TiSe2S cells
at the first cycle at 5 µV.s-1 and the second cycle at 10 µV.s-1. During the first cycle, the
Li2TiS3 cell shows one oxidative and one reductive peak from sulfur (S2-/S22-) redox
reaction at 2.69 V and 2.28 V respectively. Potentials are decreased by selenium
substitution, 2.51 V charge, 2.20 V discharge potentials are observed for Li2TiSeS2,
whereas 2.43 V charge and 2.12 V discharge potentials are observed for Li2TiSe2S. These
results indicate that new compositions have different redox activities. During
discharge, we also observe a second phenomenon that could correspond to the extra
capacity. New reductive peaks at 1.88 V appear during discharge and are more visible
at the second cycle. Besides, we observe that an oxidative peak appears at 2.14 V at the
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second cycle, indicating that the second redox process is reversible. This new peak
becomes much more defined in Li2TiSe2S cell, it is observed at 2.12 V and 1.88 V.
Figure IV - 24 : Cyclic voltammogram of (a) Li2TiS3, (b) Li2TiSeS2 and (c)Li2TiSe2S electrodes
cycled by various scan rates between 3 V – 1.5 V vs Li+/ Li; peak separation (mV) at each scan
rate (µV.s-1) to determine charge transfer kinetics of sulfur redox (d) and selenium redox (e)
processes
To understand the nature of the electrochemical process, and the electron transfer
kinetics, Li2TiS3, Li2TiSeS2 and Li2TiSe2S cells are scanned from 5 µV.s-1 to 100 µV.s-1 and
the current responses are measured and shown in Figure IV - 24a, Figure IV - 24b and
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Figure IV - 24c. Here, we observe that peak currents increase by increasing scan rates
corresponding to the relationship between peak current and scan rate expressed in
Randles-Sevcik equation (102,103). In addition, peak potentials shift to lower potentials
with increasing scan rates. This may be related to reversibility or irreversibility of the
electrochemical processes. If we describe a reversible process, peak separation (∆E) is
independent of scan rate (v) and anodic and cathodic peak currents are equal (Ipa = Ipc)
(103). Unlike reversible process, electron transfer kinetics are limited in a quasi-
reversible system, thus the peak potential is a function of scan rate (103). To distinguish
between these two processes (reversible or quasi-reversible), peak potential
separations are measured for the first (sulfide) and the second (selenium) redox
processes and the values are shown in Figure IV - 24d and Figure IV - 24e. In Figure
IV - 24d, we observe that peak separations (∆Esulfur) between 20 µV.s-1 to 100 µV.s-1 are
almost stable for Li2TiSexS3-x cells, it seems that sulfur redox process is reversible. To
confirm this, theoretical peak separations are calculated using Nernst equation that is
shown below (Eq. IV - 7) considering all the electrochemical processes are reversible
(102–104). Here, we calculate the peak potential separations based on two Li+ exchange
process for Li2TiS3, three Li+ exchange process for Li2TiSeS2 and four Li+ exchange
process for Li2TiSe2S.
∆𝐸 = 𝐸𝑝𝑎 − 𝐸𝑝𝑐 +56.5
𝑛 𝑚𝑉 at 298 K (Eq. IV - 7)
As we compare the theoretical and experimental peak potential separations, we find
that experimental peak separation values of both sulfur (∆Esulfur) and selenium
(∆Eselenium) redox processes are much higher than theoretical values (28.2 mV for
Li2TiS3, 18.8 mV for Li2TiSeS2 and 14.1 mV for Li2TiSe2S), indicating that sulfur and
selenium redox processes are not completely reversible. In Figure IV - 24e, we observe
that peak potential separations (∆Eselenium) corresponding to selenium redox process
in Li2TiSexS3-x cells increases with increasing scan rate. On the contrary, peak potential
separations (∆Esulfur) corresponding to sulfur redox process is almost stable at higher
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scan rates (from 50 µV.s-1 to 100 µV.s-1) in Figure IV - 24d. By comparing the two
different redox processes in (Li2TiSexS3-x) cells, we can conclude that electron transfer
kinetics in selenium process is much more sensitive to the scan rate.
Cyclic voltammetry can also allow determining the diffusion property of Li2TiSexS3-x
cells. The peak potentials at each scan rate can be used to calculate the lithium ion
diffusion coefficients using the relationship between measured current and scan rate,
which is expressed as i=avb (102,103). Here a and b parameters are variables and b
parameter is expressed as a sum of faradaic (diffusion controlled) and non-faradaic
(capacitive/surface controlled) currents (103). By plotting logi vs logv, b parameter can
be obtained from the corresponding slope. According to b parameter, the reaction
kinetics are either surface controlled (b=1) or diffusion-controlled process (b=0.5)
(103,105).
Figure IV - 25: a) Logarithmic and b) linear relation between peak current and scan rate of
Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes
In Figure IV - 25a, only Li2TiS3 shows a linear relation (diffusion controlled process)
between peak current and square root of scan rate (slope= 0.5, R² ≥ 0.99) among
Li2TiSexS3-x cells. The relation between peak current and scan rate in Li2TiSeS2 and
Li2TiSe2S cells goes in a nonlinear direction (slope> 0.5). For a diffusion controlled
process, peak current is related to the square root of scan rate by i=v1/2, thus it allows
applying Randles-Sevcik equation102,103 (Eq. IV - 8)
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𝐼𝑝 = (2.69 𝑥 105)𝑛3/2𝐴𝐷01/2
𝐶0𝑉1/2 (Eq. IV - 8)
where Ip is peak current in A, n the number of exchanged electrons, A the active surface
in cm2, Co the lithium concentration in mol.cm3- and V the scan rate in V.s-1.
Figure IV - 25b also shows the linear relation between peak current and scan rate with
the corresponding regression line (R²) for Li2TiS3. The concentration (Co) is calculated
as 0.06 Li+ mol.cm-3 from the unit cell volume and we consider two Li+ ions are
exchanged during charge-discharge. Using these variables in Eq. IV - 8, anodic
diffusion coefficient (Do) of Li2TiS3 is calculated as 1.99x10-9cm2.s-1. Lithium diffusion
kinetics of different sulfurous compounds are studied in the literature and the
diffusion coefficients are reported between 10-9 - 10-8 (106–108). Therefore, DLi+ that we
found is close to the reported values.
We further investigate capacity retention of Li2TiS3, Li2TiSeS2 and Li2TiSe2S cells cycled
at different scan rates (Table IV - 8). By ignoring the fact that the scan rates are quite
slow (µV.s-1), Li2TiSeS2 and Li2TiSe2S electrodes retain the capacities in a better way
than the Li2TiS3 cells at high rates. Although such a result is quite different from what
is observed in galvanostatic tests, it also supports our previous observation in which
electron transfer kinetics in Li2TiSeS2 and Li2TiSe2S compositions are much more
sensitive to scan rate.
Table IV - 8: Cyclic voltammetry test results of Li2TiS3, Li2TiSeS2 and Li2TiSe2S cells cycled at
different cycling rates between 3 V – 1.5 V vs Li+/ Li
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• Cycle life improvement with different cycling schedules
The goal of the selenium substitution was to improve the cycling stability of Li2TiS3,
therefore, cycling stabilities of Li2TiSeS2 and Li2TiSe2S should be compared with Li2TiS3.
Figure IV - 26 shows capacity retention of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes
cycled in galvanostatic mode at different discharge cut-off potentials. Cycling the cells
between 3 V - 1.5 V vs Li+/Li, Li2TiSeS2 cells show better cycling stability than Li2TiS3
cells, whereas Li2TiSe2S cells show rapid capacity fading. At higher cut-off discharge
potentials (3 V - 1.8 V) and (3 V - 2 V), all selenium substituted cells (Li2TiSeS2 and
Li2TiSe2S) show better cycling stability than Li2TiS3. To discover why Li2TiSe2S cells fail
upon cycling between (3 V - 1.5 V), we perform ex situ XRD and SEM analyses of
electrodes at different states of charge.
Figure IV - 26: Capacity retention (%) of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes cycled in
galvanostatic mode at different cycling discharge cut-off potentials: 3 V - 1.5 V, 3 V - 1.8 V
and 3 V - 2.0 V vs Li+/Li.
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Figure IV - 27: SEM images of Li2TiSe2S electrodes cycled at a rate of C/10 between (3 V - 1.5
V) and (3 V - 1.8 V)
Figure IV - 28: X-ray diffraction pattern of Li2TiSe2S electrodes cycled at a rate of C/10
between (3 V - 1.5 V) and (3 V - 1.8 V)
119
In Figure IV - 27, we observe porous morphology in the electrodes cycled at 3 V - 1.5
V. The same porous morphology is also seen in the electrodes cycled at 3 V - 1.8 V, but
it is slightly eliminated by increasing the cut-off discharge potential.
In Figure IV - 28, we detect disordered rocksalt phase containing Tix(Se-S)y type of
impurities at two different electrodes. In summary, the change in electrode
morphology and presence of Tix(Se-S)y type of impurities indicate that parasitic
reactions possibly occur during discharging the cell.
We also test a different cycling schedule in which electrodes are first discharged to 1.5
V and then charged to 3 V (Figure IV - 29). Only limited capacities are obtained upon
discharging Li2TiS3 cells, while large discharge capacities are obtained in both Li2TiSeS2
and Li2TiSe2S electrodes upon discharge. Such result indicates that Se substituted
compositions might not be pure, thus they may contain some amount of metallic
selenium that was used during the mechanochemical synthesis. This is investigated
further in the next chapter.
Figure IV - 29: Half-cell performances of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes cycled at
a rate of C/10 between 1.5 V- 3.0 V vs Li+/Li
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IV.3. Conclusion
In this chapter, Li2TiS3 powder is synthesized mechanochemically in rocksalt phase
and it shows reversible charge and discharge capacities of 339 mAh.g-1 based on two
lithium ions exchange process. To understand degradation process of Li2TiS3 in liquid
cells, various electrochemical and structural characterizations are performed. To
improve cycling stability of Li2TiS3 cells, selenium substitution is considered. Applying
the same synthesis recipe, selenium-substituted samples could not be synthesized,
thus the synthesis recipe had to be improved. After various synthesis strategies, a new
synthesis process is successfully developed and patented (100,101).
Selenium-substituted samples (Li2TiSeS2 and Li2TiSe2S) are synthesized in disordered
rocksalt phase with new synthesis process. To understand their cell performance,
cycling stability and electron transfer kinetics, galvanostatic and voltammetry tests are
conducted. Results show that Li2TiSeS2 and Li2TiSe2S provide larger discharge
capacities than Li2TiS3 cells. Cycling voltammetry test of selenium-substituted samples
at various scan rates show that oxidation and reduction potentials decrease by
selenium substitution. At high scan rates in voltammetry, selenium-substituted
samples show better capacity retention than Li2TiS3 cells. We also find that electron
transfer kinetics of selenium-substituted samples are more sensitive to the scan rate.
Afterwards, electrochemical performances of Li2TiSexS3-x electrodes are investigated at
different cycling schedules. Capacity retention of Li2TiSe2S cells is poor upon cycling
between 3 V - 1.5 V; however, we observe that it improves upon cycling between 3 V -
1.8 V and 3 V - 2.0 V. By examining the cycled electrodes with ex situ XRD and SEM
analyses, we find that Li2TiSe2S electrodes degrade during discharging the cells.
Nevertheless, increasing discharge cut-off potential might limit the parasitic reactions
that forms upon discharge.
Finally, Li2TiSexS3-x electrodes are also tested between 1.5 V – 3 V. Only limited
discharge capacities are detected in Li2TiS3 cells, whereas large discharge capacities are
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obtained in Li2TiSeS2 and Li2TiSe2S cells. Such results indicate that Se substituted
samples may not be pure, thus may contain considerable amount of metallic selenium
precursor. This is investigated further in the next chapter.
Despite detailed characterization of Li2TiSexS3-x compounds, we still do not know the
redox active species and possible redox mechanisms. To discover such properties, we
further characterize these compounds with X-ray Photoelectron Spectroscopy (XPS),
Raman Spectroscopy and neutron diffraction in the following chapter.
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Chapter V: Characterization of
Li2TiSexS3-x Compositions
In previous chapter, our new rocksalt sulfide-selenide compositions show unusual
electrochemical performances. We observe large discharge capacities in Se substituted
samples at the first cycle. This could be due to extra Li uptake at the end of the
discharge. Cubic rocksalt structure is dense and all the octahedral sites are already
filled with both Ti4+ and Li+ cations, however it is still questionable where the extra
lithium ions go in the structure. We also observe that cycling stability of Se substituted
samples is rather poor during charge-discharge. Therefore, our next target is to find
what impedes the capacity during cycling.
This chapter describes electrochemical and structural characterization of sulfur-
selenium compositions (Li2TiSexS3-x). Operando neutron diffraction and ex situ X-ray
diffraction are described to detect Li ions in the rocksalt structure. Afterwards, results
of XPS analysis are presented to elucidate the electrochemical activity in those
materials.
V.1. Operando Neutron Diffraction of Li2TiS3
Operando neutron diffraction is performed at ILL, Grenoble. We use a special cell,
“neutron cell” made of Ti2.08Zr alloy, which is transparent to neutron diffraction. The
experimental procedure is described in Chapter 2.2.b.
Before performing operando diffraction experiment, Li2TiS3 powder (≈ 120 mg) is
analyzed with ex situ neutron diffraction in ILL. The target is to collect the diffraction
patterns and to determine the quality of data before the real experiment. In ex situ
neutron diffraction pattern, we observe a large background with very small diffraction
peaks. We conclude that the powder amount is too low to obtain sufficient diffraction
peaks. Afterwards, we decide to increase powder amount (200 mg) to get sufficient
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intense diffraction peaks during the real experiment. However, we also know that
working with the cells containing large amount of powders is quite challenging, it
often results in large cell polarization with smaller capacities. Therefore, our first goal
is to optimize the electrochemical performances of the neutron cell by applying
different cell testing schedules (Figure V - 1).
Figure V - 1: Cycling of operando neutron cell cycled at a rate of C/10: a) Cell is charged to 3
V and discharged to 1.5 V b) Cell is discharged to 1.5 V and charged to 3 V, c) cell is
discharged to 1.5 V and charged to 3.5 V d) cell is charged to 3.5 V and discharged to 1.0 V
At the beginning, we prepare neutron cell as it is described in Chapter 2.b., for instance
the cell is tested at a rate of C/10 between 3 V and 1.5 V. However, it cannot be charged
to 3 V and a wide potential plateau at 2.6 V with a large polarization (0.68 V) is
observed (Figure V - 1a). We suspect that the parasitic plateau at 2.6 V might be due to
the moisture and a parafilm that is added around the nut to improve the cell tightness.
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As our main interest in this experiment is to detect extra capacity and to find out where
the extra lithium locates in the structure, we are more interested in discharge state. To
produce similar results, the neutron cell is first discharged to 1.5 V and charged to 3 V
(Figure V - 1b). Thereafter, the high cut-off potential is increased to 3.5 V and we finally
obtain a reversible cycling curve, as shown in Figure V - 1c. Nevertheless, we do not
observe the extra capacity during discharge. We further decrease the low cut-off
potential and the cell is charged first to 3.5 V and discharged to 1.0 V. After such
optimizations, we succeed to obtain reversible cycling curve and observe the extra
capacity during discharge.
Figure V - 2: Operando neutron diffraction pattern of Li2TiS3 powder: pristine powder,
discharged powder and empty neutron diffraction cell (neutron wavelength: 1.594 Å)
Figure V - 3: Simulated neutron diffraction pattern of Li2TiS3
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Figure V - 2 shows experimental and Figure V - 3 shows simulated operando neutron
diffraction patterns of Li2TiS3 powders. Based on simulated diffraction pattern, we
should see a strong peak at 31°. However, in Figure V - 2, we observe that the first peak
at 31° is lost in the background coming from the cell itself. We could only see a small
peak at 62°, but it is difficult to make a conclusion based on this result. Differently from
the simulated diffraction pattern, two additional peaks at 47 Å and 56 Å in both
pristine and discharge states are seen in Figure V - 2. Those peaks should arise from
the separator within the cell.
In conclusion, we cannot not successfully detect Li+ in the cubic rocksalt structure
because the reflections coming from the material are much weaker than the
background produced by the neutron cell, although 200 mg of powder is loaded in the
cell. Nevertheless, we decide to investigate our another question concerning what
impedes the capacity during cycling and to do so, we perform ex situ XRD study on
electrodes at different states of charge.
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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
Figure V - 4 shows cycling curve corresponding to the studied electrodes and Figure
V - 5 shows ex situ XRD of Li2TiS3 electrodes at the different states of charge.
Figure V - 4: Cell performances of Li2TiS3 electrodes at different states of charge: electrode at
the end of charge (3 V) and electrode at the end of discharge (1.6 V)
Figure V - 5: Ex situ X-Ray diffractograms of Li2TiS3 at different states of charge: pristine, at
the end of the charge (3 V) and at the end of the discharge (1.6 V)
127
Pristine Li2TiS3 electrode has a rocksalt (cubic) structure. At the end of the charge,
lithium ions deinsert from Li2TiS3 electrode and peak intensities reduce, rocksalt
structure becomes amorphous. The rocksalt structure is recovered after lithium
uptakes at the end of the discharge. Those results are consistent with the literature
where reversible structural change has been observed (62). In addition, authors indicate
that S-S bonds lengthen and shorten during charge and discharge, whereas Ti-S bonds
do not change at all (64). Such rearrangements allow flexible structural change from
crystalline to amorphous and crystalline structure.
Figure V - 6 shows SEM images of Li2TiS3 electrodes cycled at different states of charge.
In Li2TiS3 electrode at the end of the charge, expansion is obvious and some degree of
porosities are observed. However, the morphology of Li2TiS3 electrode at the end of the
discharge is quite similar to Li2TiS3 pristine electrode. Overall, we observe reversible
structural change and detect similar electrode morphologies after charge and
discharge of Li2TiS3 electrodes.
Figure V - 6: SEM images of Li2TiS3 electrodes at different states of charge: pristine electrode,
charged (3 V) and discharged (1.6 V) electrode
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V.2.b. Ex situ XRD and SEM Analyses of Li2TiSexS3-x electrodes
Figure V - 7 cycling curve corresponding to the studied electrodes and Figure V - 8
shows ex situ XRD of Li2TiSeS2 electrodes at the different states of charge.
Figure V - 7: Cell performances of Li2TiSeS2 electrodes at different states of charge: electrode at
the end of charge (3 V) and electrode at the end of discharge (1.6 V)
Figure V - 8: Ex situ X-Ray diffractograms of Li2TiSeS2 electrodes at different states of charge:
pristine, at the end of charge (3 V) and at the end of discharge (1.6 V)
As shown in Figure V - 8, reversible structural change is also detected in Li2TiSeS2
electrodes. Differently from Li2TiS3, metallic selenium is observed in Li2TiSeS2
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electrodes at the end of the charge. Thereafter, structure recrystallizes in a cubic
structure at the end of the discharge and metallic Se disappears in Li2TiSeS2 electrode.
Li2TiSe2S electrodes are also tested at different states of charge (Figure V - 9). The same
phenomenon is also observed in Li2TiSe2S electrodes, metallic selenium is detected in
charged electrodes and it disappears in discharged electrodes (Figure V - 10).
Figure V - 9: Cell performances of Li2TiSe2S electrodes at different states of charge: electrode at
the end of charge (3 V) and electrode at the end of discharge (1.6 V)
Figure V - 10: Ex situ X-Ray diffractograms of Li2TiSe2S electrodes at different states of
charge: pristine, at the end of charge (3 V) and at the end of discharge (1.6 V)
130
When Li2TiSe2S electrode at the end of the charge is studied by SEM, (Figure V - 11), a
columnar phase is observed, whereas only one phase is observed in Li2TiSe2S
electrodes at the end of the discharge. This is in accordance to the results of ex situ XRD
analysis, metallic selenium (Se0) could be formed at the end of the charge and
converted into soluble polyselenides containing Se2- at the end of the discharge.
Therefore, only disordered rocksalt structure is observed in the XRD pattern. Li2TiSe2S
electrode at the end of the discharge shows a porous structure that is different from
Li2TiS3 and Li2TiSeS2 electrodes. That might be because selenium has dissolved into the
electrolyte and left empty space in the electrode.
Figure V - 11: SEM images of Li2TiS3, Li2TiSeS2 and Li2TiSe2S electrodes at different states of
charge: pristine, at the end of charge (3 V) and at the end of discharge (1.6 V)
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Figure V - 12 shows SEM and EDX images of Li2TiSe2S electrodes. After charging the
electrodes at 3 V for 24 hours, we observe a columnar phase that has crystallized in a
different shape. EDX analysis shows that 92% of Se, 4% Ti and 4% S elements are
present in the columnar phase. This result also supports previous hypothesis
concerning formation of metallic selenium during charge process.
Figure V - 12: a) SEM images and b) EDX images of Li2TiSe2S electrodes cycled to 3V and
kept for 24 h
After those results, we have two different hypotheses:
1. Is there Ostwald ripening? We can imagine that some unreacted Se remains
at the surface of the material and it has a very small size, therefore it is not
seen by XRD. Then, when put in the electrolyte, it would dissolve and
precipitate into bigger particles following an Ostwald ripening process (109).
Therefore, diffraction peak of unreacted Se would become visible due to the
bigger size of the particles.
2. Is there a selenium redox activity? If selenium is redox active, the possible
redox process should be Se2- / Se0 that is different from sulfur redox process
(S2-/S22-). Therefore, Se0 could be produced at the end of the charge and
converted into soluble polyselenides containing Se2- at the end of the
discharge.
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To test these hypotheses, we decide to perform a high-resolution XPS study on rocksalt
sulfide-selenide pristine powders and electrodes at different states of charge.
V.3. XPS Analysis of Li2TiSexS3-x Compositions
Elucidating the electrochemical activity in those materials requires an understanding
of: what species are present in the electrodes, which species are responsible for redox
reactions and how those species change upon charge-discharge. Moreover,
electrochemical and structural characterizations of Li2TiSexS3-x show that not only
sulfur but also selenium should be redox active. To test this hypothesis, we use a
complementary technique, a high-resolution X-ray Photoelectron Spectroscopy (XPS,
introduced in Chapter 2.2.e), which is a surface sensitive technique that gives
information about both chemical states in the compositions and does semi-quantitative
analysis of the compositions. In this chapter, we present the results of a high-resolution
XPS study on disordered rocksalt Li2TiSexS3-x pristine powders and electrodes at
different states of charge. We also prepare tables that contain optimized parameters,
such as peak positions and widths at different core levels. We believe that such results
can help scientists to understand the redox processes in metal sulfides or selenides
since only few publications are present in the literature.
V.3.a. Methodology: Analysis of XPS Spectra
Peak identification/indexation: In general, the peaks obtained from the XPS analysis
are indexed based on the information in XPS Handbook110 and the online XPS
database111 that contains binding energy of species in various compositions. As we are
the first ones who designed sulfur-selenium compositions (Li2TiSexS3-x) and the
binding energy of elements differs based on the electronic environments, we not only
use the data from XPS handbook and database but we also consider similar studies in
the literature to prepare fitting models.
Strategy for peak deconvolution: To extract the hidden information from the peaks,
firstly, we take the peak fittings parameters, such as peak positions, peak ratios and
133
peak widths from the XPS handbook and literature. The latter, we fix these parameters
in the Multitap program that is used to deconvolute the peaks obtained from the High-
Resolution XPS spectra. Afterwards, we build a fitting model that is iteratively
improved until fitting parameters do not change within a specified precision (Chi<5).
The curves are fitted using 80% Gaussian, 20% Lorenzian conditions in Multitap
Program.
• Single peaks: deconvoluting Li 1s, Se 3d and Ti 3s: The binding energy of these
three core peaks are quite close (Li 1s:54-55 eV, Se 3d: 52-53 eV Ti:3s 60-61 eV)
and therefore they can be detected together at the same core level110. For
instance, Ti 3s and Se 3d can be detected at Li 1s core. Here, each core peak is
fitted and evaluated individually for quantitative elemental analysis.
• Overlapping peaks: deconvoluting Ti 2p and S 2p orbitals: For overlapping
peaks, fitting is employed based on the contributions (two peaks) originated
from spin orbit splitting of 2p cores (3/2p and 1/2p). Each peak is fitted within
the same model considering fixed peak separations110 and the ratio between two
peaks are calculated considering these two contributions at the same core level.
We follow the same procedure for fitting of S 2p core peaks.
Quantitative elemental analysis is performed by accounting peak intensities (peak area
after background removal) and the sensivity factors (Chapter 2.2.f). It gives atomic
concentrations in the analyzed samples. For elemental analysis, we account of Li+ in Li
1s core, Ti4+ in Ti 2p core, S2- (could be S22-) in S 2p core for Li2TiS3. It should be noted
that Ti4+ species can be seen at both Ti 2p and Ti 3s core levels, but the amount of
titanium species at Ti 2p core are different from those in Ti 3s core. The difference arises
from the cross-section of each elements. The number of photons obtained from Ti 2p
core is much higher than those coming from Ti 3s core; thus, we only consider one
contribution (Ti4+ in Ti in 2p core) for elemental analysis. Similarly, we account Se2-
(could be Se22-) at Se 3d core instead of Se 3p core for selenium compositions.
134
V.3.b. Survey analysis of Li2TiSexS3-x samples
A survey analysis is performed to identify the elements contained in the composition
with a survey spectrum that is recorded over a spectral range of 0−1200eV. This quick
analysis allows us determining which core of the elements should be analyzed through
XPS and what parameters should be used. Thereafter, we can optimize the quality of
signal/noise ratio for High-Resolution XPS analysis. Nevertheless, it should be noted
that such survey analysis gives general information about species in the compositions.
For detailed analysis, the samples are analyzed further with High-Resolution XPS
(penetration length of X-rays up to 5 nm).
As illustrated in Figure V - 13, O 1s (530 eV) and C 1s (284.8 eV) core peaks are detected
in powders due to sample contamination from glovebox to XPS vacuum chamber. In
Figure V - 14, similar elements are detected in electrodes as well, but differently from
the powders, F 1s core peak (687.7 eV) is detected in electrodes. This originates from
PVdF binder used during electrode preparation.
Figure V - 13: XPS analysis of powders in survey mode
135
Figure V - 14: XPS analysis of electrodes in survey mode
After survey analysis, parameters are optimized based on high signal/noise ratio and
powders are analyzed with high-resolution XPS. We decided to follow the next
procedure during the XPS analysis:
• The first target is to determine active species at specific binding energies and
distinguish them in their different electronic environments.
• Thereafter, fitting models are developed for powders and then same models are
used to analyze the electrodes of the same compositions.
At the beginning, one of our hypotheses was that there might have been a slight trace
of unreacted selenium present in the sulfur-selenium solid solutions, therefore Se
metal would have become visible in X-ray diffractograms after peak intensities
reduced at the end of the charge. To test our hypothesis, we analyze three pristine
powders and Se metal precursor through XPS.
136
Figure V - 15: XPS spectra of powders: Li2TiS3, Li2TiSeS2, Li2TiSe2S and selenium precursor
at Se 3d core
Figure V - 15 shows the Se 3d core peak of Li2TiS3, Li2TiSeS2, Li2TiSe2S and selenium
precursor. The four spectra were calibrated based on C 1s core peak at 284.8 eV. In the
figure, Li 1s and Se 3d binding energies are quite close to each other, Li 1s and Se 3d
core peaks are detected between 53 eV and 54 eV respectively. The intensity of Se2-
peak is much higher compared to intensity of Li+ peak. This could be due to the
elemental sensitivity to XPS analysis, for instance sensitivity factor of Li (0.02) is much
smaller than sensitivity factor of Se (0.67). Besides, binding energy of Se2- is rather
smaller than Se0, core peak of Se0 is detected at 56 eV. Similar binding energy of Se0 is
reported in the literature as well (112). Overall, no signature of selenium precursors is
detected in XPS spectra. Such result is not aligned with our first hypothesis,
consequently we conclude that there is no metallic selenium remaining in the powders
after the synthesis.
137
V.3.c. High Resolution XPS Analysis of Li2TiSexS3-x Powders
Figure V - 16 shows the XPS spectra and peak fittings of Li2TiS3, Li2TiSeS2 and Li2TiSe2S
powders at Li 1s, Ti 2p and S 2p core levels. The curves are fitted using 80% Gaussian,
20% Lorenzian conditions in Multitap program.
Figure V - 16: XPS spectra of Li2TiS3, Li2TiSeS2 and Li2TiSe2S powders at Ti 2p3/2-1/2
(Ti1_orange: Ti 2p3/2, Ti2_yellow: Ti 2p1/2), S 2p3/2-1/2 (S1_light blue: S 2p3/2 and S2_dark blue:
S 2p1/2), Li 1s core (navy blue), Ti 3s (purple) and Se 3d5/2-3/2 (Se1_green:Se 3d5/2, Se2_pink:Se
3d3/2)
• 72-44eV: Li 1s, Ti 3s and Se 3d cores
In Li2TiS3, we observe Li 1s core peak at 54.6 eV with a peak width of 2.1 eV. Previous
studies reported similar binding energy of Li 1s core peak in Li-S systems (113,114). In the
same energy window, we also observe Ti 3s at 60.5 eV with almost the same intensity
as Li 1s core peak (52.4%), but a much wider peak width (3.8). We use the same peak
positions, peak widths and relative intensity of Li 1s and Ti 3s core peaks for fitting
138
selenium-substituted compositions at Li 1s and Se 3d core levels. Table V - 1 shows all
optimized fittings parameters and atomic percentages of Li2TiS3, Li2TiSeS2, Li2TiSe2S
and selenium precursor at Li 1s (or Se 3d), Ti 2p3/2-1/2, and S 2p3/2-1/2.
In Li2TiSeS2, we observe two different selenides (Se1 and Se2) at Se 3d core level and
each selenide contains a doublet, the Se 3d5/2 and Se 3d3/2 peaks, with widths of 0.86 eV.
Binding energy at Se 3d5/2 core of Se1 is detected at 52.7 eV, whereas binding energy of
Se2 is found at 53.8 eV. The parameters can be found in Table V - 1. The difference
between these two energies and the presence of two Se doublets indicate that their
valence states are not the same, Se1 could be regarded as Se2- and the binding energy
of Se2 is quite close to the Se22- (115,116). In the literature, Se 3d core peaks were reported
between 53 and 54 eVs for different metal selenide systems (117,118). By semi-quantitative
analysis, we also observe that Se1 is largely present (72.4%) compared to the amount
of Se2 (27.7%) in Li2TiSeS2. Normally, we expect to detect only one type of selenide at
Se 3d, but it seems that nearly one third of the Se2- is already oxidized to Se22- at least
at the surface of the material. This means that the amount of Li in Li2TiSeS2 is less than
what we calculated (Li2-xTiSeS2).
In Li2TiSe2S, similar peak positions are detected. Besides, we observed that ratio of Se1
to Se2 increased since 81.6% of Se1 is present compared to Se2 (18.4%). The amount of
Se2- is much higher in Li2TiSe2S, as expected. Nevertheless, we observe again that some
part of Se is oxidized, thus the amount of Li in Li2TiSe2S, is less than what we calculated
(Li2-xTiSe2S).
In conclusion, we do not observe unreacted selenium in XPS analysis on powders.
Instead, we observe one type of Li+, but two different selenide types (Se2- and Se22-) in
this energy region. Besides, Ti-S and Ti-Se peaks are not easily distinguished because
their binding energy are very close (9). Therefore, they are considered as one single Ti
peak.
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• 473eV-449eV: Ti 2p core
In Li2TiS3 powder, two different Ti 2p doublets, such as Ti1 and Ti2 are observed at
455.1 eV and 457.4 eV, indicating Ti4+ present in sulfur (TiS2) and oxygen (TiO2)
environments. Presence of TiO2 peak at Ti 2p core indicates that there should be air
exposure during sample transfer from transfer box to XPS vacuum chamber. Similar
phenomenon was observed in the literature as well (119–121).
In both Li2TiSeS2 and Li2TiSe2S compositions, we observed similar binding energies
and peak widths of Ti 2p core levels. Here, Ti4+ is present in sulfur-selenium, [Ti(S-Se)2]
and oxygen (TiO2) environments. In addition, correlation between Se substitution and
Ti4+ in oxygen (TiO2) environment may exist because the more Se is present in the
composition, the greater intensity of Ti4+ in oxygen (TiO2) environment is detected.
Selenium compounds thus seem more air-sensitive.
• 171eV-150eV: S 2p and Se 3p cores
In Li2TiS3, we observe two different sulfur doublets (S1 and S2) at 159.9 eV and 160.6 eV
with widths of 1.2 eV. According to semi-quantitative analysis, S1 presents 49.8%,
while S2 presents 50.2% in Li2TiS3 composition. In the literature, many metal sulfides
contains two or three sulfur species at S 2p core and they are regarded as terminal and
bridge sulfur species (122–124). Since the energy difference between two different sulfur
species is quite small (< 1 eV), these sulfides should be regarded as S2- (S1) and S22- (S2).
Contrary to Li2TiS3, only one S 2p doublet is observed in both Li2TiSeS2 and Li2TiSe2S
compositions; however, two different selenium doublets, Se1 and Se2, are observed in
Li2TiSeS2 at 158.8 eV and 159.4 eV in percentages of 21.2% and 7.2% respectively
(Sulfur: 71.6%). As already mentioned before, Se2 might be regarded as Se22-. Similar
species are detected in Li2TiSe2S, the amount of selenium species increase in Li2TiSe2S,
for instance Se2 now reaches 20%, much higher than the one in Li2TiSeS2. On the
contrary, the percentage of Se1 found is 38.5% (Sulfur: 41.5%). In conclusion, all the
species in the powders are detected at Li 1s, Ti 2p and S 2p cores and curve fittings
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models are developed. Such models are used for curve fitting of XPS analysis on
electrodes at different states of charge.
Table V - 1 :Optimized fitting parameters of Li2TiS3, Li2TiSeS2, Li2TiSe2S and selenium
precursor at Li 1s/Se 3d, Ti 2p3/2-1/2, and S 2p3/2-1/2
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V.3.d. XPS Analysis of Li2TiSexS3-x Pristine Electrodes
After finding the best fitting models through powder analysis, we decide to continue
XPS analysis on various electrodes. Before analyzing the electrodes at different states
of charge, our next target is to investigate the surface and the bulk properties of the
pristine electrodes. One way to reach the bulk of the electrode is to scratch the electrode
with a spatula. The scratched electrodes (pristine state) are later compared with the
non-scratched electrodes and powders, XPS spectra of each compounds are shown in
Figure V - 17 (Li2TiS3), Figure V - 18 (Li2TiSeS2) and Figure V - 19 (Li2TiSe2S).
Figure V - 17: XPS spectra of Li2TiS3 at Ti 2p3/2-1/2, S 2p3/2-1/2 and Li 1s core: Li2TiS3 powder vs
non-scratched Li2TiS3 electrodes and scratched Li2TiS3 electrodes
In Figure V - 17, similar active species are observed in Li2TiS3 powder and electrodes.
We observe LiF peak at Li 1s core. That may come from reaction of sulfur with binder.
It has stronger peak intensity in non-scratched electrode, indicating that LiF is mostly
present on the electrode surface rather than in the bulk. In S 2p core, additional sulfur
species (S-C) are detected in Li2TiS3 electrodes, unlike Li2TiS3 powder. That is attributed
to sulfur-carbon bond in the literature since electrodes contain important amount of
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carbon (10%) to increase their electronic conductivity (123,124). This explains why the S-
C species was not detected before in sulfur-selenium powders. Unlike powders, we
also observe that the amount of S2- reduces in electrodes (10%), whereas the amount of
S22- increases (80%) with S-C (10%). We explain this in a new hypothesis that sulfur
may react with PVdF binder (production of LiF) and produce more oxidized sulfur
species. That also indicates that less lithium is present in the composition, the actual
formula might be Li2-xTiS3. Additionally, it might be the reason of why S-C bond is only
seen in electrodes, not in powders. Sulfur may react with carbon in the PVdF and
created S-C compound. In Ti 2p core, signal of Ti4+ in TiS2 environment becomes much
lower in electrodes, whereas signal of Ti4+ in TiO2 environment becomes stronger in
electrodes.
Same analyses are conducted for both Li2TiSeS2 and Li2TiSe2S compositions. It should
be noted that all the peak positions and peak separations are fixed at the beginning of
the XPS analysis of scratched and non-scratched electrodes, thus the results obtained
in both electrodes and powders are similar (Figure V - 18 and Figure V - 19). The main
differences between scratched and non-scratched electrodes appear in Li 1s core and S
2p core levels. Much higher selenium peak intensities at Li 1s core level are observed
in non-scratched electrodes, whereas intensities become lower in S 2p core peaks.
Similar phenomenon is observed for Li2TiSeS2 and Li2TiSe2S, LiF and S-C species are
detected in Li 1s core and S 2P core levels respectively. That also strengthens our
hypothesis that sulfur might react with PVdF binder.
Overall, peak intensity of lithium and sulfur species are much stronger in scratched
electrodes, therefore we decide to perform post mortem analysis on scratched
electrodes. Since our next target is to analyze sulfur-selenium redox process, S 2p core
levels of scratched electrodes are investigated through high-resolution XPS.
143
Figure V - 18: XPS spectra of Li2TiSeS2 at Ti 2p3/2-1/2, S 2p3/2-1/2 and Li 1s core: powder vs
scratched electrodes and non-scratched electrodes
Figure V - 19 : XPS spectra of Li2TiSe2S at Ti 2p3/2-1/2, S 2p3/2-1/2 and Li 1s core: powder vs
scratched electrodes and non-scratched electrodes
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V.3.e. XPS Analysis on Li2TiSexS3-x Electrodes at Different States of Charge
In this part, our target is to find the appropriate answer to our second hypothesis. As
we observed Se metal at the end of charge, our second hypothesis was that there
should have been a selenium redox activity and it occured from the change of Se2- into
Se0, that is why we observed Se metal at the end of the charge. According to our XPS
analysis on powders, we found out that there is no Ostwald ripening since no signature
of selenium precursor is found in XPS pattern of Li2TiSexS3-x compositions. However
we still need to verify the origin of extra capacity and elucidate redox process of Li2TiS3
composition. In this purpose, electrodes are cycled using the standard procedure
described in Chapter 2.3, dismounted and prepared for High resolution XPS as
described in Chapter 2.2.
Li2TiS3 electrodes at different states of charge
At S 2p core, three different sulfur species are observed in pristine electrode, shown in
Figure V - 20. S1 and S2 represents S2- (23.7%) and S22- (70.7) respectively, whereas S3
(S-C) (5.6%) corresponds to sulfur-carbon bond.
Figure V - 20: XPS spectra of Li2TiS3 scratched electrodes at different states of charge:
pristine, charged electrode (3 V) and discharged electrode (1.6 V)
145
At the end of charge, lithium ions are extracted from Li2TiS3 electrode, S2 (S22-) shifts to
higher binding energy and peak intensity reduces to 17%. This indicates oxidation of
S2 species. In the meantime, S1 (S2-) shifts to lower binding energy and peak intensity
increases to 43%. This indicates that sulfur might be partially reduced to maintain the
charge balance around Ti-S coordination. In electrochemical results, we assume that
sulfur oxidizes during charge and reduces during discharge. Therefore, this result is
different from what we expected. Besides, the peak at 168 eV indexed as SOx- (14%)
indicates that either Li2TiS3 might have reacted with carbonate-based electrolyte after
cycling or there is carbonate contamination on electrode surface after the cell is
dismounted. At the end of the discharge, lithium ions are inserted into Li2TiS3
electrode, all the sulfur species are found at the same binding energies as those of the
pristine state. This indicates that sulfur redox process is reversible.
Li2TiSeS2 electrodes at different states of charge
Figure V - 21 shows XPS spectra of Li2TiSeS2 scratched electrodes at different states of
charge. In pristine Li2TiSeS2 electrode Se2-(22%), Se22- (28%), S2- (45%) and S-C (5%)
species are detected at S 2p core level.
At the end of charge, lithium ions are extracted from Li2TiSeS2 electrode, sulfur
oxidizes, S2- shifts to higher binding energy and peak intensity reduces down to 22%.
On the contrary, the changes in binding energies of Se2- and Se22- are quite small and we
observe that peak intensity of Se22- increases (41%) compared to Se2- (12%). These
results seems to confirm an oxidation of the anionic network, however right now, we
are not able to confirm the exact selenium redox process because the deconvolution at
S 2p is quite difficult due to the high number of contributions with very close energies
(See Appendix).
At the end of the discharge, all the sulfur (S2-) and selenium species (Se2- and Se22-) are
detected at the similar binding energies as those of the pristine state of Li2TiSeS2
electrode. It seems that selenium redox is also reversible.
146
Figure V - 21: XPS spectra of Li2TiSeS2 and Li2TiSe2S scratched electrodes at different states
of charge: pristine, charged electrode (3 V) and discharged electrode (1.6 V)
In Figure V - 22, reversible selenium redox in Li2TiSeS2 is also seen at Se 3d core.
Similarly to Figure V - 21, shifts in binding energies are very small, however change in
the intensities is evident. It seems that Se2- converts to Se22- during charge, but it returns
to the original position during discharge. It suggests a reversible redox activity.
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Figure V - 22 : XPS spectra of Li2TiSeS2 electrodes at different states of charge: Comparison of
Se 3d and S 2p core levels
Li2TiSe2S electrodes at different states of charge
Figure V - 21 also shows Li2TiSe2S electrodes at different states of charge. In pristine
Li2TiSe2S electrode Se2-(25%), Se22- (32.5%), S2- (35%) and S-C (7.5%) species are detected
at S 2p core level. Pristine electrode now contains much more Se22- (32.5%) than in
Li2TiSeS2 (28%). At the end of charge, we observe similar sulfur and selenium redox
activities as those in Li2TiSeS2 electrodes. At the end of the discharge, the peak
positions of sulfur and selenium species are found very close to those in the pristine
state. With semi-quantitative analysis, we detect Se2-(20%), Se22- (44%), S2- (25%) and S-
C (11%) species at S 2p core.
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V.4. Discussion on the XPS Results
In pristine Li2TiS3 electrodes, we consider S2- as terminal disulfide species, whereas S22-
as central disulfide species that are bonded with Li+ and Ti4+ ions (122). Upon charging,
lithium ions are extracted from structure, central disulfide species (S22-) lose some
lithium ions and only Ti-S bonded sulfur species (S22-) remain. Meanwhile, central
disulfide species that lose neighbor lithium ions prefer creating disulfides (S2-),
therefore we observe an increase of S2- species at the end of the charge. The similar
observation is reported by A. Sakuda et al., their Monte Carlo simulation results
showed that sulfur atoms that lose at least one lithium counter made S-S bonds (S2-)
(64). With formation of disulfides, the structural volume is increased. In this way,
rocksalt Li2TiS3 has undergone smooth structural transition from crystalline to
amorphous to crystalline structure, without direct crystalline to crystalline structural
transition. At the end of the discharge process, we observe the same species at similar
binding energy, this result is consistent with what we observed in ex situ XRD results
previously.
In pristine Se-substituted electrodes (Li2TiSeS2 and LiTiSe2S), we detect central
disulfide species (S22-) and selenium species (Se22-) with diselenides (Se2-) without any
terminal disulfides (115,122,125). Like sulfide electrode (Li2TiS3), central disulfide species
(S22-) and selenium species (Se22-) lose some lithium ions and Ti-S and Ti-Se bonded
sulfur and selenium species (S22-, Se2-) remain at the end of the charge (3 V). Differently
from sulfide electrode (Li2TiS3), diselenides species tend to make bonds with central
disulfide species and produce Se-S bonds, instead of producing diselenides at the end
of the charge. That is why we observe the increase of selenium species (Se22-) instead
of diselenides (Se2-) in the XPS spectra. This also indicates that the possible redox
mechanisms between sulfur and selenium is quite different.
In previous chapter, we show that selenium redox activity appears during first
discharge at 1.88 V vs Li+/Li for both Li2TiSeS2 and Li2TiSe2S electrodes. Therefore, here
149
we only consider sulfur redox process, the shift of S22- species from low binding to high
binding energy creates available charge capacity. At the end of the discharge, we detect
similar species in selenium-subsituted samples (Li2TiSexS3-x), but it seems that more
diselenides are produced. Here, we consider that disulfides and diselenides do not
contribute to charge-discharge capacity, thus they are redox inactive species. If we
assume continuous production of diselenides in Li2TiSe2S cells at the end of the
discharge, this might also explain why we observe huge capacity fading in the
Li2TiSe2S cells previously.
According to our second hypothesis, the possible redox mechanisms of selenium
should be undergone from Se2- to Se0. With the current study, it is quite difficult to
determine this point as the actual redox process of selenium starts at the first discharge
process. Therefore, further ex situ XPS analysis on Li2TiSexS3-x electrodes at different
states of charges (2nd and 3rd cycles) are required. In sulfur 2p core level, there are high
number of contributions coming from both sulfur and selenium species. To make clear
observations on selenium redox process, it might be better to work with only selenium
species. For instance, a new composition Li2TiSe3 might be developed and electrodes
at different states of charges might be analyzed. For a comparison, high performance
pure Se electrodes might also be developed and analyzed with the same approach.
150
V.5. Conclusion
In this chapter, we focused on characterization of Li2TiSexS3-x through electrochemical,
structural and morphological studies. We observed that these materials provide extra
capacity during discharge. To understand the origin of extra capacity and the reason
of degradation process upon cycling, various structural characterizations were done.
Ex situ XRD showed that cubic rocksalt structure reversibly changed upon cycling. In
addition, metallic selenium was observed in diffraction pattern and a columnar phase
was observed in SEM images of selenium-substituted samples. After such results, we
had hypotheses, such as either unreacted selenium remains in the electrodes or
selenium redox process is present (Se2- / Se0).
To investigate our hypotheses, we performed ex situ XPS study. Our first hypothesis
failed because we did not observe unreacted selenium in powders. XPS analysis on
Li2TiS3 showed that sulfur redox process is reversible. In selenium-substituted samples,
it seems that selenium redox is also reversible, however we cannot explain and confirm
the exact Selenium redox process at this point. The deconvolution at S 2p is quite
difficult due to the high number of contributions with very close energies, thus further
analysis is necessary.
Finally, the presence of LiF at Li 1s core and more oxidized sulfur species at S 2p core
in pristine electrodes suggest that sulfur might react with PVdF binder and produce
more oxidized sulfur species. This hypothesis should be investigated by further XPS
study on Li2TiSexS3-x electrodes prepared with different binders. It may also help to
clarify the degradation process upon cycling.
151
General Conclusion and Perspectives
152
The purpose of this thesis study is to develop overlithiated disordered rocksalt
materials that provide high specific energies (> 900Wh.kg-1 at material level) and
present a cobalt-free alternative to conventional materials used in Li-ion batteries. Li-
rich cation disordered rocksalt oxides can offer large capacities (≥ 250 mAh.g-1),
however there is always irreversibility in the first cycle and origin of this irreversibility
is still unknown. Meanwhile, disordered rocksalt Li2TiS3 has attracted research interest
thanks to its excellent reversible capacities (425 mAh.g-1 at 2.3 V vs Li+/Li). However,
detailed electrochemical and structural studies of such a promising material are still
missing, especially redox properties have not been reported yet. In addition, the
question of why Li2TiS3 cells failed in traditional batteries with liquid electrolytes has
not been answered yet. Therefore, we were interested to investigate those properties
and examine the reason of degradation during charge-discharge.
In this thesis project, we have studied overlithiated disordered rocksalt family in two
different axes: disordered rocksalt oxide (Li2.2NiTi0.2Nb0.6O4) and sulfide (Li2TiSexS3-x).
In the first part of the thesis, our aim was to optimize the performance of our patented
composition, disordered rocksalt Li2.2NiTi0.2Nb0.6O4. The disordered rocksalt (cubic)
structures are generally stabilized at high temperatures (1000 oC), however the
performances are limited due to particle growth. Thus, an optimization of the synthesis
parameters of our molten salt synthesis route was carried out. The nature of the
eutectic salt, the washing solvent, as well as the duration and the temperature of the
synthesis process were investigated. An optimum was found using 0.59LiCl-0.41KCl
eutectic salt at 670°C for 1h followed by a quench, and ethylene glycol as washing
media. The material delivered 166 mAh.g-1 of charge capacity with 108 mAh.g-1 of
reversible capacity. In situ XRD was performed to investigate the irreversibility (35%)
at the first cycle. The result showed that Li2.2NiTi0.2Nb0.6O4 has undergone structural
evolutions during charge-discharge. In addition, the particle size analysis showed that
dislocations or defects might have arisen in Li2.2NiTi0.2Nb0.6O4 powder, this
phenomenon may have been the cause of the irreversibility. This result indicated that
153
purity of the disordered rocksalts systems is important to design high energy
materials.
In the second part of the thesis, disordered rocksalt Li2TiS3 powders were prepared via
mechanochemical synthesis for shorter milling duration (15h - 25h) than those in the
literature (40h – 100h). As the Li2TiS3 electrodes were tested in carbonate-based
electrolyte, they delivered a reversible of 250 mAh.g-1, however a capacity decay
starting form the second cycle was observed. Thereafter, electrode pressing was
considered to improve its electronic conductivity and we observed that both reversible
capacities (339 mAh.g-1) and cycling stability were improved. To further improve
cycling stability of Li2TiS3 cells, the partial substitution of Li2TiS3 with selenium was
considered. Our efforts to prepare Li2TiSexS3-x compositions have resulted in the
discovery of a new synthesis method and high capacity compositions (Li2TiSeS2 and
Li2TiSe2S). X-ray diffraction patterns showed that all Li2TiSexS3-x powders that were
prepared in cation disordered rocksalt phase follow a solid solution behavior.
Thereafter, electrochemical properties of new materials have been investigated in
comparison with Li2TiS3. Li2TiSeS2 and Li2TiSe2S cells showed better cycling stability
than Li2TiS3 cells at higher cut-off discharge potentials (1.8V and 2 V) while cycling
stability of Li2TiS3 cells did not improve even at higher cut-off potentials. Furthermore,
Li2TiSe2S cells failed upon cycling between 3 V and 1.5 V. To find out the reason of
degradation during cycling, we decided to investigate Li2TiSexS3-x with deep structural
characterizations.
In the last part of the thesis, we examined the lithium insertion-extraction mechanism
through disordered rocksalt Li2TiS3 with structural studies (operando neutron
diffraction, FTIR and Raman). Unfortunately, poor crystallinity of Li2TiS3 powder
made definitive conclusions difficult even though various characterization tools were
used. Afterwards, Li2TiSeS2 and Li2TiSe2S electrodes were characterized by combining
both ex situ XRD and SEM techniques. Selenium metal was observed in the charged
electrodes, but it could not be detected in both pristine and discharged electrodes.
154
Thereafter, we had a hypothesis that there might have been Oswald ripening or
Selenium redox activity during charge. To tests these hypotheses, Li2TiSexS3-x powders
and Se precursors were further analyzed with X-Ray Photoelectron Spectroscopy (XPS)
analysis. The result showed that there were no remaining selenium precursors in the
Li2TiSexS3-x powders, thus our first hypothesis related to Oswald ripening was not true.
Thereafter, we decided to discover redox active species and elucidate redox processes
in Li2TiSexS3-x by XPS. At first, we identified species (elements) present in Li2TiSexS3-x at
their core levels and developed a fitting model for both powders and electrodes.
Surprisingly, high amounts of oxidized sulfur species at 2p sulfur core and LiF species
at Li 1s core were detected in pristine electrodes, showing the possible reaction of
sulfur in the Li2TiSexS3-x with the PVdF binder during electrode preparations. Ex situ
XPS analyses of Li2TiSexS3-x electrodes at different states of charge have confirmed our
previous observations that sulfur redox process was reversible. However, it is still
quite difficult to identify exact redox process in Li2TiSexS3-x due to the presence of a
high number of contributions of selenium species. We suggest that further XPS
analyses on Se metal or Li2TiSe3 can be useful to elucidate redox process in Li2TiSexS3-x
samples.
Overall, the degradation process of Li2TiS3 in liquid cells is still a question even though
various structural characterization tools were used to understand the underlying
reason. Further XPS analysis of Li2TiS3 electrodes at different states of charge and
during cycling (after 1st, 5th, 10th and 20th cycles for instance) might be useful to clarify
this issue. The degradation might result from the Li metal anode as the lithium metal
is very reactive. Thus, testing the Li2TiS3 electrodes against graphite or silicon type
negative electrodes can be interesting. In this thesis study, we could only investigate
reversible sulfur (anionic) redox process but did not focus on cationic process as the
titanium redox is generally active at lower potentials. To determine titanium redox
activity, various ex situ XPS experiments can be performed at Li2TiS3 electrodes cycled
at lower discharge cut-off potentials. XPS could also be useful to examine the true
155
selenium redox process and explain the main differences between sulfur and selenium
redox processes. To this purpose, a new Li2TiSe3 electrode can be characterized at
different states of charge. Consequently, different high energy compounds, such as Li-
excess disordered sulfides (Li2MS3 M: Ni, Co, Fe, Mn) or chalcogenides (Li2TiX3 X: S,
Se, Te) can be prepared with our patented synthesis process and characterized with
the techniques that were reported in this thesis.
156
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Appendix Table - 1: Binding energy of species at different core levels
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Résumé
« Synthèse et caractérisation de nouveaux matériaux d’électrode positive pour des
applications Li-ion à haute énergie »
Cette thèse concerne le développement de rocksalts désordonnés surlithiés pour les systèmes
Li-Ion à haute densité d’énergie. Dans un premier volet, les paramètres de synthèse ont été
optimisés pour améliorer les performances du rocksalt désordonné Li2.2NiTi0.2Nb0.6O4. Pour
comprendre l’origine de sa forte irréversibilité au premier cycle, des analyses in situ
structurales et électrochimiques avancées montrent un changement structural avec
l’apparition d’un désordre durant la première charge. Ensuite, le rocksalt désordonné Li2TiS3
a été préparé selon notre procédé breveté. De nouvelles compositions avec une substitution au
Sélénium, Li2TiSexS3-x, ont permis d’obtenir de fortes capacités de décharge à des potentiels
inférieurs avec une meilleure cyclabilité. L’activité réversible redox du soufre a été confirmée
par électrochimie et par analyses de surface ex situ mais des caractérisations plus poussées sont
nécessaires pour élucider le procédé redox complexe du sélénium.
Abstract
« Synthesis and characterization of new positive electrode materials for Li-ion high
energy applications »
This thesis focuses on the development of overlithiated disordered rocksalts for high-energy
Li-ion systems. Firstly, synthesis parameters have been optimized to improve the
performances of the disordered rocksalt Li2.2NiTi0.2Nb0.6O4. To examine its high irreversibility
(35%) at the first cycle, in situ advanced structural and electrochemical analyses have been
performed. Results show that a structural change and disordering happen during the first
charge. In a second part, the disordered rocksalt Li2TiS3 has been prepared with our patented
process. To improve cycling stability of the cells, Li2TiS3 has been partially substituted with
selenium and new Li2TiSexS3-x compositions have been prepared. Li2TiSexS3-x cells have large
discharge capacities at slightly lower potentials. Reversible sulfur redox activity is confirmed
by electrochemistry and ex situ surface analyses, however further characterizations are
required to elucidate the relatively complex selenium redox process.