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Synthesis and characterization of inverse spinels,
intercalation materials for Li-ion batteries
Proefschrift
ter verkrijging van de graad van doctor aan de Technische Universiteit Delft,
op gezag van de Rector Magnificus prof. dr. ir. J.T. Fokkema, voorzitter van het College voor Promoties,
in het openbaar te verdedigen op maandag 16 januari om 10.30 uur
door Nitte VAN LANDSCHOOT HBO-Ingenieur in de Chemische Technologie
Dit proefschrift is goedgekeurd door de promotoor: Prof. dr. J. Schoonman
Samenstelling promotiecommissie:
Rector Magnificus, voorzitter
Prof. dr. J. Schoonman, Technische Universiteit Delft, promotor
Prof. dr. ing. D.H.A. Blank Universiteit Twente
Prof. dr. ir. S. van der Zwaag Technische Universiteit Delft Prof. dr. I.M. de Schepper Technische Universiteit Delft
Prof. dr. S.W. de Leeuw Technische Universiteit Delft
Prof. dr. A. Schmidt-Ott Technische Universiteit Delft
Dr. E.M. Kelder Technische Universiteit Delft
Contents
1
Introduction
Abstract 9
1.1 General introduction 10
1.2 Li-ion batteries 11
1.3 High voltage cathode materials 15
1.4 Lithium metal vanadates LiMVO4 (M = Ni, Co) 17
1.4.1 The inverse spinel structure 17
1.4.2 The electronic configuration 18
1.5 Scope of this thesis 19
References 21
2
Solid-state synthesis and in-situ XRD of Li(Ni,Co)VO
4Abstract 25
2.1 Introduction 26
2.2 The solid-state synthesis of LiNiVO4 and LiCoVO4 27
2.3 Experimental aspects 29
2.3.1 Powder X-ray Diffraction (XRD) 29
2.3.2 Scanning Electron Microscopy (SEM) 31
2.3.3 The electrochemical cell 31
2.3.4 Cyclic Voltammetry (CV) 32
2.3.5 Cycle test 34
2.4 The results 34
2.4.1 Li-Ni-V-O and Li-Co-V-O compositions 34
2.4.2 Morphology 38
2.4.3 The electrochemical results 39
2.5 In-situ X-ray Diffraction 41
2.5.1 Introduction 41
2.5.2 LiNiVO4 vs. Li4Ti5O12 42
2.5.3 LiCoVO4 vs Li4Ti5O12 45
2.6 Conclusions 47
3
Solid-state synthesis of doped Li(Ni, Co)VO
4Abstract 51
3.1 Introduction 52
3.2 Experimental aspects 53
3.3 Results and discussion 54
3.3.1 X-ray Diffraction 54 3.3.2 The morphology 59 3.3.3 Cyclic Voltammetry 60 3.3.4 Cycle test 63 3.4 Conclusions 65 References 65
4
Citric-acid assisted synthesis of doped LiCoVO
4Abstract 69 4.1 Introduction 70 4.2 Experimental aspects 71 4.2.1 Synthesis procedure 71 4.2.2 Phase analyses 71 4.2.3 Raman Spectroscopy 72
4.3 Results and Discussions 73
4.3.1 Analysis of the LiCoVO4 formation process 73
4.3.2 Morphology 75
4.3.3 X-ray Diffraction 77
4.3.4 Raman Spectroscopy 80
4.3.5 Electrochemical results of Cu doped LiCoVO4 83
4.3.6 Electrochemical results of Cr doped LiCoVO4 86
4.3.7 Electrochemical results of Fe doped LiCoVO4 88
4.4 In-situ Raman Spectroscopy 91
4.5 Conclusions 94
References 95
5
A structural investigation to the influence of dopants on the
electronic properties of LiCoVO
4Abstract 99
5.1 Introduction 100
5.2 Experimental aspects 101
5.2.1 X-ray Photoelectron Spectroscopy 101
5.2.2 Mossbauer Spectroscopy 103
5.3 Results and discussion 106
5.3.1 X-ray Diffraction 106
5.3.2 X-ray Photoeelectron Spectroscopy 109
5.3.3 Mössbauer Spectroscopy 114
5.4 Impedance spectroscopy 118
5.4 Conclusions 121
References 122
6
Electrochemical performance and characterisation of
Al
2O
3coated LiCo
0.94Fe
0.06VO
4Abstract 125
6.1 Introduction 126
6.2 Experimental aspects 127
6.3 Results and discussion 129
6.3.1 The dissolution phenomena 129
6.3.2 Characterization of the Al2O3 coating 131
6.4 Conclusions 142
References 142
7
Summary and outlook
145Samenvatting en vooruitzicht 151
Dankwoord 157
Curriculum Vitae 161
1
Introduction
Abstract
This chapter presents a general introduction to the field of rechargeable batteries and the current state-of-the-art. A detailed explanation will be given of the working principles of Li-ion batteries including the advantages and disadvantages.
The investigated materials LiNiVO4 and LiCoVO4 will be discussed in relationship
1.1 General introduction
Fossil fuels are a finite resource of energy and their continued use, as the world’s dominant energy supply, is damaging the environment. Future use of alternative methods of energy supply is inevitable, and nowadays a lot of research is focused on the complete range of sustainable energy technologies, i.e., solar, biomass, wind, hydropower, geothermal power, ocean-energy sources, and solar-derived hydrogen fuel. Energy storage is a key element in achieving the goals for sustainability, air quality and cost-effective, competitive goods and services. In this respect, the question arises as to which will be the future challenges and requirements for energy storage technologies?
•
Reliable and affordable electricity storage is a prerequisite for usingrenewable energy in remote locations, integration into the electricity system, and the development of a future decentralised energy supply system.
•
Energy storage has a pivotal role to play in the effort to create a future with a sustainable energy supply, that has the standard of technical services and products we are accustomed to and need.•
For both stationary and transport applications, energy storage is of growing importance as it enables the smoothing of transient and/or intermittent loads, and down-sizing of base-load capacity with substantial potential for energy and cost savings.•
Customer electronics, such as mobile phones, laptops, palmtops etc., is a growing market for which energy storage systems are a necessity.Nowadays, there are many different energy storage systems. The best known energy storage systems are: Batteries, Flywheels, Fuel cells, Capacitors, BioFuels, Compressed air, and Pumped hydro storage. Rechargeable batteries, or accumulators, are the oldest form of electricity storage and are extremely widely used. Many of today's products are powered by primary or secondary (rechargeable) batteries. Lithium-ion batteries, Ni-Cd batteries, and nickel-metal-hydride (Ni-MH) batteries are the only new battery technologies which have achieved significant market penetration in the past decades, especially in the consumer electronics [1].
weight (Whkg-1) or per unit of volume (Whl-1), that a battery is able to deliver is a
function of the cell potential (V) and capacity (Ahkg-1), both of which are linked to
the chemistry of the system. Among the various existing technologies, as shown in Fig. 1.1, Li-ion batteries currently outperform other systems, accounting for 63% of worldwide sales values in portable batteries. Japanese manufactures Sanyo, Sony and Matsushita dominate the rechargeable market, from 75% (Ni-Cd) to 70% (Li-ion batteries) of the batteries made. The Li-ion batteries dominate the Japanese cell phone market (100%) versus only 15% in the United States of America and 30% in Europe. In China one-fifth of the world market 120 million Li-ion cells are produced per year. It is expected that by 2006 the use of Ni-Cd and Ni-MH use in notebook PCs will be obsolete [2].
acid 0 50 100 150 200 250 E ner gy d e n si ty [ W h l ] -1 Energy density [W h kg ]- 1 Ni-Cd Ni-MH Li ion Li metal PL ion
Fig. 1.1. Comparison of the different battery technologies in terms of volumetric and gravimetric energy density. The use of Pb–acid batteries is restricted mainly to SLI (starting, lighting, ignition) in automobiles or standby applications, whereas Ni–Cd batteries remain the most suitable technologies for high-power applications, i.e., power tools [3].
1.2 Li-ion
batteries
such as LiClO4 in propylene carbonate [4]. The commercialisation of primary
batteries followed relatively quickly in the 1960s and 1970s. The primary cells have been used mainly for watches, calculators, medical implants and military
applications, with Zn/MnO2 being dominant in the consumer market. Not
surprisingly, the development for rechargeable batteries was far from easy and much slower. Early success with the Exxon Li/TiS2 [5] in the mid 1970s as a coin
cell for electronic watches increased the research for new types of secondary systems. Since then, large numbers of secondary systems with different cathode materials, lithium metal or different lithium compound-based anodes and electrolytes have been studied and developed. Moli Energy became the first manufacturer of commercial Li-ion batteries in the world in the late 1980s. Unfortunately, in the summer of 1989, NTT of Japan recalled all batteries for its cellular phones that used Moli Energy’s Li/MoS2 secondary cells, because of
incidents involving cells that caught fire during assembly. These accidents brought worldwide attention to the safety risks associated with primary and secondary cells. Nowadays, a lot of research is still focussed on the safety issue [6]. To overcome these safety issues, e.g. Li-dendrite formation upon charging by the use of Li-metal [7], several alternative approaches were pursued with either by varying the electrolyte, or by substituting the Li-metal for an insertion material, the so-called Li-ion or rocking-chair technology as shown in Fig. 1.2.
Fig. 1.2. Schematic representation of the operating principles of a Lithium-ion battery. During charge the electrons flow through an external circuit to the anode and the Li+ ions migrate through the electrolyte and intercalate into the anode. The
redox reactions are for charging LiMeO2 → Li1-xMeO2 + x.Li+ + x.e- and C6 +
x.Li+ + x.e- → LixC6, and for discharge Li1-xMeO2 + x.Li+ + x.e- → LiMeO2 and
LixC6 → C6 + x.Li+ + x.e-.
The discovery of the highly reversible, low voltage, Li+
C/LiCoO2 rocking-chair cell commercialised by Sony Corporation in 1991 [9].
This type of cells have a voltage of 3.6 V (three times that of alkaline-based systems) and energy densities as high as 120-150 Whkg-1 (two to three times that
of Ni-Cd). The use of a carboneous anode lowers the cell potential with 0.3 V, which is unfavourable for the power density in Wkg-1. To overcome this issue new
anode materials are desirable which have a lower intercalation voltage. New anode materials, like nano-structured alloys or intermetallics, like Mg2Si, are being
researched thoroughly.
The open-circuit voltage Voc of a lithium battery is given by the difference
in the lithium chemical potential between the cathode (µLi(c)) and the anode (µLi(a)), ( ) ( ) Li c Li a oc V F µ −µ = (1.1)
where F is the Faraday constant. Voc is determined by the energies involved in both
electron transfer and Li+ transfer. While the energy involved in electron transfer is
related to the work functions of the cathode and the anode, the Li+ transfer is
determined by the crystal structure and the coordination geometry of the site into/from which Li-ions are inserted/extracted [10]. Kinetic stability considerations require the redox energies of the cathode (Ec) and anode (Ea) to lie within the
bandgap Eg of the electrolyte, so that no unwanted reduction or oxidation of the
electrolyte occurs during charging or discharging. Thus the electrochemical stability requirement imposes a limitation of the cell voltage as given in Eq. (1.2):
( ) ( )
oc Li c Li a g
eV =µ −µ <E (1.2)
A good lithium insertion material should satisfy the following criteria:
• A lithium chemical potential to maximise the cell voltage. This implies that the transition metal should have a high oxidation state.
• A large amount of lithium should be inserted and extracted reversibly from the host material to provide a high capacity.
• The lithium insertion and extraction process should be reversible with
minimal changes in the host structure.
• The insertion host should support mixed ionic and electronic conduction
(MIEC).
• The insertion host should be chemically stable against the electrolyte
within the entire range (x) of lithium insertion and extraction.
• The insertion host should be inexpensive.
Nowadays, most of the commercial Li-ion batteries contain a polymer electrolyte instead of a liquid electrolyte. These commercial Li-ion batteries have many advantages over other rechargeable batteries as listed below [11]:
Attractive thin & flat form factor
Rather than the traditional metal can used by other small rechargeable cells, Li-ion Batteries employ a thin (110 µm), polymer-based packaging material to contain the electrochemical materials. This allows the system to have a flat and moreover thin (2 to 5 mm) form factor.
Flexible designs (footprint & thickness)
Since the case of the cell starts as a sheet of polymer-laminate, changing the footprint of the cell is very easy and cost effective. Also, when the Li-ion batteries use a 'stacked' construction, adjusting the electrode/electrolyte structure is also easy. In this way, Li-ion batteries exhibit flexibility in their mechanical properties and flexibility in their design and construction. In particular, for a stacked Li-ion battery, changing the thickness is very easy.
Excellent cycle life
Cycle life is a measure of the number of times it is possible for the required capacity (or energy) to be extracted and replaced from a cell. Li-ion batteries have excellent cycle life. Cycle life is affected by a range of parameters such as charge time, operating temperature (both during the charge and the discharge), the depth of discharge (i.e., the percentage of the total capacity used during each discharge) and the time between cycles. In general, cycle numbers of over six hundred (600) are achieved before the delivered capacity falls below 60% of the initial capacity. In addition, 300 cycles are normally achieved before the capacity falls below 80% of the initial capacity.
Low self-discharge
No memory effect
The so-called memory effect is a major drawback for NiCd technology. Although the effect is not well understood, the practical manifestation is that NiCd and NiMH cells, which have been subjected to repeated partial discharge, seem to 'remember' the electrical capacity last delivered and only then deliver this on the next discharge. The memory effect is removed only by undertaking a slow full discharge, which again causes major logistical problems. Li-ion batteries have no memory effect and can be recharged at any time without the need for a full discharge.
High charge efficiency
Li-ion batteries are almost perfectly charge efficient. This means that the number of Ah required to charge a Li-ion Polymer cell is >99.8% of the Ah delivered on the subsequent discharge. This charge efficiency allows chargers to be reduced in size and valuable energy is saved.
Excellent safety features
Li-ion batteries do not use metallic lithium or lithium alloys, and lithium is present only in the ionic form in a mixed ionic-electronic host. This greatly improves the safety of Li-ion batteries. Other safety aspects, however, are associated with the use of a separator material. The separator provides the so-called thermal shutdown during thermal runaway of the system.
Environmentally friendly
Li-ion batteries are environmentally friendly and do not contain Cd, Hg, or Pb. A lot of research is performed on the replacement of the relative expensive and toxic Co, used in LiCoO2 as cathode material in commercialised Li-ion cells, for either,
Ni or Mn.
Low-temperature operation.
Li-ion batteries are capable of delivering good capacity at low temperatures (e.g. – 20°C). Although the capacity delivered is much less than that delivered at room temperature, the fact that the use of portable electronic equipment at low temperatures is often limited in time, means that the user is still able to power its device.
1.3 High-voltage cathode materials
Currently, four high-voltage cathode materials are commercially available. These materials are based on LiNiO2, LiCoO2, LiMn2O4, and LiFePO4. These
vs. Li/Li+. An overview of different cathode materials is schematically shown in
Fig. 1.3.
Fig. 1.3. Comparison of the different intercalation materials in terms of their capacity and potential versus Li/Li+.
Recent developments of lithium batteries have shown that some three-dimensional cathode materials, Li1-xMn2-yMyO4, belong to the most promising
cathode materials for lithium batteries. Generally, compounds with a three-dimensional framework are more stable than two-three-dimensional compounds [12]. Lithium ion diffusion is easier in a three-dimensional framework than in a two-dimensional framework, because in the former structure the number of contact points of the diffusion paths for lithium ions is larger than in the latter. The spinel-phase materials are of extreme interest as the active cathode materials for their high energy density. Most of the high-voltage (> 4.5 V) cathodes have the spinel structure with the general formula Li1-xMn2-yMyO4 (M = Cr, Fe, Co, Ni, and Cu)
[13-21] or the inverse spinel structure, i.e., LiMVO4 (M = Ni, Co) [22]. An
exception is LiCoPO4, which has the olivine structure like, LiFePO4 [23]. The
voltage and structural charge and discharge voltage in the 5 V region depend on the transition metal ion M and the amount of cation substitution in Li1-xMn2-yMyO4. As
a result, the high-voltage capacity (> 4.5 V) of these manganese-based Li1-xMn 2-yMyO4 cathodes has generally been attributed to the other transition-metal ion
redox couples, i.e., Fe3+/4+, Ni2+/3+/4+, and Cu2+/3+.
The inverse-spinel materials, LiNiVO4 and LiCoVO4, have not been
studied intensively, despite their high theoretical capacities of about 148 mAh/g. Both materials are interesting candidates as cathode material for lithium-ion batteries, because of their high voltages of 4.3 - 4.9 V vs. Li/Li+. Although, their
crystal structure is well known since 1961, it is still unclear how the electrochemical properties relate to the structure.
1.4 Lithium metal vanadates LiMVO
4(M = Ni, Co)
Lithium metal vanadates were first discovered in 1961 by Jean Claude Bernier [24]. He used a solid-state method, using LiVO3 and a Ni or a Co oxide as
precursor. Two years later in 1963 Bernier et al. [25] published work on the crystallographic and magnetic study of two mixed spinel vanadate materials. The X-ray diffraction indicated a structure described as a solid solution of two spinels, e.g. the compound of LiCoVO4 comprised (5V2O5.7Li2O) + 2(V2O5.7CoO).
Magnetic susceptibilities were determined in the temperature region 80K to 1200K. The molar susceptibilities of the LiCoVO4 and LiNiVO4 obeyed the Curie-Weiss
law closely. An analysis of the combined X-Ray diffraction and magnetic data led to the compositions of Ni0.15V0.85[LiNi0.85V0.15]octO4 and Co0.3V0.7[LiCo0.7V0.3]octO4.
The use of these materials as electrodes in Li-ion batteries was discovered by Fey et al. in 1994 [22], showing that the Li+ ions in LiNiVO
4 are extracted at voltages
around 4.8 V vs. Li/Li+ and at 4.2 V for LiCoVO
4. To date, around 90 scientific
papers have been published on these materials and mainly on LiNiVO4 due to the
higher voltage vs. Li/Li+. However, the electronic conductivity of these materials is
too low. Therefore, dopants need to be explored to enhance the electronic conductivity of these materials and doping is addressed in this thesis.
1.4.1 The inverse spinel crystal structure
The family of spinels is a very large one and contains many stable and robust materials. The prototype “spinel” is the mineral MgAl2O4. Spinels have the
general formula Atet[B2]octX4, where A refers to the cations on tetrahedral (8a) sites
and B to the cations on octahedral (16d) sites of a cubic structure with space group symmetry Fd3m. The X anions, located at the 32e sites, form a cubic-close-packed array. There are 64 tetrahedral sites in a typical unit cell, one-eighth of which is occupied by the A cations, and 32 octahedral sites, one-half of which is occupied by the B cations [26]. In the inverse spinel, LiMtet[V]octO4, the octahedral sites are
occupied by the Li+ and M2+ ions (V
tet[LiM]octO4), and the tetrahedral sites are
Fig. 1.4. Schematic representation of the inverse spinel structure, LiMtet[V]octO4 (M
= Ni or Co): O2- (32e) large atoms, V5+ (8a) small dark atoms, Ni2+/ Li+ (16d) small
light atoms. A lgrey V-tetrahedron and a light Li/Co-octahedron are added to the figure.
Between the extremes of normal and inverse spinel structures, varying degrees of cation disorder can exist in which A- and B-site cations exchange positions. It would, therefore, appear at first glance that the unit cell of a spinel structure with 56 empty tetrahedral sites and 16 empty octahedral sites in the interstitial space could easily accommodate small guest ions, such as lithium ions within its interstitial spaces. The empty sites within the unit cell are the tetrahedral (8b, 48f) and octahedral (16c sites). The cation sites are all interconnected, each 8a site being surrounded by four 16c octahedral sites and 16d sites surrounding the 8b site. Due to the cubic structure, spinel materials offer a three dimensional lattice for insertion and extraction of Li+ ions [26]. The lithium-ion pathway in an inverse
spinel is 16d-8b-16d [27].
1.4.2 The electronic configuration
Both, LiNiVO4 and LiCoVO4, are transition metal oxides and have the
ability to undergo large variations in lithium concentrations, which is attributed to the valence electronic structure formed by the transition metal (Co and Ni) and oxide ions. The electronic properties of the transition metal oxide are decided to a large extent by the number of valence electrons and by the interactions of the d-orbitals of the transition metal ion with the p-d-orbitals of the oxide ion. In the inverse spinel structure, the transition metal ion is octahedrally coordinated by oxides. Ni2+ adopts (3d8, [t2g]6 [eg]2) and Co2+ adopts the high-spin configuration
(3d7, [t
2g]5 [eg]2) [28]. The d-orbitals that directly overlap the p-orbitals of the oxide
ion forming bonding and antibonding orbitals are referred to as egb and eg as shown
antibonding eg orbitals consist mainly of metal d orbitals. The d-orbitals, which do
not directly overlap with oxide p orbitals to form σ bonds, are the t2g orbitals [10].
While general trends can be rationalized with crystal field and molecular orbital models, more accurate numerical methods based on density functional theory (DFT) [29-31] are necessary to understand and predict the relationships between the electronic structure of a compound and the electrochemical properties. For the inverse spinels as studied in this thesis, no numerical studies have been reported yet.
Fig. 1.5 Schematic illustration of the bonding and antibonding orbitals that arise for a transition metal ion in an octahedral environment of oxide ions.
1.5 Scope of this thesis
This thesis describes the synthesis and the characterization of LiNiVO4 and
LiCoVO4, high voltage insertion materials for Li-ion batteries. The materials are
prepared via different synthesis routes and the influence of several dopants on the structural and electrochemical properties are discussed. A general overview of different high voltage insertion materials is given in this introductory chapter.
Chapter 2 presents the solid-state synthesis of LiNiVO4 and LiCoVO4. A
series of LixMyVzO4 (M = Ni or Co, and x+y+z = 3) compositions were synthesized
each having a different theoretical capacity. The different compositions are discussed based on the phase diagram. In-situ X-ray diffraction was used to investigate the structural changes during charging and discharging of LiNiVO4 and
LiCoVO4.
In chapter 3 the influence on the structure and electrochemical properties of substituting either Ni or Co in the inverse spinel structure for Cr, Cu, or Fe has been investigated using X-Ray diffraction, scanning electron microscopy, cyclic
voltammetry, and cycle tests. Doping levels up to 10 mol% were achieved. Neutron Diffraction experiments were also conducted on the pure and Fe-substituted inverse spinels.
In chapter 4 a citric acid-assisted complex synthesis method has been developed to prepare LiMxCo1-xVO4 with M= Fe, Cr, or Cu and x=0; 0.02; 0.04;
0.06; 0.08; 0.10. A detailed description of the formation process of LiCoVO4 using
this synthesis method is given. The structural properties are characterized, using XRD, Raman Spectroscopy, and SEM. The electrochemical properties are studied using cyclic voltammetry and cycle tests.
A more detailed interpretation of the electrochemical properties of the x=0.06 substituted LiM0.06Co0.94VO4 (M = Fe, Cr, Cu), as prepared by the citric
acid-assisted synthesis methods described in chapter 4, is presented in chapter 5. The materials are examined using X-ray Photoelectron Spectroscopy and also Mössbauer Spectroscopy for the Fe-substituted insertion material. AC-Impedance Spectroscopy is used to study the electrical conduction properties of the different materials.
In chapter 6 the cycling behaviour of Fe-doped LiCoVO4, partly coated
with Al2O3 via a wet chemical process in order to prevent the dissolution of the
different 3d-metals, is presented. The coating is analysed using TEM and X-ray Photoelectron Spectroscopy, while the influence of this coating on the electrochemical properties is being discussed.
References
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2
Solid-state synthesis
and in-situ XRD of
Li(Ni,Co)VO
4
Abstract
The lithium extraction and insertion of LiNiVO4 and LiCoVO4, prepared
via a solid-state reaction were investigated. Both cathode materials revealed a high theoretical capacity, but the capacity drops more than 50% during the first cycles.
With in-situ XRD it was found that for LiCoVO4, a second inverse spinel phase with
a lattice parameter of 8.276Å was observed, which was segregated from the initial inverse spinel phase. The segregated inverse spinel lattice parameter shifts towards
8.261Å upon Li-ion extraction. For LiNiVO4, the lattice parameter decreases from
2.1 Introduction
Since the discovery that lithium spinels [1] provide a stable three-dimensional interstitial space for unrestricted migration of lithium ions through the structure, researchers have shown great interest in exploiting this class of compounds as insertion electrodes for rechargeable lithium-ion batteries. Developments of Li-ion batteries have shown that some three-dimensional compounds are one of the most promising cathode materials for secondary lithium batteries. The hexagonal layered structure, e.g. LiNiO2, has weak bonding between
the O-Ni-O sandwich like layers and a variable c/a ratio that allows 2D Li+
transport, but constrains the volume of interstial space. Although this constraint makes the spinel structure selective for insertion of Li+ ions, it reduces ion mobility
and hence the Li+ ion conductivity σ
Li. Nevertheless, the oxospinel LiMn2O4 has a
sufficient high σLi to be used in commercially low-power cells. Diffusion of Li+
ions from a layered structure into a solvent is more difficult and the spinel structure is selective for lithium ions over solvent molecules (i.e. deleterious effects of solvent co-intercalation are avoided). The key attributes for any successful cathode material are:
•
A high free energy of reaction with lithium.•
The cathode should have ≥1/2 Li reacting per transition metal.•
A cell voltage >3.0 V.•
Rapid diffusion of lithium ions to give a high power density. Electrochemically this equates to a charge/discharge rate of 10 mA⋅cm-2 [2], or a C-rate of 3C.•
Complete chemical reversibility, which is most easily obtained if the structure does not change during insertion and reinsertion of lithium ions, i.e., Li4Ti5O12[3].
•
Low cost.•
Low toxicity.Inverse spinels possess a three-dimensional structure and, given their similar structure, the voltage difference between LiNiVO4 (4.8 V) and LiCoVO4
(4.3 V) is significant and implies that the presence and sites of nickel atoms in an inverse spinel or spinel structure may play an important role in the voltage behaviour of these materials. In Li(Co,Ni)VO4, the oxidation state of the cations is
believed to be Li 1+, Ni and Co 2+, and V 5+. It is not expected that the V atoms oxidize beyond the 5+ state so the possible extraction of Li+, is accompanied by
changes of the oxidation state of Ni and Co ions hence, from 2+ to 3+.
Several researchers have used different starting materials for preparing LiNiVO4 and LiCoVO4. Bernier et al. [4] used LiVO3 and NiCO3 as precursors,
LiNiVO4, a process not economical owing to the amount of energy and long
reaction time required. Ito et al. [5] used NiO rather than NiCO3 in the above
reaction, and increased the heating temperature up to 1000 °C. However, this heating process had to progress for as long as four days for producing LiNiVO4.
Fey et al. [6] synthesized LiNiVO4 by reacting LiNiO2 with V2O5 and effectively
reduced the required heating temperature. Nevertheless, heating the mixtures for 24 hours at 650°C was necessary. Prabaharan et al. [7] synthesized LiNiVO4 at
temperatures as low as 320 °C using the aqueous glycine-nitrate combustion
process. In order to reduce the reaction temperature and time for preparing LiNiVO4 powder, a hydrothermal process was also adopted. Orsini et al. [8]
synthesized LiNiVO4 using a new “chimie douce” method, which consisted of a
precipitation reaction occurring at a well-defined pH.
Similarly, the isostructural compound LiCoVO4 was obtained by reacting
LiCoO2 with V2O3 or V2O5 at 700 °C for one hour in air [6]. Solid-State synthesis
temperatures below 650°C are not recommended for these vanadates because they cause incomplete reactions producing CoO. A high-temperature LiCoVO4 sample
was prepared first by sintering the stoichiometric quantities of Li2CO3, Co3O4, and
V2O5 in air at 800°C for 12 hours. By using the low-temperature method LiCoVO4
was prepared by dissolving the stoichiometric quantities of LiOH ⋅ H2O, Co(NO3)2
⋅ 6H2O, and NH4VO3 in de-ionized water and dried at 150 °C for 12 hours resulting
in a brown precursor. The final product was obtained by heating the precursor at 500 °C for 48 hours [9].
The inverse spinel materials, LiNiVO4 and LiCoVO4, have been studied
intensively because of their high theoretical capacities of about 148 mAh/g and their high voltages of 4.3-4.9 V vs. Li. In addition, LiNiVO4 is also a potential
anode material [8], as it can reversibly react with seven lithium ions per formula unit when discharged below 0.2 V. This leads to a specific capacity of 800-900 mAh/g, which is approximately three times the capacity that can be obtained with commercial graphite electrodes.
2.2 The solid-state synthesis of LiNiVO
4and LiCoVO
4The synthesis method used here is the so-called solid-state synthesis. The starting materials, LiOH⋅H2O (Baker Analyzed, >99%), NH4VO3 (Acros Organics,
>98%), CH4CoO4·4H2O (Riedel de Haen, >99%), or NiCl2·6H2O (Merck, >99%)
Table 2.1. Program for the furnace for the 1st heat treatment.
Heating / Cooling rate (oC/min) Temperature (oC) Time (h) 2 110 5 5 600 10 5 RT 8
The powders were then ball-milled using a planetary micro mill (Fritsch Pulverisette 7) for 30 minutes in order to break up the agglomerates down to 5 µm. The powders were sintered again, using a different heat treatment procedure, as described in Table 2.2.
Table 2.2. Program for the furnace for the 2nd heat treatment.
Heating / Cooling rate
(oC/min) Temperature (oC) Time (h)
2 110 6
5 800 12
5 RT -
The resulting solids were cooled slowly to room temperature and afterwards pulverized by ball-milling. The ball-milling was done for one hour. The slow cooling, instead of quenching, was done to minimize the number of oxide ion vacancies and other defects, but also to minimize the amount of stress within the particles.
The solid-state synthesis method described here was used to prepare a series of different materials. The following series of LixNiyVzO4 and LixCoyVzO4
(x+y+z=3), shown in Table 2.3, was prepared each having a different composition. The series was prepared with each material having a different theoretical capacity (Qtheo) and a different occupancy of the different sites within the inverse spinel
structure. . theo n F Q M = ( 2.1) where:
Qtheo is the theoretical capacity (C.mol-1/g.mol-1=C/g)
Table 2.3. Different compositions of LixNiyVzO4 and LixCoyVzO4 (x+y+z=3). Composition of LiNiVO4 Mark Composition of LiCoVO4 Mark
Li1+aNiV1-aO4 A (a=0) Li1+hCoV1-hO4 H (h=0)
Li1+bNiV1-bO4; B (b=1/3) Li1+iCoV1-iO4 I (i=1/3)
Li1+cNi1-cVO4 C (c=1/3) Li1+jCo1-jVO4 J (j=1/3)
Li1+dNi1-2dV1+dO4 D (d=1/3) Li1+kCo1-2kV1+cO4 K (k=1/3)
Li1Ni1+eV1-eO4 E (e=1/3) Li1Co1+lV1-lO4 L (l=1/3)
Li1Ni1-fV1+fO4 F (f=1/3) Li1Co1-mV1+mO4 M (m=1/3)
2.3 Experimental
aspects
2.3.1 Powder X-ray Diffraction (XRD)
Since the wavelengths of X-ray photons are of the same dimensions as the lattice parameters of crystals, X-rays constitute a powerful tool for the structural characterization of crystal lattices, as already pointed out by von Laue in 1913 [10]. XRD is not only used for stress and texture measurements but also for quantitative phase analyses. The fundamental physical process involved, if a plane wave of X-rays impinges on a perfect crystal lattice, where atoms are arranged in the three spatial dimensions in a periodic way, is the diffraction of the primary beam by these atoms acting as diffraction centers. Constructive and destructive interference of diffracted beams result in the outgoing X-rays being constrained along certain directions, holding certain angles with the incident beam. Two equivalent analyses of the X-ray diffraction by periodic arrays of atoms were developed by von Laue and Bragg and both have been routinely used by crystallographers for the determination of crystal structures. Bragg's Law can easily be derived by considering the conditions necessary to make the phases of the beams coincide with the incident angle equals and the reflecting angle. The rays of the incident beam are always in phase and parallel up to the point at which the top beam strikes the top layer at atom z as shown in Fig. 2.1. The second beam continues to the next layer where it is scattered by atom B. The second beam must travel the extra distance AB + BC, if the two beams are to continue travelling adjacent and parallel. This extra distance must be an integral (n) multiple of the wavelength (λ) for the phases of the two beams to be the same:
nλ = AB +BC (2.2)
AB = d sinθ (2.3) Because AB = BC eq. (2.3) becomes,
nλ = 2AB (2.4)
Substituting eq. (2.3) in eq. (2.4) we have,
nλ = 2 d sinθ (2.5)
This formula is also known as Bragg’s Law.
Fig. 2.1. A schematic representation of diffraction of x-rays by a crystal.
Important material parameters, which are obtained from the diffraction pattern, are the lattice parameters. These lattice parameters can be calculated using Bragg’s Law (2.5) and the relationship between the crystal lattices (h, k, and l are the Miller induces) and the crystal constants, a, b and c, i.e..
2 2 2 2 1 hkl h k l a b c d = + + (2.6)
The inverse spinel is cubic and, therefore, a, b and c are equal. This results in only one lattice parameter. For a cubic crystal the lattice parameter (a) is, combining eqs. (2.5) and (2.6).
2 2 2 2sin( ) h k l a λ θ + + = (2.7)
2.3.2 Scanning Electron Microscopy (SEM)
Today, the scanning electron microscope (SEM) is utilized not only in material science and biology, but also in such diverse fields as medical development. The SEM is extremely useful in the characterization of materials for its range of magnification and depth of field. SEM allows imaging of surface morphology, inclusions, boundary interfaces, and general characterization of a specimen surface.
A beam of electrons is generated in the electron gun, located at the top of the microscope. This beam is attracted through the anode, condensed by a condenser lens, and focused as a very fine point on the sample by the objective lens. The scan coils are energized (by varying the voltage produced by the scan generator) and create a magnetic field, which deflects the beam back and forth in a controlled pattern. The varying voltage is also applied to the coils around the neck of the Cathode-ray tube (CRT), which produces a pattern of light deflected back and forth on the surface of the CRT. The pattern of deflection of the electron beam is the same as the pattern of deflection of the spot of light on the CRT. The electron beam hits the sample, producing secondary electrons from the sample. These electrons are collected by a secondary detector or a backscatter detector, converted to a voltage, and amplified. The amplified voltage is applied to the grid of the CRT and causes the intensity of the spot of light to change. The image consists of thousands of spots of varying intensity on the face of a CRT that correspond, to the topography of the sample. SEM measurements were performed in this work with a JEOL JSM 5800LV scanning electron microscope with acceleration voltages up to 30 kV. The SEM is equipped with an ISIS Link EDX spectrometer for element analysis. To obtain a good SEM image the samples were sputtered with gold using an Edwards Sputter Coater S150B with a current of 20 mA and a pressure between 6-8 mbar for 6 minutes.
2.3.3 The electrochemical cell
Electrochemical characterisations were performed using CR2320 coin-type cells. A schematic representation of the coin cell is shown in Fig. 2.2. Electrodes were made using a doctor blade technique on a 12µm thick aluminium foil. The electrodes contain 85-wt.% of active material, 6-wt.% of carbon black (MMM Carbon), and 9-wt.% of PVDF (Solef). From the foils, Ø15 mm discs were punched out. The cells were assembled in a helium-filled glove box with metallic Li or Li4Ti5O12 as the anode. 1M LiPF6 dissolved in EC/DMC 1:2 (Mitsubishi
on spinel electrodes. The electrochemical cells were cycled within different voltage windows and at different C-rates and are mentioned separately for each cell.
Fig. 2.2. A schematic representation of the CR2320 coin-cell type: c = The cap, d = Cu plate and spring, e = The cathode electrode foil, f = The separator, g = The anode (Li-metal), h = The can. The top SEM image is the side view of the separator material,; the lower SEM image is of a doctor-bladed LiNiVO4 electrode.
2.3.4 Cyclic
Voltammetry
Fig. 2.3. Cyclic Voltammetry: (a) cyclic potential sweep, (b) resulting cyclic voltammogram. Figure taken from ref. 11.
On the forward scan for the inverse spinel materials, the current response represents the partial oxidation from M2+ to M3+ (Co2+, Ni2+) in LiMVO
4. In the
mean time Li+ ions are extracted from LiMVO
4 to form Li1-xMVO4 (0<x<1).
During the reverse scan, the opposite process takes place and the Li+ ions are
inserted to form Li1-yMVO4 (0<y<x). Usually, several peaks are observed in a
cyclic voltammogram. They appear because the current rises when a reduction/oxidation potential of the electrode materials is reached. The current drops when a process that is diffusion limited begins, or is simply finished.
The situation in the case of lithium batteries becomes complicated by the following factors, i.e.,
•
The ohmic potential drop in the electrolyte solution increases the peakseparation in the forward and backward scan. Using several sweep rates and subsequently extrapolating to zero sweep rate can correct this.
•
There may be a charged double-layer near the cathode surface, which results in a capacitive current. This current contributes to the total current as well as to the peak separation.•
The irreversibility of the process itself causes errors.2.3.5 Cycle
test
The cycling behaviour of the cells was performed using a Maccor Battery S4000 test System that allows applying various in-house written charge/discharge programs. A standard test procedure has a 10-minute rest step. In this period the cell is expected to reach electrochemical equilibrium. Then, the cell is charged and discharged using different C-rates (C-rate is 1/Time for each (dis)charge step). This C-rate is often plotted vs. the obtained capacity to evaluate rate capability. The cells are cycled 10-20 times.
2.4 The
results
2.4.1. Li-Ni-V-O and Li-Co-V-O compositions
The Li-Ni-V-O compounds and compositions were synthesized using the starting materials as mentioned in setion 2.2 according to the compositions gathered in Table 2.3. Fig. 2.4 shows the XRD pattern of LiNiVO4 prepared via the
solid-state synthesis. The obtained powder is in good agreement with the Joint Committee on Powder Diffraction International Centre for Diffraction Data (JCPDS), now known as the ICDD File No. 73-1636 for LiNiVO4. The material is
phase pure and has a lattice parameter of 8.215 Å, which is in good agreement with the literature [6]. 10 20 30 40 50 60 70 80 90 [642] [533] [620] [440] [511] [422] [400] [311] [220] [111] In te nsity (a.u.) 2θ
Fig. 2.4. X-ray diffraction pattern of the LiNiVO4 powder synthesized at 800°C.
The results of the phase analysis of the different compositions of LixNiyVzO4 (x+y+z=3) using XRD and different JCPDS data files are listed in
Table 2.4.
Table 2.4. The nominal compositions and the obtained phases for LixNiyVzO4
(x+y+z=3) using solid-sate synthesis at 800°C in air. The letters represent the different compositions used in the ternary phase diagrams
Starting materials Compound Products Mark
LiNiVO4 LiNiVO4, A
Li4/3NiV2/3O4 LiNiVO4, Li3VO4, NiO B
Li4/3Ni2/3VO4 LiNiVO4, Li3VO4 C
Li4/3Ni1/3V4/3O4 LiNiVO4, LiVO3, Li2NiO2 D
LiNi4/3V2/3O4 LiNiVO4, Li3VO4, NiO E
LiNi2/3V4/3O4 LiNiVO4, LiVO3, Ni3V2O8 F
NiC2O4⋅ 2H2O
LiOH ⋅ H2O
NH4VO3
Li1.5NiVO4 LiNiVO4, Li3VO4, NiO G
The XRD results will be discussed as a function of the oxidation state of Li, Ni, and V. The main product LiNiVO4 and the impurities Li3VO4, LiVO4, Li2NiO2 and
NiO were identified from the diffraction patterns using a JCPDS database. The oxidation states, in all the compounds, thus including the impurities, are always the same, i.e., Li 1+, Ni 2+, V 5+. For that reason the following oxides Li2O, NiO, and
V2O5 were selected as boundary materials for constructing a phase diagram.
By drawing the concordant tie lines we have found different areas where we expect the various compositions based on the tie lines.
III
Fig 2.5. Phase diagram for LixNiyVzO4 (x+y+z=3) using a solid-sate synthesis at
800°C in air.
Sample F falls in the triangle LiNiVO4 – LiVO3 – Ni3V2O8, shown in phase
diagram I, that is in line with expectations based on the XRD results. Samples E, B, and G fall in the triangle LiNiVO4 – Li3VO4 – NiO, shown in phase diagram II,
also in line with the XRD results. Sample C should be in the triangle LiNiVO4 –
LiVO3 – Li3VO4, but probably due to a small weighing error it has shifted and lies
on the LiNiVO4 – Li3VO4 line resulting into a two-phase system. Besides, the
amount of LiVO3 could be less then 1% and, therefore, cannot be detected with
XRD. Sample D falls in the triangle LiNiVO4 – LiVO3 – Li2NiO2 (III) and this is in
accordance with the XRD results shown in Table 2.4.
The Li-Co-V-O compounds and compositions were synthesized using the starting materials as mentioned in section 2.2 according to the compositions gathered in Table 2.3. Fig. 2.6 shows the XRD pattern of LiCoVO4 using the
solid-state synthesis procedure. The obtained XRD pattern is in good agreement with the JCPDS File No. 38-1396 for LiCoVO4. LiCoVO4 was found to be phase pure with
a lattice parameter of 8.276 Å. The background/peak intensity is lower compared to that of LiNiVO4. This has also been observed by other groups [13]. The scatter
1 0 2 0 3 0 4 0 5 0 6 0 7 0 8 0 9 0 [6 4 2 ] [5 3 3 ] [6 2 0 ] [4 4 0 ] [5 1 1 ] [4 2 2 ] [4 0 0 ] [3 1 1 ] [2 2 0 ] [1 1 1 ] In ten si ty ( a .u.) 2θ
Fig. 2.6. X-ray diffraction pattern of the LiCoVO4 powder synthesised at 800°C.
The numbers in the graph represent the [hkl] Miller indices.
The results of the phase analysis of the different compositions of LixNiyVzO4 (x+y+z=3) using XRD and different JCPDS data files are listed in
Table 2.5.
Table 2.5. The nominal compositions and the obtained phases for LixCoyVzO4
(x+y+z=3) using solid-sate synthesis at 800°C in air. The letters represent the different compositions used in the ternary phase diagrams.
Starting Materials Nominal composition Products Mark LiCoVO4 LiCoVO4, H Li4/3CoV2/3O4 LiCoVO4, Li3VO4, Co3O4 I Li4/3Co2/3VO4 LiCoVO4, Li3VO4 J Li4/3Co1/3V4/3O4 LiCoVO4, LiVO3 K LiCo4/3V2/3O4 LiCoVO4, Co3O4 L
LiCo2/3V4/3O4 LiCoVO4, LiVO3, LiV3O8 M
CoCl2⋅ 6H2O
LiOH ⋅ H2O
NH4VO3
Li1.5CoVO4 LiCoVO4, Li3VO4, Co3O4 N
For the Li-Co-V-O compositions a similar phase diagram as constructed for Li-Ni-V-O is difficult to obtain, because some powders (I, L and N) contain Co3O4 as impurity and Co3O4 has a Co-oxidation state of 2.67+, therefore, these
impurities cannot be used for this phase diagram as shown in Fig. 2.7. Slobodin et al. [14] described different Li-Co-V-O compositions with Co2+ and these results are included in Fig. 2.7 (lighter coloured tie lines LiCoVO4 - Li6Co2V2O10 – Co2V2O7).
much lower [400°C - 450°C] than used in the present experiments. The area’s described in this work agree with respect to the Li-Co-V-O phase diagram. The samples that could not be included into the phase diagram shown in Fig. 2.7, were all outside the triangle Li3VO4 – Li3VO8 – LiCoVO4. It seems that in order to keep
the oxidation state of Co at 2+, a much lower synthesis temperature is required. This is explained by the CoO – Co3O4 transition temperature in air, which is
650°C. A phase diagram consisting of LiCoVO4, CoO2 and LiV3O4 as boundaries
would show that the powders I, L and N would fit in such a phase diagram.
Fig 2.7. Phase diagram for LixCoyVzO4 (x+y+z=3) using a solid-sate synthesis at
800°C in air.
2.4.2 The morphology
The results shown in the previous section show that only LiNiVO4 and
LiCoVO4 are phase pure. No other stoichiometric inverse spinel can be expected
based on Li-Ni-V-O and Li-Co-V-O. To investigate the particle size and particle size distribution, we shall only focus on the pure materials. After ball milling the powders for one hour, the particle sizes range from 1µm to 30 µm as shown in Fig. 2.8 a and b. After 5 hours ball milling the particle size is 0.5-5µm as shown in Fig. 2.8 c and d. The particle size distribution is large and seems not favourable for lithium insertion and extraction. The SEM images of LiNiVO4 and LiCoVO4 are
(a) (b)
(c) (d)
Fig. 2.8 SEM image of (a) LiCoVO4 after one hour ball milling (b) LiNiVO4 after
one hour ball milling, (c) LiCoVO4 after five hours ball milling, and (d) LiNiVO4
after five hours ball milling.
2.4.3 The electrochemical results
A LiNiVO4 vs Li4Ti5O12 cell was cycled between 2.0 V and 3.4 V for 4
cycles. Fig. 2.9 shows the cyclic voltammogram. For these experiments we used Li4Ti5O12 [3] as the anode because it lowers the cell potentials. Besides, Li4Ti5O12
2.0 2.2 2.4 2.6 2.8 3.0 3.2 3.4 3.6 -1.5x10-4 -1.0x10-4 -5.0x10-5 0.0 5.0x10-5 1.0x10-4 1.5x10-4 2.0x10-4 0 1 2 3 4 0 20 40 60 80 100 R elat ive pe ak a re a [% ] Cycle number Charge capacity Discharge capacity
Insertion of Li+ during discharging
Extraction of Li+ during charging 4th cycle 1st cycle C ur rent ( A ) Voltage (V) 2.8 3.0 3.2 3.4 3.6 3.8 -3.0x10-5 0.0 3.0x10-5 6.0x10-5 9.0x10-5 C urr en t (A) Voltage (V)
Fig. 2.9. Cyclic voltammogram of a LiNiVO4 vs Li4Ti5O12 cell run between: (a)
2.0-3.4V, (b) 2.0-3.7V with a scan rate of 0.05mV/s. The inset in figure 2.9a reflects the capacity measured by the integrated peak area as a function of the cycle number.
The cyclic voltammogram of Fig. 2.9 shows differences between the first cycle and subsequent cycles. In general, the capacity is decreasing, but no extra peaks appear. A large cathodic peak around 3.4 V appears during the cathodic scan, reflecting the deintercalation of the Li+ ions at this voltage. In the reverse scan,
there is one anodic peak around 3.1 V originating from the intercalation of the Li+
ions into LixNiVO4. The peak area, which reflects the capacity, is found to decrease
in size with increasing cycle number. From this figure it is also observed that the capacity of the LiNiVO4 vs Li4Ti5O12 degrades over four cycles. When the cell is
observed in the anodic scan, due to the instability of the cathode material at this cut-off voltage. Hence, a parasitic reaction must occurs at this potential.
A LiCoVO4 vs Li4Ti5O12 cell was cycled between 1.5 V and 3.0 V for 5
cycles as shown in Fig 2.10. The CV’s from Fig. 2.10 reveal the reversibility of the LiCoVO4 compound for Li+ insertion/extractions. In the anodic scan a large anodic
peak around 2.3 V is observed. The decrease in sizes of the cathodic and of the anodic peaks with the cycle number shows the loss of capacity upon cycling, as shown in the inset in the figure.
1.6 1.8 2.0 2.2 2.4 2.6 2.8 3.0 -6.0x10-4 -4.0x10-4 -2.0x10-4 0.0 2.0x10-4 4.0x10-4 6.0x10-4 8.0x10-4 1.0x10-3 0 1 2 3 4 5 0 20 40 60 80 100 Rel at iv e pe ak ar ea [% ] Cycle number Charge capacity Discharge capacity Insertion of Li+ during discharging Extraction of Li+ during charging 5th cycle 1st cycle Cur re n t [A ] Voltage [V]
Fig. 2.10. Cyclic voltammogram of a LiCoVO4 vs. Li4Ti5O12 cell, run between
2.0-3.4V with a scan rate of 0.05mV/s.
2.5 In-situ X-ray diffraction
2.5.1 Introduction
In Li-ion batteries, both the negative and positive electrode, as shown in Fig. 2.2, contain the active material, an electronic conductive powder (Carbon) and a polymer binder. These porous electrodes are impregnated with electrolyte during the assembly to reduce polarization. These materials are able to accommodate variable quantities of lithium ions. As a rule, the lithium insertion or de-insertion into the electro-active material results in lattice changes of the insertion host. These changes may be relatively minor, e.g., a small expansion or contraction of the lattice, i.e., Li+ insertion into Li
4Ti5O12 [3], but can also become major, involving a
changes can be conveniently studied by means of in-situ XRD. However, great care must be exercised in designing the cell. This is because XRD probes the bulk of the particles, while their potential is determined mainly by the lithium concentration at the particle surfaces [17]. Thus, rapid equilibration of the entire insertion electrode should be achieved in a well-designed cell. In Fig. 2.11 the schematic drawing of the in-house built in-situ cell is presented, which was used to study the structural changes in the Li(Ni,Co)VO4 cathode materials. The cell consisted of two stainless
steel parts separated by a PTFE ring. The aluminium foil was used for sealing and as a current collector. A Bruker D8 Advance X-Ray Diffractometer with Cu-Kα radiation was used, while the cell was cycled using a MACCOR battery tester. The cells, LiNiVO4 vs. Li4Ti5O12 and LiCoVO4 vs. Li4Ti5O12, were charged at
0.2mA/cm2 and discharged at 0.1mA/cm2. LiPF
6 dissolved in EC/DMC 1:2
(Mitsubishi Chem.) was used as the electrolyte and Solupor (DSM Solutech) as the separator. The selected 2θ range was between 52° and 68°, thus showing the [422], [511], and [440] Bragg peaks. The scan time was one hour. The in-situ experiments were carried out for the first three cycles. An aluminium peak arising from the current collector is observed at 65.3° and was used as an internal reference. 1 2 3 4 5 6 7 8
Fig. 2.11. A schematic presentation of the in-house built in-situ cell. 1 = Stainless steel current collector, 2 = Stainless steel top cover, 3 = PTFE ring, 4 and 5 = Rubber O-ring, 6 = Teflon ring on which the cell is assembled, 7 = Stainless steel pin,⑧8 = Two screws.
2.5.2 LiNiVO
4vs. Li
4Ti
5O
12of the peaks are from the Cu-Kα2 reflections. A monochromator to filter the Cu-Kα2 reflections was not used in this measurement. Fig. 2.14 shows a peak shift during the charging of the cell resulting in a lattice parameter decrease followed by an increase upon subsequent discharging. This change in the lattice parameter from peak fitting is shown in Fig. 2.14 as a function of the cycle time.
0 5 10 15 20 25 30 2.0 2.5 3.0 3.5 4.0
. .
.
.. .
.
..
.
.
.
.
..
.
.
.
D C B A V oltage [V ] Time [hr]Fig. 2.12. Charge/discharge curve of LiNiVO4 vs. Li4Ti5O12. Icharge = 0.2mA/cm2,
Idischarge = 0.1mA/cm2. A, B, C and D indicate the time an XRD pattern was
recorded for cycles 2 and 3.
54 56 58 60 62 64 2.20 V 3.25 V 3.29 V 3.34 V 3.38 V 3.42 V 3.60 V 3.52 V 3.50 V 3.47 V 3.45 V 3.40 V 3.31 V OCV 3.60V A 3.60V C 2.20V B 2.20V D In ten si ty [a .u.] 2θ [°]
Fig. 2.13. In-situ X-ray diffraction patterns of the LiNiVO4 electrode for 2θ =
0 5 10 15 20 25 30 8.18 8.19 8.20 8.21 8.221 0.53 0.80 0.52 0.66 0.74 0.63 x in LixNiVO4 C D C D 1st Discharge (D) 1st Charge (C) La tt ic e par amet er [ A ] Time [hr]
Fig. 2.14. Calculated lattice parameters of LiNiVO4 versus charge/discharge time
as derived from Fig. 2.13.
The ex-situ XRD patterns as shown in Fig 2.15 show that the inverse spinel structure stays intact with a change in the lattice parameter, as seen in Fig 2.14. The intensity of the diffraction peaks does not change during cycling indicating a stable host structure, i.e., no cracking or internal stress as a result of the extraction or insertion. No phase transitions were observed during extraction, which occurs for LiMn2O4 [18]. The capacity fading is not due to the structural degradation of the
inverse spinel. Another reason for the capacity fading is the diffusion of the different cations, like Ni or V, to the empty octahedral sites. It is known that Ni diffuses within LiNiO2 during charging [20, 21]. Vanadium is also known to
diffuse through the spinel LiV2O4 [22] during extraction. Although, the ex-situ
XRD pattern did not show a change in the peak intensity ratio between the different diffraction peaks, it is still very likely that also in LiNiVO4 the cations diffuse to
10 20 30 40 50 60 70 80 90
*
* reflection of the Al-foil
* *
Pristine sample after 5 cycles in the discharged state
In te ns ity [a .u .] 2θ [°]
Fig. 2.15. Ex-situ XRD patterns of LiNiVO4 for a fresh electrode and an electrode
after 5 cycles measured between 10°-90° in the discharge state. The peaks indicated with an asterisk represent the aluminium current collector reflections.
2.5.3 LiCoVO
4vs Li
4Ti
5O
12Fig. 2.16 shows the first three cycles and shows a similar capacity fading as observed for LiNiVO4. According to the XRD pattern, the unit cell parameter of
LiCoVO4 is calculated to be 8.276Å A monochromator was used in this
experiment to filter out the Cu-Kα2 reflections.
0 5 1 0 1 5 2 0 2 5 1 .0 1 .5 2 .0 2 .5 3 .0 3 .5
. .
.
.
.
.
.
.
.
.
.
.
. .
.
..
D C B A Vo ltage [ V ] T im e [h r]Fig. 2.16. Charge/discharge curve of LiCoVO4 vs. Li4Ti5O12. Icharge = 0.2mA/cm2,
Idischarge = 0.1mA/cm2. The dots indicate when an XRD pattern was recorded and A,
Furthermore, Fig. 2.17 reveals the appearance of a new inverse spinel compound upon Li-extraction of LiCoVO4. The lattice parameter of this segregated
spinel phase shifts to lower values upon Li-extraction. Moreover, Fig. 2.18 shows the ex-situ XRD pattern of a fresh LiCoVO4 foil and one after the 5th cycle. It is
clearly seen that the second inverse spinel phase is present.
52 54 56 58 60 62 64 66 68 0 200 400 600 800 1000 1200 3.40 V D 2.00 V C 3.40 V B 2.00 V A In te n sit y [a .u .] 2θ
Fig. 2.17. In-situ X-ray diffraction patterns of a LiCoVO4 electrode for 2θ = 52-68°
over the potential range of 2.0 to3.40 V. The voltages are the end voltages, where A, B, C and D correspond to the same voltages as shown in Fig. 2.15.
In summary, the initial inverse spinel exhibits a constant lattice parameter, i.e., 8.276Å, while a new phase reveals a lattice parameter shift towards the value of 8.261Å. This gradual change in the lattice parameter is calculated from the peak shift. The line broadening is due to the presence of a thin shell at the interface of the particle of depleted Li+ ions. This shell formation is attributed to the low bulk
diffusion coefficient of Li+ ions. This is then reason for the partial Li-ion extraction
and hence, limits the utilization of LiCoVO4, as a relatively large amount of
28 32 36 40 44 48 52 56 60 0 50 100 150 200 After 5thdischarge OCV Aluminium
*
*
*
In ten sit y [a .u.] 2θFig. 2.18. Ex-situ XRD patterns of LiCoVO4 for a fresh electrode and an electrode
after 5 cycles, measured between 28°-60°in the discharge state. The shoulders of the new peaks are not from the Cu-Kα2 reflections. The peaks indicated with an asterisk represent the aluminium current collector reflections.
2.6 Conclusions
Using a solid-state technique, a pure phase of LiCoVO4 and LiNiVO4 is
formed at 800°C. SEM analyses indicate that the particle size of the powders is 6-30µm and after ball milling the powder for 5 hours the particle size decreases to 0.5-5µm. Using a lower synthesis temperature, i.e., <700°C, LiVO3 is formed and
at higher temperatures, >900°C, NiO or CoO is formed. The synthesis of different
Li-Ni-V-O compounds at 800°C was only possible for LiNiVO4. The other
compounds exhibit different impurities which, all can be explained in a phase diagram with NiO, V2O5, and Li2O as boundary materials. The oxidation state of Li
1+, V 5+, and Ni 2+ did not change during the synthesis. For the Li-Co-V-O compounds, a phase diagram having CoO, Li2O, and V2O5 as boundaries, is
constructed. When a sample is outside the triangle LiCoVO4-Li3VO4-LiV3O8 the
oxidation state of the Co 2+ changes to Co 2.67+. The CoO-Co3O4 transition
temperature in air, i.e., 650°C, is the most rationall reason for that change. Both LiNiVO4 and LiCoVO4 exhibit a low initial charge capacity, which
materials is different as observed from the in-situ XRD measurements. The capacity fading of LiNiVO4 and LiCoVO4 was investigated by in-situ X-ray
diffraction. The results clearly show that the lattice parameter of LiNiVO4
decreases from 8.215Å to 8.185Å during Li-extraction and increases to 8.205Å during Li-insertion. During the following cycles the initial lattice parameter is not reached. No new phases were observed. The intensity of the diffraction peaks does not change during cycling indicating a stable host structure, i.e., no cracking or internal stress occur as a result of the Li-ion extraction or insertion. The in-situ
measurements of LiCoVO4 reveal that two phases are formed during the
Li-extraction. This phase transformation takes place during the first cycle. This new phase has a lattice parameter of 8.261Å and its lattice parameter increases to 8.276 Å during Li-insertion. Furthermore, the in-situ results show that the lattice parameter of this new phase shifts during insertion and extraction of the Li-ions and that the lattice parameter of the initial phase does not shift. This is explained by a surface layer of shell formation, attributed to the low diffusion coefficient of the Li+ ions. This then accounts for the partial Li-ion extraction and, thus, limits the
utilization of LiCoVO4, because a relatively large amount of Li-ions will remain in
the bulk.
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