Download Low-Cost Iron-Based Cathode Materials for Large-Scale

Digital Comprehensive Summaries of Uppsala Dissertations
from the Faculty of Science and Technology 179
Low-Cost Iron-Based Cathode
Materials for Large-Scale Battery
Applications
ANTON NYTÉN
ACTA
UNIVERSITATIS
UPSALIENSIS
UPPSALA
2006
ISSN 1651-6214
ISBN 91-554-6559-5
urn:nbn:se:uu:diva-6842
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List of papers
This thesis comprises a summary based on the following papers, referred to
in the text by their Roman numerals:
I.
The surface chemistry of carbon-treated LiFePO4 particles for
Li-ion battery cathodes studied by PES,
M. Herstedt, M. Stjerndahl, A. Nytén, T. Gustafsson, H. Rensmo,
H. Siegbahn, N. Ravet, M. Armand, J.O. Thomas and K. Edström,
Electrochem. Solid-State Lett., 6 (2003) A202.
II.
A neutron powder diffraction study of LiCoxFe1-xPO4 for
x=0, 0.25, 0.40, 0.60 and 0.75,
A. Nytén and J.O. Thomas,
Accepted for publication in Solid State Ionics (2006).
III.
Electrochemical performance of Li2FeSiO4 as a new Li-battery
cathode material,
A. Nytén, A. Abouimrane, M. Armand, T. Gustafsson and
J.O. Thomas,
Electrochem. Comm., 7 (2005) 156.
IV.
The Lithium Extraction/Insertion Mechanism in Li2FeSiO4,
A. Nytén, S. Kamali, L. Häggström, T. Gustafsson and
J.O. Thomas,
J. Mater. Chem. (2006) DOI:10.1039/b601184e (in press).
V.
An ab initio study of the Li-ion battery cathode material
Li2FeSiO4,
P. Larsson, R. Ahuja, A. Nytén and J.O. Thomas,
Accepted for publication in Electrochem. Comm.(2006) (in press).
VI.
Surface characterization and stability phenomena in Li2FeSiO4
studied by PES/XPS,
A. Nytén, M. Stjerndahl, H. Rensmo, H. Siegbahn, M. Armand,
T. Gustafsson, K. Edström and J.O. Thomas,
Submitted to J. Mater. Chem.
Comments on my own contribution to the papers in the thesis:
I.
Synthesis and electrochemical measurements. Participation in the
discussion of the results.
II.
Sample preparation and data analysis. Main author of the paper.
III.
The majority of the experimental work. Main author of the paper.
IV.
The majority of the experimental work. Main author of the paper.
V.
Planning the calculations and participation in the evaluation of the
results. Source of the experimental data.
VI.
Synthesis and electrochemical measurements. Planning the
experiments and full participation in the evaluation of the results.
Main author of the paper.
Contents
The scope of this thesis .................................................................................11
1. Introduction...............................................................................................12
1.1 Lithium battery history in brief ..........................................................12
1.2 The Li-ion battery...............................................................................13
1.3 Large-scale battery applications.........................................................15
2. Cathode materials for Li-ion batteries ......................................................16
2.1 Commercial cathode materials ...........................................................16
2.2 LiFePO4 ..............................................................................................17
2.3 LiCoxFe1-xPO4.....................................................................................19
2.4 Li2FeSiO4............................................................................................19
3. Methods ....................................................................................................21
3.1 Materials synthesis .............................................................................21
3.2 Electrode and battery fabrication .......................................................21
3.3 Electrochemical testing ......................................................................22
3.4 X-ray diffraction.................................................................................22
3.5 Neutron diffraction .............................................................................24
3.6 Photoelectron spectroscopy................................................................25
3.7 Mössbauer spectroscopy.....................................................................26
3.8 Scanning electron microscopy............................................................27
3.9 Theoretical calculations......................................................................27
3.10 Other techniques...............................................................................28
4. LiFePO4 as cathode material.....................................................................29
4.1 Surface chemistry of carbon-treated LiFePO4 ....................................29
4.2 Transition-metal substitution in LiFePO4 ...........................................32
5. Li2FeSiO4 as cathode material ..................................................................34
5.1 Synthesis and phase purity .................................................................34
5.2 Electrochemical performance.............................................................35
5.3 Structure .............................................................................................37
5.4 Redox mechanism ..............................................................................39
5.5 Surface chemistry ...............................................................................41
6. Conclusions and future work ....................................................................45
Acknowledgements.......................................................................................48
Summary in Swedish ....................................................................................49
References.....................................................................................................52
Abbreviations
AES
DEC
DFT
DOS
EC
EPDM
EV
GGA
HEV
HR-SEM
ICP
LiBOB
LiTFSI
MS
NiCd
NiMH
PC
PES
PSD
SEI
SEM
SR
TMO
XPS
XRD
Atomic emission spectroscopy
Diethyl carbonate
Density functional theory
Density of states
Ethylene carbonate
Ethylene propylene diene copolymer
Electric vehicle
General gradient approximation
Hybrid electric vehicle
High resolution scanning electron microscope
Inductive coupled plasma
Lithium bis(oxalato)borate
Lithium bis(trifluoromethylsulfonyl)imide
Mössbauer spectroscopy
Nickel cadmium
Nickel metal hydrid
Propylene carbonate
Photoelectron spectroscopy
Position sensitive detector
Solid electrolyte interphase
Scanning electron microscope
Synchrotron radiation
Transition-metal oxide
X-ray photoelectron spectroscopy
X-ray diffraction
The scope of this thesis
It would now appear relatively certain that the Li-ion battery concept currently used worldwide in mobile phones, lap-tops and recently in power
tools, is also destined to become the battery concept of choice for large-scale
applications such as electric and hybrid electric vehicles (EVs and HEVs) in
the near future. This will lead, directly or indirectly, to reductions in the use
of fossil fuels and thereby to an overall improvement in our environment
worldwide. This development has motivated our search for new low-cost
active cathode materials with retained high energy- and power-density. A
commercial Li-ion battery today generally incorporates LiCoO2 or
Li(Co,Ni)O2 as cathode material, but these will prove too expensive for large-scale applications; a much cheaper transition metal, preferably the very
cheapest - iron (Fe) - is to be preferred, combined with a (cheap) graphitebased anode. The quantities needed in a large-scale battery will be huge;
while a mobile-phone battery consumes only a few grams of some relatively
expensive active cathode material (like LiCoO2), an HEV battery could need
kilogram quantities to store sufficient energy. So far, no such material has
emerged on the commercial market, although a number of possible candidates are being developed.
The goal of this thesis work has therefore been to investigate and improve
Fe-based cathode materials in terms of their structural and especially their
electrochemical properties. Compounds with the general formulation
LiyMXO4, for M=Fe, Co and X=P, Si have been investigated using a variety
of techniques. Photoelectron spectroscopy (PES) using both laboratory- and
synchrotron-based radiation has been used to study the surface chemistry for
carbon-coated LiFePO4 (paper I), since it is vital to understand the reactions
occurring at the electrochemical interfaces within the lithium-ion battery to
be able to improve both their safety and performance. Neutron powder diffraction was used in the structural study of LiCoxFe1-xPO4 (paper II). A
combination of X-ray diffraction (XRD), Mössbauer spectroscopy (MS),
photoelectron spectroscopy (PES), electrochemical measurements and DFT
calculations were then used in a set of studies of the new cathode material
Li2FeSiO4 (papers III-VI).
11
1. Introduction
1.1 Lithium battery history in brief
Battery technology has come a long way since the Italian physicist Alessandro Volta in 1800 described the first electrochemical cell, which came to
be known as the Volta pile [1]. Today, the main focus in the battery World is
on the Li-ion battery, which has the great advantage over other battery concepts of providing a high gravimetric and volumetric energy density
(Fig. 1.1) [2]. G.N. Lewis initiated the first work on lithium-based batteries
already in 1912, but it was not until the early 1970s that the first commercial
primary lithium battery reached the market [3]. At the same time, intercalation compounds were proposed as cathode materials for secondary
(rechargeable) batteries. In these early batteries, TiS2 was used as active
cathode material, lithium foil as anode and lithium perchlorate in dioxolane
as electrolyte [4,5]. The combination of lithium metal and a liquid electrolyte
was not ideal, however.
Figure 1.1 Comparison of different battery technologies in terms of volumetric
(Wh/l) and gravimetric (Wh /kg) energy density; figure taken from [2].
12
Particularly, dendrite formation (leading to short-circuiting, thermal runaway
and fire) [6] and the fact that lithium reacts violently with moisture led to
metallic lithium ultimately being replaced by a second intercalation compound which could also host Li+ [7,8]. By the end of the 1980s and the
beginning of the 1990s this had led to the so-called Li-ion or “rocking-chair”
technology. In 1991, the Sony Corporation was the first company to introduce rechargeable Li-ion batteries onto the market for mobile-phone applications [9].
Another way to increase the intrinsic safety was to replace the liquid
electrolyte by a dry polymer electrolyte [10]. This concept still tends to be
limited to elevated-temperature applications, however, due to the low lithium
conductivity of dry polymer electrolytes at ambient temperatures.
1.2 The Li-ion battery
A Li-ion battery basically comprises three vital components: a cathode, an
anode and a separator containing some liquid electrolyte between. The cathode, which provides the lithium ions, is generally a transition-metal oxide
(TMO) based compound such as LiMO2 (M= Co, Ni or Mn,) [11-13] or
LiMn2O4 [14]. These types of material are attractive since they exhibit high
redox potentials (ca. 4V) with respect to Li/Li+. LiCoO2 and various doped
compositions such as LiNi1-x-yCoxMyO2 (M=Mg, Al, Ga or Ti) are commonly
used in today’s commercial Li-ion batteries. In this connection, there is considerable current interest in LiNi1/3Co1/3Mn1/3O2 [15,16].
A carbon-based (graphite) anode [17] is used almost exclusively in
today‘s commercial Li-ion batteries due to its low-cost and good cycling
performance. Other anode materials under current investigation involve
lithium-metal alloys, LixM, where M= Al, Sn, Si, Sb, etc. [18,19]. Although
these materials generally display high volumetric capacities, they tend to
suffer from high volume expansion/contraction upon cycling, which causes
structural degradation with a resulting reduction in cycle-life. Intermetallic
compounds like Cu6Sn5 [20], InSb [21] and Cu2Sb [22] have a somewhat
reduced volume change on lithium insertion but, since one of the component
metal atoms is electrochemically inactive (e.g., Cu in Cu2Sb), their
gravimetric capacity is lower.
The liquid electrolyte in the battery serves as conductor for Li-ions and
comprises some lithium salt dissolved in a mixture of liquid organic carbonates. Commonly used carbonates are: ethylene carbonate (EC), diethyl
carbonate (DEC), dimethyl carbonate (DMC) and propylene carbonate (PC).
Lithium hexafluorophosphate (LiPF6) and lithium bis(trifluoromethane
sulfonyl)imide (LiTFSI) are two of the more commonly used lithium salts.
13
On charging, lithium ions are extracted from the LiMOx positive electrode
and inserted into the graphite negative electrode, while electrons move
through the outer circuit (Fig. 1.2). On discharge, the reverse process takes
place.
Figure 1.2 A schematic representation of the Li-ion battery.
A number of parameters and terms are used to characterize the performance of a battery; typically:
gravimetric capacity (mAh/g) – the charge involved in the cell reaction per
mass unit; gravimetric energy density (mWh/g) – the energy per mass unit,
given by the operating potential (in V) u the gravimetric capacity (in
mAh/g); volumetric energy density (mWh/l) – the energy per volume unit;
charge and discharge rate (C/t) – t being the time in hours for one complete
charge or discharge.
For the cell reaction LinA o A + nLi+ + ne-, the theoretical gravimetric
capacity (Tcap in mAh/g) is given by:
Tcap
1 / M u nF
3 .6
where M is the molecular weight of LinA and F is Faraday’s constant.
14
1.3 Large-scale battery applications
There are today a number of standard viability requirements from the car
industry for batteries for electric and hybrid electric vehicles (EVs and
HEVs). Their price should be in the range 100-150 USD/kWh, energy density: 500-1000 Wh/kg, with a 5-10 year lifetime. The total World demand for
light hybrid electric vehicles is estimated to be ca. 4.5 million units by 2013.
This increase from ca. 310 000 HEV sold in 2005 is motivated by rising fuel
costs and increased emission regulations. The most important markets are
expected to be in the United States, Western Europe and Japan, although the
gains will be greater in the United States and Japan; HEV growth is projected to be much slower in Europe, where diesel-fueled vehicles already
exceed 40% of the total passenger car market. Another interesting market is
likely to be China [23].
Approximately 310 000 HEVs were sold worldwide during 2005 - around
205 000 of these in USA [24]. The most attractive vehicle in this context is
the widely-acclaimed Toyota PRIUS, which represents 55% of the total sales
in 2005. Toyota have recently (Dec. 2005) announced that they plan to
replace the NiMH batteries which they use today by Li-ion batteries in their
next-generation PRIUS, which is scheduled for 2008.
The Li-ion battery concept can also find other large-scale applications
outside the car industry. The most important of these is the storage of electrical energy produced by renewable power sources such as solar, wind and
wave energy. A practical example is the incorporation of Li-ion batteries
into today’s commercially available residential solar-heating systems.
To make a real breakthrough in these fields, a low-cost cathode material
with a high intrinsic capacity which is also environmentally acceptable is an
essential precondition. In this context, iron-based cathode materials, such as
LiFePO4, are of particular interest, since the raw materials involved are
potentially cheap. We must be aware, however, that it is equally important to
arrive at a corresponding low-cost synthesis process, which can be up-scaled
for mass production.
15
2. Cathode materials for Li-ion batteries
2.1 Commercial cathode materials
The most widely used cathode material in today’s commercial Li-ion batteries ars the layered lithium transition-metal oxides, Li(Co,Ni)O2. Although
LiNiO2 has a higher specific capacity and is potentially cheaper than LiCoO2, it is both chemically and electrochemically unstable, leading to both
synthesis and safety problems. LiCoO2 is superior in this connection and
was, indeed, the cathode used in Sony’s very first Li-ion batteries. The deintercalation and re-intercalation of Li+ into LiCoO2 takes place around
4 V, which should result in very high gravimetric energy densities for full Li
extraction, giving voltages up to 4.7 V). However, when such large amounts
of lithium are removed (at high potentials), irreversible structural changes
occur. In LiCoO2 and cobalt-rich LiNi1-xCoxO2 phases, this is caused mainly
by exothermic oxygen reaction with the electrolyte [25-27]. Such effects are
even more pronounced in the less thermally stable LiNiO2. The best
compromise is reached in nickel-rich LiNi1-xCoxO2, but even here other
structural instabilities arise in the highly delithiated state, where nickel ions
migrate into the lithium sites in the structure [28], thereby limiting lithium
diffusion and reducing storage capacity. For this reason, only 0.5 Li is
removed in commercial applications, corresponding to a cut-off at 4.2 V and
a theoretical capacity of 130 mAh/g.
A number of electrochemically inert substituents for nickel or cobalt have
been tested in attempts to stabilize the structure, e.g., Al, Ga, Mg and Ti. The
LiNi1-y-zCoyAlzO2 system has been found to give the best performance, and is
today used in many commercial applications. SAFT has constructed cells
with this substituted nickel oxide that have been cycled 1000 times at 80%
depth-of-discharge with an energy density of 120-130 Wh/kg [29].
Mixed metal oxides of nickel, cobalt and manganese, especially the
LiNi1/3Co1/3Mn1/3O2 system, also offer materials with lower cost and higher
stability compared to the pure LiCoO2. Such materials are currently under
careful scrutiny, both regarding the complex behaviour of the transition metals and their optimum composition. In the LiNi1/3Co1/3Mn1/3O2 system, it has
been shown that high capacities of 150 mAh/g for a 4.2 V cut-off and 200
mAh/g when charging up to 4.6 and 5.0 V [15,16] are attainable, and that a
16
certain amount of nickel in the lithium layer seems to stabilize the structure
and improve capacity retention.
2.2 LiFePO4
Intercalation into phosphate-based compounds has been known ever since
Delmas et al. reported the reversible intercalation of Na+ into the
NASICON-type phase NaTi2(PO4)3 to yield Na3Ti2(PO4)3 [30]. Following
the discovery of Li+ insertion into the insulator Fe2(SO4)3 [31], research focused on the electrochemical behaviour of other NASICON-related
compounds, e.g., Li3Fe2(PO4)3 [32]. At this stage, LiFePO4 itself was
considered as an impurity that was to be avoided in the synthesis of other
iron phosphate compounds. A year later, Goodenough first described a new
class of intercalation compounds which he called the phospho-olivines,
LiMPO4 (M=Fe, Mn, Co, Ni) [33] (still believing more in the NASICONrelated compounds [34]), but it was some time before LiFePO4 was actually
recognized as one of the most promising cathode materials for Li-ion
batteries. Its poor electronic conductivity was initially seen as an
insurmountable obstacle.
LiFePO4 has an orthorhombic unit cell (space-group: Pnma;
a= 10.3290(3) Å, b= 6.0065(2) Å, c= 4.6908(2) Å, V= 291.02(1) Å3 and
Z= 4) with the oxygen atoms arranged in a slightly distorted, hexagonal
close-packed arrangement (Fig. 2.1). The iron and lithium atoms occupy
octahedral sites; phosphorous atoms occupy tetrahedral sites. Corner-shared
FeO6 octahedra are linked together in the bc-plane, while LiO6 octahedra
form edge-sharing chains along the b-axis. The tetrahedral PO4 groups
bridges neighbouring layers of FeO6 octahedra by sharing a common edge
with an FeO6 octahedron and two edges with LiO6 octahedra. The strong
P–O bonds cause a polarization of the oxygen atom towards the phosphorous
atom in the Fe–O–P linkage, causing the covalent contribution to the bonding strength in Fe–O to decrease through the induction effect. This has the
effect of lowering the Fermi level of the Fe2+/Fe3+ redox couple and thereby
increasing the cell potential compared to the corresponding oxide LiFeO2.
17
Figure 2.1 The crystal structure of LiFePO4 viewed along the c-axis; dark and light
grey polyhedra represent PO4 and FeO6, respectively.
The fact that LiFePO4 is non-toxic, low-cost and has a high thermal and
mechanical stability makes it a major candidate as electrode material in
future lithium-ion batteries for large-scale applications (cf. Section 1.3). Under electrochemical cycling, it shows a very flat voltage plateau at 3.45V vs.
Li/Li+ with a theoretical capacity of 170 mAh/g based on the reaction
LiFePO4 o FePO4 + Li+ + e- [33]. Lithium extraction in this reaction was
long considered to be a topotactic two-phase process. It has been shown very
recently, however, that the situation is somewhat more complicated, involving Li-rich (0.9 d x d 1.0) and Li-poor (0 d x d 0.1) LixFePO4 phases,
separated by a broad immiscibility gap [35].
The major drawback with LiFePO4 is its low intrinsic electronic
conductivity. Various synthetic and processing strategies are used today in
efforts to overcome this problem. Decreasing the particle size [36-38] and
coating the particles with some highly conducting material (typically carbon)
[39-41] are two realistic solutions to the problem. Both approaches have
since been confirmed to go a long way towards giving full capacity
utilization [42]. The problems of low electronic conductivity would also
seem to be less serious at elevated temperatures [43].
18
2.3 LiCoxFe1-xPO4
Another possible means of increasing the degree of utilization of the
intrinsic capacity of LiFePO4 could be the partial replacement of Fe by some
other transition metal (Mn, Co or Ni) to increase the electronic conductivity
[44]. Improved capacity retention and better cycling stability at very high
rates (8-20C) have been shown for LiFe0.9M0.1PO4 (M=Ni, Co, Mg)
compared to both pure and carbon-coated LiFePO4 [45].
Electrochemical cycling of pure LiCoPO4 shows a voltage plateau at
4.8 V for the Co2+/Co3+ redox couple [46]. The high irreversible capacity
observed on cycling is probably not only an effect of poor electronic conductivity, but could also be due to electrolyte-salt decomposition occurring at
the high potential used. In a study of the electrochemical properties of the
solid solution series LiCoxFe1-xPO4, Penazzi et al. reported an increasing
capacity with respect to LiFePO4 at low Co levels (LiCo0.2Fe0.8PO4) [47]. In
a similar study of the same system, a continuous lowering of the Co2+/Co3+
plateau was observed with increasing amounts of Fe, which generated a better capacity retention [48]. The structure characterization of this solidsolution series was in both cases, however, only based on powder XRD
measurements. A structural study of the LiCoxFe1-xPO4 system for x= 0,
0.25, 0.40, 0.60 and 0.75 was therefore undertaken using neutron diffraction
to focus more closely on the structural detail of the Co/Fe substitution
(paper II).
2.4 Li2FeSiO4
In the search for new, low-cost cathode materials, the tremendous variety
offered by silicate chemistry motivated an effort to exploring the Li-Fe-Si-O
system. It was anticipated that we could find the same type of lattice stabilization effect (“induction effect”) as in LiFePO4 (cf. Section 2.1) as a result of
the strong Si-O bonds in the system. The lower electronegativity of Si vs. P
will result in a lowering of the Fe2+ œ Fe3+ couple. Such materials should
therefore have a lower electronic band gap, and therefore a higher electronic
conductivity. A material based on Fe and Si would also be potentially lowcost in terms of raw material costs; indeed, Fe- and Si-oxides represent
>10% of the Earth’s crust.
One of the first electrochemical investigations of this type of silicatebased material was conducted in 1994, when Armand attempted to
delithiate Li2MnSiO4 [49]. The observation of an irreversible oxidation peak
at 3.8-4.0 V was not encouraging and the investigations were not pursued
further until the discovery of the phospho-olivines. However, attempts to
synthesize LiFe3+SiO4 ended up with a multi phase product containing
mainly iron-based spodumene (LiFeSi2O6). Similar efforts to synthesize
19
LiCr3+SiO4 again resulted in the spodumene-type compound LiCrSi2O6. A
number of different lithium iron silicates were thereafter synthesised, but
none of these exhibited any electrochemical activity until the successful synthesis of Li2FeSiO4 in 2003 [50]. Although the final product contained an
appreciable amount of unreacted material (Li2SiO3), reversible electrochemical activity was observed (Fig. 2.2). Tarte et al. [51] had, in fact, described
this compound 30 years earlier as being isostructural with Li3PO4 (spacegroup: Pmn21) and determined its lattice parameters; its detailed structure
had not been probed more deeply. In order to achieve electrochemical
activity in Li2FeSiO4, the use of different carbon-coating techniques has
been of great importance.
The very first paper describing the electrochemical behaviour of
Li2FeSiO4 is included in this thesis (paper III). More detailed studies using a
combination of XRD, Mössbauer spectroscopy (MS) and electrochemical
cycling have probed the lithium extraction/insertion mechanism (paper IV);
DFT calculations are reported in paper V; and an analysis of the surface
chemistry is made in paper VI. The most significant results of these studies
are summarized in Chapter 5.
Figure 2.2 An early cyclic voltammogram for an impure Li2FeSiO4 sample at 80qC
and 20 mV/h. Cell configuration: < Li2FeSiO4 _ polymer electrolyte _ Li-foil >.
20
3. Methods
3.1 Materials synthesis
All cathode material studied in this thesis have been prepared by solidstate synthesis. In general, the starting materials were dispersed in acetone
and then mixed and ground thoroughly, either by planetary ball-milling or by
milling with marble balls in a plastic container on a roller-band. For the
Li2FeSiO4 compound, this latter type of mixing was carried out for at least
two months. After evaporating the acetone, the resulting powder was
annealed in a tube furnace at temperatures ranging from 500 to 700qC in a
N2-gas flow; CO/CO2 (50:50) gas was used for the case of Li2FeSiO4. In all
cases, it was necessary to ensure an oxygen-free atmosphere to suppress the
formation of Fe3+ species. More detailed information about the different
starting materials and synthesis conditions can be found in paper I
(LiFePO4), paper II (LiCoxFe1-xPO4) and paper III (Li2FeSiO4).
3.2 Electrode and battery fabrication
A cathode slurry was prepared by mixing the active material with Super-P
carbon powder and ethylene propylene diene terpolymer (EPDM) rubber
binder in an excess of solvent cyclohexane. The homogeneous slurry was
then spread thinly onto an Al-foil current collector. Circular discs with a
diameter of 2 cm were punched out and dried under reduced pressure at
120qC in an Ar(g)-filled glove-box (<5 ppm O2, <1 ppm H2O) for at least
6 h. Half-cells of configuration < electrode with active material _ glass wool
or Solipur” separator soaked in electrolyte _ Li-foil > were assembled and
packed in a flexible polymer-coated aluminium foil (a so-called “coffee-bag”
cell). The electrolytes used were either 1M LiTFSI salt in a 1:1 solution of
ethylene carbonate (EC) and propylene carbonate (PC) or 1M LiPF6 salt in a
2:1 solution of ethylene carbonate (EC) and diethyl carbonate (DEC).
A transparent polymer film with no aluminium content was used as packing material for in situ XRD measurements in transmission mode to minimize the amount of unwanted diffraction peaks.
21
The electrochemically cycled electrodes used for surface characterization
were dismantled in an Ar(g)-filled glove-box and thereafter transported in a
specially designed, sealed sample-holder to the instrument. In this way, the
electrodes were protected from exposure of air and moisture, which is
crucial since many of the surface species in the SEI layer react rapidly on
contact with O2, CO2 and H2O.
3.3 Electrochemical testing
Galvanostatic cycling (or cyclic chronopotentiometry), performed on a
Digatron BTS-600 battery tester, has been the most widely used electrochemical technique during the course of this work. A constant current (I) is
applied to the cell and the change in potential versus time (t) is monitored.
The gravimetric capacity (cap) is calculated from the total amount of charge
passed per unit mass of active electrode material (m) for a complete charge
(or discharge) according to Eq. 3.1:
cap = I · t / m
(3.1)
The cycling rates are commonly presented as C/t, where t is the time in
hours for a full charge or discharge. For example, C/25 means one charge (or
discharge) cycle in 25 hours.
Potentiostatic cycling (cyclic voltammetry) was used in paper IV to obtain
information about the complex electrode reactions in Li2FeSiO4 during the
first cycles. Here, the potential is varied linearly (when low current limits are
used) with time between two cut-off voltages, and the change in current is
recorded. Measurements were carried out on a Bio-Logic VMP2 cycling
system.
3.4 X-ray diffraction
When an X-ray beam interacts with the electrons of an atom, it is scattered in all directions. A constructive interference occurs in crystalline materials, where Bragg’s law is satisfied:
2dsinT = nO
(3.2)
where n is an integer, O is the wavelength, d is the distance between equivalent atomic planes and T is the angle between the incident beam and these
planes. In an X-ray diffraction (XRD) experiment, the scattered intensity is
measured as a function of scattering angle 2T. Analysis of the diffraction
pattern obtained provides information about the crystal structure e.g., bond
22
lengths and bond angles; it is also a powerful method to determine the phasecontent of a sample.
The Rietveld method [52,53] has been used to perform crystal structure
refinements of both X-ray and neutron powder diffraction data. A leastsquares refinement is carried out until the best fit is obtained between the
entire observed powder diffraction pattern and a calculated pattern. This
procedure involves the minimization of the expression:
¦ w i y i ,obs y i , calc 2
i
where yi,obs and are the observed and calculated intensities for the ith step,
respectively, and wi is the weighting assigned to an individual observed net
intensity yi,obs. Profile and structure parameters are varied to obtain an optimal fit. The calculated intensities, yi,calc, are determined by summing the
calculated contributions from overlapping Bragg reflections and adding the
background, yi,b:
y i , calc ( 24 i )
2
s ¦ L K FK ) 24 i 24 K PK A y i , b ( 24 i )
K
where s is the scale factor, K represents the Miller indices, (h k l), for a
Bragg reflection, LK contains the Lorentz, polarization and multiplicity factors, FK is the structure factor, ) is the reflection profile function, PK is the
preferred orientation function and A is an absorption factor.
In quantitative phase analysis, the weight percentage (Zp) of a phase p in
a mixture of m phases is given by:
m
Zp = np(ZMV)p / ¦ ni(ZMV)i
i 1
where Mi is the formula weight of the phase i involving Zi formula units/cell
of volume Vi [54]; ni is the refined scale factor for the ith phase. A quantitative phase analysis was performed in paper IV.
The quality of the refinement is given numerically by a number of agreement or reliability factors, such as the Bragg R-value:
R Bragg
¦k I k (obs) I k ( calc) / ¦k I k (obs)
where Ik(obs) and Ik(calc) are the observed and calculated integrated intensities for reflection k.
The Rietveld refinements in this work have been performed with the program FULLPROF [55].
23
XRD in direct combination with electrochemical cycling (in situ XRD) is
a very efficient way to study Li-ion insertion/extraction mechanisms in electrode materials. A schematic of the experimental set-up is shown in Fig 3.1.
Lithium battery
Positionsensitive
detector
Monochromator
X-ray tube
Potentiostat
Figure 3.1 A schematic illustration of the in situ XRD set-up.
Use of thin “coffee-bag” type cells (described in Section 3.2) makes it possible to perform powder diffraction measurements in transmission mode [56].
The advantage of this is that the measurement can be performed directly on
the cell in which the electrochemical cycling is taking place. Moreover, the
entire bulk of the electrode material is probed during the measurement, in
contrast to reflection-mode studies, where an often unknown depth below the
surface of an electrode is probed. A disadvantage, however, is that all
components in the cell contribute to the final diffractogram. To partly solve
this problem, the electrode of interest can be removed from the cell after
electrochemical cycling and repacked on its own in a “coffee-bag” type (cf.
Section 3.2), thereby avoiding the contribution from the lithium foil and the
separator.
Conventional X-ray powder diffraction measurements were made
in-house on an automatic STOE & Cie GmbH STADI powder diffractometer
equipped with a small straight-wire position-sensitive detector (PSD)
covering ~7q in 2ș. Exposures were made in transmission mode using Gemonochromatized Cu KĮ1 radiation (O = 1.54059 Å). In situ measurements
were conducted on beam line I711 (O = 1.2519 Å) at the Swedish National
Synchrotron Radiation Laboratory MAX in Lund.
3.5 Neutron diffraction
Neutron radiation for materials investigations is produced either in a
nuclear reactor or in a spallation source; this makes their availability quite
low. Neutron diffraction is therefore used mainly as a complementary technique to XRD. In a neutron diffraction experiment, the scattering process
24
takes place when neutrons interact with the atomic nuclei. As a consequence,
and in contrast to XRD, neutron scattering cross-sections are independent of
scattering angle and vary in an irrational way with mass number A. This can
make it possible to distinguish adjacent atoms in the periodic table (Fig. 3.2);
a fact which was used in paper II to study structural variations in the solid
solution series LiCoxFe1-xPO4.
Figure 3.2 Irregular variation of neutron scattering amplitude with atomic weight.
(Taken from Research, London, 7, 257, 1954).
The neutron powder diffraction measurements were performed at the
steady-state medium-flux research reactor R2 in Studsvik, Sweden. A monochromator system with two copper single-crystals (220-reflection) in parallel
alignment gave the wavelength O=1.470(1) Å. The samples were contained
in vanadium tubes during the data collection; vanadium being virtually transparent to neutrons. A detector bank of 35 independent 3He detectors was
used, and data were collected in the 2T-range 4-140o. Refinement of the data
was again performed using the Rietveld technique [52,53].
3.6 Photoelectron spectroscopy
The photoelectron spectroscopy (PES) measurements performed in the
context of this thesis exploited both synchrotron radiation (SR) and in-house
monochromatized AlKD radiation (1486.6 eV) (papers I and VI). Experiments using SR were performed on beam line I411 at the Swedish National
Synchrotron Radiation Laboratory MAX in Lund [57]. The in-house monochromatic measurements, also referred to as X-ray Photoelectron Spectroscopy (XPS), were carried out on a PHI 5500 spectrometer.
25
In PES, a sample under ultra-high vacuum is irradiated by photons with a
specific energy (hQ). If these photons have energies larger than the binding
energy (EB) of a certain state, photoelectrons will be ejected from the occupied energy levels. The measured kinetic energy (Ek) of these ejected electrons are used to calculate EB according to the expression EB = hQ - Ek – Ȃ;
where Ɏ is the potential difference between the Fermi level of the sample
and the vacuum level of the spectrometer (see Fig. 3.3).
Figure 3.3 The principal of photoelectron spectroscopy (PES).
The intrinsic surface sensitivity of PES is a result of the shallow escape
depth of the ejected electrons. Although the impinging photons can penetrate
a few micrometers into the sample, the photoelectrons can only travel 50100 Å before losing their energy. Only photoelectrons from the outermost
surface will therefore reach the detector and contribute to the photoemission
spectrum.
Depth-profiling has been performed by varying the incoming excitation
photon energy (using SR) for a given core level. This will result in different
kinetic energies of the emitted electrons, which means that they have been
ejected from different depths in the sample (paper I). This type of depth
profiling has the advantage of being non-destructive and more surfacesensitive compared to the Ar+-ion sputtering technique, which is normally
used in combination with monochromatized XPS analysis.
3.7 Mössbauer spectroscopy
There are about 50 known Mössbauer isotopes; the most widely used is
Fe representing more than 90% of all publications. The technique is based
on the recoil-free emission and resonant absorption of nuclear J-rays in
solids, referred to as the Mössbauer effect after Rudolph Mössbauer who was
awarded the Nobel Prize in 1961 for his discovery. The experimental set-up
for a Mössbauer experiment can be seen in Fig. 3.4. The source is a radioactive isotope (57CoRh was used in paper IV) which is moved back and fore
with a constant acceleration. In this study, the absorber was an 57Fe-
57
26
containing electrode in a Li-ion battery, which was held at different states of
charge or discharge in the cycling process.
Figure 3.4 Experimental set-up for the Mössbauer measurement.
From an analysis of the spectrum obtained, the isomer shift and quadrupole splitting parameters of different assigned doublets can provide definitive information about the redox states of Fe. The relative amounts of each
of these states can be estimated from a comparison of the relative intensities
of the doublets in a spectrum. To eliminate the effect of texture in a sample,
measurements can be performed at the magic angle (54.7q) [58].
3.8 Scanning electron microscopy
Surface studies of a solid sample can be made by Scanning Electron
Microscopy (SEM). The surface is scanned with a focused beam of electrons
of a certain energy. By detecting the electrons emitted, an image of the surface can be created. A Leo 1550 HR-SEM has been used to produce micrographs of powder samples and electrode surfaces studied in this thesis.
3.9 Theoretical calculations
Quantum mechanical calculations can be used to compute the properties
of molecules and their interactions with one another. The ground-state geometrical structure and electronic energies of a molecule can, for example, be
calculated theoretically for a given set of nuclei and electron parameters. All
relevant information describing the interaction of nuclei and electrons is
contained within the Schrödinger equation. However, since an exact solution
of this equation is not possible for a many-electron system, certain approximations have to be made. In Density Functional Theory (DFT), the calculated total energy of a system is directly correlated to its electron density.
Starting with an initial set of trial parameters, and performing the calculations iteratively until convergence is obtained, gives the totally relaxed state
of a system. The achieved minimum value of the total-energy functional is
the ground-state energy of the system, and the electron density that yields
27
this minimum value is the exact single-particle ground-state electron density.
It should be noted, however, that all DFT calculations are performed at 0 K.
Ab initio level calculations are reported in paper V; these use DFT with a
plane-wave basis set, as implemented in the VASP software package [59],
using the generalized gradient approximation (GGA) of Perdew and Wang
[60]. Using structural data taken from papers III and IV, a study was made
of Li2FeSiO4 and the three possible local lithium configurations within a
single crystallographic unit cell for the delithiated phase LiFeSiO4. Parameters from paper I were used to obtain the crystal structure, atomic fractional
coordinates and crystallographic cell shape for the relaxed state of
Li2FeSiO4. The three starting crystal structures for LiFeSiO4 were obtained
by removing two Li ions from the optimized Li2FeSiO4 unit cell before
letting the system relax. From these relaxation processes, the final total energies and the Density of States (DOS) were derived for the two materials.
A number of electrochemical properties were derived directly from the
difference in DFT total energies – a procedure offering improved accuracy
through the cancellation of errors. Typically, the cathode discharge process
assumes the half-cell reaction vs. Li/Li+:
LiFeSiO 4 x Li
xe
o Li 1 x FeSiO 4
The open circuit voltage (OCV) for an intercalation reaction involving xLi+
can then be calculated from the energy difference if volume and entropy
effects are neglected:
OCV |
E( LiFeSiO 4 ) x E( Li ) E( Li 1 x FeSiO 4 )
x
In this work, all OCVs were calculated for the case of x=1.
3.10 Other techniques
The elemental analysis technique used in paper II was Inductively Coupled
Plasma - Atomic Emission Spectroscopy (ICP-AES). The interested reader is
directed, for example, to reference [61]. Elemental analysis was made for
carbon using a combustion method. This technique is described in paper IV.
28
4. LiFePO4 as cathode material
LiFePO4 is today a rather well described cathode material, which is close
to being implemented in commercial applications. Its initial drawback of low
electronic conductivity has now been overcome by various synthetic and
processing strategies, as described in Chapter 2. One approach has been to
attach amorphous carbon to the LiFePO4 particles during the synthesis
process [39]. High capacity retention can be achieved in this way for
extended electrochemical cycling (140 mAhg-1 at C/2 rate for more than 700
cycles) [62].
4.1 Surface chemistry of carbon-treated LiFePO4
In the study reported in paper I, a polypropylene powder was added in the
synthesis to enhance the electrochemical performance. This carbontreatment of the particles resulted in a carbon content of 0.56 wt% and an
electronic conductivity of about 2u10-5 Scm-1.
The capacity obtained at room temperature for this carbon-treated
LiFePO4 was higher than that for untreated LiFePO4 reported earlier [43].
On raising the cycling temperature to 40qC, the capacity increased and the
irreversible capacity decreased compared to the electrochemical measurements conducted at room temperature. This might suggest that surface-film
formation at higher temperatures does not limit electrochemical performance
as it does for other cathode materials such as LiNi0.8Co0.2O2 or LiMn2O4
[26,63].
Surface characterization of pristine carbon-treated LiFePO4 electrodes
showed no evidence for the presence of Li2CO3. This is contrary to results
obtained for other cathode materials. The absence of Li2CO3 could have a
bearing on long-term cycling and electrode impedance, since it has been
suggested that initially present Li2CO3 could react with the electrolyte in an
assembled cell and contribute to surface-film formation [25,26].
Reference measurements on a carbon-treated LiFePO4 electrode exposed
to electrolyte (1 M LiPF6 in EC/DEC) were performed to determine the
binding energies of the salt species present on the surface. It could be
noticed that no trace of the electrolyte solvent was detected.
29
PES spectra of C, F, O and P were measured on carbon-treated LiFePO4
electrodes cycled at 23 and 40qC, using photon energies corresponding to a
kinetic energy of 500 eV for the detected electrons, which facilitated detection of all elements at the same depth (Fig. 4.1). For both electrodes, the only
C1s peak corresponds to carbon black, the EPDM binder and the carbon on
the LiFePO4 particles. The absence of solvent reaction products (like polycarbonates) on the carbon-treated LiFePO4 electrode is again in direct contrast with the results of surface studies conducted on LiNiO2, LiNi0.8Co0.2O2
and LiMn2O4 [25,26,63,64], and suggests that the phosphate group in
LiFePO4 does not participate in solvent reactions as do the oxides. The salt
compounds detected (LixPFy, LixPOxFy and LiF) are in agreement with
previous results for other cathode materials. Larger quantities of salt-based
species and more of the oxygenated species were detected on the surface of
the electrode cycled at 40qC. Peaks from the phosphate group are clearly
seen in both the O1s and P2p spectra, indicating that the thickness of the
salt-based surface compounds is less than the penetration depth (~50 Å) or
that the surface compounds are not evenly distributed.
Depth-profiling using SR-excited PES was achieved by varying the
photon energy. The only C1s peak detected at all photon energies (i.e., at
different penetration depths) originates from carbon black, the EPDM binder
and the carbon on the LiFePO4 particles, which clearly shows that no solvent
reaction products are present at the electrode surface. The detection of peaks
from the phosphate group in the O1s and P2p spectra indicates that the saltbased surface film does not completely cover the electrode surface. The
surface layer formed cannot therefore be described as an “SEI layer”, but
rather as a partial coating of a salt-based film.
The results from the PES study suggest that the formation of a surface
film does not limit the cell performance of carbon-treated LiFePO4 as it does
other cathode materials. It can also be concluded that electrochemical
cycling of the material at higher temperatures enhances the capacity and
reduces the irreversible capacity compared to cycling at 23qC, which again
confirms earlier results for untreated LiFePO4 [43]. However, these results
are based on “half-cell” measurements, i.e., where the cell-anode is metallic
lithium. In a recent study of high-temperature stored and cycled carboncoated <LiFePO4_LiPF6-electrolyte_graphite> cells, a significant capacity
fade was observed for cells cycled at 37qC and 55qC [65]. The carbon-coated
LiFePO4 electrode was found to release Fe ions into the electrolyte on ageing
at these temperatures, and the presence of Fe could be confirmed on the
graphite electrode. Replacing LiPF6 by LiB(C2O4)2 (LiBOB) significantly
improved cycling stability. Further surface analyses of LiFePO4 electrodes
are clearly needed.
30
Figure 4.1 C1s, F1s, O1s and P2p PES spectra for carbon-treated LiFePO4 electrodes cycled at 23 qC (a-d) and at 40 qC (e-h) recorded at photon energies of 790,
1194, 1040 and 643 eV, respectively. This results in a photoelectron kinetic energy
of 500 eV. F1s, O1s and P2p spectra are shown together with fitted peaks to clarify
peak assignments. Pairs of Gaussian peaks were used to fit the P2p doublets (P21/2,
P23/2) to a spin-orbit splitting of 0.8 eV and an area ratio of 1:2.
31
4.2 Transition-metal substitution in LiFePO4
As mentioned earlier, LiFePO4 suffers from low electronic conductivity,
which gives poor rate-performance and limits its use in high power-density
applications. It has been shown both experimentally and by theoretical
calculations that a possible means of increasing the intrinsic electronic conductivity in LiFePO4 can be the partial replacement of Fe by some other
transition metal, such as Ni or Co [44,45]. In earlier studies, the preparation
of the solid solution series LiCoxFe1-xPO4 was reported and the electrochemical performance evaluated. It was shown that an increase in the
capacity with respect to LiFePO4 and better capacity retention were achieved
for low levels of Co substitution [46,47]. From a structural point of view, it
is generally assumed that the Co substitutes into the Fe-site in this solid solution series, but this has never been confirmed. To do so, we need neutron
diffraction, since X-ray diffraction cannot distinguish adjacent transition
metals in the periodic table.
In this study, LiCoxFe1-xPO4 samples were prepared for x = 0. 0.25, 0.40,
0.60 and 0.75 and their structures investigated by both XRD and neutron
diffraction. All samples were confirmed to be phase-pure, with the same
orthorhombic space-group as LiFePO4: Pnma. The a and b lattice parameters
decrease, while c increases linearly with the amount of Co2+ substituted,
giving an overall decrease in the unit-cell volume.
From Rietveld refinements of the neutron diffraction data (c.f., Section
3.4), it was indeed confirmed that Co substitutes for Fe in the 4c-site for all
compositions. A typical Rietveld fit for the neutron powder diffraction data
for one of the compositions (LiCo0.6Fe0.4 PO4) is shown in Fig. 4.2.
The mean M–O (M=Co, Fe) distance in the LiCo0.75Fe0.25PO4 composition
is, on average, slightly shorter than in the other compositions, implying
somewhat stronger M–O bonding for x=0.75. This is supported by the XPS
results of Wang et al., where the highest O1s binding energy (suggesting
shorter Fe/Co–O distances) was observed for the similar composition
LiCo0.80Fe0.20PO4 [47]. A distortion of the phosphate tetrahedra is also
observed for the LiCo0.75Fe0.25PO4 composition, where one P–O distance is
considerably shorter then the rest. Wang et al. also observe the highest P2p½
binding energy in their XPS measurements for the LiCo0.80Fe0.20PO4
composition. These observations taken together would suggest that a certain
amount of Co (x0.75) can be accommodated at the Fe-site without
distorting the structure but, beyond this Co-content, the structure tends to
become less stable. This type of feature is reminiscent of a feature recently
discussed by Yamada et al. [35], where they propose the notion of two stable
LixFePO4 phases (a Li-rich and a Li-poor phase) rather than the commonly
assumed x=1 and x=0 phases LiFePO4 and FePO4. It could well be that the
relevant phase diagram for the LiCoxFe1-xPO4 system also involves this type
of subtlety.
32
Figure 4.2 Observed, calculated and difference neutron powder diffraction profile
for LiCo0.60Fe0.40PO4 (O=1.47 Å). The row of short vertical lines marks the calculated positions of the reflections.
33
5. Li2FeSiO4 as cathode material
5.1 Synthesis and phase purity
Li2FeSiO4 was prepared using a solid-state synthesis technique in which
Li2SiO3 and FeC2O4˜2H2O were mixed and ground together thoroughly with
10 wt % of a carbon precursor in the form of a carbon gel, formed by the
polymerisation of resorcinol-formaldehyde [66,67] (paper III) or polyethylene-poly(ethylene glycol) (paper IV). The carbon precursor was added to
form a “carbon nano-coating” on the Li2FeSiO4 particles in order to increase
the electronic contact between the electrochemically active particles. In the
following heat treatment, the annealing temperature of 700qC was kept for
20 h. To suppress the formation of Fe3+ compounds, the heat treatment was
conducted in a flow of CO/CO2 gas (50:50). These conditions were critical;
lowering the annealing temperature below 650qC or exchanging the CO/CO2
gas with an inert gas (N2) resulted in traces of unreacted starting material in
the final product. However, the key factor to obtain an optimal product from
an electrochemical activity point of view was the mixing and grinding
procedure of the starting materials prior to the heat treatment. Several weeks
of mixing/grinding were needed.
Even then, the as synthesised powder was not phase pure. The presence of
wüstite, Fe2+1-xO and LiFe3+Si2O6 was confirmed both from Mössbauer
spectroscopy and X-ray diffraction measurements. From multi-phase
Rietveld refinements of XRD data (c.f., Section 3.4), the amount of Fe1-xO
and LiFeSiO2O6 was determined to be 5(1) and 10(1) wt%, respectively. The
Fe1-xO composition was established to be Fe0.91O from extrapolation of the
variation of lattice constant versus wt% Fe in Fe1-xO, as obtained by Jette et
al. [68].
The carbon content in the as synthesised powder was determined to be
4.68(5) wt%. From analysis of SEM pictures, the particle size was found to
be in the range 50-500 nm, with an average value of 150-200 nm. The small
particle sizes are most certainly a result of the extended mixing/grinding
time.
34
5.2 Electrochemical performance
Electrochemical cycling of Li2FeSiO4 at 60qC in galvanostatic mode
shows a flat voltage plateau at 3.10 V in the first charge-cycle (Fig. 5.1). On
the following discharge cycle, the observed shift in the potential plateau to
2.80 V indicates some rearrangement of the structure. The potential plateaus
under further charge and discharge cycles are flat and stable at 2.76 V and
2.80 V, respectively, suggesting little or no subsequent change in structure.
This suggests that some type of electrochemical annealing process is
occurring during the first cycle, which could involve a phase transition to a
more stable structure.
Roughly 75% of the extractable Li+ ions, corresponding to a specific capacity of 125 mAh/g (at a cycling rate of C/25), could be withdrawn on the
first charge. In the following cycles, no irreversible capacity is observed and
the overall capacity loss over 120 cycles is only ~3%; in fact, a capacity gain
of ~6% occurs during the first cycles.
Figure 5.1 The galvanostatic voltage curve for Li2FeSiO4.
An electrochemical study using cyclic voltammetry (Fig. 5.2) of this same
material resulted in the appearance of two prominent features not seen earlier: firstly, the oxidation peak appears at higher voltages and the reduction
peak at lower voltages, which is explained by the much faster
charge/discharge rate used here (sweep-rate: 72 mV/h) compared to the
galvanostatic experiment. Since a higher current is passing through the
battery, the polarization increases. More importantly, the structural changes
occurring upon cycling are clearly more complex than as suggested by the
galvanostatic cycling. The shape of the cyclic voltammogram represents a
classic ECE reaction, where E represents an electron transfer at the electrode
35
surface, and C represents a homogeneous chemical (“structural”) reaction in
the bulk or at the surface. The first oxidation peak (Q at 3.19 V in Fig. 5.2)
corresponds to the first electron transfer in the ECE reaction, in which Fe2+ is
oxidised to Fe3+. This is followed by a “chemical” step, giving rise to the
first reduction peak (R) at 2.96 V. The more precise nature of this step is
difficult to define on the basis of electrochemical data alone, and will therefore be addressed later on the basis of structural data. The second reduction
peak (R' at 2.72 V) represents the second electron transfer process (where
Fe3+ is reduced to Fe2+). Only a slight remnant is seen of the reduction peak
at 2.96 V on the second discharge, whereas the second reduction peak (T) is
clearly dominant and has shifted to 2.68 V. This peak continues to shift
towards lower potential on the subsequent ten discharges, ending up at 2.66
V. The same behaviour, with a shift of the oxidation peak (Q at 3.19 V)
towards lower voltage (S at 2.93 V), is seen during the charge cycles, again
implying some type of structural rearrangement which stabilizes over the
first 10 cycles, with the main rearrangement actually occurring during the
first 2 cycles.
Figure 5.2 Cyclic voltammogram of the first four cycles for Li2FeSiO4; the letters
mark the states at which Mössbauer (P-U) and X-ray diffraction measurements
( P-T) were made; sweep-rate: 72 mV/h.
36
5.3 Structure
The crystal system for Li2FeSiO4 was confirmed to be orthorhombic, with
space-group: Pmn21, and refined cell parameters: a=6.2667(5),
b=5.3296(5), c = 5.0147(4) Å, V=167.49(1) Å3; Z=2. This cell agrees well
with the very early study of Tarte and Cahay [51], where Li2FeSiO4 was
proposed to be isostructural with Li3PO4. The crystal structure was refined
routinely by the Rietveld method (c.f., Section 3.4) using atomic startcoordinates taken from the low temperature Li3PO4 structure [69]. The Li, Fe
and Si atoms were all confirmed to be tetrahedrally coordinated to oxygen
atoms (Fig. 5.3).
DFT calculations of Li2FeSiO4 and the three possible local lithium configurations in the assumed same unit cell for the delithiated phase LiFeSiO4
(denoted by A, B and C in Fig. 5.4) were performed in an attempt to gain
further insights into the complicated structural rearrangement occurring
during the first electrochemical cycle. From these calculations, the optimized
structures and their lattice parameters were obtained, which were compared
with the experimental values given in papers III and IV (Table 5.1).
The final cell volume for Li2FeSiO4 was found to be ca. 1% larger than
the experimental value, which is consistent with the tendency to underestimate bond strengths using GGA functionals. The optimized cell volume
increased by 3% on going from Li2FeSiO4 to the lowest-energy configuration of LiFeSiO4; this compares to 1.2% from experiment. The calculated
average voltage vs. Li/Li+ was 2.66 V, 2.79 V and 2.88 V for the three
different Li configurations, with the lowest voltage corresponding to the
lowest-energy structure.
Table 5.1 Optimized lattice parameters. Values in brackets are experimental
values; standard deviations given in parentheses. For LiFeSiO4, small deviations (<0.02Å) from perfect orthogonality have been neglected.
Structure
Li2FeSiO4
LiFeSiO4 (A)
LiFeSiO4 (B)
LiFeSiO4 (C)
a /Å
6.313
[6.266(1)
b /Å
5.393
5.329(1)
c /Å
4.979
5.015(1)
V /Å3
169.5
167.5(1)]
6.11
6.12
6.08
[6.507(3)
5.56
5.57
5.64
5.213(2)
5.03
5.07
5.07
5.001(2)
170.9
172.8
173.9
169.6(1)]
37
Figure 5.3 The crystal structure of Li2FeSiO4. Four Li ions forming a plane have
been highlighted in the structure.
Figure 5.4 Three different ways (denoted by A, B and C) of removing two lithiums
out of four from the optimized Li2FeSiO4 unit cell to form LiFeSiO4. The common
plane of the Li ions is the same as in Figure 6.1.
We can note that the experimentally determined voltage vs. Li/Li+ on the
second and subsequent charge and discharge cycles is 2.80 V and 2.76 V,
respectively (Fig. 5.1). It could be that the fully-optimized theoretical
structure actually corresponds to the stable structure obtained experimentally. The higher-energy theoretical structures could therefore correspond to
38
some intermediate metastable phase giving the higher experimentally
observed voltage on the first cycle. Alternatively, the experimental value of
2.80 V vs. Li/Li+ observed on the second charge-cycles could, of course,
derive from a mixture of possible LiFeSiO4 phases, while the higher firstcycle experimental voltage would then correspond to a higher-energy
metastable structure being formed in the synthesis. There is also
experimental evidence (paper IV) to suggest a decrease in the short-rangeorder correlation length on cycling, corresponding to Li- and Fe-site
occupations progressively becoming more randomized. This would clearly
influence the resulting voltage on successive cycling – as is indeed observed
experimentally.
The DOS calculations also confirm that Li2FeSiO4 is a semiconductor; the
band gap of 0.15 eV being small enough to allow some conduction at room
temperature. One should bear in mind, however, that band-gaps are often
underestimated in DFT calculations.The DOS for LiFeSiO4 corresponds to
an insulator with a band-gap of 1.1 eV. Unlike Li2FeSiO4, this is probably
too large to allow conduction at room temperature, which would clearly
imply poor battery cyclability. This is not especially apparent, however,
from the successful experimental cycling observed, although much of the
experimental work has been performed at 60oC to improve cyclability and
material utilization, while DFT calculations are performed at 0K!
5.4 Redox mechanism
A special effort was made to investigate the complex structural changes
and the lithium extraction/insertion mechanism occurring on cycling the
material. In this study, a combination of electrochemical cycling, synchrotron in situ diffraction and Mössbauer spectroscopy was used. The points at
which X-ray and Mössbauer experiments were performed are indicated in
the cycling curve shown in Fig. 5.2.
Multiphase Rietveld refinements (c.f., Section 3.4) were performed on the
XRD data (Fig. 5.5) recorded at the different charge and discharge points
(Q, R, S and T in Fig. 5.2) to quantify the amounts of the various phases
present [54]. The relative amounts of crystalline Li2FeSiO4 and LiFeSiO4
deduced from these refinements are presented in Table 5.2.
Table 5.2 The relative amounts (in %) of Li2FeSiO4 and LiFeSiO4
obtained from multiphase Rietveld refinements of the XRD data.
Cycle #
0
½
1
1½
2
Li2FeSiO4
LiFeSiO4
100
0
5.2(6)
94.8(6)
90.4(3)
9.6(3)
3.6(5)
96.4(5)
91.5(4)
8.5(4)
39
The volume of the unit-cell for the delithiated form (LiFeSiO4) is slightly
(~1%) larger than for Li2FeSiO4; this difference is significantly less than that
occurring in the LiFePO4-FePO4 system (~7%) [43], not to mention the
much larger volume changes occurring in intermetallic anode materials, e.g.,
42% for the transformation of Cu2Sb to Li3Sb [22]. This small volume
change for Li2FeSiO4 indicates that the mechanical stress created in the
material is rather low during electrochemical cycling. This can contribute, at
least in part, to the excellent observed long-term cyclability (cf., Ch. 5.2).
Figure 5.5 X-ray diffractograms taken at fully charged and fully discharged states
during the first two charge/discharge cycles between 2.0 and 3.7 V: P) 0 cycles; Q)
½ cycle; R) 1 cycle; S) 1½ cycle; T) 2 cycles. Peaks corresponding to Li2FeSiO4 and
LiFeSiO4 are marked with their indices. Peaks marked with an asterisk (*) originate
from the “coffee-bag” polymer.
A self-consistent picture begins to emerge in which Li2FeSiO4 is the
dominating phase in the reduced state, while LiFeSiO4 dominates in the oxidised state during electrochemical cycling. This is supported both by XRD
and Mössbauer measurements. The XRD results indicate around 95%
delithiation/lithiation of the material, while the Mössbauer results suggest
90% Fe2+/Fe3+ conversion during charge/discharge. The electrochemical
measurements, however, suggest only 75% utilisation of the active material.
40
The existence of two impurity phases (LiFeSi2O6 and Fe0.91O; confirmed by
XRD) could partly explain the lower utilisation achieved in the electrochemical measurements; the mass of the electrode was not corrected for
these impurities. If a correction is made, assuming that none of the
impurities contribute to the total capacity and using the total amount of impurities derived from multiphase refinement of the diffraction data (15 wt%),
a value of 89% utilisation is obtained, which is now in excellent agreement
with XRD and Mössbauer results.
The observed shift of the potential plateau from 3.10 to 2.80 V clearly
suggests some structural phase transition in Li2FeSiO4. From electrochemical measurements conducted at moderate rates (>C/10), it was found that this
structural rearrangement is rather slow, evolving over the first 2-3 cycles. A
clear indication of the nature of this structural rearrangement is seen from
the Rietveld refinements of the in situ XRD measurements (batteries cycled
at the low rate of C/25) in terms of Li and Fe ion-site exchange. The most
pronounced change in the Li:Fe ratio at the 4b-site (from 96:4 to 40:60) is
observed during the first charge-cycle. If this ion rearrangement corresponds
to structural stabilization, it would cause a lowering of the potential plateau,
which is indeed what is observed; on subsequent cycling, the Li:Fe ratio at
the 4b-site shifts between 60:40 in the lithiated state and 40:60 in the
delithiated state. Since the surroundings of the Fe ions, whether in 4b- or 2asites, are rather similar, it is hard to draw any more definitive conclusions on
the basis of the supporting Mössbauer results.
It should be noted in this context that an X-ray powder diffraction study
has recently been reported by Dominko et al. for the related material
Li2MnSiO4 [70], where partial occupation was observed in alternate
tetrahedral sites by Li and Mn ions. Inter-migration of Li and Fe/Mn ions is
clearly a critical factor in determining the Li ion extraction/insertion
mechanism in both materials. More precise structural data is needed,
however, preferably from single-crystals, to better resolve these issues.
5.5 Surface chemistry
Surface reactivity of pristine electrodes in contact with air, and surfacefilm formation after electrochemical cycling at 60qC using a LiN(SO2CF3)2based electrolyte were investigated in paper VI. The PES spectra for pristine
carbon-coated Li2FeSiO4 electrodes are shown in Fig. 5.6. Pristine, uncycled,
air-exposed electrodes showed a significantly greater amount of Li2CO3 on
their surfaces compared to electrodes stored under inert atmosphere. This is
in agreement with the observation made on the surface of pristine electrodes
for transition-metal oxides like LiMn2O4, LiNiO2 and LiNi1-x-yCoxAlyO2
[25,26,71], but contrary to the observation for LiFePO4 (paper I). The formation of carbonates on the Li2FeSiO4 surface on exposure to air probably
41
occurs through the reaction of atmospheric oxygen and CO2 with the lithium
in Li2FeSiO4. Since the carbon-black peak at 284.4 eV is not significantly
weaker in the C1s spectra for the air-exposed electrode, this would suggest
that the coverage of the carbonate-based film is only partial.
Figure 5.6 PES spectra recorded for pristine unexposed (left column) and airexposed (right column) carbon-coated Li2FeSiO4 electrodes.
The electrochemical cycling behaviour for an unexposed and an airexposed sample is shown in Fig 5.7. The low specific capacity for the airexposed (13 mAh/g) compared to the unexposed sample (113 mAh/g) on the
first charge-cycle suggests that lithium has been withdrawn from the
structure on exposure to air. However, the capacity on the first dischargecycle is the same for the two samples. From complementary electrochemical
measurements (not shown) on cells with a graphite anode (instead of
lithium), it was confirmed that the lithium withdrawn on exposure to air does
not participate in the subsequent intercalation/deintercalation process.
42
(a)
(b)
Figure 5.7 First-cycle performance of: (a) an unexposed, and (b) air-exposed sample
of carbon-coated Li2FeSiO4 at 60 qC; cycle rate C/25 in 1M LiTFSI EC:PC (1:1).
PES spectra were obtained for Li2FeSiO4 electrodes electrochemically
cycled at 60qC. Only small differences were observed compared to reference
measurements of an uncycled electrode stored in electrolyte. No degradation
products of the LiTFSI salt (such as LiF) were detected on the surface of the
two electrodes, which indicates that the salt is stable under cycling. Observations of additional peaks between 285 and 289 eV in the C1s spectra, as
compared to the pristine electrode, were therefore attributed to solvent
reaction products. Surface studies of LiMn2O4, LiNiO2 and LiNi0.8Co0.2O2
show that the solvent reaction products have been mainly hydrocarbons,
polycarbonates and lithium alkyl carbonates [25,26]. However, no peak is
detected at 290-291 eV in our spectra, which implies a total absence of
carbonate-based compounds. This also leads to the conclusion that the
Li2CO3 initially present had already disappeared from the surface in the
electrolyte-exposed electrode before electrochemical cycling.
In a surface study by Zhuang and Ross [72] of the SEI layer formed on
aged graphite electrodes, the dominant species were characterized to be
lithium succinate (LiO2CCH2CH2CO2Li), lithium oxalate (Li2C2O4) and
lithium methoxide (LiOCH3). They suggest lithium succinate as a model
compound for lithium carboxylates such as Li-formate, -acetate and
-propionate. Such compounds could explain the presence of peaks at 286 and
288-289 eV in our C1s spectra. These types of surface species should accumulate on cycling at 60qC.
The major part of the (albeit very thin) surface film formed would seem to
originate from the salt in its original state, although more solvent reaction
products are present after cycling compared to the reference electrode.
However, the C1s peak from the carbon black, EPDM and the carbon
coating (284.4 eV) is still the strongest for both electrodes, thus supporting
the conclusion also drawn for LiFePO4 that the surface film formed does not
completely cover the electrode surface. The occurrence of solvent reaction
43
products at the surface is in direct contrast with the situation found in
LiFePO4 (paper I). The different surface-layer components present on the
electrode surface of LiFePO4 and Li2FeSiO4, respectively, are illustrated
schematically in Fig. 5.8.
(a)
(b)
Figure 5.8 Schematic models of the surface-layer components on the electrode
surfaces of: (a) LiFePO4 and (b) Li2FeSiO4.
44
6. Conclusions and future work
This work has provided new insights into the surface chemistry of carbontreated LiFePO4, along with a deeper understanding of structural features in
the doped solid solution system LiCoxFe1-xPO4, in the context of their rôle as
cathode materials. Its major contribution, however, has been the introduction
and investigation of Li2FeSiO4 as a new cathode material for Li-ion batteries.
Studies have been made of structural and electrochemical processes occurring both at the surface and in the bulk of this material. A summary of the
more important conclusions drawn from these studies, together with some
suggestions for future work, are presented below.
Surface analysis of carbon-treated LiFePO4 has revealed that only saltbased reaction products are present on the electrode after electrochemical
cycling (using lithium metal as anode) at temperatures up to 40qC; no products of any solvent reactions are detected. This is contrary to results obtained
for other cathode materials, and is an indication that the oxygen in the phosphate group does not participate in any surface reactions, as does the oxygen
in today’s transition-metal oxide based cathode materials. The absence of
Li2CO3 on pristine electrodes is also contrary to results obtained for these
currently used cathode materials, and could possibly explain the different
surface chemistry of LiFePO4. The observed stability of carbon-treated
LiFePO4 results in most promising electrochemical performance, especially
at higher temperatures.
However, subsequent high-temperature cycling performance tests in cells
with the general configuration <LiFePO4/C_LiPF6-electrolyte_graphite> have
shown a rather dramatic capacity loss compared to cycling performance at
room temperature [57]. It is found that replacing the LiPF6-salt by
LiB(C2O4)2 (“LiBOB”) can partly solve this problem. In the light of these
observations, some future studies can be suggested:
-
-
Investigation of the surface film formed on both the cathode and the
anode after cycling an <LiFePO4/C_LiPF6-based electrolyte_graphite>
“whole-cell” at elevated temperatures, especially to probe the phenomenon of metal-ion (here, Fe) cross-over from cathode to anode.
A comparative study in which LiB(C2O4)2 is used as salt instead of
LiPF6.
45
Our X-ray and neutron powder diffraction investigations of the solid solution series LiCoxFe1-xPO4 for x=0.0, 0.25, 0.40, 0.60 and 0.75 confirm that
Co substitutes for Fe in the octahedral 4c-site for all compositions. The a and
b lattice parameters decrease linearly with increasing Co2+ substitution,
while the c parameter increases, resulting an overall decrease in the unit-cell
volume. The observation of a slightly shorter mean MO (M=Fe,Co)
distance and a distortion of the phosphate tetrahedra for the LiCo0.75Fe0.25PO4
composition could suggest that a certain degree of phase instability arises for
high Co-content. Based on these results, a suggestion for some future work
could be:
-
A careful study of the complete phase diagram for the solid solution
series LiyCoxFe1-xPO4 to confirm (or otherwise) this proposed phase
instability at high levels of Co substitution, and to ascertain the optimal composition both from an electrochemical and electronic conductivity viewpoint.
Investigations of the electrochemical performance of Li2FeSiO4 have
confirmed that the material cycles between Li2FeSiO4 and LiFeSiO4.
Excellent capacity retention (< 3% capacity loss over 120 cycles) and minimal irreversible capacity on the first cycle has been achieved for cells cycled
at 60qC with an LiTFSI-based electrolyte. The potential plateau observed at
3.10 V on the first charge-cycle is lowered to 2.80 V on subsequent chargecycles. This observation is suggested to involve some type of structural
rearrangement in Li2FeSiO4, with a transformation to a more stable structure.
A proposal for the nature of this structural rearrangement is given in terms of
Li and Fe ion-site exchange.
DFT calculations for Li2FeSiO4 and its delithiated form LiFeSiO4 show
rather good agreement with experiment for the relaxed atomic structures and
cell parameters derived. Three possible local Li arrangements within a single
crystallographic unit cell were proposed for LiFeSiO4. Their total energies
and related open circuit voltages (OCVs) suggest that the remaining Li ions
are randomly distributed; which is indeed supported experimentally.
Investigations of the surface chemistry of Li2FeSiO4 show that Li2CO3 is
formed on the surface of pristine electrodes exposed to air. This is contrary
to the case of LiFePO4 electrodes. Furthermore, the surface-film formed on
electrochemical cycling of Li2FeSiO4 electrodes at 60qC using a
LiN(SO2CF3)2-based electrolyte suggests high salt stability and only very
small amounts of solvent reaction products. These are mainly of
Li-carboxylate type; neither carbonates nor LiF are found. This excellent
capacity retention and minimal irreversible capacity during the first cycle is
probably a direct result of this very thin surface film.
These results clearly suggest LiFeSiO4 to be a highly promising positiveelectrode material for large-scale Li-ion battery applications. To further
46
improve and develop this material, a number of important issues should be
focused upon:
-
It is vital to obtain phase-pure Li2FeSiO4. The variable parameters in
the solid-state synthesis process should therefore be optimized
systematically. Other synthesis techniques, such as sol-gel, hydrothermal and precipitation methods, should also be probed in efforts to
achieve phase-purity.
-
Reducing particle size and optimizing its distribution will certainly
improve electrochemical performance. Sol-gel or precipitation
methods are certainly two possible synthesis techniques to achieve
this.
-
Alternative carbon-coating techniques should also be tested and
optimized to improve contact between the active particles in the
electrode.
-
The complex structural rearrangement on electrochemical cycling
should be explored further – preferably using single-crystal XRD
techniques.
-
Substitution of other transition metals into the Li2FeSiO4-system
should be studied further. The most obvious choice would be manganese (Mn) by virtue of its 2+ and 4+ oxidation states; the system
Li2(Fe1-xMnx)SiO4 thus holds the promise of a “>1-electron reaction”
of type:
Li+2(Fe2+1-xMn2+x) SiO4 Ÿ Li+1-x(Fe3+1-xMn4+x)SiO4 + (1+x)Li+ + (1+x)e-
Substitution of only 20% of the Fe sites by Mn (x=0.20) would result
in a theoretical capacity well in excess of 200 mAh/g.
47
Acknowledgements
First of all, I would like to thank my supervisor Professor Josh Thomas.
Your never-ending support, generosity and deep understanding of Science
have been invaluable for me to accomplish this thesis and to gain a broader
perspective on the World of (and around) Chemistry. I am deeply grateful
for the opportunity you have given me to develop as a scientific researcher.
Thank you Torbjörn Gustafsson for your support, all scientific
discussions and help with technical stuff (especially concerning the PSD!).
I send a special “thank you” to Professor Michel Armand, who took such
good care of me during my visits to Montréal, Canada, and who introduced
me to the iron-silicates. Without you sharing your ideas and knowledge with
me, this thesis would have had a much less exiting content.
Saeed Kamali and Lennart Häggström are acknowledged for all their help
with Mössbauer measurements, Håkan Rensmo for our many fruitful
discussions regarding those photoelectrons and Peter Larsson and Rajeev
Ahuja for the skill you display in performing the calculations.
Thanks to all the people who have made the Department of Materials
Chemistry such a nice and friendly place to work at during all these years. A
special “thanks” to Daniel (whom I have “followed” for more than 30 years),
my room-mate Hanna for all scientific discussions and laughs, Igor for
being a good friend (and never saying no to a party) since the very first
chemistry course, Martin for just being a cool dude and Mårten S for all our
joint trips to Max Lab. A big “thank you” to the whole technical and
administrative staff: your assistance has been truly indispensable.
Tack också till alla kompisar utanför labbet som fått mig att tänka på annat än kemi. Till Nicklas (min resekompanjon på andra sidan klotet), Nea
och Per för alla vilda västkustresor, Janne för allt fågelskåderi och till min
bror Erik för våran starka vänskap. Varmt tack till mamma och pappa för att
ni alltid har trott på mig och till all annan släkt, såväl ny som gammal, för
stöd och uppmuntran.
Till sist, tack till Annika och Nelly för all kärlek – utan er är jag ingenting.
Älskar Er båda mer än mest.
Uppsala – 10 April 2006
48
Summary in Swedish
I dagens samhälle används uppladdningsbara litiumjonbatterier i ett flertal
småskaliga applikationer som t.ex. mobiltelefoner, bärbara datorer och elverktyg. Jämfört med andra batterikoncept, som t.ex. nickel metallhydrid
(NiMH) och nickel-kadmium (Ni-Cd) batterier, är litiumjonbatteriet ett miljövänligare alternativ, med lägre vikt, förhållandevis låga råmaterial kostnader och en större flexibilitet när det gäller dess design. Framförallt erbjuder
det också en högre energitäthet. Energitäthet är ett mått på hur mycket energi
som kan lagras per vikt- eller volymsenhet och är av stor betydelse när när
man vill skapa små, lätta batterier med hög effekt.
Figur 1 Schematisk bild av ett litiumjonbatteri.
Ett uppladdningsbart litiumjonbatteri består av en katod (positiv pol), en
anod (negativ pol) samt en elektrolyt mellan dessa som kan vara fast, flytande eller gelartad (Fig.1). Gemensamt för alla elektrolyter är att de innehåller
ett litiumsalt (t.ex. LiPF6 eller LiN(SO2CF3)2) som bidrar till jonledningsförmågan i elektrolyten. I de batterier som finns på den kommersiella marknaden idag består katoden vanligtvis av en metalloxid, såsom LiCoO2 eller
49
dopade varianter av LiNiO2, medan anoden utgörs av grafit (kol). Vid uppladdning vandrar Li-jonerna från katoden, genom elektrolyten, till anoden,
medan elektroner frigör sig från katoden och överförs via en yttre krets till
anoden. Under urladdning, vilket är den spontana processen, sker den motsatta reaktionen. Detta upp- och urladdningsförfarande ska i ett optimalt
batteri kunna utföras upp emot 1000 gånger utan att någon större mängd
kapacitet går förlorad.
Litiumjonbatterier kan även tänkas användas i storskaliga applikationer som
t.ex. elbilar eller elhybridbilar. Detta kommer att realiseras inom en mycket
snar framtid, då Toyota deklarerat att litiumjonbatterier ersätter NiMH batterier i nästkommande generation av deras elhybridbil Toyota PRIUS.
När det gäller utveckling av batterier för storskaliga applikationer spelar
priset för råmaterialen till katoden en avgörande roll. I ett mobiltelefonbatteri
där endast ett par gram av det aktiva katodmaterialet behövs är detta inte av
lika avgörande betydelse, vilket medför att ett förhållandevis dyrt material
som LiCoO2 används. I ett batteri till t.ex. en elhybridbil kommer däremot
flera kilogram av katodmaterial att användas, vilket motiverar att ersätta den
dyra kobolten (Co) med en mycket billigare metall – helst med den allra
billigaste, nämligen järn (Fe). Denna avhandling har därför fokuserat på att
studera och utveckla järn-baserade katodmaterial som framförallt lämpar sig
i Li-jonbatterier för storskaliga applikationer.
Litiumjärnfosfat (LiFePO4) är i dagsläget det mest lovande järn-baserade
katodmaterialet. För att kunna anpassa och utveckla materialet på bästa sätt
är det viktigt att undersöka och förstå de kemiska processer som sker i materialet under upp- och urladdning. En studie genomfördes därför för att undersöka vilka reaktioner som sker i gränsskiktet mellan en LiFePO4 elektrod
och elektrolyten under cykling (en cykel = en upp- och urladdning) vid
rumstemperatur och 40qC, samt hur dessa reaktioner påverkar batteriets prestanda. Resultatet påvisade att ganska lite produkter bildadas på ytan av LiFePO4 elektroden, vilket skiljer sig från liknande studier av andra katodmaterial. I dessa tidigare studier har det föreslagits att en större mängd reaktionsprodukter på ytan av elektroden medför en ökad resistans, vilket leder
till försämrade cyklingsegenskaper och en lägre erhållen kapacitet. Det har
också fastslagits att mer produkter bildas på ytan när temperaturen höjs.
Även om en något större mängd produkter kunde detekteras på elektroder
cyklade vid 40qC, konstaterades att den erhållna kapaciteten vid den elektrokemiska cyklingen var högre jämfört med elektroderna cyklade vid rumstemperatur. Dessa resultat förstärker bilden av LiFePO4 som ett lovande
katodmaterial för storskaliga applikationer, där arbetstemperaturen förväntas
ligga över rumstemperatur.
50
I sökandet efter nya järn-baserade katodmaterial för Li-jonbatterier har litiumjärnsilikat (Li2FeSiO4) visat sig vara ett mycket lovande alternativ. Materialets kapacitet och cyklingsegenskaper har utvärderats med olika elektrokemiska mätmetoder, vilka innefattar både galvanostatisk och potentiastatisk
upp- och urladdning. De strukturella förändringarna på atomär nivå under
den elektrokemiska cyklingen vid 60qC har studerats genom en kombination
av in situ röntgendiffraktion (XRD) och Mössbauer spektroskopi. Dessa
mätningar avslöjade att materialet cyklar mellan Li2FeSiO4 och LiFeSiO4,
men att stora strukturella förändringar sker i materialet under den första cykeln. Det faktum att den initiala potentialen på 3.10 V sjunker till 2.80 V
under denna första cykeln, talar för att en mer stabil struktur bildas. Denna
bibehålls under fortsatt cykling. Endast små förluster i kapacitet på mindre
än 3% noterades för batterier cycklade mer än 100 gånger.
Ytstudier liknande de som gjordes för LiFePO4, utfördes också på
Li2FeSiO4. Dessa påvisade att endast ett tunt skikt av produkter, till största
del bestående av oreagerat litiumsalt, bildas på ytan vid elektrokemisk cykling. Den goda kapacitetsbehållningen vid långtidscykling är säkerligen
kopplad till att endast en tunn ytfilm bildas och att litiumsaltet är stabilt under upp- och urladdning.
Från de resultat som erhållits för Li2FeSiO4 kan man dra slutssatsen att
materialet är mycket lovande för litiumjonbatterier i storskaliga applikationer.
51
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Acta Universitatis Upsaliensis
Digital Comprehensive Summaries of Uppsala Dissertations
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