Towards a Lithium-ion Fiber Battery
by
MASSACHUSETS NTTE
OF TECHNOLOGY
Benjamin Grena
JUL 0 12013
B.Sc. Physics,
Ecole Polytechnique, France, 2010
LIBRARIES
Submitted to the Department of Materials Science and Engineering
in partial fulfillment of the requirements for the degree of
Masters of Science in Materials Science and Engineering
at the
MASSACHUSETTS INSTITUTE OF TECHNOLOGY
September 2013
@ Massachusetts Institute of Technology 2013. All rights reserved.
Signature redacted
Author .............
Departme&AfMaterials Science and Engineering
August 15, 2013
Signature redacted
Certified by........
/
Yoel Fink
Professor of Materials Science
Professor of Electrical Engineering and Computer Science
Thesis Supervisor
Accepted by...................
Signature redacted,
G,-
Gerbrand Ceder
Chair, Departmental Committee on Graduate Students
Towards a Lithium-ion Fiber Battery
by
Benjamin Grena
Submitted to the Department of Materials Science and Engineering
on August 15, 2013, in partial fulfillment of the
requirements for the degree of
Masters of Science in Materials Science and Engineering
Abstract
One of the key objectives in the realm of flexible electronics and flexible power sources
is to achieve large-area, low-cost, scalable production of flexible systems. In this thesis
we propose a new Li-ion battery architecture in a fiber form that could be the building
block to large-area, conformal, flexible power sources, achieved through fiber thermal
drawing. This architecture is based on the key-finding of using thermally induced
phase separation as a method to introduce porous structures inside thermally drawn
fibers for the very first time. This new versatile process allows us to incorporate
ionically conductive gel-polymer electrolytes in fiber cores in a very simple way, with
ionic conductivities suitable for a battery application. The rest of our proposed infiber battery architecture is composed of composite electrodes, which we fabricate
and characterize. A model system is tested and a detailed pathway towards the first
successful fabrication of a Li-ion fiber battery is given.
Thesis Supervisor: Yoel Fink
Title: Professor of Materials Science
Professor of Electrical Engineering and Computer Science
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Acknowledgments
!
I first would like to thank my advisor Professor Yoel Fink for giving me the opportunity to conduct research in his group. His inspiring visions, contagious enthusiasm
and profound advice have always allowed me to move forward, even at times where
I felt cornered by the difficulties of life and research. I owe him a lot.
I am extremely thankful and indebted to my colleagues and friends in the group.
Guillaume Lestoquoy, who has always found the time to answer my numerous emails
about grad school and MIT before my arrival here. I think him, as I would not be
working in this wonderful lab without him. Dr. Sasha Stolyarov who has given me
the inspiration for this great project, and has always shared his ideas and excitement.
I also thank the rest of the group, No6mie, Chong, Alexander, Lei, Xiaoting, Andres,
Tara, Jeff and Michael. Through our daily interactions and many discussions, you
have helped me become a better scientist.
I thank the rest of the MIT community with whom I have had the chance to
interact. Professor Samuel Allen, for his eye-opening lectures and fruitful discussions.
Mark Belanger for his patience at the Edgerton machine shop. Thanks as well to
the many DMSE students that have made my experience here much richer and more
fun than it would have been without them.
I thank my friends in family in France, whom I miss more than anything else (even
more than "saucisson"). I am furthermore deeply grateful to share my appartment
with my roommates Hadrien, Pierre and Matthieu, among whom I have to include
the free-electron Julien. Thanks for being awesome, guys!
Lastly, I thank my sweet sweet love Eh6onore for being the sunshine of my life,
and for her constant support and affection (and rare grouchiness). Merci
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)
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Contents
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Introduction
Flexible electronics . . . . . . . . . . . .
Flexible power sources . . . . . . . . . .
Why a Lithium-ion fiber battery? . . . .
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Background
1.1 Lithium-ion batteries: fundamentals
1.1.1 Working mechanism . . . .
1.1.2 Anode and cathode materials
1.1.3 Electrolyte materials . . . .
1.2 Thermal drawing process . . . . . .
1.2.1 Process description . . . . .
1.2.2 Advantages and constraints
1.3 Challenges and proposed architecture
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Case of the electrolyte
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1.3.2
Case of the electrodes
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1.3.1
Fabrication of a porous structure inside a fiber
2.1 Thermally-induced phase separation of a polymer solution
2.1.1 The TIPS process . . . . . . . . . . . . . . . . . .
2.1.2 Materials selection . . . . . . . . . . . . . . . . .
2.2 Fiber preparation . . . . . . . . . . . . . . . . . . . . . .
2.2.1 Preform preparation . . . . . . . . . . . . . . . .
2.2.2 Fiber drawing . . . . . . . . . . . . . . . . . . . .
2.3 Fiber characterization . . . . . . . . . . . . . . . . . . .
2.3.1 SEM imaging and porosity . . . . . . . . . . . . .
2.3.2 Spherulite size distribution . . . . . . . . . . . . .
2.4 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . .
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Ionic conduction inside a fiber
3.1 The two birds/one stone process
3.1.1 Description of a gel-polymer electrolyte
3.1.2 TBOS process . . . . . . . . . . . . . .
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3.1.3 TBOS in a film geometry ...........
Fiber drawing and characterization . . . . . .
3.2.1 Preform preparation and fiber drawing
3.2.2 SEM characterization . . . . . . . . . .
3.2.3 Ionic conductivity of the fibers . . . . .
Conclusions . . . . . . . . . . . . . . . . . . .
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Conclusion
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4 Pathway towards a Lithium-ion fiber battery
4.1 Electrode fabrication and characterization . . . .
4.1.1 Fabrication method . . . . . . . . . . . . .
4.1.2 Electrical conductivity . . . . . . . . . . .
4.1.3 SEM imaging . . . . . . . . . . . . . . . .
4.2 Battery assembly and testing in a coin-cell . . . .
4.2.1 Method . . . . . . . . . . . . . . . . . . .
4.2.2 Results and discussion . . . . . . . . . . .
4.3 Next steps: preform preparation and fiber drawing.
4.4 Conclusions . . . . . . . . . . . . . . . . . . . . .
8
Introduction
Flexible electronics
The domain of flexible electronics has been attracting considerable research and
public interest over the past few years. Flexible electronics potentially offer a scope
of applications much larger than that enabled by traditional wafer-based electronics.
Devices such as flexible organic light-emitting diodes displays [1, 2], or flexible solar
cells [3, 4] have already emerged from research groups all around the world and are
slowly transitioning to the market.
Flexible power sources
Among all flexible electronic devices, flexible power sources occupy a crucial position.
Indeed, flexible power sources are a key element in order to achieve fully flexible
electronic systems. So far, Li-ion batteries [5, 6, 7] as well as supercapacitors [8, 9]
have seen the day in a flexible form, from various research groups. Typical fabrication
methods for such devices include polymer nanocomposite fabrication and stacking,
or microbattery array fabrication by electrodeposition and lithography. However
these methods are limited in terms of device sizes and scalability. One of the key
research objectives in the realm of flexible power sources - and more generally flexible
electronics - is thus to achieve large area, low-cost, and versatile production of flexible
systems.
Why a Lithium-ion fiber battery?
In this work, we propose and present a fabrication scheme based on the thermal
drawing process leading to flexible Li-ion fiber battery devices. Motivations for this
work are three-fold:
" Firstly, fibers can be extremely easily integrated into flexible objects of any
shape and size through the millenial technique of weaving. Thus, we believe
a Li-ion fiber battery could be the building block to long-sought large-area
conformal flexible power sources.
" Secondly, fibers are ubiquitous. Transforming traditional fibers into potential
energy storage devices could make a big difference in our daily lives, but also in
9
the design habits of consumer electronics. A back of the enveloppe calculation
shows that a pair of jeans could withhold enough energy to charge an iPhone,
assuming it's made of fiber batteries with an energetic capacity as low as one
tenth of a traditional Li-ion battery.
* Lastly, the thermal drawing process is a cheap and intrinsically scalable fabrication method, which is not the case of current Li-ion flexible batteries fabrication
methods.
In order to propose a suitable architecture and materials' processing, one needs to
grasp the mechanism at play in Li-ion batteries, in addition to the specificities of the
thermal drawing process and the constraints it may impose on materials selection.
Thus we will start by giving a quick review of the fundamentals of Li-ion batteries
as well as of the thermal drawing process. We will actually see how a potential
fabrication scheme emerges from these constraints - particularly through the idea of
using phase separation to introduce porous domains inside fibers.
We will then go into greater details in the idea of resorting to phase separation of a
solution during the draw to introduce porous structures in a fiber, by first explaining
the scientific basis and evolving to more complex and functional developments.
With these structures in hand, we will see how one can tune the fabrication
process to fabricate gel-polymer electrolyte core fibers in a single fabrication step,
with no post-processing needed.
Lastly, we will present work on the electrode front. We will detail the fabrication
method and present the properties of our fabricated battery electrodes. Finally we
will give the pathway towards the first functional Li-ion fiber battery by giving the
steps of a relatively simple fabrication method.
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Background
1
Lithium-ion batteries: fundamentals
1.1
In order to design a Li-ion fiber battery fabrication scheme, one needs to grasp
the mechanisms at play in a traditional Li-ion battery as well as the "catalog" of
materials used. We will thus quickly review the fundamentals of Li-ion polymer
batteries, which is the common architecture to all flexible Li-ion batteries so far.
1.1.1
Working mechanism
The working mechanism for such rechargeable batteries is often reffered to as "the
rocking chair" mechanism, or intercalation/deintercalation mechanism [10]. Upon
discharge, for instance, lithium ions flow through the electrolyte from the anode to
the cathode where they intercalate directly into the crystal structure of the cathode
material (cf. Fig. 1). The opposite occurs upon charging the battery. In each case the
driving force is an electrochemical potential difference between anode and cathode,
and the flow of lithium ions is accompanied by a counter flow of electrons in the
external circuit.
Electrons
--
+
oa
A deCathode
collector
Current
collector
Polymer electrolyte
Figure 1: Schematic representation of a Li-ion polymer battery under discharge.
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1.1.2
Anode and cathode materials
Good electrode materials should exhibit high gravimetric capacities to maximize
the energy storage, small volume changes between lithiated and unlithiated states
to limit strains on the system, and high diffusivity of lithium ions to ensure rapid
charging.
A very typical material for the anode is graphite. Upon charging, Li intercalates
in between the graphene layers of the graphitic structure to eventually form the structure LiC6 with a small volume change, and a theoretical capacity of 370 mAh.g-1
[10]. Other materials are also used, such as lithium metal (mostly for research purposes due the hazards of lithium), or more recently tin and silicon nanoparticles
[11, 12].
Concerning the cathode, the most commonly used material is LiCoO 2 , a layered
lithium oxide with a theoretical gravimetric capacity of about 300 mAh.g- 1 . Many
other materials have been studied and used, such as spinel lithium metal oxides, or
more recently lithium iron phosphate [10].
Active anode and cathode material are generally introduced into a polymer composite electrode, by mixing the particles with a suitable polymer binder and adding
carbon black for improved electrical conduction. The choice of binder is dictated by
its electrochemical stability and also its affinity to the particles [10]. PVdF is again
very generally the preferred-choice as a polymeric binder. The role of the added carbon black is to ensure a sufficient electrical conductivity. Indeed, it is important that
the electrodes have a high enough conductivity such that the electrons generated as
lithium ions exit the anode upon discharging, for instance, and travel through the
electrodes up to the current collector where they are to be swept away. Electrodes
are processed so as to maximize the amount of active material per weight of binder,
to achieve a good conductivity (generall around a few S.m-1) and capacity, without
sacrifying the cohesivity of the final structure.
-
Both anode and cathode are always in contact with a current collector whose role
is to ensure conduction of the electrons from the electrode to the external circuit.
Generally, copper and aluminum are used as anode and cathode current collectors
but quite generally any metal with a low reactivity towards the electrode is a good
choice for a current collector.
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1.1.3
Electrolyte materials
Electrolyte materials for Li-ion batteries can be of three types: liquid nonaqueous
solutions, polymer-based, or solid electrolytes.
Liquid nonaqueous electrolytes In the first case, a lithium salt is dissolved into
a nonaqueous solvent. Driven by an electrochemical potential difference, the lithium
ions migrate directly in the liquid phase and yield ionic conductivity. This type
of electrolyte requires the use of an insulating membrane in between the electrodes
called a separator, blocking the transport of electrons through the solution, which
would cause the battery to short. Many lithium salts have been studied, but the
two most common are lithium hexafluorophosphate LiPF6 and lithium perchlorate
LiClO 4 . Organic solvents from the carbonate family are typically used such as ethylene carbonate, dimethyl carbonate or propylene carbonate. This kind of electrolytes
exhibit relatively high ionic conductivity, on the order of 10 mS.cm-' [10]. However
because of safety issues due to solvent leakage and short lifetime due to dendrite
growth from anode to cathode shorting the battery from the inside [10], liquid electrolytes are no longer used and have been replaced by polymer electrolytes.
Gel-polymer electrolytes and polyelectrolytes Polymer electrolytes are generally based on a porous polymeric structure filled with liquid electrolyte. Such structures might often be referred to as gel-polymer electrolytes, due to the swelling of
the polymer by the liquid electrolyte. Several host polymers have been investigated
such as poly(acrylonitrile), poly(methyl methacrylate) or poly(vinylidene fluoride)
[13]. Performances are measured based on ionic conductivity, electrochemical stability towards the electrodes, or cycling behavior. Despite a somewhat lower ionic
conductivity in general (on the order of 1-5 mS.cm- 1 ), gel-polymer electrolytes display many advantages over their liquid counter parts such as increased safety due
to lower leakage and lower reactivity, longer battery lifetimes thanks to suppression
of dendrite growth, as well as better flexibility [13]. For these reasons, gel-polymer
electrolytes are so far the best electrolytes for Li-ion batteries, and are found in
industrial batteries. They are furthermore almost ubiquitously found in flexible battery systems, owing to their flexibility and increased safety. For these reasons, we
shall focus on incorporating such an electrolyte in our fibers.
Recently, systems in which lithium salts are dissolved directly into a polymer
network have seen the day. Such systems are known as polyelectrolytes, and are
13
particularly compelling due to the absence of liquid phase and thus of toxic solvent. However such polyelectrolytes so far display low conductivity rendering them
unsuitable for room temperature applications [10].
Solid electrolytes Lastly, we mention solid electrolytes which are ionically conductive solid materials in which ion migration occurs by diffusion in a crystal lattice.
Solid electrolytes are a relatively novel kind of electrolytes and are currently attracting a growing research interest [10], however they are not suited for flexible battery
applications and are thus not the focus of the present work.
1.2
1.2.1
Thermal drawing process
Process description
The thermal drawing process starts by the fabrication of a multimaterial macroscopic object called a preform, identical in its geometry and composition to the final
fiber however much larger in its cross sectional dimensions and shorter in length. A
preform is usually comprised of a cladding material and functional materials, in a
well-chosen architecture. The cladding material can either be an amorphous thermoplastic or silica glass, depending on the temperature regime sought. Its function
is mostly to support the stresses arising in the fiber during the draw process. Examples of thermoplastic materials used as cladding include polycarbonate, polysulfone,
polyetherimide, or poly(methy methacrylate). In polymer preforms, cladding and
active materials are assembled by either film-rolling of film stacking, followed by
a consolidation step under vacuum to remove trapped gases and fuse the different
materials into one solid part. Because the cladding material represents usually up
to 95% of the preform in weight, it dictates the regime of draw temperatures. A
general rule of thumb is that the appropriate draw temperature is about 80*C above
the cladding polymer's glass transition temperature.
Once assembled and consolidated the preform is taken to a draw tower and preheated in a vertical furnace to an elevated temperature. The furnace temperature is
controllable in three points corresponding to the top, bottom and middle zone. The
middle zone temperature is usually set to be the hottest, and thus we often refer to
the middle zone as the hot zone. During preheating the hot zone temperature is set
to around 100 to 120'C above the cladding material's glass transition temperature.
The aim of this preheating step is to locally soften the preform in the region of the
14
The preform is fed in the furnace
at a controlled feed speed
Draw tower furnace
Necking zone
Hot zone)
The fiber is pulled on by a capstan
at a controlled speed
Figure 2: Schematic representationof the thermal drawing process. The preform is loaded
into the furnace and fed downwards at a controlled feed speed. The cladding is
represented in yellow and the "active material" in purple. The preform necks
down in the hot zone and extends into a fiber, which is pulled on at a controlled
speed. Figure courtesy of G. Lestoquoy.
hot zone to a point where it necks down and starts falling under the effect of grav-
ity. We call this event the "bait-off". Following the bait-off, the bottom part of the
preform is cut and thrown away, and the proper drawing of the fiber then starts.
The fiber is fed into a capstan which pulls on the fiber at a constant controllable
speed vcapstan, meanwhile the preform is fed into the furnace at another prescribed
feed speed Vdownfeed (cf. Fig. 2). Conservation of volume dictates that the drawdown ratio v, relating the cross-sectional dimensions of features in the preform to
_
Lpreform
-
dimensions in the fiber, is set by the processing speeds as:
Vcapstan
Vdownfeed
Lfiber
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A set of lasers measures the cross-sectional dimensions of the fiber, and a tensiometer measures the tension in the fiber, which is then used to compute the stress
in real time. By monitoring the process with a LabView program and properly
controlling the different control knobs (mostly capstan speed and furnace temperatures), one can draw a fiber extending hundreds of meters in length. Provided the
different materials in the preform are able to codraw at the draw temperature, the
fiber conserves the geometry and composition of the preform, but its cross-sectional
dimensions are reduced by a factor v of 10 to as large as 500.
This technique enables fabrication of features down at the submicron-scale, extending kilometers in length. Examples of functional devices achieved by this fabrication method include photonic bandgap transmission fibers [14], photodetecting
fibers [15], fiber capacitors [16], or even acoustic fiber transducers [17].
1.2.2
Advantages and constraints
The advantages of the thermal drawing process are numerous: the method allows for
production of fiber devices with features down at the nanoscale in a relatively cheap
and highly scalable fashion. Furthermore, once a fiber is drawn, it can be weaved and
integrated into large-area multicomponent architectures of various shapes [18, 19, 20].
Thus, the fiber is an elementary building block to large-area conformal and flexible
electronic systems.
This method however imposes constraints on the materials that can be used. First
of all, it relies on the application of heat and relatively large temperatures. This can
cause some materials to decompose or degrade. Second, the materials composing a
preform should be able to co-flow, in other words they should have similar rheological
properties in the regime of draw temperatures. If a crystalline material is to be
introduced in the fiber, it needs to have a melt point below the draw temperature.
All theses constraints limit the potential combinations of materials in preforms.
1.3
Challenges and proposed architecture
We've seen how a Li-ion fiber battery could be a very compelling device due to
its ease of integration into large-area flexible systems, and low-tech and scalable
production. However, based on the constraints previously mentioned, the thermal
drawing process seems to be somewhat unsuited to the different battery elements,
particularly to the gel-polymer electrolyte, relying on a porous polymer structure.
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1.3.1
Case of the electrolyte
-
A naive idea would be to incorporate porous PVdF films directly into a preform and
attempt to draw at low stress in order to avoid pore collapse. An intuitive image
backing that idea up is that of lava flow; the lava's high viscosity and the flow's low
shear rate traps air bubbles and the structure resulting from solidification is highly
porous. However we never obtained such structures when trying to draw porous
PVdF. It is likely that the viscosity of PVdF was too low at the attempted draw
temperatures, and that the extensional shear at the necking point was too intense
to preserve the shape of the air pockets.
Instead of trying to work against those very conditions, we devised a process
making use of them. We achieved the incorporation of a porous PVdF structure
the backbone of the polymer electrolyte - by phase separating a PVdF solution in
the draw process. Details on this process are given in Sections 2 and 3.
1.3.2
Case of the electrodes
In addition to the electrolyte issue, it would be impossible to draw active electrode
materials as such. Graphite sublimes before melting at atmospheric pressure, and
lithium cobalt oxide decomposes before melting.
Instead, the solution we chose was to incorporate active material nanoparticles
in a polymeric matrix, and adjust composition and host polymer type in order to
tune the rheology of the composite to reach drawability. Such composites had in
fact already been used in a number of thermally drawn fiber devices, usually in
the form of carbon black composites as viscous electrodes [20, 16]. Fabrication and
characterization of such electrodes are detailed in Section 4 of this thesis.
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2
Fabrication of a porous structure inside a fiber
2.1
Thermally-induced phase separation of a polymer solution
Phase separation is a phenomenon by which a homogeneous solution of multiple
components undergoes separation into two distinct phases. Phase separation can
be observed in any given class of materials, and in either solid or liquid solutions.
Depending on the specific system, phase separation can be induced by the change of
an intensive parameter such as temperature, composition, magnetic field, etc. We
will focus on temperature-induced phase separation as a way to generate porosity
inside our fibers. In fact, the so-called TIPS process is a very common method for
polymer microporous membrane fabrication [21, 22], including in a fiber geometry
[23]. However these fibers are generally single-material hollow-core porous fibers,
processed by solution extrusion, and cannot be part of more complex multimaterial
fiber architectures.
2.1.1
The TIPS process
Process description The TIPS process for microporous polymer membrane fabrication quite generally follows the following protocol [21]:
1. One first forms a homogeneous solution of a polymer with a high-boiling point
liquid known as the solvent, at an elevated temperature
2. The solution is then cast into a film or poured into a mold
3. The solution is subsequently allowed to cool down more or less rapidly in order
to induce phase separation and solidification (or vitrification) of the polymer
phase
4. The solvent is removed by vacuum evaporation or solvent extraction
The TIPS process is a fairly easy and versatile way of fabricating microporous
membranes. Compared to other methods, such as diffusion-induced phase separations of polymer solutions, it is simpler and more controllable. Indeed, once a
polymer-solvent system has been chosen, there are only two parameters one can act
upon: volume fraction and cooling rate.
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+
T
binodal
curve
spinodal
curve
0
C
Figure 3: Schematic phase diagram of a UCST polymer solution.
Theory
The thermodynamics of the TIPS process is well described by the Flory-
Huggins theory of polymer solutions, which gives the free energy of a polymer solution
in terms of temperature, polymer volume fraction, molecular weight, and interaction
parameter x between the polymer and solvent molecules.
From this free energy it
is possible to construct phase diagrams for polymer solutions, provided one has a
measure of the interaction parameter.
A typical phase diagram of upper critical
temperature solution is given in Fig. 3. A homogeneous solution formed in region 1
thus phase separates into two phases either by a binodal or spinodal decomposition
when cooled down. The polymer rich phase eventually forms the continuous polymer
matrix, and the solvent-rich phase is subsequently removed to leave voids. Depending
on the crystallinity or amorphicity of the polymer and its interaction parameter
with the solvent, the phase diagram can be more complex and exhibit a polymer
crystallization line [24], and the phase separation may either be liquid-liquid or liquid-
solid.
The specific microstructure, including porosity, pore size and size distribution,
pore shape, or spherulite size in the case of a semi-crystalline polymer, is mostly dictated by the interplay between the thermodynamics and the kinetics of the transformation. Many experimental and theoretical studies discussing the effects of various
parameters are available [21, 22, 24]. We will only quote the general result that the
faster the cooling rate, the smaller and more numerous the pores, whereas for slower
cooling rates, pores tend to be less numerous and larger [21].
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2.1.2
Materials selection
Constraints and options In order to achieve a temperature-induced phase separation inside a fiber during the draw process, we needed to choose a polymer-solvent
system as well as a suitable cladding material. The choice of these different components is obviously interdependent. Specifically, the chosen solvent should:
- have a boiling point higher than the draw temperature dictated by the cladding
- not dissolve the cladding in any way
- allow phase separation of the polymer, ie. dissolve it at temperature close to
the drawing temperature, but not at room temperature
Additionally, the core polymer should be a suitable choice for a polymer-electrolyte
host polymer, as this is eventually our end goal.
We naturally turned to PVdF, being one of the most used polymers for polymerelectrolytes mostly because of its high chemical stability with respect to electrode
materials [10]. It followed from materials data that, among others, the choice of
propylene carbonate as a solvent and cyclic olefin copolymer (hereafter, COC) as
a cladding material was suitable. COC is furthermore an excellent moisture barrier [25], which would be extremely important in a battery application particularly
because of the electrolyte's sensitivity to moisture.
Film experiments We went on to validate the choice of these materials by experiments at the preform-level. Films of porous PVdF were fabricated following the
TIPS method. PVdF (Kynar 761, Arkema) was dissolved into propylene carbonate
(Alfa Aesar) at 2000C in a reaction bottle, in a 20:80 weight/weight ratio. We then
cast the solution on a preheated glass plate and allowed the films to cool down under different conditons. One film was cooled down in a oven by slowly decreasing
the temperature overnight, another film was cooled by simply placing it at room
temperature. The remaining solvent Was then leached out with ethanol. The membranes were subsequently dried under vacuum, freeze-fractured in liquid nitrogen to
ensure a clean cross-section, and gold sputtered for electron conduction in the SEM.
We then observed them using a scanning electron microscope (JSM-6010LA, JEOL).
Micrographs are shown in Fig. 4.
From the micrographs it is visible that the rapidly-cooled sample exhibits a porous
microstructure whereas the slowly-cooled sample seems dense at the micro-scale.
20
(a)
(b)
Figure 4: Micrographs of PVdF films obtained by thermally-induced phase separation at
(a) high cooling rate and (b) low cooling rate.
This can be explained considering that, at the limit of very low cooling rates, the
structure has enough time at elevated temperatures to rearrange itself to minimize
its surface area through coarsening of the phase separated structure. In the rapidlycooled samples, the structure is quenched into a phase separated state and the polymer chains do not have enough mobility to rearrange.
Furthermore, in the porous
sample the polymer was in the form of spherulites. This was expected as PVdF is a
highly crystalline polymer, known to form such microstructures upon solidification
[21].
Rheology measurements
In the context of an eventual fiber drawing, we could
not limit ourselves to knowing that our system exhibited a phase separation between
200'C and room temperature. We needed to know at what temperature did the phase
separation occur, for a given cooling rate. The proper way of obtaining this phase
separation temperature would be by using a Differential Scanning Calorimeter. For
technical reasons, we could not use such an equipment. Instead, we resorted to an
indirect measurement using a rheometer (AR 20001, TA Instruments). We measured
the viscosity of a PVdF/propylene carbonate in a 20:80 weight ratio, from 200'C to
25 C, under a constant shear rate of 1 s-1. We obtained the viscosity curve presented
in Fig. 5
We see that the solution's viscosity slowly increases from 200'C to about 40'C
from 1 Pa.s to 10 Pa.s. At 40'C, the viscosity abruptly jumps by around 3 orders of
magnitude. We interpret this large increase as the signature of the phase separation.
21
10
10
10+
10
20 1 16 180 200
20
40
80
(60
1
Temperature (in *C)
.
Figure 5: Viscosity as a function of temperature for a PVdF/propylene carbonate in a
20:80 weight ratio, from 200' C to 250 C, under a constant shear rate of 1 s
Indeed, upon phase separating, the system goes from a homogeneous solution to a
mixed phase system composed of a nearly solid PVdF network swollen by solvent. We
thus anticipate a much higher viscosity in the phase separated state due to the solid
polymer contribution. Furthermore the observed temperature is in good agreement
with values reported by other groups performing DSC measurements on the same
system [26]. Thus our system in fact phase separates at a rather low temperature,
and in practice we will always be above that temperature during the draw.
2.2
Fiber preparation
2.2.1
Preform preparation
The first attempt to incorporate a porous structure inside a fiber was made with
a very simple core/cladding structure. COC rods were acquired (TOPAS 6015S,
Boedeker Plastics) and machined using a lathe to 9" long cylinder, 1.25" in diameter.
A 5/16" hole was drilled over about 5" into the rod. These pure COC preforms were
then baked in a vacuum oven at 110'C for at least two weeks to eliminate moisture.
We then prepared a PVdF/propylene carbonate solution 20:80 weight ratio in a
reaction bottle, on a hot plate at 150'C. The homogeneous solution was subsequently
poured into the core of the COC preform. We then took the preform immediately
to the draw tower to begin the draw process.
22
2.2.2
Fiber drawing
As we have seen, the draw process starts with a preheating step prior to bait-off where
the preform is locally heated to up to 100 to 120'C above the cladding material's
glass transition temperature. In the case of a COC cladding with a T9 of 150 C, this
corresponds to locally heating the preform at about 260'C. This is above the boiling
point of propylene carbonate, and thus called for a buffer region at the bottom of the
preform made of pure COC. We purposely placed the bait-off location in the bottom
of that buffer region, and subsequently turned the temperature down during the
draw in order to reach a temperature below the boiling point of propylene carbonate
before the solution penetrated the hot zone. This allowed for the temperature of
the solution to always remain below the boiling point of the solvent, and thus avoid
substantial solvent loss and change of concentration. In order to limit evaporation,
a PTFE plug was furthermore placed into the core hole.
2.3
2.3.1
Fiber characterization
SEM imaging and porosity
Method In order to observe the microstructure of our fibers and conclude on their
porosity, we prepared samples for SEM imaging. Because of the vacuum needed in
scanning electron microscopy, we had to dry our samples from any trace of propylene
carbonate. To do so, we chose to expel the core from the fiber cladding uSing a thin
steel wire, and leach the extracted RVdF core with ethanol to remove any trace of
propylene carbonate, prior to vacuum drying the samples to get rid of the ethanol.
We then freeze-fractured the samples in liquid nitrogen, as well as gold-sputtered
them to enable conduction of the electrons in the SEM.
Results and discussion The micrographs clearly show a porous structure constituted of interconnected polymer spherulites and voids, very similar to that obtained
in film experiments.
We estimated the porosity q, defined by the ratio of void volume over total
volume, by carefully weighing the dried core samples and measuring their radius r
and length 1. Knowing the density PPVdF of the PVdF, the porosity follows from our
measurements by:
= 1 mmeas
2
PPVdF
23
.
irr
Figure 6: SEM micrograph of the fiber cross section after drying and freeze fracturing.
The interconnected spherulite structure leaves voids and the core of the fiber is
thus porous. The overall diameter of this sample was around 150 pm
24
We got a porosity varying from 15% to 25%, with a large deviation from sample
to sample due to their very small sizes and weights, close to our microbalance's
precision. However this porosity range is in good visual agreement with the SEM
micrographs.
2.3.2
Spherulite size distribution
Objective The overall porosity is an extremely important parameter of porous
structures, especially for use as gel-polymer electrolyte where it is tied to the liquid
electrolyte intake. However the precise microstructure is also a crucial factor. Therefore, we wanted to study the microstructure in more details, and specifically devise
a method to obtain quantitative informations on the spherulite size distribution. We
could then relate this distribution to the different fabrication parameters (polymer
concentration, draw speed, furnace temperature,...). This could be interesting both
to fine tune the process such that the microstructure meets our needs, but also on a
scientific point of view to study the effect of these parameters on the solution stability
and polymer crystallization.
Method Instead of extracting the size distribution by hand, we implemented an
image analysis algorithm which does so automatically. The algorithm is based on the
use of the Hough transform [27]. The principle is as follows. To find the spherulites
of radius R in the image:
" We create a so-called accumulator space which is made up of a cell per each
pixel of the image, and we initiate all cells to 0.
" We detect the edges of the picture (by a color gradient for example).
" For each edge point, we increment all the cells of the accumulator space that
are on a circle of radius R of center the current edge point. Those points are
possible candidate for being the center of a circle of radius R passing by that
edge point.
" Once we have done so for all edge points of the picture, we search for local
maxima in the accumulator space, ie. we search for the cells whose value is
greater than every other cell in its neighborhood. These local maxima are the
points with the highest probability of being the location of the center of the
circle that we are trying to locate.
25
2<q
f
I2d
Ra"k (oi1mo
Figure 7: Typical outcome of the algorithm. The red circles correspond to the detected
spherulites. The data are used to compute the histogram on the right.
We can then increment R to R+dR, and keep track of the local maxima and their
values. Eventually we will find the centers and radii of the circles which most probably correspond to a spherulite in the image.
The algorithm was implemented in
C++, and a typical outcome is presented on Fig. 7.
Results and discussion
We see that the algorithm is relatively efficient in finding
spheres, as it manages to find both small and large spherulites. The program does not
generate many defects, the number of "fake" spherulites or the number of erroneous
size estimates is rather low.
By tuning the different arbitrary thresholds in the
program (particularly the edge detection, but also the threshold for local maxima),
one can further improve the program's efficiency. Therefore it allows for a good size
distribution estimate, on a statistical sample.
Future work will focus on precisely
relating this microstructure to fabrication parameters.
2.4
Conclusions
Incorporating porous domains inside a fiber is difficult, as the draw conditions tend
to destroy any porous microstructure incorporated in a preform. However, we have
shown in this section that one can in fact use those very draw conditions, and specifically the high to low temperature transition, to induce phase separation of a polymer
solution leading to porous microstructures inside fibers. We first gave evidence of
this at the so-called preform level by carrying out experiments with films. We then
took the concept to the draw tower and achieved a porous core fiber in a very novel
and simple fashion. This structures may be interesting for many applications as they
may be used as drug delivery fiber vehicles, chromotography colons inside fibers, or
- as we will focus on next - gel-polymer electrolytes for use in battery applications.
26
3
Ionic conduction inside a fiber
3.1
The two birds/one stone process
We have achieved the incorporation of a porous structure inside our fibers, which
has the potential to act as a host polymeric structure for a gel-polymer electrolyte.
However our fibers are so far not functional as such. Indeed, let us quickly review
the materials, structure and conduction mechanisms of a gel-polymer electrolyte.
3.1.1
Description of a gel-polymer electrolyte
Gel-polymer electrolytes, or sometimes called plasticized polymer electrolytes, are
ionically conductive materials composed of a polymer or polymer blend, a solvent
or plasticizer dissolving a lithium salt, and potentially some additives. Gel-polymer
electrolytes are halfway between liquid electrolytes and polyelectrolytes, in that they
display both the diffusive properties of liquid electrolytes, and the mechanical stability of polyelectrolytes or solid electrolytes [13].
A gel-polymer electrolyte is typically made according to this 3-step method:
1. Construct a polymer porous film. Typically, phase separation methods will be
used such as diffusion-induced or temperature-induced phase separation.
2. Rinse and dry the membrane under vacuum to eliminate any leftover traces of
solvent used for the phase separation process.
3. Load the membrane by immersing it in a liquid electrolyte bath.
The ionic conductivity in gel-polymer electrolytes is due to lithium ion diffusion
both in the liquid electrolyte in the percolating porous network, but also in the
swollen polymer phase. Therefore the final ionic conductivity is essentially a function
of three factors [28]:
- the conductivity of the chosen solvent and lithium salt system
- the interaction between the polymer and solvent and particularly the ability
for the solvent to swell the polymer
- the specific porous polymer microstructure, including porosity, pores' shapes
and sizes, etc.
27
PVdF
e-
solution
SLi+
Homogeneous solution at high T
1000
Phase-separated system at
room temperature
M
lonically conductive material
Figure 8: Schematic description of the two birds/one stone process for film electrolyte
preparation.
The membrane's porosity should be high enough to allow for sufficient liquid electrolyte intake - source of ionic conductivity - but not too high so as to guarantee the
mechanical stability and cohesivity of the final system.
As mentioned earlier, many polymers have already been used as gel-polymer
electrolyte host polymers, however PVdF so far is the preferred choice due to its high
electrochemical stability [13]. Concerning liquid electrolyte, it has been shown that
the use of molar solutions of lithium perchlorate or lithium hexafluorophosphate in
ethylene carbonate, dimethyl carbonate, propylene carbonate, or mixtures of those,
leads to gel-polymer electrolytes with ionic conductivities on the order of a few
mS.cm-
1
3.1.2
TBOS process
at room temperature [10], which is a suitable value for a functional battery.
In order to make our fibers functional, we would therefore need to replace the solvent
occupying the pores after phase separation has occured with a lithium battery electrolyte. Technically, we would thus need to dry the fiber under vacuum to remove the
solvent, and subsequently replace it with a liquid lithium battery electrolyte, through
capillary rise for instance. Both of these processes (solvent drying and liquid electrolyte rise) are vi-processes, and therefore the time it would take to accomplish
these post-processing steps scales as the square of the fiber's length. Thus, resorting
to this 3-step method would completely go against the scalability of the thermal
drawing process.
Instead of resorting to such time-consuming post-processing steps, we have devised a single-step method by which our fiber's core immediately behaves as a ionically conductive medium upon exiting the draw tower, with no post-processing nec-
28
essary. To achieve this, we simply changed the composition of the polymer solution
inserted in the core of the preform and added a lithium salt, lithium perchlorate.
Provided the lithium salt does not have a substantial effect on the stability of the
formed solution, we should still observe a phase separation by going from draw temperature to room temperature - the difference being that the pores would now be
already filled with a liquid lithium ion-conducting electrolyte. Thereby we would
obtain an ionically conductive medium in one fabrication step, with no drying and
activation necessary, which we Kke to refer to as the Two Birds/One Stone process,
or TBOS process.
The choice of lithium perchlorate was guided by its ability to be dissolved in
propylene carbonate, as well as its high thermal stability. The decomposition temperature of lithium perchlorate was found to be around 505 C, way above typical
draw temperatures of around 200 C and comparable to other measures [29]. A
schematic description of the TBOS process for film electrolyte fabrication is given in
Fig. 8.
3.1.3
TBOS in a film geometry
Film preparation To valide this TBOS process idea, we first carried out experiments outside of the fiber, in film geometries. We first prepared a 1M solution of
lithium perchlorate in propylene carbonate, and homogenized it by mixing at 70'C
for 12 hours in a glovebox. Then, we mixed PVdF and this liquid electrolyte solution
at 200*C in a 20:80 weight ratio for 2 hours. The obtained homogeneous solutions
were then cast and spread with a doctor blade on a preheated glass plate, and cooled
to room temperature in the glovebox.
Film characterization The films were first rinsed with ethanol and dried under
vacuum in order to be observed in the SEM. Dried films were freeze-fractured in
liquid nitrogen prior to being gold-sputtered. A typical SEM micrograph is shown in
Fig. 9. The structure is obviously porous, and is very similar to those obtained with
solutions of PVdF in pure propylene carbonate. This is a sign that phase separation
still occurs even when lithium perchlorate is added to the system.
29
Figure 9: SEM micrograph of the cross section of a film prepared by phase-separationof
a PVdF + liquid electrolyte solution.
3.2
3.2.1
Fiber drawing and characterization
Preform preparation and fiber drawing
Preforms were prepared in the same fashion as previously. COC rods were machined
to make a 9" long cylinder of 1.25" oustide diameter with a 5/16" hole drilled over
5" in length so as to keep a pure COC part of about 4" at the bottom of the preform.
A 20:80 weight ratio of PVdF in a 1M solution of LiClO 4 in propylene carbonate
was prepared in a glovebox, and homogenized at 200'C for 3 hours.
The solution
was then poured into the core of the COC rod and the preform was sealed at its end
with a PTFE plug. The preform was then drawn in the draw tower following the
same precautions than detailed in Section 2.2.2. We obtained over 70 meters of fiber
from the first draw, with at least 50 meters containing a PVdF and liquid electrolyte
mixture in its core.
3.2.2
SEM characterization
Method
After this first successful draw, we prepared our fibers for SEM imaging.
The objective was to confirm the presence of pores and thus of a phase separation
during the draw process, but also to probe the effect of the presence of lithium salt
in the system. To do this, we prepared samples in two different ways. Some samples
were expelled from the core using a thin steel wire, and rinsed with ethanol so as to
remove any trace of propylene carbonate or lithium salt. They were then dried under
vacuum for 48 hours. The other samples were simply directly put under vacuum for a
30
Figure 10: SEM micrographs of the fiber cross sections. a) Sample prepared by core extraction, ethanol leaching and drying. b) Higher magnification view of core
displaying "rice-ball" structure due to LiClO4 precipitation on the polymer
spherulites. c) Sample prepared by vacuum drying for 48 hours. d) Higher
magnification view of core displaying spherulitic structure.
longer period of time to completely dry out the propylene carbonate. In both cases,
samples were freeze-fractured after the drying step, and gold-sputtered to enable
conduction of the electrons in the SEM. The micrographs are shown in Fig. 10.
Results and discussion
The first obvious observation is that the structures are
porous, and composed of interconnected polymer spherulites, as was the case previously in the absence of lithium perchlorate. We can thus conclude that qualitatively,
the presence of lithium perchlorate does not inhibit the phase separation. However
we note a major difference in microstructure between the samples rinsed and the
ones that were simply dried.
On the one hand, in the dried samples the lithium
perchlorate has precipitated onto the polymer spherulites so as to form a kind of
31
"rice-ball" structure. On the other hand, the samples that were rinsed show no
sign of crystallites and the microstructure closely resembles the one obtained previously. This was to be expected as ethanol is miscible with propylene carbonate and
also dissolves lithium perchlorate. The rinsing step thus sweeps away any trace of
LiClO 4 . This rice-ball microstructure could potentially be interesting to achieve large
surface-to-volume objects of different crystalline solids precipitated onto a polymer
scaffold.
3.2.3
Ionic conductivity of the fibers
Method We proceeded to measure the ionic conductivity of our fibers, by the
AC impedance spectroscopy method [30]. In this method, a test cell is fabricated
by connecting an electrolyte of known dimensions to two metal electrodes. The
cell is than run into an impedance analyzer to measure the impedance over a wide
frequency range, and the results are generally presented in the form of a Nyquist plot
(or complex-plane plot). One than maps these results onto a physically-motivated
equivalent circuit in order to extract the electrolyte bulk resistance and thus the
conductivity.
In our case, we have achieved this test cell by cutting samples of fibers at various
lengths, and measuring their diameter with an optical microscope. In parallel we
polished the ends of two copper wires, that we then plugged into the core of the fiber
at both ends of the samples to contact the core, as seen in Fig. 11.b). A simple
equivalent circuit for this system (based on [30]) is shown on Fig. 11.c).
Equivalent circuit analysis The resistor Rb corresponds to the bulk resistance to
the lithium ions' migration in the electrolyte. The capacitance Cb corresponds to the
dielectric polarization of the electrolyte and in a first approximation can be expressed
as COb= cEoA/l, where E is the dielectric constant of the polymer/solvent composite,
Eo the vacuum permittivity, and A and 1 are the area and length of the electrolyte.
Because the migration resistance and polarization physically happen in parallel, it's
sound to place these elements in parallel in the equivalent circuit. The capacitances
C, correspond to the electrical double-layers at the electrode-electrolyte interfaces.
Indeed, during each half-cycle of the alternating field, there is an ionic charge buildup
in the electrolyte close the electrode compensated by an equal and opposite charge
on the electrode, and associated with a second diffusive layer of charge further away
32
x 10
14
a)
12
0
length = 3 mm
length = 10 mm
X
length = 16 mm
0
length = 28 mm
A
length = 58 mm
C)
Con
conmta
R
10
N
Cbulk
d)
-z30
Rb
0
2
4
6
8
10
Z' (in 9)
12
Z'
14
X 10r
Figure 11: a) Experimental Nyquist plots for fiber samples of different lengths containing
phase-separated electrolyte. b) Picture of the measurement setup. c) Model
equivalent circuit of the test cell in picture above. d) Theoretical ideal Nyquist
plot for equivalent circuit.
in the electrolyte - thus creating an electrical double-layer acting like a capacitor.
Because of the small sizes of the electrical double-layers, we expect that C, > C.
The complex impedance of the above circuit can be written:
______wRbCb
- j
R
2
t)2C)
2
+
Z = Z' + jZ"= Rb[
W Ce
At high frequencies, because Cc > Cb, the behavior is dominated by the RbCb combination in parallel. Thus, in the complex plane we expect a semicircle (characteristic
of RC parallel impedances) at high frequencies, centered on [Rb/2, 0] and with radius Rb/2. At low frequencies, the impedance 1/jCew of the contact capacitances
will dominate over Rb and thus the bulk capacitance makes no contribution to the
impedance. The circuit behaves like a RbCe combination in series and should lead
to a straight vertical line in the complex plane, located at Z' = Rb. A schematic
representation of the expected theoretical Nyquist plot is given in Fig. 11.d).
33
Bulk resistance
(in kQ)
5.2 101
1.6 102
3.0 102
5.4 102
1.1 103
Conductivity
(in mS.cm- 1
2.2
2.3
2.0
2.0
1.9
)
Sample length
(in mm)
3.0
10
16
28
58
Table 1: Values of measured bulk resistance and computed conductivity for different fiber
sample lengths, with a core diameter of 580 pm
Results and discussion The test cell shown in Fig. 11.b) was run into an
impedance analyzer (Solartron 1287A, Solartron Analytical) between 1Hz and 1Mhz,
for different fiber lengths. The resulting Nyquist plots (Z" as a function of Z') are
shown in Fig. 11.a).
One first notices a relatively good qualitative correspondence between the experimental graphs and the theoretical plot. Particulary, the bulk resistance of the
electrolyte, Rb, scales linearly with the fiber length. From those values, and knowing
that the core radius for these samples was 290 pm, we can compute conductivities
of the samples, as reported in Tab. 1. The conductivities do not vary much with the
fiber length, which is expected. The average computed value is 2.1 mS.cm 1 , which
is a reasonable value, comparable to similar systems prepared in film forms [31] and
suitable for a battery application. The variations from this average value are within
experimental error.
Nonetheless we notice a discrepancy from the model at high frequencies, in the
presence of an inductance-like semicircle (close to the origin). This inductive behavior
at high frequencies seems not to depend on sample size and we believe it might be due
to cabling - however further experiments should be conducted. Another unexpected
behavior is that at low frequencies. Indeed, we observe a finite slope feature whereas
we expected a straight vertical line in the Nyquist plane. Such a behavior has been
reported and is a classical discrepancy from the simple model presented previously,
generally attributed to contact effects and added resistance between electrolyte and
electrodes [30]. It does not affect the reading of the bulk resistance Rb.
34
3.3
Conclusions
Even though promising, the porous microstructures obtained in Section 2 were not
functional as gel-polymer electrolytes as such, and it would have taken substantial
post-processing to make them so. Instead, we have devised a single-step process we
refer to as the TBOS process, upon which fibers exiting the draw tower contain a gelpolymer electrolyte core with a high conductivity. To achieve this, we simply added
lithium perchlorate to the solution of PVdF and propylene carbonate. Experiments
in films showed that the lithium salt seemed not to have a substantial effect on the
phase separation. This was confirmed in the fibers we made, as we observed both
porosity and high ionic conductivity, suitable for a battery application.
35
4
Pathway towards a Lithium-ion fiber battery
In the previous sections we have shown how to incorporate an ionically conductive gelpolymer electrolyte in a fiber core with conductivity suitable for a battery application.
The next step towards a functional lithium-ion battery in a fiber is thus to fabricate
and incorporate electrodes in the fiber, adjacent to the core. Due to the rheological
constraints of the drawing process, we have turned to polymer compositing as a
fabrication scheme. In this section we detail the fabrication process as well as the
preform-level characterization of the electrodes, and we give the pathway towards
incorporating the electrodes in a preform and drawing the final structure.
4.1
Electrode fabrication and characterization
4.1.1
Fabrication method
Materials selection Before even choosing a processing method, we first needed to
choose the proper materials to work with, ie. both the polymer host and the active
materials for cathode and anode.
For the polymer binder, because we intended to eventually draw the composite
electrodes adjacent to the electrolyte-filled core, we needed a polymer that:
- is resistant to propylene carbonate, the solvent used for the phase separation
- has a rheological behavior close to that of the cladding in order for the two to
to co-flow
- has a large electrochemical window so as be stable during the battery's operation
We chose to use COC as a way to get by the first two constraints. However no
data exists on the use of COC as an electrode binder material and on its potential
window - therefore experiments are needed to assess the electrochemical behavior of
this material.
Concerning the anode and cathode active material, we went for well-established
materials. As the anode active material we chose graphite, and as the cathode active
material we chose lithium cobalt dioxide LiCoO 2 . Both exhibit suitable thermal
stability and are reasonably safe to work with.
36
LiCoO 2
-
Carbon black
4%
COC
96%
Conductivity
Non conductive
36%
10%
14%
20%
9%
90%
86%
80%
55%
3.0 S.m- 1
8.0 S.m~
30 S.m-'
13 S.m-
Table 2: Average values of measured electrical conductivity for films of different compositions, in weight ratios.
Processing method The final goal being to achieve submicron particles polymer
composites, we first started by grinding the LiCoO 2 powder (Alfa Aeasar) using a
ball-mill to reduce the inital particle size which was on the order of 100 /im. The
carbon black particles (Alfa Aesar) were already below micron sized upon reception.
We fabricated electrode films by a solution-casting method. We dissolved COC
(T9 =1360 C, Polysciences) in toluene in a 10:90 weight ratio in a capped reaction
bottle, and stirred on a hot plate at 60'C for several hours until a homogeneous
solution was obtained. We then introduced a certain amount of carbon black for the
anodes, and carbon black and LiCoO 2 for the cathodes, and stirred at high shear
rates for several hours. The resulting slurry was then poured onto a cleaned glass
plate and spread evenly using a doctor blade. The films were placed in an oven at
60*C at ambient pressure for 4 hours, and for 20 additional hours under vacuum.
We thus obtained self-standing films for a wide range of concentrations.
4.1.2
Electrical conductivity
Method Self-standing anode and cathode films were tested for their electrical conductivity. The method went as follows. Samples of precise geometry and known
compositions were cut out of larger films, and the thickness was measured using a
micrometer. We then applied copper tape on both ends of the samples and contacted them with metal probes in order to obtain the U-I curve using a home-made
setup. A typical outcome of the measurements is shown in Fig. 12. Subsequently,
we extracted the values of the sample's resistance and thus the conductivity. Values
are reported in Tab. 2.
Results and discussion Electrical conduction is uniquely due to carbon black,
because both COC and LiCoO 2 are insulators. The mechanism for conduction is
37
R6=086
40h
Rf
R
(58nA
.M
20
o~o
--
-
Ar
o~os
oi11
0.005
0~A
0.01
/
1oo-F-
-0.01
-0__.005
0
0.015
Figure 12: Typical U-I curve obtained from resistance measurements on composite electrodes, with a nearly perfect linear behavior. The inset shows a photograph of
the measurement setup.
essentially percolation of the carbon black particle network [32]. There should be no
conduction for carbon content lower than the percolation threshold, and conduction
above it.
Above the threshold, conductivity scales with the carbon black volume
fraction. This is indeed what we observe from the measurements. Specifically samples
with 4% carbon for 96% COC in weight were absolutely not conductive, whereas we
measured a finite conductivity for samples containing 9% carbon. We also see that
carbon concentrations around and above 10% in weight are sufficient to ensure a
high-enough electrical conductivity of the electrodes.
4.1.3
SEM imaging
Method
We proceeded to imaging our samples using a SEM in order to assess the
good dispersion of active material, but also qualitatively probe the porosity of the
fabricated structures.
To do so, we prepared samples by freeze-fracturing them in
liquid nitrogen to preserve a clean cross section, and gold-sputtered them to ensure
high electrical conduction. Typical cross sections are shown in Fig. 13.
38
Figure 13: a) Cross-section a composite cathode with composition 36:9:55 weight ratio
in LiCoO2 /C/COC. b) High magnification image of composite cathode crosssection in a region containing non LiCoO2 particles. c) EDS analysis of a
LiCoO2 particle, with proof of cobalt content. d) Cross-section images of
a composite anode with composition 10:90 weight ratio in C/COC. e) High
magnification image of composite anode cross-section.
39
Results and discussion Concerning the cathode (cf. Fig. 13.a., b., and c.), we
confirm the presence of LiCoO 2 by performing an EDS analysis (cf. Fig. 13.c.). We
qualitatively see that the LiCoO 2 particles are well dispersed, with only few large
aggregates observed over the different samples that we imaged. The particle size
goes from a few microns to tens of microns, thus hinting that additional grinding
might be useful. Imaging regions devoid of active material particle reveals a porouslike structure composed of interconnected polymer filaments. This hints at a possible
phase separation which might occur upon removal of toluene from the solutions. The
polymer morphology (exhibiting no spherulite or related structure) is typical of an
amorphous polymer [21] - which is the case of COC. Further DSC experiments to
quantify this transition should be carried out. Carbon black agglomerates on the
order of 100 nm can be observed. Anodes (cf. Fig. 13.d. and e.) as well display a
qualitatively good dispersion of carbon black particles.
4.2
Battery assembly and testing in a coin-cell
Before incorporating the electrodes in a preform and drawing the battery structure,
we wanted to take a systems perspective. The objective was to validate whether or
not our proposed battery was functional, as well as test the electrochemical stability
of the COC as binder material.
4.2.1
Method
To do this we assembled a coin-cell battery in a glovebox by stacking and crimping
lithium metal as the anode, our TBOS electrolyte, and our composite LiCoO 2 /C/COC
of composition 36:9:55 weight ratio as the cathode, in a CR 2016 coin-cell case. The
reason we used a lithium metal anode was to decouple the uncertainty of a disfunctional anode or cathode, by focusing only on our home-made cathode. We then took
the coin-cell batteries to a battery cycler and performed 100 cycles with C/10 rate.
We deduced the corresponding current needed by measuring the weight in active
material in the cathode films. We obtained cycling curves similar to that presented
in Fig. 14.
40
Charge
Discharge
4.5-
4
3.5
3
-
2.5
2-
0
2000
4
6000
8000
10000
12
Time (s)
Figure 14: Voltage as a function of time for the C/10 cycling experiment on a
Li/TBOS/Composite LCO coin-cell. The positive slope regions correspond
to charging and negative slopes to discharging. Each peak corresponds to one
cycle.
4.2.2
Results and discussion
We see that the assembled systems do indeed behave as rechargeable batteries, however stop functioning after a relatively small number of cycles, close to 6.
The
capacity on the first cycle was measured to be only 7 mAh, which is close to 20
times smaller than the theoretical capacity based on active material content. This
discrepancy can be due to a non-percolating LiCoO 2 network, causing the lithium
intercalation to occur only in the particles at the electrode/electrolyte interface. Furthermore, the electrode exhibits a low-porosity based on the SEM imaging, and this
too might limit the amount of active material accessible to the lithium ions. The
extremely rapid capacity fade can also be accounted for by these same remarks.
Indeed, if lithium always intercalate in the same particles, it can induce structural
changes which will limit the capacity [33].
As a conclusion, our cathode and electrolyte were functional and allowed for
a Li-ion battery assembly.
The resulting batteries nonetheless performed rather
poorly, with a low initial capacity and rapid capacity-fade.
41
However there is room
for improvement through higher LiCoO 2 content or increased porosity.
4.3
Next steps: preform preparation and fiber drawing
The next step is to co-draw the electrodes and electrolyte in a COC cladding. For
this, we propose a preform fabrication method in multiple steps, illustrated in Fig.
15.
Start from
COC rod
Machine
hole and
slots
Insert
composite
anode and
cathode
Roll COC films
around
preform and
insert current
collectors
Consolidate
Figure 15: Li-ion fiber battery preform fabrication plan.
We propose to start from a pure COC rod in which we will drill a hole to contain
the gel-polymer electrolyte, as well as machine two slots along the length of the
rod. We will then insert our composite anode and cathode of precise geometry into
those slots, and wrap COC film around the structure. After a few layers we will
consolidate the structure to fuse the materials together, and machine two small slots
adjacent to the electrodes that will contain our current collectors. We propose to use
the eutectic alloy Bi5 2Sn 48 which is a good metal conductor with a melting point of
138 C (as reported by the manufacturer, Indium Corporation). We will then proceed
to wrapping more COC film around the preform and consolidate one last time.
This method would lead to electrodes adjacent to the core of the preform, and
current collectors in contact with the electrodes. We would then undergo the draw
process at high stress to maintain the shape of the electrodes and current collector
and limit bulging or flowing of the electrodes inside the core.
4.4
Conclusions
In this section we have presented results on the fabrication and characterization of
composite anodes and cathodes. Our electrodes are made or COC as a binder ma42
terial and should not be damaged by propylene carbonate during the draw process.
Furthermore the electrodes exhibit a high enough electrical conductivity so as to be
used in batteries. Electrochemical testing of coin-cell batteries made of composite
cathodes, TBOS electrolyte and lithium metal anodes validated the use of the cathode and electrolyte as elements of a battery. The next steps are obviously to include
these electrodes in preforms and attempt a draw of the final structure.
43
Conclusion
A Li-ion fiber battery could be an extremely compelling device because of its ease of
integration into large-area conformal and flexible electronic systems, scalable production, and the ability to distribute energy sources rather than localize them. However
fabricating such a device comes with a number of challenges, the main one being the
issue of incorporation of a porous polymeric domain inside a fiber as host material
for a gel-polymer electrolyte.
In this thesis we presented a process that enables the incorporation of porous
domains inside fibers through the phase separation of a polymer solution. This a
very versatile method which allows to fabricate porous structures in geometries not
previously possible, such as a porous core/dense shell fiber structures. These porous
domains may furthermore be functionalized. In particular we have shown how the
process can be tuned to produce ionically conductive gel-polymer electrolytes inside
fibers, with high ionic conductivities, thus paving the way towards a functional Li-ion
battery in a fiber. This achievement also relies on composite electrodes with suitable
properties, which we have fabricated by a solution casting method. The next - and
potential final - step would be to codraw these different elements into a fiber battery,
with electrodes adjacent to the electrolyte core.
44
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