ACTAUNIVERSITATIS

UPSALIENSISUPPSALA

2018

Digital Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology 1630

Block Copolymer Electrolytes

Polymers for Solid-State Lithium Batteries

ANDREAS BERGFELT

ISSN 1651-6214ISBN 978-91-513-0233-1urn:nbn:se:uu:diva-340856

Dissertation presented at Uppsala University to be publicly examined in Häggsalen,Ångströmlaboratoriet, Uppsala, Friday, 23 March 2018 at 09:00 for the degree of Doctorof Philosophy. The examination will be conducted in English. Faculty examiner: AssociateProfessor Yoichi Tominaga (Tokyo University of Agriculture & Technology).

AbstractBergfelt, A. 2018. Block Copolymer Electrolytes. Polymers for Solid-State Lithium Batteries.Digital Comprehensive Summaries of Uppsala Dissertations from the Faculty of Science andTechnology 1630. 68 pp. Uppsala: Acta Universitatis Upsaliensis. ISBN 978-91-513-0233-1.

The use of solid polymer electrolytes (SPEs) for lithium battery devices is a rapidly growingresearch area. The liquid electrolytes that are used today are inflammable and harmful towardsthe battery components. The adoption of SPEs could drastically improve this situation, but theystill suffer from a too low performance at ambient temperatures for most practical applications.However, by increasing the operating temperature to between 60 °C and 90 °C, the electrolyteperformance can be drastically increased. The drawback of this approach, partly, is that parasiticside reactions become noticeable at these elevated temperatures, thus affecting battery lifetimeand performance. Furthermore, the ionically conductive polymer loses its mechanical integrity,thus triggering a need for an external separator in the battery device.

One way of combining both mechanical properties and electrochemical performance is todesign block copolymer (BCP) electrolytes, that is, polymers that are tailored to combine oneionic conductive block with a mechanical block, into one polymer. The hypothesis is thatthe BCP electrolytes should self-assemble into well-defined microphase separated regions inorder to maximize the block properties. By varying monomer composition and structure of theBCP, it is possible to design electrolytes with different battery device performance. In Paper Iand Paper II two types of methacrylate-based triblock copolymers with different mechanicalblocks were synthesized, in order to evaluate morphology, electrochemical performance,and battery performance. In Paper III and Paper IV a different strategy was adopted,with a focus on diblock copolymers. In this strategy, the ethylene oxide was replaced bypoly(e-caprolactone) and poly(trimethylene carbonate) as the lithium-ion dissolving group.The investigated mechanical blocks in these studies were poly(benzyl methacrylate) andpolystyrene. The battery performance for these electrolytes was superior to the methacrylate-based battery devices, thus resulting in stable battery cycling at 40 °C and 30 °C.

Andreas Bergfelt, Department of Chemistry - Ångström, Polymer Chemistry, Box 538,Uppsala University, SE-751 21 Uppsala, Sweden.

© Andreas Bergfelt 2018

ISSN 1651-6214ISBN 978-91-513-0233-1urn:nbn:se:uu:diva-340856 (http://urn.kb.se/resolve?urn=urn:nbn:se:uu:diva-340856)

To Martina and my family

List of Papers

This thesis is based on the following papers, which are referred to in the text by their Roman numerals.

I d8-Poly(methyl methacrylate)-Poly[(oligo ethylene glycol)

methyl ether methacrylate] Tri-block Copolymer Electro-lytes: Morphology, Conductivity and Battery Performance A. Bergfelt, L. Rubatat, R. Mogensen, D. Brandell and T. Bow-den Polymer 131 (2017) 234 - 242

II Poly(benzyl methacrylate)-Poly[(oligo ethylene glycol) me-thyl ether methacrylate] Triblock Copolymers as Solid Elec-trolyte for Lithium Batteries A. Bergfelt, L. Rubatat, D. Brandell and T. Bowden Submitted

III εε-Caprolactone-based Solid Polymer Electrolytes for Lith-

ium-Ion Batteries: Synthesis, Electrochemical Characteriza-tion and Mechanical Stabilization by Block Copolymeriza-tion A. Bergfelt, M. J. Lacey, J. Hedman, Christofer Sångeland, D. Brandell and T. Bowden Submitted

IV A Mechanical Robust yet highly Conductive Di-block Copol-

ymer based Solid Polymer Electrolyte for Room Tempera-ture Battery Applications A. Bergfelt, R. Mogensen, M. J. Lacey, D. Brandell and T. Bow-den Manuscript

Publications not included in this thesis, but with general interest to the reader:

V Graft copolymer electrolytes for high temperature Li-battery applications, using poly(methyl methacrylate) grafted poly(ethylene glycol)methyl ether methacrylate and lithium bis(trifluoromethanesulfonimide) M. Bergman, A. Bergfelt, B. Sun, T. Bowden, D. Brandell and P. Johansson Electrochimica Acta, 175 (2015) 96 - 103

VI A Robust Water-Based, Functional Binder Framework for High-Energy Lithium-Sulfur Batteries M. J. Lacey, V. Österlund, A. Bergfelt, F. Jeschull, T. Bowden and D. Brandell ChemSusChem, 10 (2017) 2758 – 2766

VII A Self-Assembly of Cholesterol End-Capped Micelles for Drug Release M. Gao, A. Bergfelt, T. E. Larsson and T. Bowden Manuscript

Reprints were made with permission from the respective publishers.

Contents

Scope of the thesis .................................................................................... 11

Introduction ............................................................................................... 13Lithium batteries .................................................................................. 13Block copolymers ................................................................................. 16

Atom transfer radical polymerization .............................................. 18Ring-opening polymerization .......................................................... 19

Solid polymer electrolytes ........................................................................ 20Ionic conductivity ............................................................................ 21Lithium-ion transference number .................................................... 23

Block copolymer electrolytes ................................................................... 24Triblock copolymers ............................................................................ 24Diblock copolymers ............................................................................. 26Polyesters and polycarbonates ............................................................. 29Summary .............................................................................................. 31

Results and discussion .............................................................................. 32Triblock copolymers ............................................................................ 32

Thermal properties and morphology ............................................... 33Ionic conductivity ............................................................................ 37Electrochemical performance .......................................................... 38Conclusion triblock copolymers ...................................................... 40

Polyester-based electrolytes ................................................................. 41∂-Valerolactone-based electrolytes .................................................. 41

ε-Caprolactone-based electrolytes ........................................................ 43Synthesis .......................................................................................... 43Ionic conductivity ............................................................................ 43Thermal properties and morphology ............................................... 48Mechanical properties ...................................................................... 48Diblock copolymer via controlled synthesis ................................... 49Electrochemical performance .......................................................... 52Conclusion diblock copolymers ...................................................... 55

Conclusion ................................................................................................ 56

Svensk sammanfattning ............................................................................ 58

Acknowledgements ................................................................................... 60

References ................................................................................................. 61

Abbreviations

AFM ATRP BCP BCT CV DMA DMF DSC

Atomic force microscopy Atom transfer radical polymerization Block copolymer Poly(benzyl methacrylate)-b-poly(ε-caprolactone-r-trimethylene car-bonate) Cyclic voltammetry Dynamic mechanical analysis Dimethylformamide Differential scanning calorimetry

GPC LiTFSI LFP PBnMA PCL PDI PEO PMMA POEGMA PS PTMC ROP SANS SAXS SC SCT SEI SPE TMC VTF

Gel permeation chromatography Lithium bis(trifluoromethanesul-fonyl)imide LiFePO4 Poly(benzyl methacrylate) Poly(ε-caprolactone) Poly dispersity index Poly(ethylene oxide) Poly(methyl methacrylate) Poly([oligo ethylene methyl ether]methyl methacrylate) Polystyrene Poly(trimethylene carbonate) Ring-opening polymerization Small-angle neutron scattering Small-angle X-ray scattering Poly(styrene)-b-poly(ε-caprolactone) Poly(styrene)-b-poly(ε-caprolactone-r-trimethylene carbonate) Solid electrolyte interphase Solid polymer electrolyte Trimethylene carbonate Vogel-Tammann-Fulcher

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Scope of the thesis

To properly characterize and evaluate solid polymer electrolytes (SPEs), sev-eral points of view have to be adopted. The general requirements for a SPE include high ionic conductivity, lithium transport number close to 1, high me-chanical integrity, and electrochemical stability over a broad voltage and tem-perature window. The hypothesis is that block copolymer (BCP) electrolytes have the ability to combine ionic conductivity and mechanical properties when phase separated into well-defined microstructures. By using different block constituents, the electrochemical stability and mechanical properties could be altered.

The overall objective of this thesis lies on the synthesis of BCPs; the for-mulation of SPEs; the study of their morphology, mechanical properties, and electrochemical properties; the analysis of their performance in solid-state lithium battery devices. To do this methacrylate-based triblocks and polyester-based diblocks have been synthesized by atom transfer radical polymerization and ring-opening polymerization. In an attempt to narrow the topic and make it possible to compare the results, the electrolytes have been prepared with LiTFSI as salt, and the cell design throughout the studies is Li | SPE | LiFePO4, that is, half-cells utilizing lithium metal as the anode and LiFePO4 as the active material in the cathode.

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13

Introduction

Lithium batteries A lithium (or lithium-ion) battery is made up of several components. Its main components are the anode (negative electrode), cathode (positive electrode), and the electrolyte (see Figure 1). It also has important supporting compo-nents, such as the current collectors, which are usually copper for the anode and aluminum for the cathode side. [1, 2] When a liquid or gel electrolyte is used, a separator is needed to prevent the electrodes from coming into contact and causing a short circuit. A casing is also needed to prevent the electrolyte to leak out. When a battery is discharging, electrons flow from the negative electrode to the device we want to power, and then returns to the cathode of the battery. While this is happening, ions are moving in the electrolyte be-tween the electrodes, thus closing the circuit. This reaction could be exempli-fied for graphite- and LiFePO4-based batteries by a redox scheme, which com-bines the electrodes reduction and oxidation:

Negative electrode: LiC6 C6 + Li+ + e- Positive electrode: FeIIIPO4 + Li+ + e- LiFeIIPO4 Cell reaction: LiC6 + FeIIIPO4 C6 + LiFeIIPO4

This reaction describes a so-called lithium-ion full cell. When the anodes are made of solid lithium, the cells are called lithium half-cells. Depending on the material that is used in the electrodes, the battery cell will have different ca-pabilities to achieve a cell potential, and the higher the cell potential - or the higher the specific capacity of the active material - the higher the energy out-put. One of the most commonly used anode materials in commercial cells is graphite, which has a theoretical capacity of 372 mAh g-1, whereas lithium has a specific capacity of 3,680 mAh g-1. [3, 4] The main reason for using graphite anodes is to avoid lithium dendrite growth, which will penetrate the cell mem-brane, thus causing thermal runaway that may lead to fires and explosions as a result. This sequence of events can be prevented by using graphite anodes wherein lithium ions are intercalated into the graphite structure, thus resulting in the confinement of lithium ions in a non-metal state, hence the term “lith-ium-ion battery” (LIB). [5] The LIBs used today utilizes a liquid electrolyte, which will generate a high diffusion rate of the ions; the rate determines the

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speed in which energy is released from the battery. These electrolytes are typ-ically made up of organic solvents such as ethylene carbonate, diethyl car-bonate, dimethyl carbonate, and propylene carbonate. [6, 7] As stated above, these solvents are rather volatile and inflammable. [8] Cycling within the elec-trochemical stability window of the electrolyte is important, otherwise the electrolyte will be reduced or oxidized, thus possibly leading to a clogging of the separator membrane pores causing a negative effect on the battery device. [8, 9] However, the electrolyte degradation could actually be beneficial if it forms a robust solid electrolyte interphase (SEI) on the anode. [10] The SEI will protect the electrolyte and anode from further degradation, and at the same time shuffle ions from electrode to electrolyte. However, the volume expansion of the electrode during cycling will generate cracks in the SEI,

Figure 1. Illustration of a battery device during discharge, where the electrons move from the negative anode to the positive cathode while powering a device. During the same time, lithium-ions are moving from the electrolyte to the cathode to close the cir-cuit.

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which will open up new sites that will react with the electrolyte. In this way, the amount of lithium will be consumed over time, thus resulting in a gradual decrease in capacity utilization. [7] Considerable research and engineering ef-forts have been dedicated to optimizing each component, such as cell design, control of particle size of active materials, and LIB managing systems, to max-imize the capacity output of the LIB. [11-18] In this way, the components have been designed such that it today is possible to power cell phones, laptops, and even electrical vehicles. However, there is a demand for batteries with higher energy and power outputs to enable the operation and increase the driving range of electrical cars, busses, and heavy vehicles, such as trucks and heavy-duty construction vehicles. Thus, research has focused on designing new types of electrode materials that could incorporate more capacity and generate higher energy density by increasing the cell voltage.

The electrolyte is an essential part of a battery and needs to be the focus of study to further the development of the lithium batteries. By using a solid pol-ymer electrolyte (SPE), several of the drawbacks associated with liquid elec-trolytes could be overcome. [19, 20]. The benefits of SPEs are potentially sev-eral: a solid electrolyte would be less flammable and less aggressive against the battery components. [1, 20, 21] In addition to safety improvements, the use of SPEs can enable the production of a denser battery stack because no separator or casing would be necessary to keep the liquid electrolyte from leaking out of the battery pack, thereby generating higher energy density. [22] The requirements for a competitive SPE include high ionic conductivity, pos-itive transference number close to 1, thermodynamical and/or kinetical stabil-ity towards electrodes, and high moduli. [23] A major challenge in developing such an SPE is that polymers with high moduli are generally poor ion conduc-tors at room temperature. One of the most used and studied SPE host polymer is poly(ethylene glycol) (PEO), which renders thermally and electrochemi-cally stable electrolytes. It is characterized by its low glass transition temper-ature (Tg) (ca. -30 °C), electrolytes with comparatively high ion conductivities (ca. 10-4 S cm-1 at elevated temperatures), and good ability to dissolve lithium salts. [23] The drawbacks of PEO are that its ion conductivity is less satisfying than its mechanical properties and that its transport number is low (ca. 0.1 – 0.3). [24] At high molecular weights, PEO shows mechanical integrity but becomes semicrystalline below 60 °C, thus negatively affecting the ionic con-ductivity. This is generally true for most systems: crystallinity lowers the seg-mental motion of the chains, thus lowering the ionic conductivity. However, a too low modulus will allow for lithium dendrite growth and will be a safety concern in the long run. [25, 26]

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Block copolymers Engineered polymers find their advantage in several areas, like batteries, sur-factants, semiconductors, solar cells and many more. Precision and perfection in polymer synthesis is important when tailoring the properties of the resulting polymeric material. This means that there are several interesting characteristic properties of macromolecules that are important. These properties include the length, structure, and size of the polymer, backbone and side-group motion and functionality; supra molecular structure; molecular conformation and or-ganization at interfaces; morphology, crystallinity and self-assembly, just to mention some. The fabrication of polymers has for many decades been per-formed by techniques such as free radical polymerization, which provide lim-ited possibilities for BCP design, but in 1956 living polymerization was made possible by Szwarc, making it possible to polymerize different types of alkene monomers and cyclic monomers. [27] If performed under the right conditions no termination or chain transfer will occur, making it possible to polymerize BCPs with well-defined block lengths. The main drawback with this synthesis method is the sensitivity towards impurities. However, given the recent pro-gress made in the field of polymer chemistry, a useful toolbox has been devel-oped, thus broadening the design possibilities with the use of robust polymer-ization techniques, such as atom transfer radical polymerization (ATRP), re-versible addition/fragmentation chain-transfer (RAFT) polymerization, ni-troxide-mediated polymerization (NMP), and different types of controlled ring-opening-polymerization (ROP). With these techniques a broad range of molecular architectures and compositions have been made available, making it possible to design linear multi blocks with arbitrary block sequences, star and graft BCPs, bottle brush polymers and many more. These controlled meth-ods makes it possible access defined molecular weights and low polydisper-sity indexes (PDIs) (i.e., <1.1), and at the same time control the topology, functionality and end-group fidelity. The end group fidelity is highly im-portant, since it enables that the just prepared polymer could be used as a ma-croinitiator for growing a new block. Although not there yet, the toolkits pro-vided today is narrowing the gap between natural and synthetic polymers. [28-31] In the current study, primarily ATRP and ROP have been used to synthe-size the SPE host materials.

With the access and advance in controlled polymerization techniques it is possible to target the topics precision and perfection in polymer chemistry, but what is also important is prediction, where morphology and macromolecular properties are of interest. [32, 33] Generally, the more advanced the architec-ture and composition is, the harder it is to predict morphology and polymer properties.

The theoretical thermodynamical foundation for BCPs was made by Ka-wai, Hashimoto and coworkers, and Meier in the 1950-1960s and later on by Helfand in the 1970s. [34-36] In 1980 Leibler published a paper based on the

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random phase approximation, which explained the laws for diblock copoly-mers. [37] A diblock copolymer is a sequence of monomer A coupled with a sequence of monomer B. Polymers that are made up of different monomers usually have low compatibility and will not mix but will phase separate into domains of A and B. However, in a BCP, the blocks are chemically linked and will not have the possibility to macrophase separate; instead, they will mi-crophase separate. If the conditions are right, the BCPs could arrange them-selves such that they create a long-range order (see Figure 2). [38, 39] On a macro level, the material will act as one, however, but on a micro level, the material has the individual properties of the block components. [39] The mi-crophase separation into discrete domains depends on several parameters. The microphase separation for an AB diblock copolymer depends on the degree of polymerization of the blocks (NA, NB), the volume fractions of the components (φA, φB), and the Flory-Huggins interaction parameter χAB, which is a thermo-dynamical interaction parameter that will dictate the compatibility between A and B. [37] [38, 39]. A wide range of instruments and techniques has been developed in order to study these materials, such as small angle neutron- and X-ray scattering (bulk structure), differential scanning calorimetry (thermal properties), transition electron microscopy and atomic force microscopy (im-aging), just to mention some tools for polymer characterization.

Figure 2. Transmission electron microscopy (TEM) image showing a microphase sep-arated PBnMA-POEGMA-PBnMA BCP electrolyte from Paper II.

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Atom transfer radical polymerization ATRP was developed in 1994 by Sawamoto and co-workers and in 1995 by Matyjaszewski and co-workers, based on the principles of atom transfer radi-cal addition. [40-42] ATRP is based on the concept of controlling the equilib-rium between dormant and propagating radical species, where the key concept is to keep the number of propagating radicals as low as possible, in order to avoid any radical termination. By doing so, the growth time for a polymer chain is increased from about a second to several hours or days. The activation step consists of a halogenated initiator which is homolytically cleaved. The halogen is then transferred to a metal complex (usually copper with a nitrogen-based ligand). The initiator has to be halogenated, but can also bear desirable types of functionalities that could be used for later chemical manipulations. The metal complex will dictate the equilibrium of propagating and dormant species, and together with the monomer concentration and growing radicals, this will dictate the rate of propagation (see Figure 3). To achieve a high con-trol the initiation has to be fast; there should be no chain termination or chain transfer together with a high chain deactivation rate. By manipulating the con-centration of initiators, monomers, solvent, temperature, type of metal com-plex, etc., it is possible to find the right conditions for a controlled reaction. However, it should be said that there always will be a fraction of terminated chains, typically between 1 – 10%. When the polymer has been recovered it is possible to use it as a macroinitiator for further reactions; depending on monomer functionality, different possibilities for end-group fidelity and func-tionality can exist. There are several different types of ATRP techniques, such as ARGET-ATRP, SET-LRP, SARA-ATRP, eATRP and photo induced ATRP. These techniques all aim to minimize the amount of catalyst and ena-ble the production of BCPs from a wide variety of monomers with a high de-gree of control. [28]

Figure 3. ATRP reaction scheme. On the left-hand side is the dormant P-Br polymer. When activated by the catalyst complex the polymer starts propagating, attaching monomers to the polymer until it is deactivated by the Cu-complex.

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Ring-opening polymerization Several catalytic systems exist, and they are categorized according to the polymerization of cyclic carbonates, functionalized carbonates, cyclic esters, etc., via ROP. [43-45] Many of the monomers that could be polymerized are considered to be biodegradable and/or derived from natural products, thus making them desirable for use because the degradation products that are formed in some cases are harmless. [46, 47] Therefore, a lot of effort has been devoted to finding polymerization techniques that can control the synthesis of this broad family of monomers. [48] Given that the size of the ring, and the substituents, will affect the polymerization thermodynamics and kinetics, it could be hard to copolymerize a range of these monomers. Several different catalytic systems exist for ROP synthesis, but tin(II) 2-ethylhexanoate (SnOct2) is one of the most classically used organo-metallic catalysts because it is highly efficient, soluble in most commonly used solvents and low toxicity. However, the synthesis has to be performed at 110 °C to 130 °C for several days until the reaction is finished. [49] When using the state of the art organic catalyst systems of today, the reaction could be conducted at ambient condi-tions (30 °C) at a polymerization duration of a couple of hours. [44, 50] How-ever, as stated above, the choice of catalytic system will depend on the type of monomers that are used, which is the reason for why SnOct2 was used in Paper III and Paper IV. The SnOct2 reaction mechanism is debated, how-ever, it is believed that it follows a coordination-insertion mechanism, where the metal complex and the alcohol initiator in equilibrium forms a metal oxide alkoxide. The reaction then proceeds via coordination of the monomer and catalyst via bond rearrangements (see Figure 4). The metal alkoxide remains the reactive center throughout the reaction. [43, 51-53]

Figure 4. Preinitiation and coordination-insertion steps for the polymerization of ε-caprolactone with SnOct2 to form poly(ε-caprolactone).

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Solid polymer electrolytes

PEO was the first polymer to be used in LIBs, and has been the focus of re-search since about 1970, when it was discovered that PEO has the ability to dissociate lithium salts. [23] The PEO electrolyte can be considered an inter-mediate between a liquid electrolyte, where the ions are solvated and moves with a coordinating shell, and a solid electrolyte, where the ions move through a fixed crystal matrix. The ionic mobility in PEO is believed to be assisted by the segmental motion of the polymer chains and by ion hopping between the coordination sites in and between the PEO chains. [54] Therefore, it is im-portant that the electrolyte is well above its glass transition temperature (Tg), which is approximately -50 °C for pure PEO, and that the electrolyte is amor-phous, because crystallinity has a crucial effect on the ionic conductivity. However, considering that PEO is a non-branched and symmetrical polymer with ether oxygen in the repeating unit, PEO crystallizes to a high degree and has a melting temperature (Tm) of approximately 60 °C. Its ionic conductivity is several magnitudes in order lower than an equivalent but amorphous elec-trolyte. [55, 56] Depending on the composition and type of salt, complex phase diagrams could be found, thus indicating that the crystallization is rather diverse to its nature. [57, 58] From a battery device point of view, crystallinity is a disadvantage, because it eliminates any possibility for a room-tempera-ture-based PEO electrolyte. However, when the operating temperature in-creases above the Tm of the electrolyte (ca. 60 °C), the crystals melt, and the ionic conductivity increases; the drawbacks of this process include increased parasitic side reactions and decreased mechanical properties. Therefore, this is a rather complicated matter to solve, and several strategies have been eval-uated. One approach is the modification of the polymer structure by attaching pendant groups to the main chain. [59] Ion coordinating side groups are of special interest because they are dynamically detached from the main chain; therefore, they are less affected by the molecular weight and stiffness of the main chain. The major drawback with such an approach is the decrease in mechanical properties, with electrolyte resembling viscous liquids rather than solid materials. [24, 55, 60] Another approach is the use of fillers or different types of salts to lower the degree of crystallinity or to lower the melting point of the different phases. [61] A rather different approach is the hypothesis that ion conductivity could occur in the crystalline parts of the polymer. This hy-pothesis is based on a decoupling of the ion conductivity from the polymer segmental motion, which was realized in a crystalline PEO6:LiXF6 system,

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where X = P, As, Sb. The same system was later doped with N(SO2CF3)- ions, which improved the ionic conductivity and allowed the system to reach values of 10-6 S cm-1 at room temperature. [62-64] It has been speculated that the ion conductivity is confined within the chain fold regions, and that the Li+ are guided by the crystalline lamella structure. [65] Moreover, several studies have been conducted to understand the segmental motion and structural relax-ation time in SPEs. A hypothesis is that if the polymer is designed in such a way that it is loosely packed but still “frustrated”, the lithium ions will slip through the polymer matrix even if the structure is frozen. [66, 67] In an at-tempt to design such a structure, an oxanorborene polymer backbone with short PEO chains as pending side groups were synthesized and investigated as an SPE. The ionic conductivity was below that of PEO at room temperature, but the method of constructing a rigid and frustrated chain packing system lead to a decoupling of five to six order of magnitude improvement compared to a reference system. This finding showed that it is possible to decouple the ion mobility from the main chain motion. [68]

Ionic conductivity A fundamental property of SPEs is ionic conductivity, which is expressed as:

where is the concentration of charge carriers (Li+), is the charge of the charge carriers, and is the mobility of charge carriers, which is a function of temperature and pressure. [69] When dissolving a lithium salt into a poly-mer host, the salt will be dissociated by the ether oxygen in the repeating units in the PEO backbone. Typically, three to six ethylene oxides will coordinate a Li+ via ion coordination to the polymer (see Figure 5). [57, 70] When in-creasing the amount of free charge carriers, the ionic conductivity increases, but a common phenomenon is that the salt concentration versus ionic conduc-tivity reaches a peak value. Furthermore, increasing the salt concentration af-ter this point will only decrease the overall ionic conductivity. This is an out-come of the trade-off between the free charge-carrier concentration and the ion-ion interaction starts to become significant as the ion concentration in-creases, thus creating long-range interaction resulting in a more rigid system, which lowers the segmental motion of the polymer. Furthermore, the ion-ion interaction will result in ion paring, that is, the amount of free charge carriers will actually decrease. [71] Furthermore, the anion could have a dramatic im-pact on the salts coordination strength to the polymer, which will affect the segmental motion of the polymer chains, and in the end affect the ionic con-ductivity and lithium-ion transport number. [72, 73] The conclusion is that the

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polymer should allow for a high degree of ion dissociation and segmental mo-tion to reach a high ionic conductivity. Given that ion mobility relies on chain motion, decreasing the Tg of the SPE will increase the chain motion and the ionic conductivity. It is common to fit the ionic conductivity data from AC spectroscopy with the Vogel-Tammann-Fulcher (VTF) equation, which is an empirical relationship that holds for polymer electrolytes above their Tg, and is expressed as follows:

σ σσ

where σ is the ionic conductivity, Eσ is a pseudo-activation energy that typ-

ically is related to the segmental motion of the polymer chains, kB is the Boltz-mann constant, Aσ is a prefactor that is proportional to the number of carrier ions and To is a reference temperature that is usually associated with the ideal glass transition temperature at which the configurational entropy becomes zero (usually said to be ca 50 K below Tg). [74, 75] The VTF relationship is commonly observed for dry polymer electrolytes, and is based on the concept of free volume and on an ionic transport that is coupled with the segmental motion of the polymer chains. By fitting the equation to the data, it becomes possible to evaluate whether the electrolyte is amorphous, as well as under-stand the salt dissociation and segmental motion of the chains and how these correlate to the overall ionic conductivity. [76, 77] Polymer electrolytes are a special case wherein the ionic conductivity depends not only on the lithium-ion transport, but also on chain relaxation. This relationship is described by the following:

OO

O

O

O

O

OO

Li+

TFSI-

Figure 5. Schematic illustration of PEO:s coordination to LiTFSI.

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where is the decoupling index between the chain relaxation time and the ion conductivity relaxation time . These two properties are closely related and are hard to separate. The most commonly used approach to reach high ionic conductivity is to decrease the viscosity, that is, lower the Tg. However, as mentioned, when the viscosity decreases the mechanical properties are lost. This is one of the reasons for investigating the possibility of designing a new type of polymeric electrolyte with a high decoupling index.

Lithium-ion transference number Another important property is the lithium-ion transport number, t+. The mass transport in an electrolyte, that is, the motion of lithium ions from an electrode to the other, is said to occur in three ways: (1) by migration, where the ions move in response to an electric field, but where the total ion conduction is divided between different ions; (2) by diffusion, where species of all kinds move in response to a concentration gradient; and (3) by convection, which covers heat transfer and mechanical effects (this could be ignored for SPEs). Therefore, cell polarization will drive the reactions at the electrodes, and the ions will move in the electrolyte by migration and diffusion. In this context, the transport number t+ is the fraction of charge carried by the cationic species, and the corresponding anion transport number t- will add up such that t+ + t- = 1. This means that the transport number cannot have negative values. The transference number T+ is the number of moles of cationic species transferred per Faraday of charge, and T+ may be less than or more than 1. The difference between t+ and T+ is that in real electrolytes, clusters and neutral pairs may form at high salt concentrations, thus making T+ the more commonly reported value. There are some different methods for performing this analysis, but the commonly used method by Bruce and Vincent is easy and elegant and will therefore be explained. [78] A bias voltage of 10-20 mV is applied to a sym-metrical Li | SPE | Li cell, and the current response is followed until the current reaches steady state. An impedance spectrum is recorded both before and after the voltage step and T+ is calculated as follows:

where Iss is the current at steady state, I0 is the initial current, Rp,0 is the

initial resistance, Rp,ss is the resistance at steady state and ∆V is the voltage step. The key assumptions are low polarization potentials, negligible ion-ion interaction, and no consideration of ion clusters and ion pairs. [78-80]

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Block copolymer electrolytes

By combining a block that has the ability to solvate lithium salts with a block that has mechanical integrity and can microphase separate such as that like in Figure 6, it should be possible to create BCP electrolytes that combines the need for electrochemical and mechanical properties. [81] [20] A fully phase separated BCP electrolyte is believed to maximize the ionic conductivity and mechanical properties. If the blocks are mixed, the ionic conductivity and the mechanical properties will be lowered. As already stated, the microphase sep-aration for an AB diblock copolymer depends on the degree of polymerization of the blocks (NA, NB), the volume fractions of the components (φA, φB), and the Flory-Huggins interaction parameter χAB. However, when mixing salt and polymer, the thermodynamics of mixing will be complicated, thus hindering the ability to predict the microphase separation.

Triblock copolymers Gray et al. at St Andrews University, United Kingdom, conducted an early study in 1987 on the use of the BCP concept when synthesizing a styrene-butadiene-styrene triblock copolymer with pendant short-chain PEO grafted to the backbone. They found that the ionic conductivity did depended both on

Figure 6. Transmission electron microscopy (TEM) image of a microphase separated triblock copolymer electrolyte (Paper II), where the black stripes represents one block, and the white stripes represent the second block. The inset scale bar represents 1 µm.

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the casting technique, because it affected the degree of phase separation, and the amount of salt. The ionic conductivity reached high values but only at el-evated temperatures. [82, 83] In 2002, Jannasch synthesized a triblock copol-ymer with the general structure polyethylene-polyether-polyethylene. Differ-ential scanning calorimetry (DSC) analysis showed that the polymer was sem-icrystalline, with a high melting temperature (Tm above 100 °C), and a Tg of approximately -50 °C. The ionic conductivity reached 10-6 S cm-1 at room temperature, but no mechanical, morphological or battery testing was con-ducted on the material. [84] Many works have since been focused on using polystyrene (PS) as the mechanical block, which has a high Tg and is an amor-phous polymer, and a high modulus. Kanamura et al. investigated a comb triblock SPE with the structure PS-poly[(oligo ethylene glycol) methyl ether methacrylate] (POEGMA)-PS. The polymer was fully amorphous, and the ionic conductivity reached 10−4 S cm-1 at room temperature for weight frac-tions of POEGMA above 85%. A lithium battery operating with LiCoO2 as cathode material was cycled with a discharge capacity of 100 mAh g-1 at C/10 and 30 °C. TEM images clearly showed that the BCP electrolyte did mi-crophase separate, but the tensile strength was rather low at ~3 MPa. [85] This is far from the 6 GPa that has been calculated to stop dendrite formation, as proposed by Monroe and Newman. [25, 26] In 2003, Wang et al. grafted PEO to a PS block to form a PS-(PS-g-PEO)-PS triblock copolymer with a Mw of 340,000 g mol-1, where the PEO part made up 60 wt% of the polymer matrix. TEM imaging showed that the SPE microphase separated into a well-ordered pattern, with the PS phase forming long-range rods in a matrix of PS. The storage modulus E´ for the BCP electrolyte (20:1 EO:Li+) was about 0.1 GPa between 20 °C to 100 °C, as determined by dynamic mechanical analysis (DMA). The SPE was fully amorphous, and the ionic conductivity was close to that of PEO, reaching almost 10-4 S cm-1 at 30 °C. A Li | SPE | LixMnO2 battery device was cycled at 40 °C with approximately 100 mAh g-1 of LixMnO2 between 2.0 and 3.5 V at C/6. However, the cathode binder material was prepared with 30 wt% of low molecular weight PEO (Mw = 2,000 g mol-

1), and not the BCP electrolyte. When a liquid is used, the porous active ma-terial is well wetted, allowing for high ion diffusion and a high rate perfor-mance of the cell. A bulky or high molecular weight SPE might have difficul-ties filling the pores in the active material, which will reduce the contact area between the active material and the SPE, resulting in large impedance and capacity loss during cycling. This is usually addressed via either solvent cast-ing the SPE onto the electrodes, or by using the SPE both as a binder and electrolyte. A low molecular weight PEO could wet the surface of the active material and allow for ion transport between the active material and the elec-trolyte. However, it was not stated in why this approach was used. These re-sults are still competitive, both from a mechanical and battery device point of view, combining a modulus with high ionic conductivity and excellent battery

26

performance at low temperatures. [86] In an attempt to address the low lith-ium-ion transference number of PEO, Ryu et al. synthesized a single-ion triblock copolymer of poly(lauryl methacrylate) (PLMA)-b-poly(lithium methacrylate)-b-poy(oxyethylene methacrylate) with molecular weights span-ning from approximately 30,000 g mol-1 to 80,000 g mol-1 with PDI:s close to 1.20. The BCP electrolyte underwent microphase separation, as proven by SANS and AFM imaging. The transference number was estimated to unity, but the ionic conductivity was in the order of 10-7 S cm-1 at room temperature, and no mechanical testing was made or a battery device was constructed to prove the capability of the electrolyte in a battery device. [87] Bouchet et al. performed an in-depth study of the different architectural types of linear and comb BCPs based on PS as the mechanically supporting block and POEGMA or PEO as the conductive block. When comparing linear and comb BCPs, they found that the linear PS-PEO-PS structure provided a suitable balance be-tween ionic conductivity and mechanical properties, and that the supporting block should comprise up to 30-40 wt% of the matrix. Although the methac-rylate comb SPE PS-POEGMA-PS did not show any crystallization (and high ionic conductivity at room temperature), cyclic voltammetry indicated poor electrochemical stability at 80 °C, thus resulting in poor battery performance and suggesting that the maximum operating temperature for methacrylate-based SPEs is up to 70 °C. They also investigated a series of PS−PEO−PS copolymers with different molecular weights and block compositions. The analysis showed that the ionic conductivity increased with the PEO molecular weight; this finding is counterintuitive to what had been published before. The explanation to these findings was that there are two conductive regions: one region is the bulk PEO, and the second region is a “dead zone” located at the interface of the two blocks where block mixing occurs. Owing to this mixing, the chain flexibility decreases, thus lowering the ionic conductivity in this zone compared to that of bulk PEO. The volume was estimated to be 4-5 EO units wide ( 1.6 nm). This phenomenon was also presented by Balsara and Kanamura in their work on diblock copolymer PS-PEO electrolytes. [88, 89]

Diblock copolymers The most investigated type of BCP structure for electrolyte usage is diblocks. Although the mechanical properties of diblock copolymers are generally good, those of triblock copolymers are considered to be better because they can form more anchoring points, thus forming a more crosslinked network. [88] How-ever, the synthesis of diblock copolymers is often more straightforward and they are thus easier to synthesize. Some early BCP electrolyte systems were investigated by Mayes et al., who focused on the acrylates PMMA, PLMA, and POEGMA. In their first paper, where PLMA was used as the mechanical

27

block, they found that the ionic conductivity was enhanced when the conduct-ing block was attached to a soft rubbery block rather than to a hard-glassy block. Their diblock copolymer exhibited an ionic conductivity one order of magnitude higher than that of a randomly copolymerized PLMA/POEGMA polymer mixed with lithium triflate (LiTf). The diblocks microphase sepa-rated, as proven by DSC, which showed two distinct Tg:s for all compositions. TEM images also confirmed the phase separation and indicated a lamella-type of microphase separation. However, rheology testing showed a modest storage modulus G´ of 105 Pa, which is far from the 6 GPa that has been reported to be needed to prevent dendrite growth. [25, 26] Battery cycling tests for the PLMA-POEGMA diblock copolymer indicated good reversible capacity and capacity retention for at least 25 cycles. [90] In their second study, the authors shifted focus to PMMA-POEGMA diblock electrolytes and performed small-angle neutron scattering (SANS) measurements in order to investigate the bulk properties of the electrolyte and how its morphology is affected by changes in temperature. It has been concluded in earlier papers that PEO and PMMA are compatible in the amorphous state and that blends of these two components are homogenous down to a length scale of 50 nm. However, the properties of these systems drastically change upon the addition of a salt (depending on the type of salt). The SANS profiles showed that it is possible to increase the or-der-disorder transition temperature (TODT) of the polymer system from below room temperature to between 100 °C and 200 °C, depending on salt concen-tration. The authors used non-deuterated polymers, and they employed LiTf to trigger the changes in the TODT. [91] These studies demonstrated the ad-vantages of using BCPs rather than randomly copolymerized polymers.

Many theoretical and experimental efforts have been dedicated to under-standing how diblock copolymer electrolytes undergo microphase separation, and the main focus has been devoted to the PS-PEO system, where most of the PEO is linear. PS is often used to create fully phase-separated BCPs with ordered, nanostructured morphologies at length scales of up to 100 nm, thus yielding long-range ordered structures that typically have lamellar, spherical, or cylindrical domain structures. [85, 92-94] The Balsara group has in a series of articles investigated how molecular weight, composition, salt doping, and morphology dictate the ionic conductivity and mechanical strength of PS-PEO. In an early paper, they studied the effect of molecular weight on the mechanical and electrochemical properties. By changing the molecular weight and composition of the diblocks they managed to create either lamella or hex-agonal morphologies, as proven by both small-angle X-ray scattering (SAXS) and TEM. The ionic conductivity was rather low, and they managed to reach 10-4 S cm-1 only at high temperatures. Rheology gave a G´ of approximately 100 MPa, but the measurements were performed at room temperature, and not at the operating temperature of the battery device. [95] In a later study, they focused on different PS-PEO molecular weights, ranging from 2.7 to 13.7 kg mol−1. An interesting phenomenon was that the ionic conductivity decreased

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with an increase in molecular weight, however, at a certain molecular weight, the ionic conductivity increased again, when exceeding approximately 10 kg mol−1. The parameters that controlled this phenomenon were the Tg of the PS block and the width of the PEO domains. As stated earlier, the explanation to this behavior was the “dead zone” at the block interface, where a mixing of the blocks lowers the ionic conductivity. [89] The distribution of lithium ions has been shown to be located mostly in the bulk PEO phase because the ability to coordinate Li+ near the PS-PEO interphase is lowered by the polymer chain immobility. [96, 97] PS is often used to create fully phase-separated BCPs with ordered, nanostructured morphologies at a length scale of up to 100 nm, thus yielding long-range ordered structures that typically have lamellar, spher-ical, or cylindrical domain structures. [85, 92-94] In a very elegant manner, way Jo et al. synthesized PS-PEO BCPs with −OH, −SO3H, or −SO3Li PEO end groups. Depending on only the end group and salt content, the SPEs formed different types of microphase morphologies, spanning from lamella to hexagonal and gyroid domain structures, as proven by SAXS and TEM imag-ing. [94] This points at how subtle the changes in molecular structure need to be in order to manipulate the morphology.

The microphase-ordered structures could form grains, and several studies have been performed to investigate and predict how the ionic conductivity is affected by the grain size. It is believed that ionic conduction occurs not only in the PEO domains, but also at the grain boundaries. Thus, to maximize the ionic conductivity, it is important to determine whether the size and direction of the grain boundaries can be controlled. [76, 98] Balsara and co-workers showed that the grain size is an important parameter to take into consideration, because not only the microphase separation has to be controlled, but also the structural alignment and grain orientation. The ionic conductivity will in turn be affected by these parameters but in a somewhat counterintuitive manner. The ionic conductivity actually decreased as the grain size increased, thus leading to the conclusion that electrolytes with long-range order and large structures are less favorable than electrolytes with small and disordered grains. [98] The relationship between structure and ionic conductivity in ordered and disordered SPEs has also been studied by Lodge et al., who found that mi-crophase separated but disordered morphologies provide better ionic conduc-tivity than polycrystalline systems with long-range order and grain bounda-ries. Thus, the mixed regions between conducting and insulating blocks are hence important to consider because these interfaces are believed to limit the overall ionic conductivity. [99]

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Polyesters and polycarbonates Alternative electrolyte host materials to ether-based polymers are currently being evaluated. One such branch is based on polyesters and polycarbonates (see Figure 7 for exemplifying polymer structures). Wei and Shriver con-ducted some pioneering work in 1998 when they investigated two types of homopolymers, poly(vinylene carbonate) and poly(1,3-dioxolan-2-one-4,5-diyloxalate) mixed with LiTf. The SPEs were semicrystalline, and the ionic conductivity reached approximately 10-7 S cm-1 at 40 °C, which could have been an outcome of the crystallinity, with a Tm above 100 °C for both systems. The authors concluded that the ion transport was decoupled from the segmen-tal motion because the polymer formed a rigid backbone structure. [100] In-spired by this work, Smith et al. investigated a series of high molecular weight poly(trimethylene carbonate) (PTMC) (300,000 g mol-1) electrolytes by using LiTf and lithium perchlorate (LiClO4) as salts. The electrolytes were fully amorphous, with a Tg of approximately -15 °C. The ionic conductivity was the highest for LiClO4 with a molar concentration of 1:2 Li+:CO (lithium ion to carbonyl oxygen), reaching approximately 10-6 S cm-1 at 30 °C. [101] In a later study they investigated the PTMC system with lithium tetrafluoroborate, and cyclic voltammetry showed that the SPE showed good electrochemical stabil-ity up to about 4.5 V, however, no battery cycling was performed. [102] Munch Elmér and Jannasch investigated a poly(ethylene oxide-co-ethylene carbonate) macromonomer with a molecular weight of 2,650 g mol-1, and an ethylene carbonate content of 28 mol%, before it was mixed with LiTFSI and then UV-polymerized. The final SPE was characterized with AC impedance, and the ionic conductivity was high, ranging from about 10-5 S cm-1 at 30 °C, up to 10-3 S cm-1 at 100 °C. [103] Tominaga and co-workers have in a series of papers in depth investigated poly(ethylene carbonate) (PEC)-based electro-lytes. When comparing salt and salt content they found that the transference number reached 0.63 when using 80 mol% of LiTFSI, in contrast with approx-imately 0.1 when using 60 mol% LiBF4 at 80 °C. [73] In their following papers they used dielectric spectroscopy to investigate the ionic conductivity and the correlation between solvation and salt concentration. The conclusion was that the segmental motion of the polymer chains at high salt concentrations are plasticized due to salt aggregates and a intramolecular screening phenomenon. [72, 104] It was also found that PCE with 80 wt% LiTFSI showed high elec-trochemical stability, and a Li ⏐ SPE ⏐ LiFePO4 device was cycled at 30 °C with relatively high capacity (ca. 80 – 90 mAh g-1 of LFP) at C/5. [105] They

Figure 7. Polymer structures for PCL and PTMC compared with that for PEO.

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have also investigated PEC copolymers with ether based side groups and PEC-PEO copolymers, which all show interesting electrochemical properties at high salt concentrations. [106-109] PTMC was investigated again as electro-lyte material in 2014 by Sun et al., with an in-depth characterization of the system. PTMC with a molecular weight of 368,000 g mol-1 was synthesized and mixed with different amounts of LiTFSI, and a Li | SPE | LiFePO4 half-cell was constructed that cycled at C/13 at 60 °C. The PTMC electrolyte showed good electrochemical stability, but the initial capacity was rather low. However, the capacity increased with cycling, and reached a plateau of ap-proximately 150 mAh g-1 of LiFePO4 after about 70 days. The reason for this was that the electrolyte hade to be infiltrated into the cathode core in order to maximize the full capacity of the cathode. [110] In 2005, Fonseca and Neves presented an SPE based on poly(ε-caprolactone) (PCL) and LiClO4. A Li | SPE | LiNiCoO2 battery device was constructed and cycled for approximately 50 cycles, showing an initial specific capacity of about 150 to 160 mAh g-1 of LiNiCoO2. [111] In 2015 Mindemark et al. prepared a series of copolymers of PTMC and PCL, and by doing so they combined PCL which has a low Tg (-60 °C), with PTMC that have a Tg of approximately -15°C. The benefit with combining these two monomers is that TMC breaks up the symmetry and re-duce the ability of PCL to crystallize. The PCL content ranged from 10 to 40 wt%. The best performing electrolyte was an electrolyte with 40 wt% PCL, with 10-5 S cm-1 at room temperature. A Li | SPE | LiFePO4 half-cell was con-structed and cycled at 60 °C, showing almost 150 mAh g-1 of LiFePO4 at C/5. [112] These very encouraging results made them look into a new series of PTMC and PCL copolymers mixed with LiTFSI, but with a lower amount of PTMC. It was found that for an 80:20 monomer proportion between PCL and PTMC, and with 36 wt% LiTFSI, the ionic conductivity reached approxi-mately 10-4 S cm-1 at room temperature. The lithium-ion transference number was calculated to be 0.62, compared with that of PEO at 0.12. By combining both high ionic conductivity with a high transference number, it was possible to cycle a Li | SPE | LiFePO4 half-cell at room temperature with high capacity at a reasonable rate. The main drawback with the PCL:PTMC electrolyte is that it is has poor mechanical properties. This was addressed in a later paper where gamma irradiation was used to crosslink the electrolyte in an assembled device, however, this strategy showed a limited success. [113, 114]

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Summary It is possible to create BCP electrolytes with good mechanical properties, but the ionic conductivity and the lithium-ion transference number are not high enough to cycle a PEO-based electrolyte at room temperature. The solution is to increase the operating temperature to 60 °C or 100 °C with the penalty of low mechanical properties and parasitic side reactions, which in turn will af-fect battery performance and lifetime. However, when the PEO is grafted to a polymer backbone, it is possible to create low Tg BCPs. The main focus of previous research has been to use PS as the mechanical block, but it could well be of interest to study the morphology and electrochemical behavior of sys-tems with other mechanical blocks because it is clear that the morphology af-fects the ionic conductivity, and thus ultimately the overall battery perfor-mance.

The use of polyesters and polycarbonates, especially the combination of the two, is an interesting alternative to PEO as a polymer electrolyte host material. They combine electrochemical stability with high ionic conductivity and a high lithium-ion transference number. An interesting option would be to add me-chanical stability to these systems and at the same time find a regime where the crystallinity is mitigated in order to maximize the ionic conductivity.

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Results and discussion

Triblock copolymers Papers I and II focus on triblock copolymers, where the midblock was POEGMA, which is an ionic conductive comb polymer, flanked with mechan-ical blocks of either deuterated PMMA (d8-PMMA) or poly(benzyl methacry-late) (PBnMA), to form the polymer structures d8-PMMA-POEGMA-d8-PMMA and PBnMA-POEGMA-PBnMA (see Figure 8). ATRP was used to synthesize the POEGMA midblock, which was purified and dried before be-ing used as a macroinitiator for the second ATRP step where d8-PMMA or PBnMA blocks were polymerized. The pure POEGMA block formed a low-viscous polymer, with a Tg close to -60 °C (Mn ca. 30,000 g mol-1). A series of d8-PMMA based BCP with different molecular weights was synthesized. The PBnMA BCP was synthesized to give one molecular weight, resulting in 35 wt% PBnMA. The final BCPs, as seen in Table 1, had low PDI values, thus indicating well-controlled polymerization conditions. Films made from the POEGMA macroinitiators alone did not form any self-standing films, but the resulting BCPs formed transparent and rubbery-like materials.

Figure 8. Schematic illustration of the triblock copolymers synthesized in Paper I (left) and II (right), utilizing POEGMA as the ionic conductive block and d8-PMMA and PBnMA as mechanical blocks. I symbolizes the initiator.

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Table 1. Molecular weights and PDI values from GPC analysis, where the number af-ter d8-BAB means wt% of POEGMA.

Entry Mn (g mol-1) PDI

d8-MMA

POEGMA 33,145 1.13

d8-BAB75 44,194 1.13

d8-BAB69 48,300 1.13

d8-BAB60 55,510 1.21

d8-BAB48 68,900 1.15

d8-BAB43 76,400 1.17

PBnMA

POEGMA 31,085 1.13

b-BAB35 47,830 1.22

Thermal properties and morphology When mixed with LiTFSI, the formed d8-PMMA and PBnMA electrolytes were transparent, non-sticky and mechanically stable, making the films easy to handle. As anticipated, differential scanning calorimetry (DSC) analysis showed that all samples were fully amorphous. None of the electrolytes in the d8-PMMA-based series showed a second Tg, thus indicating a lack of mi-crophase separation (see Table 2). Furthermore, d8-BAB60 was analyzed with different salt concentrations, ranging from 0 to 23 wt% LiTFSI to evaluate the effect of salt content on the phase separation, but no second Tg was noticed. However, the PBnMA based electrolytes 8:1, 10:1, and 20:1 (ethylene oxide to Li+ ratio, EO:Li+), showed a second clear Tg when LiTFSI was added, thus indicating that the phase separation was triggered by the addition of salt be-cause the BCP without salt was in the mixed state. DSC analysis showed that the phase separation was temperature and time dependent because the second Tg was smeared out on the second heat ramp. To clarify these findings, SANS was used for the d8-PMMA electrolytes and SAXS was used for the PBnMA-based electrolytes in order to study the bulk morphology. The SANS peaks were modelled with the random phase approximation (RPA) model, which is used for polymers in the mixed state. As can be seen in Figure 9, the RPA model fits the scattering curves, regardless of composition. The Flory-Hug-gins interaction parameter was determined with the RPA model, which gave negative values - meaning that the interaction parameter favors mixing. This

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finding, together with the lack of a second Tg in the DSC analysis and TEM imaging where no phase separation was to be found, supports the conclusion that the d8-PMMA-based electrolytes are in the mixed state. Table 2. DSC results for the d8-PMMA and PBnMA based triblock copolymers.

Entry wt% LiTFSI Tg,1 (°C) Tg,2 (°C)

d8-PMMA

d8-BAB75 30.1 -52.2 -

d8-BAB69 28.1 -50.5 -

d8-BAB48 21.6 -38.5 -

d8-BAB43 19.8 -40.5 -

d8-BAB60 0 -54.2 -

d8-BAB60 8.9 -49.4 -

d8-BAB60 12.8 -44.9 -

d8-BAB60 22.8 -37.2 -

PBnMA

b-BAB35 0 -43.9 -

20:1 14.7 -52.5 55.5

10:1 29.5 -48.3 55.8

8:1 36.8 -42.7 59.1

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Figure 9. SANS data for the d8-PMMA based triblock copolymer electrolytes, fitted with the RPA model (solid black lines).

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In Figure 10 the SAXS profiles for the PBnMA-based electrolytes are pre-sented for the heating and cooling ramps from 30 °C to 150 °C and back to 30 °C. The heating ramp shows several broad scattering peaks at 30 °C. However, when heating the sample to 150 °C the scattering peaks become well defined, and the first-order peak q* shifts to a smaller q. The peak position versus tem-perature is not fully reversible, as in agreement with DSC analysis. For all samples three secondary order peaks are observed at √4q*, √7q*, and √14q*, thus indicating a self-assembly into a hexagonal packing of cylinders, which is in agreement with the expected theoretical equilibrium morphology given by the triblock composition. The cylinder-to-cylinder distance, calculated from d = (4*π)/(√3q*), gave that sample 8:1 is 45.4 nm at 30 °C before and 49.2 nm at 30 °C after the annealing step, the distance for sample 10:1 is 48.0 nm and 53.7 nm, and 46.5 and 51.1 nm for sample 20:1. In the temperature range explored, no TODT was reached, thus indicating that the annealing at 150 °C only allowed the structure to release any kinetically hindered locking of the structure.

The TEM images show mainly elongated objects with few regions present-ing isotropic objects (see Figure 11). These observations can be interpreted as cylinders forming a limited hexagonal packing, which could support the broad peaks seen in the SAXS curves collected at 30 °C before the annealing step.

Figure 10. SAXS profiles for sample 8:1, 10:1 and 20:1, measured at 30 °C up to 150 °C, and back to 30 °C.

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Ionic conductivity The temperature dependence of the ionic conductivity for the solvent-cast SPEs was evaluated by AC impedance spectroscopy, and the results are shown in Figure 12. The SPEs exhibit VTF behavior, which is expected for PEO-based systems. The low-temperature ionic conductivity can perhaps be con-sidered comparatively high, being in the order of 10-5 S cm-1, which can be explained by the rather limited molecular weight of the BCPs. The Tg in-creases with the amount of d8-PMMA, thus resulting in that d8-BAB75 is showing the highest ionic conductivity for the d8-PMMA-based electrolytes. In the PBnMA-based electrolyte series, the ionic conductivity increases with the addition of LiTFSI, and there seems to be no dominant long-range ion-ion interactions, which would have caused peak ionic conductivity. This may be a result of a more pronounced microphase separation, mitigating any mixing of the blocks. Furthermore, this means that more salt could be added before a possible peak ionic conductivity is reached. It is also clear that the electrolytes are amorphous even though they are PEO based. This is an outcome of the BCP comb design. Previous studies have shown that there is an optimum chain length of ca. 10 - 20 EO units before the side groups starts to crystallize. [59]

Figure 11. TEM images of sample 8:1 before the annealing step. The images clearly show that the structure before the annealing step is lamella, but mixed with regions that could be cylindrical, see circle. The left scale bar is 200 nm, and the right is 500 nm.

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Electrochemical performance The d8-BAB75 and 8:1 electrolytes were evaluated in Li | SPE | LiFePO4 half-cell devices (see Figure 13). The d8-BAB75 based half-cell did not cycle for more than about 30 cycles before it failed; this failure is most probably as a result of lithium dendrite growth. The limited specific capacity observed could be a consequence of wetting issues between electrolyte and active material, a common problem for SPEs. The PBnMA 8:1 electrolyte cycled at a lower specific capacity; however, it managed approximately 140 cycles at C/10 as a trade off before the cell failed. Several cells were cycled for both systems, and they all showed similar behavior: the 8:1-based electrolyte managed to cycle consistently over 100 cycles, whereas the d8-BAB75 based batteries failed much earlier. Cyclic voltammetry (CV) analysis showed that the electrolytes showed high electrochemical instability, which could be one explanation to the rather limited battery cycling success, see Figure 14. The electrochemical instability of comb methacrylate-based electrolytes have been noticed in an earlier publication where random copolymerized PMMA and POEGMA was used. [115] The same electrochemical phenomenon have also been observed by Devaux et al. when evaluating linear and comb BCP electrolytes, and PS as the mechanical block and NMP as the synthesis method. [88]

Figure 12. Ionic conductivity for the d8-PMMA (d8-BAB)- and PBnMA-based electro-lytes (8:1, 10:1, and 20:1). Dashed lines represent VTF-fits.

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Figure 13. Cyclic voltammogram of electrolyte 8:1 in a lithium vs. stainless steel cell configuration operating at 60 °C, cycled at 1 mV s-1.

Figure 14. Cycling of Li | SPE | LiFePO4 half-cells at 60 °C and at C/10.

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Conclusion triblock copolymers It is known that even a small addition of salt could have a significant effect on the phase separation in BCP electrolytes, and thus also on the ion transport and mechanical properties. Rather surprisingly the d8-PMMA-based electro-lytes were found to be in the mixed state, for a wide range of molecular weights and salt concentrations, as proved by fitting the data with the RPA model. The single broad peak is thus a correlation hole scattering peak. But when PBnMA was used the phase separation was triggered by the addition of salt, which indicates that the phase separation have to be related not only to the chain length but also to the Flory-Huggins interaction parameter, which not necessarily needs to show a linear relationship. However, if the overall BCPs were larger, both the A and B blocks, it is very much likely that the phase separation would be more pronounced, which points at that this is in-deed a delicate matter to predict, since it is known that the Flory-Huggins in-teraction parameter actually depends on N, and the salt concentration, and that the relationship is non-linear. [81]

The ionic conductivity was relative high for these systems, which could be a consequence of the high mobility of the comb-PEO and the relatively low overall molecular weight of the BCPs. Sample 8:1 showed the highest ionic conductivity, which is attributed to that it microphase separated. However, it might be so that the high ionic conductivity depends both on an increase in charge carrier concentration and a low degree of block mixing.

The methacrylate BCPs that were investigated here gave electrolytes with good mechanical robustness and relatively high ionic conductivity. However, CV analysis showed low stability versus the cell components. The elevated cycling temperature may allow for parasitic side reactions to be pronounced, which could be an explanation for the fading capacity in the 8:1 PBnMA-based electrolyte. Although the mechanical properties were not quantitatively analyzed, the d8-BAB75 electrolyte was softer as compared with the 8:1 elec-trolyte, and dendrite growth may be the reason for the rapid cell failure. It was rather unexpected that the methacrylate based electrolytes are so electrochem-ically reactive because gel electrolytes based on PMMA has been frequently used in battery devices, where they do not show the same reactivity. However, the materials studied here are dry systems, which will experience rather dif-ferent impedance and mass transport properties compared with gel electro-lytes. [116, 117] The poor performance of comb methacrylate-based electro-lytes was also observed by Devaux et al., who used PS as the mechanical block and NMP as the ynthesis method. [88] The conclusion so far is that these methacrylate based comb copolymers are of limited success for lithium bat-tery applications. Although having relatively high ionic conductivity, the elec-trochemical stability and cycling capability is poor.

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Polyester-based electrolytes In an attempt to evaluate polyesters and polycarbonates as SPEs, a series of BCPs were synthesized and evaluated (see Figure 15). It is of interest to study these type of SPEs since they show high lithium-ion transference numbers together with promising mechanical properties and ion conduction. [107, 109, 113] Two different polyester based polymers were evaluated, poly(∂-valerolactone) (PVL) and poly(ε-caprolactone) (PCL). The general approach was to synthesize a mechanical block via ATRP and use it as macroinitiator for the following ring opening polymerization (ROP) step. The results for the PCL-based polymers are presented in Papers III and Paper IV.

∂-Valerolactone-based electrolytes PVL with approximately 25,000 g mol-1 was synthesized and mixed with LiTFSI to form a salt series. The films were mechanically stable at low salt concentrations but quickly lost mechanical integrity at higher salt loadings, thus making it impossible to analyze the entire series. However, the ionic con-ductivity was high for the electrolytes that could be measured. In an attempt to increase the mechanical integrity, two different diblock copolymers were synthesized: a PS-PVL (SV) diblock copolymer and a PS-poly(∂-valerolac-

Figure 15. Synthesized polyester-based polymers, synthesized with ATRP and ROP.

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tone-r-trimethylene carbonate) (SVT) diblock. The hypothesis is that by add-ing PS and TMC the mechanical integrity should increase and the degree of crystallinity should decrease, respectively, thus allowing for both mechanical integrity and high ionic conductivity. The PS macroinitiator was polymerized via ATRP, thus generating a macroinitiator with narrow PDI; but generally, the control of the polymerization was lost in the ROP step, thus resulting in BCPs with broad PDIs (see Table 3). Moreover, even when a mechanical block was attached, the mechanical integrity was poor, and it was not possible to analyze a full series from any of the prepared polymers. Some conductivity data can be seen in Figure 16.

Table 3. Molecular weights from GPC analysis.

Entry Mn(g mol-1) PDI wt% A-block PVL 25,500 2.58 - PS-OH 27,900 1.13 - SV 45,800 2.00 61 SVT 53,300 1.77 52

Figure 16. Ionic conductivity for ∂-valerolactone based electrolytes.

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ε-Caprolactone-based electrolytes

Synthesis A PS-PCL (SC) diblock was synthesized via a two-step ATRP and ROP syn-thesis, with the aim to increase the mechanical properties. However, given that PCL has a high degree of crystallinity, the following BCPs were copolymer-ized with TMC, R,L-lactide or ∂-decalactone with ε-caprolactone to form the BCPs PS-poly(ε-caprolactone-r-trimethylene carbonate) (SCT), PS-poly(ε-caprolactone–r-∂-decalactone) (SCD) and PS-poly(ε-caprolactone–r-D,L-lac-tide) (SCL), with the idea to lower the symmetry and generate a lower degree of crystallinity. In the last BCP the mechanical A-block was altered by using PBnMA as a mechanical A-block, thus forming the BCP poly(benzyl methac-rylate)-poly(ε-caprolactone-r-trimethylene carbonate) (BCT). As for the ∂-valerolactone based BCPs, the control of the polymerization was lost in the ROP step, particularly when TMC was used, thus resulting in BCPs with broad PDIs (see Table 4).

Table 4. Molecular weights from GPC analysis.

Entry Mn (g mol-1) PDI wt% A-block PCL 24,400 1.86 - PS-OH 27,900 1.13 - SC 49,200 1.43 57 SCD 53,700 1.47 52 SCL 51,700 1.57 54 SCT 51,200 2.33 54 PBnMA-OH 26,000 1.13 - BCT 79,000 2.00 33

Ionic conductivity The difference between PCL and PVL is that PCL have one more carbon in the repeating unit, but the effect on the mechanical properties is noticeable. Where PVL lost all mechanical integrity when LiTFSI was added, the PCL electrolytes formed robust, hard, non-sticky and rubbery-like electrolytes, thus making them easy to handle. The temperature dependence of the ionic con-ductivity for the PCL electrolytes with 9 and 17 wt% LiTFSI shows a clear effect of crystallinity. However, when increasing the salt concentration to 29 wt% LiTFSI or increasing the temperature to 60 °C (above the Tm), the crys-tallinity decreased, thus allowing for higher ionic conductivity when the elec-

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trolyte becomes fully amorphous (see Table 5 and Figure 17). The ionic con-ductivity for the SC electrolyte with 9 and 17 wt% LiTFSI also shows a clear effect of crystallinity (see Figure 18). The addition of ∂-decalactone greatly reduced the crystallinity, but the ionic conductivity suffered (see Figure 19). Electrolytes based on the homopolymer of D,L-lactide do not form any stable films when adding LiTFSI; the mechanical integrity is low, and the resulting films easily fall apart. However, SCL formed a though polymer with a limited ionic conductivity compared with the other electrolytes (see Figure 20). The SCT electrolytes utilizing TMC forms a linear polymer without bulky side groups, which could interfere with the ion motion throughout the polymer ma-trix. This could be one of the explanations for the high ionic conductivity for this series (see Figure 21).

When switching the mechanical block from PS to PBnMA, the ionic con-ductivity was somewhat increased, with BCT17 showing the highest ionic conductivity for all diblock electrolytes (see Figure 22 and Table 5). The peak in ionic conductivity for the electrolytes is attributed to the optimum free ionic-charge-carrier concentration. Excessive salt results in a crosslinking ef-fect, ion pairing, and a net decrease in ionic conductivity. An interesting ob-servation was that the BCT-based SPEs films shimmered with a clear blue color. Table 5. Highest ionic conductivity for diblock SPEs at 30 °C.

Entry Ionic conductivity at 30 °C (S cm-1)

PCL 1.68E-05 SC 3.84E-06

SCD 1.43E-06

SCL 8.93E-07

SCT 4.82E-06

BCT 9.06E-06

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Figure 17. Ionic conductivity for PCL series.

Figure 18. Ionic conductivity for SC series.

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Figure 19. Ionic conductivity for SCD series.

Figure 20. Ionic conductivity for SCL series.

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Figure 21. Ionic conductivity for SCT series.

Figure 22. Ionic conductivity for BCT series.

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Thermal properties and morphology DSC analysis showed that the LiTFSI-free polymers were highly crystalline; however, with the addition of LiTFSI, the crystallinity decreased (see Table 6). By adding TMC to the BCP, the symmetry of PCL was broken, thus further decreasing the degree of crystallinity. With the addition of LiTFSI, the SCT and BCT electrolytes became fully amorphous at lower salt concentrations than compared to the PCL and SC electrolytes. However, no second Tg was found for the BCP electrolytes, thus indicating that no phase separation oc-curs. It should also be pointed out that these values come from the second heat scan, meaning that the thermal history have been erased. Several of the sam-ples showed crystallinity on the first scan but not on the second scan.

Table 6. DSC values for PCL, SC, SCT and BCT SPEs.

Entry Tg [°C] Tc [°C] Tm [°C] PCL0 −34.3 27.9 55.5 PCL9 −38.2 10.9 51.3 PCL17 −42.9 14.4 51.2 PCL23 −44.1 −1.5 45.0 PCL29 −40.2 - - SC0 −36.7 26.8 55.1 SC9 −37.5 21.0 53.4 SC17 −40.1 −8.5 45.6 SC23 −41.8 - - SC29 −35.4 - - SCT0 −63.0 −10.4 42.2 SCT9 −42.1 −14.4 36.7 SCT17 −42.9 - - SCT23 −37.1 - - SCT29 −24.8 - - BCT0 −42.3 1.8 49.7 BCT9 −41.0 -6.1 44.4 BCT17 −48.9 2.7 39.3 BCT23 −41.3 - - BCT29 −35.1 - -

Mechanical properties DMA was performed to evaluate the moduli as a function of temperature for SC23, SCT17, and BCT17 (see Figure 23). The results show that the storage modulus E´ is highest for BCT17, and that the modulus is not far from that of PS, which is added as a reference material. The drop in modulus for all the

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electrolytes occurs at a rather low temperature, which may relate to the Tm. The addition of TMC to SCT17 shows a dramatic effect on the storage mod-ulus E´, which decreases by almost one order of magnitude compared with SC23, even though it is using PS as a mechanical block, which is rather sur-prising. The rheological properties of BCT17 where analyzed with oscillation frequency sweeps at 30 °C and 60 °C, corresponding to the battery cycling temperatures. Rheology data at 60 °C showed that the moduli of BCT17 had a rapid decrease and a crossover at ca. 1 Hz, thus indicating that the material loses its mechanical stability and starts to flow; this finding correlates with the DMA results where a crossover at approximately 40 °C can be observed. At 30 °C the G´ and G´´ were well separated, thus indicating that the electrolyte will maintain its mechanical integrity over time.

Diblock copolymer via controlled synthesis It was found that good control of the ROP step was possible to achieve when the ATRP reaction was quenched with dry THF and filtered through basic Al2O3 directly into the ROP reaction vessel. In the earlier protocol, the ma-croinitiator was precipitated into methanol, before being dried in a vacuum oven for a week; thereafter, it was used in the ROP step in bulk at 130 °C for approximately 2 – 3 days. By adopting the more controlled synthesis method, the PDI values decreased from 2.0 to 1.30 (see Table 7). The difference in

Figure 23. DMA analysis of SC23, SCT17 and BCT17 with PS, added as reference material. The crossover at ca. 40 °C correlates with the Tm of the diblocks. Only E´ is plotted for the sake of clarity.

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synthesis method is observed in the DSC analysis, where BCT17 shows a clear second Tg at 32 °C and SCT17 a Tg at 100 °C, when adopting the optimized method. However, the DSC thermogram was much more complex for the SCT17 electrolyte, thus poosible indicating a complex morphological structure.

A corresponding electrolyte series with a deuterated SCT BCP (d8-styrene) was analyzed with SANS, which showed multi-peak high intensity scattering, see Figure 24. The microstructure of the SCT electrolytes are not yet fully understood, but preliminary AFM and TEM imaging indicates that the elec-trolytes lack long-range order. Although the microphase structure (lamella, hexagonal, etc.) could be predicted by equilibrium theory, the actual structure will depend on film preparation, and parameters such as temperature, pressure, and solvent. The solvent will play a very crucial role for the final microphase structure because the solvent to some degree will be selective to one of the blocks. This will lead to the swelling of one of the block domains, thus causing an altering off the effective volume and result in films that to some degree will be in a state of non-equilibrium. [118, 119] The DSC thermogram for BCT17 showed two clear Tg:s, and AFM imaging showed a clear microphase sepa-rated structure, indicating a hexagonal packing morphology (see Figure 25).

However, the ionic conductivity was not that affected by this phenomenon, according to a comparison of the ionic conductivity between the two synthe-sized BCT polymers. This also holds for the mechanical properties, where only a slight increase in G´ was observed in the rheological. This is rather unexpected results because it was assumed that a microphase separated elec-trolyte would have higher ionic conductivity and stronger mechanical proper-ties than a mixed electrolyte system. This indicates that good control of the polymer architecture may not be that important for this system, or that the annealing time in the AC impedance and rheological set up was too short to allow for full microphase separation to occur. Table 7. Molecular weights from GPC analysis.

Entry Mn (g mol-1) PDI wt% A-block PS-OH 11,000 1.16 - SCT 35,400 1.21 31 PBnMA-OH 7,200 1.15 - BCT 22,900 1.30 31

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Figure 24.SANS profiles for the SCT electrolytes on the cooling ramp. The curves are combined from three angles from the NGB30 beamline at NIST. As the salt increased, the scattering peaks become more pronounced.

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Electrochemical performance PCL, SC, SCT, and BCT electrolytes were evaluated in Li | SPE | LiFePO4 half-cells. The PCL based half-cells could not be operated for more than a few cycles, with the best performing cell managing only approximately 16 cycles before failing. The cells were annealed for 6 h at the operating temperature before cycling to maximize the interfacial contact between the electrolyte and the lithium anode. During this time, the open circuit voltage (OCV) fluctuated significantly, thus indicating reactions between the electrolyte and one of the electrodes, presumably the lithium anode. This may be supported by the fact that it was hard to obtain a cyclic voltammogram, and it was impossible to perform a transference number analysis because no steady current was ob-tained. However, full-cells with the configuration graphite | PCL23 | LiFePO4 were possible to cycle at 40 °C at a rate of C/20. Also these cells showed an

Figure 25. AFM images of BCT17 collected using phase contrast and tapping mode tips.

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initially high instability in the OCV, indicating electrolyte degradation. The-secells were able to cycle with ca. 80 – 100 mAh g-1 of LiFePO4.

However, these analysis methods were possible to conduct on the SC, SCT, and BCT electrolytes. The general voltammogram, in a stainless steel vs. lith-ium cell configuration, showed a passivation of the lithium electrode and a relatively high current peak at 1 V. It is yet unclear what this current peak is an outcome of, but it may be a redox activity effect of the ester groups. The transference number (T+) were estimated with the Bruce–Vincent method to be 0.56 for SC23, 0.68 for SCT17 and 0.64 for BCT17. These T+ values are far higher than those usually reported for PEO, with usual T+:s of ca. 0.1 to 0.3. [24, 78, 120] All the diblocks showed a flat OCV during the 6 h annealing at the operating temperature for the Li | SPE | LiFePO4 half-cells, thus indi-cating that the electrolytes were sufficiently chemically and mechanically sta-ble with the cell components. Upon cycling at 60 °C, the capacity decreased significantly during the first two cycles for the SC23-based cell. After this initial rapid drop, the capacity levelled out but continuously faded upon fur-ther cycling. The cell operated for over 250 cycles until it ultimately failed (see Figure 26, battery C). The rapid decrease in capacity and the poor cycling stability is hard to explain, and whether this is an effect of SEI on the anode side or the solid permeable interphase (SPI) on the cathode side is hard to determine. However, the overall battery performance and lifetime under prac-tical operating conditions depend strongly on these interphases.

SCT17 formed a smooth and hard electrolyte surface when solvent-casted onto the cathodes. It can be concluded that the SCT17 electrolyte greatly im-proves the cycling performance as compared with the PCL- and SC-based

Figure 26. Cycling of Li | SPE | LiFePO4 half-cells. Battery B is utilizing SCT17 as a binder material and electrolyte, cycling at C/5. Both batteries were operated at 40 °C.Battery C is using SC23 as electrolyte, operating at C/10 and at 60 °C and D and Eis using BCT17 as electrolyte.

54

SPEs. A typical cycling behavior is shown in Figure 26 (battery A). The ca-pacity was initially high (the theoretical value is 170 mAh g-1, wheras the prac-tical is close to 150 mAh g-1 of LiFePO4), and the half-cell easily operated for more than 300 cycles. To increase the mass transport of lithium ions through the cathode and generate good compatibility with the electrolyte and the active material, the electrolyte was used both as a binder (instead of polyvinylidene difluoride) (PVdF) and electrolyte. The cycling stability was greatly improved with this approach and it was possible to cycle the cell at C/5 at high capacity at 40 °C; see Figure 26, battery B.

To understand the effect of using SCT17 as a binder in the cathode, an intermittent current interruption technique was applied to follow the resistance in the Li | SCT17 | LiFePO4 half-cell. In this experiment, the resistance was measured from short current interruptions during the cycling of the cell. [121, 122] From Figure 27, it is clear that the polarization and resistance during cycling are much higher in the cells using PVdF as binder than in those using SCT17, probably because of mass transport limitation of ions in the inner parts of the cathode.

The Li | BCT17 | LiFePO4 half-cells were cycled at 30 °C at C/10, and at 60 °C at C/5. The half-cell cycling at 60 °C continued with a constant capacity fade, with a starting capacity of ca. 115 mAh g−1 of LFP until it reaches ca. 80 mAh g−1, at which point the cell was turned off. The cell cycling at 30 °C shows an initially low capacity that increases over time. This is believed to be

an outcome of poor interfacial contact between the electrolyte and the active

Figure 27. A two-electrode resistance analysis during cycles 1, 10, and 20 of LiFePO4 half-cells cycling with SCT17 as electrolyte (upper) and as binder (lower) at 40 °C.

55

material. Thus, a second cell was cycled using BCT both as binder material (together with LiTFSI) and as electrolyte. Although partly overcoming the in-itial low capacity, the overall specific capacity is still low but with a stable cycling profile, see battery E in Figure 26.

Conclusion diblock copolymers It was seen that PCL based electrolytes showed high ionic conductivity, but a limited electrochemical stability, which made it impossible to cycle lithium half- or full cells for any extended time. On the other hand, when using SC as electrolyte, the electrochemical stability increased. Although cycling at low capacity and at a modest rate of C/10, the addition of PS block indicates that some kind of electrochemical stabilization occurs. The SCT electrolytes showed superior steady cycling behavior compared with the PCL and SC elec-trolytes, cycling close to full capacity at 40 °C and at C/5. Moreover, the bat-tery cycling performance was greatly improved by incorporating the electro-lyte as a cathode binder material, thus showing that the wetting of the active material is crucial for minimizing the overpotential and resistance in the cell. However, although SCT17 showed the best electrochemical performance, it also showed the lowest mechanical properties compared with the SC and BCT electrolytes. When switching the mechanical block from PS to PBnMA the mechanical properties increased, thus making the BCT material a good candi-date for robust all-solid-state batteries because it combines high ionic conduc-tivity, a high T+, electrochemical stability and good mechanical properties. However, the capacity utilization of the active material with this electrolyte is not yet as high as for SCT17. It is not clear whether this is due to the wetting issues of the active material, or an inherent issue of BCT17 because the mate-rial otherwise displays a highly promising combination of properties.

The synthesis of the diblocks was improved when care was taken to avoid the precipitation of the macroinitiators into a solvent. DSC, AFM, and SANS analyses showed that when prepared under these conditions the electrolytes became microphase separated. This resulted in a slight increase in ionic con-ductivity and G´ modulus. AFM imaging suggests that the BCT17 electrolyte forms a hexagonal cylindrical perforated structure. The SCT electrolytes showed strong scattering, but AFM and TEM imaging indicate that the sys-tems lacks long-range order. Whether this is an outcome of sample prepara-tion, choice of solvent, annealing time, or any other parameter is still difficult to say.

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Conclusion

The field of SPEs is indeed multidisciplinary, where synthesis, battery assem-bly, and advanced characterization of polymer, electrolyte, and battery have to be mastered in order to fully study the ability of the SPE. This thesis work started with synthesizing a series of randomly copolymerized PMMA and POEGMA polymers. It was found that the ionic conductivity was rather high for this series; thus, the natural next step was to synthesize BCPs in an attempt to further increase the ionic conductivity with the hypothesis that microphase separated blocks would allow for both high ionic conductivity and mechanical integrity. However, as presented in Paper I, it turned out that the SPEs did not microphase separate. However, by replacing PMMA with PBnMA as me-chanical blocks in Paper II, the electrolytes did microphase separate, and the ionic conductivity increased as a result. On the other hand, the methacrylate-based SPEs were not electrochemical stable, and the battery performance for these types of SPEs was limited.

By switching to the polyester and polycarbonate-based electrolytes, not only high ionic conductivity was achieved, but also high lithium-ion transfer-ence numbers, good electrochemical compatibility, and promising mechanical properties. The promising mechanical properties, together with the fact that the SCT and BCT diblocks seems highly reliable electrolyte materials (i.e., they do not suffer from any “sudden death”), makes these BCPs to interesting candidates for further investigations as all-solid-state battery electrolytes. However, the interphase between electrolyte and electrode have not been in-vestigated in the scope of this thesis, and most of the experiences so far points at that this interphase being the elucidating parameter that needs to be ad-dressed in order to truly fulfill the potential of the SCT and BCT electrolytes.

It has not been studied, but it is most likely that the bulk properties of BCT17 behaves rather differently than when used as a binder in the cathode, where the electrolyte forms interfaces with several different components. It is well known that when the thickness decreases the surface and interface of a polymer becomes increasingly important, which could affect ion and electron transport. Depending on the interaction orientation the phase separation could be altered. The molecular weight is therefore important because when the thickness approaches the radius of gyration the properties of the polymer will change drastically, such as the morphology, density and the Tg. This preferen-tial configuration could create a ion barrier at the interface, but which pro-

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motes electron transport, an artificial SEI layer. Therefore, the interface inter-action could potentially dictate several important SPE performance parame-ters. With this said, maybe it is not optimal to use BCT17 as both electrolyte and binder, but rather use a two component system, where BCT17 is the bulk electrolyte and another polymer electrolyte is used as binder material in order to maximize ion and electron transport at the interface with the active material.

The other parameters for optimization these diblocks are more easily ad-dressed. High molecular weight block copolymers of ca. 100,000 g mol-1 should be used in order to allow for molecular entanglements, thus preventing the material from creep. By optimizing the AB block ratios and the amount of TMC it should be possible to maximize the mechanical robustness and ionic conductivity. Hopefully, this could realize materials that combine high ionic conductivity, lithium-ion transference numbers, good electrochemical com-patibility, and high mechanical integrity.

Furthermore, it should be of interest to investigate different lithium salts and active materials because changing the anion could drastically alter several properties of the electrolyte like the ion conductivity, lithium transport num-ber, Tg and mechanical properties. For all SPEs, one of the main bottle-necks for utilizing the full capacity of the active material is the ionic conductivity. The manner in which to solve this problem is still an open question. Although still struggling with increasing the ionic conductivity, initial studies showed that SCT17 was able to be cycled in full cells with commercial NMC and graphite electrodes, where the electrodes were infiltration casted, sandwiched together, and cycled in coin cells at 40 °C. This “plug and play” approach worked well, although with a low specific capacity as a result of the poor in-filtration of the electrolyte into the densely packed electrodes. These experi-ments also indicate that the NMC cathodes perform better than the LFP-based cathodes. Another option would be to use a low-viscosity SPE as cathode binder, while using the high-viscosity polymer as the bulk electrolyte between the electrodes.

What has only been touched upon in this thesis is the role of morphology and its influence on battery performance. In the literature, it has been shown that when the alignment was parallel to the electrodes the ionic conductivity of a PEO-based BCP increased from 10-9 S cm-1 to 10-6 S cm-1, for a BCP with only 23 wt% of PEO. [123] Other approaches for creating nanoscale control include solvent vapor annealing, solvent casting and addition of fillers, but further studies needs to be conducted to make these technologies more acces-sible for large-scale ordering. [38, 124-127] These very exiting finding opens up for further developments in the field of SPEs, by combining several differ-ent interesting scientific topics.

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Svensk sammanfattning

Fasta polymera elektrolyter (FPE) är av stort intresse att använda i litiumbat-terier, då flytande elektrolyter är både brandfarliga, giftiga samt sliter hårt på batteriets samtliga komponenter, vilket i sin tur påverkar batteriets livslängd och prestanda. FPE skulle även kunna tillföra mekaniska egenskaper, vilket potentiellt skulle medföra nya möjligheter när man designar batterier i forma av att man inte behöver en separator mellan elektroderna, och på sätt ökar batteriets energidensitet. Det är flera parametrar som måste tas hänsyn till när man designar en ny elektrolyt, mekaniska egenskaper, jonledning, jonselekti-vitet samt elektrokemisk stabilitet är de viktigaste. Det största problemet med polymera elektrolyter är deras dåliga jonledning i förhållande till deras meka-niska egenskaper; ofta måste man välja mellan dessa två egenskaper. FPE li-der emellertid fortfarande av för låg prestanda för praktiska tillämpningar vid rumstemperatur, och temperaturen måste upp till 60 °C-90 °C för att få till en praktisk prestanda. Nackdelen med att höja temperaturen är dock att olika ty-per av sidreaktioner kan accelerera och bli märkbara, vilket påverkar batteriets livslängd och prestanda på ett negativt sätt. Ett sätt att kombinera både meka-niska egenskaper och elektrokemisk prestanda är att polymerisera blocksam-polymerer (BSP), det vill säga polymerer som kombinerar ett joniskt ledande block med ett mekaniskt block i en polymer. Genom att variera monomerkom-positionen och strukturen hos BSP är det möjligt att tillverka elektrolyter med olika batteriprestanda. I artikel I och artikel II syntetiserades två typer av metakrylatbaserade BSP:er med olika mekaniska block för att utvärdera mor-fologin, den elektrokemiska stabiliteten och batteriförmågan. Denna avhand-ling började med att syntetisera en serie sampolymerer av metylmetakrylat (MMA) och metylmetakrylat med etylenoxid som sidogrupper. Det visade sig att den jonledande förmågan var hög, så nästa naturliga steg var att polymeri-sera BSP:er, i ett försök att ytterligare öka den jonledande förmågan med hy-potesen att mikrofasseparerade BSP:er kan kombinera både hög jonisk kon-duktivitet och mekanisk integritet. Men som framgår i artikel I visade det sig att elektrolyterna inte fasseparerade. I artikel II ersättas MMA:n med ben-zylmetakrylat (BMA), och då fasseparerade elektrolyterna. Tyvärr visade det sig att de metakrylatbaserade elektrolyerna inte är elektrokemiskt stabila, och batteriprestanda var begränsad.

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I artikel III och artikel IV antogs en ny strategi, och fokus förflyttades till diblocksampolymerer, i vilka den jonkoordinerande gruppen byttes från ety-lenoxid till ester- och karbonatgrupper. Det visade sig att batteriprestandan för dessa elektrolyter var överlägsen de metakrylatbaserade elektrolyterna. Det var möjligt att cykla halvceller med konfigurationen Li | elektrolyt | LiFePO4 vid 40 °C och med hög kapacitet under en längre tid. När BMA användes som mekaniskt block i kombination med kaprolakton och trimetylkarbonat erhölls en BSP som visade sig vara ett attraktivt alternativ till etylenoxid som poly-merelektrolyt, då elektrolyten kombinerade elektrokemisk stabilitet med hög jonkonduktivitet och transporttal, samt mekanisk integritet. Genom att växla till polyester- och polykarbonatbaserade elektrolyter uppnåddes alltså inte bara hög jonledningsförmåga utan även höga transporttal, god elektrokemisk kompatibilitet och lovande mekaniska egenskaper. De lovande mekaniska egenskaperna tillsammans med att BSP:erna verkar mycket pålitliga som elektrolytmaterial, de drabbas inte av några plötsliga kortslutningar, gör dessa BSP:er till intressanta kandidater för fortsatta studier som elektrolyter för strukturella batterier.

Figur 12. De två studerade blocksampolymererna i artikel I (vänster) och II (höger).

Figur 2. Bilden visar blocksampolymeren som studerades i artikel IV, som är en kom-bination av benzylmetakrylat som mekaniskt block samt ett jonledande block av kapro-lakton och trimetylenkarbonat.

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Acknowledgements

As Ramiro said when a were about to start at Uppsala, enjoy the ride! Truly it has been a ride, both emotionally and in life. Not to mention the ride between Stockholm and Uppsala... I guess I have to acknowledge SJ for (mostly) taking me home to Stockholm late evenings, and the fact that this thesis actually was written (mostly) on the train.

To my supervisors: thank you for giving me an amazing once in a lifetime opportunity to freely explore the world of polymers, I believe that this freedom will be hard to find again. Laurent, thank you for taking your time helping me with the SAXS analysis and all the time you spent on discussing science with me. I’m grateful to have had the opportunity to meet such a friendly and help-ful person. Adrian! Thank you for introducing me to the world of neutrons, a truly amazing but at the same time daunting world. The trip to NIST was an eye-opener for me, it was great to meet so many friendly and enthusiastic peo-ple that really wants to contribute. I will remember this trip for the rest of my life.

All members of the polymer chemistry and PUB-group throughout the years, thank you for the friendly atmosphere and being the people you are! Special thanks to Fabian, Ronnie and Matthew who helped me with the battery chemistry part, without you guys life had been really hard. I really appreciate your friendliness and your enthusiasm to science (and beer drinking), thank you for these years.

I have had a couple of project workers throughout the years helping me with various stuff in the lab. Maybe everything wasn’t that organized and planned in detail, but it was great fun and I’m truly grateful for the help, thank you so much! Special thanks to Jonas Hedman and Guiomar Hernández, with-out your help in the end I would not have made it, at the time I was too tired to do all the lab-work myself.

Daniel Bermejo and Fabian Jeschull, you have been great friends and nice companions in the lab, so well as at the pub. I wish you the best of luck what-ever you do in life.

To my family and friends, it is hard to explain what I’m been doing in Upp-sala, but it doesn’t matter, thank you for supporting me anyways and balanced my life, with football, orchestra, and just having fun! And Martina, I hope that we will have more time for each other and our small family know, exciting new times lie ahead.

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