University of Groningen Ferroelectric block copolymers: from self-assembly towards potential application Terzic, Ivan

Ferroelectric block copolymers: from self-assembly towards potential application

Terzic, Ivan

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Ferroelectric Block Copolymers:

From Self-Assembly Towards

Ivan Terzić PhD thesis

University of Groningen

April 2019

Zernike Institute PhD thesis series 2019-13 ISSN: 1570-1530

ISBN: 978-94-034-1529-1 (printed version) ISBN: 978-94-034-1528-4 (electronic version)

The work described in this thesis was performed in the research group Macromolecular Chemistry and New Polymeric Materials of the Zernike Institute for Advanced Materials at the University of Groningen, The Nether-lands. This work was supported via the Innovational Research Incentives Scheme of The Netherlands Organisa-tion for Scientific Research NWO (VICI Grant 724.013.001).

Cover design: Dina Maniar, dina.maniar09@gmail.com Lay-out: RON Graphic Power, www.ron.nu

Printing: ProefschriftMaken || www.proefschriftmaken.nl

Ferroelectric Block Copolymers:

From Self-Assembly Towards

Potential Application

PhD thesis

to obtain the degree of PhD at the University of Groningen

on the authority of the Rector Magnificus prof. E. Sterken

and in accordance with the decision by the College of Deans. This thesis will be defended in public on

Monday 1 April 2019 at 12.45 hours

by

Co-supervisor Em. prof. G. ten Brinke Assessment Committee Prof. B. Ameduri

Prof. A.J.H.M. Rijnders Prof. F. Picchioni

Table of contents

1 Introduction 9

1.1 Piezoelectricity and ferroelectricity 10

1.2 Poly(vinylidene fluoride), crystalline structure, co-, terpolymers and their application 12

1.3 Polymer-based multiferroic materials 18

1.4 Polymer-based dielectric materials for capacitive energy storage 20 1.5 Block copolymers: self-assembly, confined crystallization and selective

dispersion of nano-objects 28

1.6 Synthesis of PVDF-based block copolymers 33

1.7 Self-assembly of PVDF-based block copolymers 38

1.8 Ferroelectric properties under nanoconfinement 41

1.9 Scope of this thesis 44

1.10 References 46

2 CuAAC Click Chemistry: A Versatile Approach Towards PVDF-based Block Copolymers 53

2.1 Introduction 55

2.2 Experimental section 56

2.2.1 Materials 56

2.2.2 Synthesis of 4-(chloromethyl)benzoyl peroxide 56 2.2.3 Synthesis of 4-((trimethylsilyl)ethynyl)benzoyl peroxide 57

2.2.4 Synthesis of chlorine-terminated PVDF 57

2.2.5 Synthesis of azide-terminated PVDF 57

2.2.6 Synthesis of (TMS-alkyne)-terminated PVDF 58

copolymers 60 2.2.12 Purification of the PVDF based block copolymers 60

2.2.13 Characterization 60

2.3 Results and Discussion 61

2.4 Conclusion 69

2.5 References 71

3 Electroactive Materials with Tunable Response Based on Block Copolymer Self-Assembly 75

3.1 Introduction 77

3.2 Experimental section 78

3.2.1 Synthesis of block copolymers 78

3.2.2 Polymer film preparation 81

3.2.3 Polymer characterization 82

3.2.4 Hysteresis loop measurements 82

3.3 Results and Discussion 83

3.4 Conclusion 92

3.5 References 93

4 Tailored Self-Assembled Ferroelectric Polymer Nanostructures with Tunable

Response 97

4.1 Introduction 99

4.2 Experimental section 101

4.2.1 Materials 101

4.2.2 Synthesis of alkyne-terminated P(VDF-TrFE) 101

4.2.3 Synthesis of alkyne terminated P2VP 102

4.2.4 Synthesis of block copolymers 102

4.2.5 Polymer film preparation 103

4.2.6 Characterization 103

4.2.7 Hysteresis loop measurements 104

4.3 Results and Discussion 104

4.4 Conclusion 116

5 Polymer-based Multiferroic Nanocomposites via Block Copolymer Self-Assembly 121 5.1 Introduction 123 5.2 Experimental section 125 5.2.1 Materials 125 5.2.2 Synthesis of (Co2+ _Fe 23+)–oleate precursor 125

5.2.3 Synthesis of cobalt ferrite (CoFe2O4) nanoparticles 125

5.2.4 Ligand exchange and preparation of hydrophilic nanoparticles 126

5.2.5 Preparation of polymer films 126

5.2.6 Characterization 126

5.3 Results and Discussion 127

5.4 Conclusion 136

5.5 References 137

6 Highly Defined Ferroelectric Block Copolymer-based Dielectric

Nanocomposites 141

6.1 Introduction 143

6.2 Experimental section 145

6.2.1 Materials 145

6.2.2 Synthesis of hafnium oxide nanorods 145

6.2.3 Ligand exchange 146

6.2.4 Preparation of Polymer Films 146

6.2.5 Characterization 146

6.3 Results and Discussion 147

6.4 Conclusion 157 6.5 References 159 Summary 162 Samenvatting 166 List of publications 170 Acknowledgements 171

1.1 Piezoelectricity and ferroelectricity

Piezoelectricity describes the ability of a material to undergo a change in electrical polarization and to generate an electric potential under applied mechanical stress.[1]

The electric response produced after applying stress on the material increases linearly with mechanical load. The reverse effect also applies; piezoelectric materials experience mechanical deformation as a response to an applied electric field. This electromechanical property was first observed in quartz in 1880 by the brothers Pierre and Jacques Curie, although it was not widely used before World War I. Nowadays, piezoelectric materials have found their applications in sensors and actuators, photonics, energy harvesting and self-powered electronics, to name a few.[2–6]

The origin of piezoelectricity is strongly related to the crystalline structure of a material. It was found that the crystalline lattice of piezoelectric materials is non-centrosymmetric, meaning that applying mechanical stress can induce a change in the position of atoms which leads to the formation of net dipole moments, causing a polarization of the material. The existence of net polarization induces the formation of a charge at the surface of materials, as depicted in Figure 1.1. When no stress is applied, the electric current will not flow through the external electric circuit connected with the material (Figure 1.1a). Depending on the crystalline structure, compressive stress can cause an increase or decrease of the polarization. In both cases, the current starts to flow to balance the surface charge that has been changed after the material was subjected to the mechanical stress (Figure 1.1b,c). Conversely, when a piezoelectric material is electrically loaded, the dipole moment is produced, which results in deformation.

Certain piezoelectric materials, called ferroelectric materials, demonstrate a spontaneous polarization that can be reversed by the application of an external electric field.[7] _{Ferroelectric}

materials consist of regions in which dipole moments have the same orientation, called ferroelectric domains. A newly prepared ferroelectric contains many domains with randomly oriented polarizations, causing a zero overall polarization. Subjecting the material to an electric field induces the orientation of dipoles in the direction of the field and creates a spontaneous polarization (Figure 1.2a). After the field removal, in spite of some dipoles being able to change their orientation, the material preserves the net polarization, called remanent polarization (P_r). The direction of the spontaneous polarization can be inverted by changing the direction of the electric field, which yields the characteristic hysteresis loop behavior of a material (Figure 1.2b). Ferroelectric materials have a coercive field (Ec), defined

as a minimum field necessary to switch the full remanent polarization. At low electric fields, the dipoles inside the material are not affected by the change of the electric field, yielding a linear dielectric behavior. A material starts to polarize when the field value approaches the

Introduction | Chapter 1

1

coercive field, while a further increase in the electric field leads to polarization saturation. Ferroelectric materials, however, demonstrate this property bellow a particular temperature, called the Curie temperature (T_c). Above this temperature, the spontaneous polarization and hysteresis behavior disappear and the material becomes paraelectric.

The ferroelectric materials most widely used so far have been lead zirconate titanate (PZT), barium titanate and similar ceramic materials.[8,9] _{However, a high content of lead in PZT}

restrains the application of this material, while low a Curie temperature, together with a reduced ferroelectric response compared to PZT, limits further implementation of barium titanate. Their application is additionally limited due to their high fragility and hardness, high cost and increased thermal and electric conductivity. In contrast to this, organic Figure 1.1 Piezoelectric effect. (a) No mechanical load, (b) compressive load leading to a decrease in

polarization relative to the no load condition, (c) tensile load leading to an increase in polarization relative to the no load condition.

Figure 1.2 (a) Domain structure of ferroelectric materials, (b) Hysteresis loop of ferroelectric

material. Dashed line corresponds to the poling of the unpoled ferroelectric material with zero net polarization. - - - - - - + + + + + + + + + + + + + + + + + + P z - - - - - - - - - - - + + + + + + + + + + + z - - - - - - + + + + + + + + + + + + + + + + z I=0 P e- I + - + - + - + - P e- I

a

b

c

electric field field removal after

E E P Pr Psat Ec a b

nanoscale.[10,11] _{Their physical properties can be efficiently controlled by proper chemical}

modifications. Furthermore, most of them are biocompatible, which allows their integration inside the human body. Among all ferroelectric polymers, poly(vinylidene fluoride), PVDF, and its copolymers are the most used and have the highest piezoelectric and ferroelectric activity.[12,13]

1.2 Poly(vinylidene fluoride), crystalline structure, co-, terpolymers and

their application

PVDF is the polymer of vinylidene fluoride, prepared by free-radical polymerization of this monomer, mostly performed in aqueous dispersions.[13] _{It has a high content of fluorine}

(contains 59.4 wt. % of fluorine and 3 wt. % of hydrogen atoms), which makes it inert to various solvents and acids and causes its high thermal resistance. Additionally, the presence of fluorine atoms with large electronegativity leads to a large dipole moment perpendicular to the chain. PVDF is a semicrystalline polymer, with 50-70 % crystallinity and a melting temperature of 155-190 °C (depending on the crystalline polymorph, molecular weight and crystallization temperature). PVDF crystals are arranged in typical spherulitic aggregates consisting of radially arranged crystalline lamellae and intervening amorphous regions with a disordered conformation.[14]

PVDF exists in five different crystalline phases, designated as α-, β-, γ-, δ- and ε- phases, depending on molecular chain conformations (Figure 1.3).[15] _{The α-phase is hexagonal and}

consists of polymer chains in the trans-gauche-trans-gauche’ (TGTG’) conformation.[16] _The

unit crystalline cell of the α-phase contains two aligned chains antiparallel to each other causing a compensation of dipoles and the non-polar nature of the α-phase. Chains in the β-phase have an all-trans planar zigzag conformation with all dipoles oriented in one direction, which yields the highest dipole moment of all crystalline phases of PVDF (2.10 D) (Figure 1.3d).[17] _{The presence of a highly polar β-phase is demonstrated to lead to excellent}

piezo- and ferroelectric activity.[18] _{The polar γ- phase can be described as a mixture of the α-}

and β-phase with TTTGTTTG’ chains conformation, whereas the δ-phase is a polar analogue of the α-phase. Both phases, despite their lower dipole moments and polarization, are accountable for the ferroelectric property of PVDF.[19,20]

The paraelectric α-phase is a kinetically favorable phase and is generally obtained from the melt by cooling the polymer at a normal rate or by solution crystallization at a temperature above 70 °C.[21,22] _{The polar analogue, the δ-phase, is obtained by applying high electric}

fields (≥ 170 MV m-1_{) at room temperature on the α-phase.}[19] _{In addition, Martín et al. have}

Introduction | Chapter 1

1

step solid-state processing of a non-polar phase, by applying moderate pressures below the melting temperature.[23] _{The γ- phase is mostly obtained via solvent casting from polar}

solvents at low temperatures, in the presence of poly (ionic liquids) or crystallization at high temperatures.[24–26] _{Since the β-phase exhibits piezo- and ferroelectric properties, special}

attention is dedicated to obtaining or increasing the content of this phase. The β-phase is frequently obtained by mechanical stretching of the α-phase in the temperature range of 70-100 °C[27]_{, from the melt under high pressure}[28] _{or ultrafast cooling}[29]_{, heat-controlled spin}

coating[30]_{, from solution casting at low temperatures from solvents like dimethylformamide}

(DMF) and dimethylacetamide (DMAc)[31,32] _{or by the addition of nucleating agents, such as}

nanoparticles[33]_{, ionic salts}[34] _{or polymers like PMMA}[35] _{(Figure 1.4). One recent application}

of nanofabrication techniques for the preparation of PVDF nanostructures has proven to be also beneficial for the β-phase formation: The crystallization confinement in 3D spherical structures, obtained either by embedding PVDF nanoparticles in another amorphous polymer, or self-assembly of PVDF-based block copolymers yielded the homogeneous nucleation and exclusive formation of the β-crystalline phase.[36,37]

The addition of a comonomer bulkier than VDF in some cases promotes the formation of the β-phase after cooling down from the melt. For example, the incorporation of trifluoroethylene (TrFE) with one extra fluorine atom, weakens intermolecular interactions, expands the inter-chain distance and favors the alignment of polymer chains in an all-trans conformation, independently of processing conditions.[38,39] _{Contrary to TrFE, monomers} Figure 1.3 Unit cells of (a) α, (b) δ, (c) γ, and (d) β forms of PVDF crystals viewed along the c-axes and

schematic chain conformations for (e) TGTG′ (α/δ), (f) TTTGTTTG′ (γ), and (g) all-trans (β) rotational sequences. Red, cyan, and blue spheres represent F, C, and H atoms. Reprinted with permission from Ref. [17]. Copyright 2012 American Chemical Society.

a

b

c

d

This resulted only in the reduction of crystallinity, while the crystalline polymorph obtained from the melt was the same as for PVDF. The preferential crystallization of poly(vinylidene fluoride-co-trifluoroethylene) (P(VDF-TrFE)) into the ferroelectric phase and its easier processing (smooth dense films are obtained from spin coating[42]_{, broader solvent range}

for solvent casting, possibility of ink-jet printing[43]_{) explain its wider application compared}

to PVDF.

P(VDF-TrFE) is a semicrystalline polymer in the whole compositional range (most used and studied copolymers are with VDF compositions in the range of 50 to 80 mol%).[44,45] _Unlike

PVDF, this copolymer demonstrates a Curie temperature (Tc) below the melting point, where

the ferroelectric-to-paraelectric transition occurs (Figure 1.5).[46] _{The crystalline polymorphs}

present in this copolymer have a similar conformation as the phases of PVDF, which often causes confusion in literature. Tashiro et al. introduced a new classification of the P(VDF-TrFE) crystalline phases.[47] _{According to this classification, P(VDF-TrFE), depending on the}

processing conditions and temperature, can exist in three different crystalline phases: the low temperature ferroelectric phase (LTFE) with an all-trans conformation similar to the β-phase of PVDF; the high temperature paraelectric phase (HTPE) with a conformationally disordered chain structure; and the cooled ferroelectric phase (CLFE) that consists mostly of trans sequences with some gauche defects.[48] _{Upon cooling a polymer from the melt,}

it first crystalizes in the HTPE phase and further cooling down results in a paraelectric-to-ferroelectric solid-state phase transition at the Curie temperature. The same transition, just at a higher temperature, takes place during the heating of P(VDF-TrFE) in the LTFE phase. An increase in the amount of TrFE in the copolymer increases the lattice spacing and consequently reduces the T_c and coercive field of the polymer.[44,49] _{The incorporation} Figure 1.4 Methods for obtaining crystalline phases of PVDF.

MELT SOLUTION

α

β

γ

δ

poling high T drawing low T ultradrawing high T annealing very high p

poling very high E

annealing very high T poling at high E solid-state-processing (high p)

Introduction | Chapter 1

1

of different stereoirregular TrFE units leads to different regiodeffects and can cause the transition from the all-trans to the 3/1-helix conformation. Hence, the compositional region 49 mol% ≤ C_VDF ≤ 55 mol% is characterized by the presence of both the ferroelectric all-trans and the 3/1-helix conformation and an antiferroelectric-like double loops hysteresis that occurs at an even lower VDF content.[50]

The switching mechanism and shape of the hysteresis loop of P(VDF-TrFE) can be easily altered by controlling the lattice spacing and size of the crystals. This is achieved either by electron-beam irradiation or incorporation of a third bulky less polar monomer such as CFE and CTFE.[51–54] _{Both methods limit the length of P(VDF-TrFE) crystalline segments and}

increase the lattice spacing, causing weaker coupling forces between ferroelectric domains and easier dipole switching. The incorporation of CFE inside P(VDF-TrFE) increases the interchain distance due to the incorporation of the bulkier monomer. Even though in P(VDF-CFE) copolymer CFE is excluded from the crystalline phase due to its bulkiness, the presence of TrFE and the expansion of the interchain distance grant the successful incorporation of the bulky monomer inside the crystalline structure. The CFE here provides a temporary physical pinning of the P(VDF-TrFE) chains and reduces the size of the crystals and degree of crystallinity.[55,56] _{In this way, polymer chains can move freely and induce the so called relaxor}

ferroelectric behavior. Upon applying the electric field, the dipoles from the P(VDF-TrFE) units align firstly with the direction of the field. At higher field values small CFE defects with a large dipole moment (1.8 D) also show a response and orient themselves in the direction of the field. This altered switching mechanism results in a double hysteresis behavior (DHL) as Figure 1.5 (a) DSC thermograms of P(VDF-TrFE). FE, PE, and M denote ferroelectric, paraelectric,

and melt phases, respectively. (b) WAXD profiles for the uniaxially stretched P(VDF-TrFE) film at different temperatures. Reprinted with permission from Ref. [39]. Copyright 2013 Elsevier.

that shows no response to the applied electric field, causing a narrow single hysteresis loop (SHL) behavior (Figure 1.6b).[52] _{The addition of any other bulkier monomer such as the HFP}

results in the exclusion of HFP units from the crystalline regions and in a reduced degree of crystallinity and coercive field, without the formation of a relaxor ferroelectric.[57] _However,

recently, Li et al. have demonstrated that mechanical stretching could pull HFP units into the crystalline domains, forming strong pinning spots and causing relaxor ferroelectric behavior.[48]

On the other hand, the electron beam irradiation transforms the TrFE units into >CH-CF₃ units and additionally causes the internal crosslinking.[51] _{These new groups present defects}

that not only increase the interchain distance and allow easier dipole switching, but also cause the formation of polar nanodomains. The internal crosslinking points play the role of a strong chemical pinning of P(VDF-TrFE) chains and result in narrow hysteresis loops (Figure 1.6c). All of the previous methods for reducing and tuning the shape of hysteresis loops induce a reduction of the Tc to near room temperature.[54,58] Having in mind that the

ferroelectric (relaxor ferroelectric)-to-paraelectric phase transition is accompanied with the large jump in a dielectric constant, these materials are the polymers with the highest ever measured dielectric constant, offering considerable potential for the application as dielectrics for capacitive energy storage (See 1.4).

Figure 1.6 Schematics of nanodomain formation and corresponding hysteresis loop shapes in

(a) P(VDF-TrFE-CFE), (b) P(VDF-TrFE-CTFE) and (c) e-beam-irradiated P(VDF-TrFE). Reprinted with permission from Ref. [87]. Copyright 2014 American Chemical Society.

Introduction | Chapter 1

1

Figure 1.7 Schematic representation of (a) ferroelectric field-effect transistor, (b) ferroelectric

capacitor (c) ferroelectric diode. Ferroelectric and semiconductor are denoted as FE and SC, respectively.

The existence of spontaneous polarization and bi-stabile dipole switching in ferroelectric polymers makes them valuable functional materials for memory storage. To store information, memory devices have to be able to achieve at least two different states-bits upon applying voltage, and ideally the states should stay stable after voltage removal. Different architectures of memory devices containing ferroelectric polymers have been developed: ferroelectric field-effect transistors (FeFETs), capacitors and resistive memories with diode architecture.[59] _{The working principle of FeFETs relies on the use of ferroelectric}

polymers as gate dielectrics (Figure 1.7a).[60] _{The polarization state of ferroelectrics can lead}

either to the accumulation or depletion of electrons in/from the channel, which results in an increase or decrease in conductivity. By changing the gate voltage, the polarization state of the ferroelectric polymer can be altered, whereas only a small source-drain potential is necessary to read out the polarization state. In this way, an easy and nondestructive reading is achieved. Contrary to this, ferroelectric capacitors consist of a thin ferroelectric film sandwiched between two electrodes (Figure 1.7b).[10,61] _{The application of an electric}

field higher than the coercive field causes the polarization of the ferroelectric polymer. Depending on the polarization direction, two different states can exist and be detected by measuring the switching current. Even though the capacitor device is much simpler than FeFETs, the original state can be changed in a read-out process since the electric field is again applied in order to detect the switching current. Asadi et al. developed a novel architecture for memory storage based on the phase separated blend of the ferroelectric P(VDF-TrFE) and a semiconducting polymer poly(3-hexyl thiophene) (P3HT) (Figure 1.7c).

[61] _{After the phase separation, the active layer consists of a conductive continuous pillar}

embedded in the ferroelectric polymer matrix. The charge injection from the electrode to semiconducting polymers is limited without the ferroelectric polarization, which leads to the low conductivity of the device. After the poling the P(VDF-TrFE) in the right direction, the injection barrier is lowered and a higher current is measured. Therefore, the ferroelectric poling can give either a low or a high current depending on the poling direction.

FE FE

SC

FE SC FE SC FE

mechanical energy, sound and vibration energy into electricity. Additionally, the pyroelectric nature of ferroelectric polymers, where polarization changes with temperature, allows them to be applied as flexible temperature sensors.[63] _{The hybridization of ferroelectric materials} with different inorganic, either magnetic or high dielectric constant, components can furthermore expand the application range of ferroelectric polymers, either as multiferroic devices or dielectrics for capacitive energy storage.

1.3 Polymer-based multiferroic materials

Multiferroics (MF) are materials that simultaneously display ferroelectricity and ferromagnetism (Figure 1.8a).[64] _{This means that they can be both electrically polarized}

and magnetically ordered at the same time. The interest in these materials is mostly driven by the possibility of coupling between these two ferroic orders - the magnetoelectric (ME) effect.[65,66] _{In this way, the electrical polarization of a material can be switched with}

application of a magnetic field and the opposite, the magnetization of the material, can be tailored with the applied electric field. This highly valuable behavior for the application has been found to exist in some single-phase oxides (e.g. BiFeO₃, BiMnO₃), where the coupling between magnetic moments and electric dipoles takes place.[67,68] _{However, the values of ME}

coupling coefficients are not sufficient for practical applications. In addition, the presence of two orders with different phase transition temperatures causes just a finite number of materials to exhibit coupling at room temperature. In order to fulfil practical requirements, multiferroic composites with a large ME coupling coefficient at room temperature are developed.[69–71] _{Compared to single phase multiferroics, MF composites offer wider design}

flexibility by allowing a wide range of both ferroelectric and ferromagnetic materials to be employed for their preparation.

The ME effect in MF composites is not an intrinsic property of the material, but a tensor property mediated by strain, as depicted in Figure 1.8b.[72,73] _{Similar to piezoelectric}

materials, the presence of the magnetic field induces strain in the ferromagnetic phase of the composite. The close contact between the phases allows the transfer of the strain from the magnetic to the piezoelectric component, which produces a change in its electrical polarization. In the same way, the presence of an electric field can induce a change in the magnetization of the ferromagnetic phase.

Even though ceramic-based composites have demonstrated the highest measured ME coupling coefficient, they still have a limited application, mostly due to their brittleness, high cost, and an increased conductivity combined with significant losses. To solve these problems, a lot of attention is transferred to the preparation of lightweight and flexible

Introduction | Chapter 1

1

ferroelectric polymer-based composites.[66] _{For the largest number of composites, the}

copolymer P(VDF-TrFE) is used as a ferroelectric component.

Two types of MF composites have been described in the literature: laminated composites and particulate nanocomposites. Despite much higher values of ME coupling coefficients obtained for laminated composites, the simple fabrication and possibilities for miniaturization and large-scale manufacturing, together with the absence of degradation at the interface, present obvious advantages for the practical application of nanocomposites. MF nanocomposites are predominantly prepared by the mixing of magnetic nano-objects (in most cases ferrite-type nanoparticles) and P(VDF-TrFE) in a good solvent for both components (often dimethylformamide) with the help of some external stimulus (e.g. ultrasonication, heating) and spreading on a glass substrate.[74,75] _{After drying and further}

crystallization of the polymer matrix, nanocomposite films are electrically poled by applying a high electric field to obtain a piezoelectric response of the material.

The interface quality between the two phases is demonstrated to be, mostly on the example of ceramic nanocomposites, one of the most determining factors for the creation of nanocomposites with a large ME response.[76] _{Nevertheless, little research is done on the}

quantitative relation between the ME response of polymer-based nanocomposites and interface effects and is expected to be further explored. The ferroelectric-ferromagnetic interphase architecture is responsible for ME coupling coefficient values and can either increase or reduce them, depending on the proper design. For example, ceramic particulate nanocomposites with a properly tuned and adjusted interface demonstrated a five times higher ME response as compared to microcomposites.[77] _{However, the presence of an} Figure 1.8 (a) Multiferroics combine the properties of ferroelectrics and magnets in single material.

(b) Strain induced magnetoelectric coupling between the ferroelectric and the magnetic phase. Magnetostrictive

H

Piezoelectric

- - -

+ + + + + +

Magnetic field Strain Strain Polarization

- - -

+ + + + + +

H

Magnetoelectric effect

(Strain/Magnetic field ) Χ (Polarization/Strain)

composites with fillers at the nanoscale due to the large interfacial area of these materials. It is reported that the addition of silane surfactant on the surface of nanoparticles reduces the piezoelectric response of the polymer matrix, which leads to a reduced magnetoelectric coupling.[78]

The preparation of multiferroic nanocomposites paved the way for the development of a wide range of advanced materials. For example, the existence of two ferroic orders provides extra degrees of freedom for creating next-generation memory devices. As both ferroelectric and ferromagnetic materials are used for storage of binary information (ferroelectric random-access memories, FeRAMs, and magnetic random-access memories, MRAMs) the coexistence of these two phases grants the realization of the four-state logic in a single device.[79] _{Additionally, the electrical control of magnetization based on the}

magnetoelectric coupling proposes the chance for combining the advantages of both FeRAMs and MRAMs. In MRAMs magnetic bits are written by using a high magnetic field which is slow and requires high writing energy. A possible solution for this problem lies in an easymagnetic bit switching by applying an electric field.[80] _{Next to memory devices,}

magnetoelectric coupling has also found its application in energy harvesting technologies, especially in the coupled magnetic and vibration energy harvesting[81]_{, or for the preparation}

of self-powered magnetic field sensors that transfer magnetic into electrical signals.[82]

1.4 Polymer-based dielectric materials for capacitive energy storage

Dielectric capacitors have attracted both academic and commercial interest as one of the most promising compact technologies for the storage and delivery of stored energy.[83,84]

Compared to other energy storage capacitors (electrochemical or supercapacitors), they offer lower losses and can withstand higher operation voltages, while releasing stored energy in an extremely short period of time creating intense power pulses. These properties allow their application in pulsed power systems, such as defibrillators or lasers, in electronic elements for the conversion of direct to alternating current energy for hybrid electric vehicles or wind turbine generators.[85,86]

Capacitors consist of two metal electrodes separated by an insulating dielectric material (Figure 1.9a). The electric energy inside the material is stored by keeping the opposite charges on the electrodes separated by the dielectric material. The voltage difference between the electrodes forms the electric field across the dielectric. Increasing the electric field results in more energy accumulation inside the capacitor, while decreasing it results in the discharge of the stored energy. As illustrated in Figure 1.9b the stored energy inside the material can be calculated as

Introduction | Chapter 1

21

1

(1)

where E is the applied electric field and D is induced electric displacement. Electric displacement values are related to the polarization of the dielectrics D = P + ε₀E. For linear

dielectrics the following equation can be written as

(2)

Here ε0 = 8.85 x 10-12 F m-1 is the vacuum permittivity, εr is the relative permittivity of a

dielectric material and E_b is the breakdown strength. From this equation it becomes obvious that the energy density of the capacitor is determined by the values of relative permittivity and breakdown of the dielectric. The dielectric constant of the material is a macroscopic property highly related to its polarizability. There are in total five types of polarization: electronic, atomic, orientational, ionic and interfacial polarization.[17] _{Out of all of them, only}

the orientational polarization based on a response of permanent dipoles to the applied electric field is shown to be suitable for fast energy discharge in high-power energy storage devices.[87] _{On the other side, breakdown strength represents the critical field where the}

electrically conductive path is formed in a material, turning it from an insulator to a conductor, and in case of polymers is mostly affected by free volume, mechanical strength and thermal conductivity of a material.[88] _{Both a high relative permittivity and breakdown strength are}

necessary to achieve high energy density of dielectrics. However, these two elements are in most cases mutually exclusive. Giving the quadratic relation between breakdown strength and energy density, the high breakdown strength is crucial for the improvement of energy density values.

𝑈𝑈

= %

𝐸𝐸 𝑑𝑑𝐷𝐷

𝑈𝑈

=

1

_{2 𝜀𝜀}

𝜀𝜀

𝐸𝐸

(1)

where E is the applied electric field and D is induced electric displacement. Electric displacement values

are related to the polarization of the dielectrics D = P + ε

E. For linear dielectrics the following equation

can be written as

(2)

Here ε

= 8.85 x 10

F m

is the vacuum permittivity, ε

is the relative permittivity of a dielectric

material and E

is the breakdown strength. From this equation it becomes obvious that the energy

density of the capacitor is determined by the values of relative permittivity and breakdown of the

dielectric. The dielectric constant of the material is a macroscopic property highly related to its

polarizability. There are in total five types of polarization: electronic, atomic, orientational, ionic and

interfacial polarization.

_{Out of all of them, only the orientational polarization based on a response}

of permanent dipoles to the applied electric field is shown to be suitable for fast energy discharge in

high-power energy storage devices.

_{On the other side, breakdown strength represents the critical}

field where the electrically conductive path is formed in a material, turning it from an insulator to a

conductor, and in case of polymers is mostly affected by free volume, mechanical strength and thermal

conductivity of a material.

_{Both a high relative permittivity and breakdown strength are necessary}

to achieve high energy density of dielectrics. However, these two elements are in most cases mutually

exclusive. Giving the quadratic relation between breakdown strength and energy density, the high

breakdown strength is crucial for the improvement of energy density values.

𝑈𝑈

= %

𝐸𝐸 𝑑𝑑𝐷𝐷

𝑈𝑈

=

1

_{2 𝜀𝜀}

𝜀𝜀

𝐸𝐸

(1)

where E is the applied electric field and D is induced electric displacement. Electric displacement values

are related to the polarization of the dielectrics D = P + ε

E. For linear dielectrics the following equation

can be written as

(2)

Here ε

= 8.85 x 10

F m

is the vacuum permittivity, ε

is the relative permittivity of a dielectric

material and E

is the breakdown strength. From this equation it becomes obvious that the energy

density of the capacitor is determined by the values of relative permittivity and breakdown of the

dielectric. The dielectric constant of the material is a macroscopic property highly related to its

polarizability. There are in total five types of polarization: electronic, atomic, orientational, ionic and

interfacial polarization.

_{Out of all of them, only the orientational polarization based on a response}

of permanent dipoles to the applied electric field is shown to be suitable for fast energy discharge in

high-power energy storage devices.

_{On the other side, breakdown strength represents the critical}

field where the electrically conductive path is formed in a material, turning it from an insulator to a

conductor, and in case of polymers is mostly affected by free volume, mechanical strength and thermal

conductivity of a material.

_{Both a high relative permittivity and breakdown strength are necessary}

to achieve high energy density of dielectrics. However, these two elements are in most cases mutually

exclusive. Giving the quadratic relation between breakdown strength and energy density, the high

breakdown strength is crucial for the improvement of energy density values.

Figure 1.9 (a) Schematic representation of electrostatic capacitor that consist of two metal

electrodes separated by an insulating dielectric material. (b) D−E hysteresis loops for both linear and nonlinear dielectric polymers under high-field switching.

- - - + + + + + + + + V - + + - + - + - + - + - + - + - E P Electric field Elect ric d isp lace me nt _{released energy} a b

Among all candidates, polymeric materials are particularly favored, mostly due to their large-area, low cost production of high-quality flexible films using straightforward polymer processing methods.[89,90]] _{Additionally, polymer materials possess a significantly higher}

breakdown strength and a unique self-healing behavior. After the breakdown, an open circuit instead of a short circuit is formed at the failure point, since the part of metal electrodes evaporates from the breakdown spot due to the heat generated during breakdown.[91]

At the moment, current technology commonly uses metalized biaxially oriented polypropylene (BOPP) films. This nonpolar polymer has an exceptionally high breakdown strength (>700 MV m-1_{) and a low dielectric loss (<0.02 %), as well as good mechanical strength.} [92] _{Nevertheless, the nonpolar nature of BOPP is responsible for its low dielectric constant}

(εr = 2.2) and consequently low energy density values even at extremely high electric fields

(≈1-2 J cm-3_).[93] _{This definitely clarifies the critical requirement for the preparation of novel} dielectric materials with improved energy density.

The incorporation of polar groups into polymer chains commonly induces an increase in the dielectric constant compared to BOPP.[94,95] _{Especially ferroelectric polymers, like PVDF}

and co- or terpolymers thereof exhibit a high polarization and an intrinsically high dielectric constant (ε_r of PVDF is around 10), due to the high dipole moment and packing density of C-F

bonds, and are currently the most studied polymers for high energy storage applications.[17]

Pristine PVDF and P(VDF-TrFE), despite displaying a significantly higher dielectric constant and displacement than BOPP, have a ferroelectric nature that is linked with high energy losses. As demonstrated in Figure 1.10a a small part of the stored energy can be used after the discharge. Around half of the stored energy actually produces heat inside the material and causes the premature breakdown and reduction in the durability of the capacitor. As already discussed in 1.2 the reduction of the size and the increase in the distance between ferroelectric domains, achieved by incorporating defects into P(VDF-TrFE) in the form of CFE or CTFE units or crosslinking spots, leads to relaxor ferroelectric behavior with slimmer hysteresis loops (Figure 1.10b,c). For example, P(VDF-TrFE-CFE) relaxor ferroelectric manifests the highest dielectric constant among all known dielectric polymers at room temperature (ε_r > 50), and the discharged energy density of this polymer can reach values close to 10 J cm-3_.[96]

The main drawback of relaxor ferroelectric polymers is related to the saturation of polarization at low electric fields and any further increase in the electric field produces only a small change in the displacement and stored energy. Therefore, researchers shifted their focus to copolymers of VDF with a large amount of bulky CTFE or HFP used to reduce the hysteresis and dielectric constant, so that the saturation can take place at higher electric fields. These copolymers have demonstrated enviable values of discharged energy density

Introduction | Chapter 1

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Figure 1.10 (a) Schematic D−E hysteresis loops for (a) normal ferroelectrics, (b) relaxor ferroelectrics,

and (c) antiferroelectrics.

as 25 J cm-3 _{for P(VDF-CTFE) and even 30 J cm}-1 _{for P(VDF-HFP).}[90,97] _{It is worth noting that the}

high breakdown strength of the polymer is necessary to achieve these high values of energy density. Some of these polymers, such as the high energy density P(VDF-HFP), have already been produced on large-scale in the form of 2 µm thick capacitor films. Again, the efficiency of these polymers is reduced compared to BOPP and other linear dielectrics, particularly at a higher electric field caused by the poling and growth of polar ferroelectric domains. Conduction losses also contribute to the reduction of the charge-discharge efficiency of polymers.

Several approaches have been developed to reduce both ferroelectric and conducting losses and to obtain the ultimate polymer material for high energy density. Crosslinking of P(VDF-CTFE) with triallyl isocyanurate resulted in reduction of the hysteresis and an increase in the polarization at the same time, which is opposite to all so far reported examples of ferroelectric polymer networks in which crosslinking reduces crystallinity and causes a decrease in the polarization (Figure 1.11a,b).[98] _{Optimized crosslinked films showed a}

charge-discharge efficiency of ≈83 % at 400 MV m-1_{, compared to melt-stretched P(VDF-CTFE) with}

65%. Dielectric losses can also be reduced by incorporating lower dielectric polymers in the form of either blends or graft copolymers. The formation of the confining low dielectric constant polystyrene layer at the periphery of PVDF crystals suppressed the hysteresis loop in P(VDF-CTFE)-g-PS graft copolymers and resulted in antiferroelectric-like or even linear dielectric behavior with obviously reduced losses (Figure 1.11c-e).[99,100] _{The mechanism of}

the nanoconfinement effect on the switching behavior of ferroelectric polymers will be explained in detail in 1.8. Another particularly attractive method refers to the application of coextruded multilayer films consisting of alternating layers of a ferroelectric polymer with a high energy density, low polarity, high breakdown strength polymer such as polycarbonate

a

_b

c

d

e

f

_g

Figure 1.11 (a) Discharged energy density, (b) charge–discharge efficiency of ferroelectric polymer

networks made via crosslinking of P(VDF-CTFE) with triallyl isocyanurate. Adapted with permission from Ref. [98]. Copyright 2013 Nature Group. (c) Synthesis of P(VDF-TrFE-CTFE)-g-PS graft copolymer using ATRP. D-E loops of (d) P(VDF-TrFE) 93/70 and (e) P(VDF-TrFE-CTFE)-g-PS (14 %) demonstrating change in the switching nature after incorporation of the insulating polymer chains in the structure of P(VDF-TrFE). Adapted with permission from Ref. [99]. Copyright 2011 American Chemical Society. (f) Schematic of charge migration and charge buildup in multilayered films depending on the layer thickness. (g) Average breakdown strength of 32-layer (solid circles) and 256-layer (open circles) PC-P(VDF-HFP) multilayers as a function of volume concentration. Adapted with permission from Ref. [101]. Copyright 2012 American Chemical Society.

Introduction | Chapter 1

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the breakdown strength of the films with ultra-low layer thickness of only 50 nm gives considerably higher breakdown strength than obtained for either constituent (Figure 1.11g). In addition to the development of all-polymer dielectric materials, much attention is devoted to the preparation of high energy density nanocomposites consisting of inorganic nano-objects inside the ferroelectric polymer matrix.[103] _{The whole concept is based on}

combining two components–polymer with high flexibility and breakdown strength and ceramics with a high dielectric constant and polarization. A combination of these two components should in theory give higher energy density than it is possible to achieve with individual materials. However, this task still represents a real challenge. The behavior of nanocomposites is influenced by various factors, such as the interface between fillers and the polymer matrix, the shape, orientation and spatial distribution of the fillers.[104]

Additionally, the large mismatch in the dielectric constant between the ceramic fillers and the matrix causes a great increase in the local electric field around the fillers, reducing the breakdown strength of nanocomposites.[105,106]

Interfaces in nanocomposites are defined as a transitional region of a finite thickness between nanoparticles and the polymer matrix where the properties differ from either the nanofiller or the polymer matrix. The reduction of filler sizes to nanometer dimensions causes an increase in the volume fraction of the interfaces to over 50 %, showing the dominant role of the interface in the dielectric response of nanocomposites.

After the addition of nanoparticles in the polymer matrix, the surface of the particles becomes charged because of different chemical potentials between the matrix and particles. Figure 1.12 shows the distribution of the charges in the interfacial region for positively charged nanoparticles. The charges in the nanoparticles cause the absorption of negatively charged counterions on the particle surface to form a Stern-Helmholtz double layer. Next to this layer is the Gouy-Chapman diffused layer with both positive and negative ions. Since this layer is filled with mobile charges, it plays a predominant role for the dispersion of nanoparticles and also dielectric properties of nanocomposites. Therefore, the manipulation of the interface became the most investigated approach for improving the energy density of nanocomposites. Current manipulations of the interface are mostly related to improving dispersions of nanoparticles and manipulating the dielectric constant mismatch.

The dispersion of nanoparticles with a large surface-to-volume ratio is often accompanied by the agglomeration of nanoparticles during the preparation of nanocomposites.[107]

compatibility between nanoparticles and the polymer matrix should be improved. The most common method to improve compatibility is the surface modification of nanoparticles. The ligands used for the modification should contain a functional group that strongly interacts with the surface of nanoparticles and the part that grants good miscibility with the polymer matrix. A large variety of ligands, e.g. silane coupling agents, dopamine-based ligands, phosphonic acid-based ligands, ethylenediamine have been studied in order to improve the dispersion of nanoparticles.[107–109] _{Furthermore, the simple hydroxylation of the nano-object}

surface is used to increase the concentration of hydroxyl functional groups that can interact with fluorine atoms making hydrogen bonds.[108] _{Another approach, consisting of grafting}

a polymer to the surface of nanoparticles, is proven to not only contribute to the strong interaction with the polymer matrix, but also that the polymer can act as a matrix itself. For example, grafting P(VDF-HFP)-g-poly(glycidyl methacrylate) to the surface of BaTiO3

nanoparticles yields the formation of core-shell nanocomposites with a homogeneous dispersion and enhanced dielectric properties.[109] _{Similar results are obtained after grafting}

fluorinated acrylate monomers from the surface of BaTiO₃ nanoparticles.[110]

The second problem associated with dielectric polymer nanocomposites is the existence of the dielectric mismatch at the interface caused by the difference in the dielectric constant between the filler and the polymer matrix (Figure 1.13). As a consequence, it causes the presence of regions with a significantly increased nominal electric field values that are specifically sensitive to an electrical breakdown.[111] _{To avoid this phenomenon, a careful}

choice of polymer and ceramic fillers with a comparable dielectric constant is necessary. Li et

al. demonstrated that the incorporation of titanium dioxide nanoparticles with a dielectric Figure 1.12 (a) Diffuse electrical double layer produced by a positively charged particle in a polymer

matrix with the resulting electrical potential distribution. (b) Conduction via diffuse double layers in nanocomposites. Adapted with permission from Ref. [103]. Copyright 2016 American Chemical Society. - - - - - - - - - + + + + + + + + +

a

b

Introduction | Chapter 1

1

constant of 47 in P(VDF-TrFE-CFE) with εr = 42, even though it does not show an increase

in the dielectric constant, leads to a significantly improved discharged energy density and breakdown strength.[112] _{The other way to address this issue is to form a gradient in the}

dielectric constant between the phase in the form of core-shell fillers. Liu and co-workers showed that the formation of a silica layer on the surface of the BaTiO₃ decreases the dielectric loss and improves the discharged energy density of PVDF-based nanocomposites compared to the neat BaTiO₃. [113]

In addition to the polymer-filler interface, the dielectric anisotropy strongly influences the dielectric response of nanocomposites, and is associated with the shape, orientation and spatial distribution of fillers (Figure 1.14). The phase field modelling of nanocomposites demonstrates that the alignment of nanorods perpendicular to the electric field reduces the values of the local electric field, which results in an increase of the breakdown strength.[105]

An even larger improvement of dielectric properties is achieved using 2D nanosheets, since they make a strong barrier to a charge migration in the direction of the field.[114] _{In addition,}

it has been observed that the formation of a multilayered structure with alternation layers consisting of different types of fillers can significantly increase the discharged energy density and breakdown strength. For example, Liu and co-workers incorporated boron nitride nanosheets as outer layers into PVDF to obtain high breakdown strength, while the PVDF inner layer filled with barium strontium titanate nanowires grants a high dielectric constant. After optimization, corresponding nanocomposites display a discharged energy density of 20.5 J cm-1 _{with breakdown strength of 588 MV m}−1_{, significantly better than those}

of single layered films.[115]

Figure 1.13 (a) Schematic illustration of modeled nanocomposite structure. (b) Space charge and

(c) conductivity distribution inside the nanocomposite. Reprinted from [110], with the permission of AIP Publishing.

1.5 Block copolymers: self-assembly, confined crystallization and selective

dispersion of nano-objects

Previous research on ferroelectric polymers and their devices have demonstrated that their ferroelectric and piezoelectric properties can be easily improved by preparing ferroelectric polymer nanostructures using different nanofabrication techniques.[10,116,117]

The confinement of crystallization inside nanostructures causes a preferential orientation of polymer chains, which grants an easier switching of ferroelectric dipoles. Different procedures have been used for the preparation of ferroelectric nanostructures, such as electrospinning, nanoimprinting or using templates as anodized alumina.[116,118,119] _However,

these methods have different drawbacks that limit their wider application. Some of them demand long fabrication times, while others are still expensive and complicated for mass production. In addition, the use of ferroelectric polymers is often accompanied by high dielectric losses and reduced breakdown strength compared to other nonpolar polymers.

[101,120] _{They also lack the functionality that can be used for the improved dispersion of}

nano-objects crucial for the preparation of various nanocomposites or for obtaining new valuable properties (i.e. better adhesion to substrates or electrodes, better thin film formation).[107,121] Figure 1.14 3D simulations of microstructure effects on breakdown: (a) breakdown phase

morphology in the nanocomposites with different microstructures, (b) evolutions of the breakdown phase volume fraction under applied electric fields, and (c) extracted breakdown strengths for corresponding nanocomposites. Reprinted with permission from Ref. [105]. Copyright 2018 John Wiley & Sons Inc.

Introduction | Chapter 1

1

For these reasons, it would be of great interest to develop an alternative method for the preparation of ferroelectric polymer nanostructures that can simultaneously resolve all the above-mentioned issues.

Block copolymers consist of two or more chemically different polymer chains attached at their ends. In general, the small mixing entropy and unfavorable mixing enthalpy cause a low miscibility between different polymer chains. The connectivity between the blocks, however, prevents their macrophase separation and allows them to microphase separate on the nanometer scale. The balance between the interaction enthalpy, that induces the stretching of polymer chains, and elastic entropy that resists the stretching, controls the self-assembly of block copolymers. The microphase separation of block copolymers can be regulated by the total degree of polymerization (N), the volume fractions of the blocks, architecture and the Flory-Higgins segment-segment interaction parameter (χ) that measures the incompatibility between the blocks.[122–124] _{The degree of microphase separation of block}

copolymers is determined by the product χN. With increasing temperature, the product χN decreases and the system becomes homogeneous above a certain temperature (T_ODT). At

different ratios between the blocks, different morphologies can be obtained for diblock AB and triblock ABA copolymers: from lamellar, to hexagonally packed cylinders, body-centered spheres, closed-packed spheres and bicontinuous gyroid phase (Figure 1.15).[125]

More complex structures can be formed by incorporating an additional polymer chain in the form of triblock ABC copolymers or by changing the architecture of block copolymers (star or branched instead of linear).[126]

The incorporation of a crystallizable block into the structure of block copolymers complicates the self-assembly process due to the strong influence of crystallization on the morphology formation. The thickness of crystals is comparable to the size of the nanodomains obtained via microphase separation. Therefore, the interplay between crystallization and microphase separation determines the type of the obtained morphology. The behavior of crystalline-amorphous block copolymers above the melting temperature (T_m) is the same as for amorphous-amorphous block copolymers. Upon cooling down, the crystalline block undergoes a crystallization process. The outcome of the crystallization depends strongly on the molecular characteristics of both blocks.[127,128] _{In the case of weakly segregated or}

homogeneous block copolymer systems the structure preformed in the melt is not resistant enough against the crystallization, and the microphase separated structure is destroyed. Instead, the asymmetric lamellar morphology, consisting of alternating crystalline-amorphous layers, is formed. This mode of crystallization is called break-out crystallization (Figure 1.16).[129] _{In contrast to this, when the microphase separation is a stronger driving}

morphology obtained in the melt is preserved after the crystallization (Figure 1.16).[130]

Two types of confinement exist: hard and soft. When the glass transition temperature (T_g) of the amorphous block is significantly higher than the crystallization temperature of the crystalline block, the crystallization is confined to the hard domains formed after the vitrification of the amorphous block. The hard domains are difficult to deform during the crystallization and consequently bad quality crystals with a low degree of crystallinity are formed.[131] _{The crystallization can be confined even if the amorphous block is not vitrified}

before the crystallization, in case the segregation strength between the block is high enough. Here, the soft domains can slightly deform during the crystallization, which results in the increased degree of crystallization.[132]

The crystallization behavior of polymer chains confined within isolated block copolymer nanodomains is substantially different from pristine polymers. The confined crystallization inside spherical domains (in case of hard confinement also often cylindrical) is always associated with a strong reduction in the crystallization temperature. This can be explained by the change in the mechanism of the crystal nucleation-from heterogeneous to Figure 1.15 Schematic representation of the bulk morphologies observed in AB diblock copolymers

as function of volume fraction of block B (f_B). Reproduced from Ref. [124] with permission from The Royal Society of Chemistry.

Introduction | Chapter 1

1

Figure 1.16 Possible crystallization mechanisms of microphase-separated crystalline–amorphous

diblock copolymers. Reprinted with permission from Ref. [133]. Copyright 2015 Elsevier.

homogeneous nucleation.[127] _{The number of spherical domains highly exceeds the amount}

of nucleus forming impurities present in the block copolymer, causing the homogeneous nucleation. Additionally, the rate of crystallization is increased for the crystallization inside nanodomains. The crystal growth is almost instantaneously finished inside small dimension structures formed using the block copolymer self-assembly. This is completely different from the crystallization of the pristine homopolymer, in which a sigmoidal increase in the crystallinity is observed over time.[133]

Well-ordered block copolymer nanostructures found their application for the preparation of thermoplastic elastomers, nanopatterning applications, the preparation of nanostructured materials, orientation of the crystalline stems, in membrane science and for optoelectronic devices.[134–138] _{Block copolymers also present a promising platform for the dispersion of}

nano-objects (NOs) inside the polymer matrix and the preparation of nanocomposites.[139,140]

The use of block copolymers for the dispersion of NOs offers various advantages compared to the dispersion of NOs in homopolymers. The self-assembly of block copolymers allows

copolymers is a consequence of the relationship between the conformational entropy of polymer chains, translational entropy of NOs and enthalpy caused by the creation of the polymer-NO interface.[141] _{The entropy loss related to the NOs insertion is dictated by the}

size of NOs and block copolymer domains, while the interaction between polymer chains and NOs determines the enthalpy of insertion.[142] _{The important feature of block copolymer}

nanocomposites is the selective dispersion of NOs, where nanoparticles tend to localize in one of the domains formed by the self-assembly. The localization of NOs is mostly governed by the size, concentration and surface functionalization of the NOs. Small nanoparticles often localize at the interface between blocks, since the gain in translational entropy prevails over the loss in the conformational entropy of polymer chains after the nanoparticle incorporation. Conversely, the loss in conformational entropy is a dominant factor for the dispersion of larger nanoparticles, which leads to their localization in the interior of polymer domains (Figure 1.17a).[143] _{The increase in NO concentration can also strongly control the}

NOs location. With increasing the concentration of NOs in block copolymers, the amount of NOs inside polymer domains tends to increase (Figure 1.17b).[144] _{However, after a certain}

concentration has been reached, a loss in conformational entropy becomes dominant and prevents the uniform dispersion of NOs. Therefore, in some cases, NOs either start to macrophase separate or to localize inside the other block. Nevertheless, the surface chemistry of NOs has the strongest influence on the specific localization of NOs inside block copolymer nanodomains, and can be tuned by the type, grafting density and molecular weight of the ligand.[139]

The addition of NOs inside block copolymers can significantly influence the morphology of block copolymers. The incorporation of large nanoparticles often results in the formation of a macrophase separated morphology with NO-rich and poor regions. It is also observed that a change in the concentration of NOs inside block copolymers can result in order-to-order and order-to-order-to-disorder-to-order phase transitions. For example, the selective incorporation of mercaptoacetic acid-modified CdS nanoparticles inside poly(4-vinylpyridine) domains of poly(styrene-b -4-vinylpyridine) induces the transition from a hexagonally packed cylinder into a lamellar structure (Figure 1.17c). [145]_{Further increase in the nanoparticle concentration}

causes an order-to-disorder phase transition. Conversely, a favorable interaction between NOs and one of the blocks can induce the disorder-to-order phase transition (Figure 1.17d,e).

[146] _{This is often observed in the systems where NOs form selective hydrogen bonding with}

one of the segments of a block copolymer.

Block copolymers have different influence on the organization of NOs of different shape. While they can only cause the selective dispersion of nanoparticles, the addition of anisotropic 1D nanosheets or 2D nanorods inside block copolymer domains can results in the preferred orientation of the fillers. In this way, the conformational entropy loss of

Introduction | Chapter 1

1

polymer chains is minimized. The aspect ratio of fillers plays a major role in their orientation. Shorter nanorods tend to orient to a lesser extent, while larger rods have a strong preference to orient parallel to the domain interface.[147]

1.6 Synthesis of PVDF-based block copolymers

To achieve the self-assembly of block copolymers into well-ordered morphologies, the

a b

c d

e

Figure 1.17 (a) Influence of NP size on particle localization: small nonselective nanoparticles locate

in the interfacial region, while large selective particles locate in the interior of the domain. Reprinted with permission from Ref. [143]. Copyright 2003 American Chemical Society. (b) A concentration-dependent localization of nanoparticles inside block copolymer nanodomains: at lower concentration nanoparticles are locates in the center of lamellar domains, while the amount of nanoparticles inside polymer domains tends to increase with their content. Adapted with permission from Ref. [144]. Copyright 2009 Nature Group. (c) Cylindrical-to-lamellar order-to-order phase transition due to CdSe nanoparticles incorporation in PS-P4VP. Adapted with permission from Ref. [142]. Copyright 2005 American Chemical Society. (d) and (e) Nanoparticle-driven assembly of block copolymers. Adapted with permission from Ref. [146]. Copyright 2011 American Chemical Society.

Even though various synthetic methods for the preparation of well-defined block copolymers have been advanced in recent years, especially in the field of controlled radical polymerizations, the synthesis of fluorinated PVDF-based block copolymers still presents a very difficult task. Many fluorinated monomers propagate with very reactive radicals that often undergo side reactions and are in the gaseous state under ambient conditions. This adds more complexity into the synthetic procedure, requiring the use of high temperature high pressure equipment. Additionally, the small difference in the size of substituents in the VDF monomer allows not only the head-tail addition of the monomers in the propagation step, but also a significant number of head-head units that have negative influence on the ability of the VDF to be polymerized in a controlled fashion with satisfying molecular weights.[13,148] _{The preparation of PVDF-based block copolymers has so far been achieved}

using three main methods: preparation of functional PVDF telomers followed by chain extension, using reversible deactivation radical polymerization (RDRP) techniques, such as iodine transfer polymerization (ITP) and reversible addition-fragmentation chain transfer (RAFT) polymerization and using functional benzoyl peroxide initiators.

The free-radical polymerization of the VDF and other fluorinated monomers, in the presence of halogen containing chain transfer agents that can undergo the C-X bond cleavage, results in the preparation of PVDF-X telomers that can be further used in the preparation of block copolymers (Figure 1.18). For example, Destarac et al. performed the atom transfer radical polymerization (ATRP) of different monomers (styrene, methyl acrylate) starting from the trichloromethyl-terminated PVDF prepared via a telomerization reaction using chloroform as a chain transfer agent.[149] _{A Linear increase in molecular}

weight with a monomer conversion and narrow molecular weight distribution proved the controlled nature of the ATRP reaction. However, the telomers used in the reaction had low molecular weight and a degree of polymerization between 1 and 16. On the other side, the emulsion copolymerization of VDF and HFP in the presence of the same chain transfer agent gave polymers with significantly higher molecular weight (up to 24 kg mol -1_).[150] _{The subsequent ATRP reaction led to a formation of P(VDFr HFP)bPS and P(VDFr}

--HFP)-b-PMMA block copolymers with relatively low dispersities of 1.2-1.5. Another route to synthesize block copolymers containing PVDF starting form PVDF-based telomers was explored by Holdcroft’s group.[151,152] _{Here, they performed the coupling of low molecular}

weight dibromo-terminated PVDF and α,ω-bis(dihydroxy)- poly(arylene ether sulfone), PAES, using a condensation reaction in the presence of a strong base NaH. The prepared multiblock copolymers demonstrated a relatively broad dispersity around 2, characteristic for a polycondensation. The subsequent sulfonation of the PAES-b-PVDF block copolymer was successfully achieved and the materials demonstrated great potential for fuel cell membrane applications.