BZT Barium zirconium titanate C Concentration C * Critical overlap concentration CFE Chlorofluoroethylene Cp Volume heat capacity CTFE

Keywords:

i Engineering the solid-state structure of vinylidene fluoridebased systems for efficient energy storage and refrigeration

Abstract:

Vinylidene fluoride (VDF)-based semi-crystalline polymers comprise different polar and apolar polymorphs in their solid-state structure, which induce a wide range of properties from ferroelectric and relaxor ferroelectric. This thesis concerns the engineering the solid-state structure of these VDF-based polymers, more specifically VDF-based terpolymers containing Trifluoroethylene (TrFE), chlorofluoroethylene (CFE), such as P(VDF-TrFE-CFE) and/or chlorotrifluoroethylene (CTFE) in their compositions, either in neat polymer, but also blended with poly methyl methacrylate (PMMA) and in organic/inorganic nanocomposite systems containing barium titanate-based perovskite oxide nanoparticles.

It has been shown that the solid-state structure formation in the VDF-based polymers is highly dependent on parameters governing the solidification process, including the casting temperature, casting time, and entanglement density of the polymer in solution. By changing these parameters, various solid-state structures, and resulting ferroelectric and ferroelectric relaxor behaviors, were obtained in VDF-based polymers with ~70/30 VDF/TrFE ratio in their compositions. The formation of these solid-state structures is also affected by the entanglement density of the chains in solution, which affects the distribution of CFE or CTFE termonomer, and trans-gauche and all-trans conformations, in the solid-state structure, and thus induce different ferroelectric and relaxor ferroelectric properties. It has been also shown that the induced tg conformations in the structure are not stable during the storage at room temperature and they evolve into all-trans conformations after a while. This instability is minimized by increasing the casting temperature to 105°C for the concentrated solutions.

It has been shown that in the PMMA:P(VDF-TrFE-CFE) blend systems, PMMA with the dipoledipole interactions with VDF chains, arrange these chains in more polar structures. PMMA is well-miscible with VDF-based polymers in the amorphous phase; thus the blending decreases the overall crystallinity and lamellae thickness of the VDF-based polymers. However, this decrease in overall crystallinity is accompanied with an increase in a-b lateral extension, which includes the polar direction. More importantly, it has been proven that PMMA can increase the breakdown strength of the system, which is highly desirable for the energy storage applications.

ii In order to produce polymer nano-composites, ferroelectric barium titanate 100-200nm-sized nanoparticles were synthesized via two hydrothermal and polyol methods. To manipulate the ferroelectric behavior of these nanoparticles, various amounts of Zr +4 ions were introduced into the structure through the synthesis of barium zirconium titanate (BaZrxTi(1-x)O3) with three different Zr contents (x=0.12, 0.2, and 0.3). Pseudo-cubic to cubic phase transitions are observed in 115°C for the synthesized barium titanate nanoparticles, which are shifted to lower temperature when the Zr +4 ions are added to the structure. Broad polar-to-apolar phase transitions are observed at 70-80°C for nanoparticles with 12 mol.% Zr with respect to Ti, while broader polar-to-apolar phase transitions at 20-30°C are detected for the nanoparticles with 20 mol.% Zr. It has proven that organic/inorganic nanocomposite systems comprising VDF-based polymers and barium titanate-based nanoparticles with polar-to-apolar phase transitions in the same temperature range can show enhanced polarizability.

Finally, energy storage ability and the electrocaloric effect of the designed systems have been investigated, and improvement in these applications has been detected in the selected systems.

Keywords: VDF-based polymers, ferroelectricity, nanocomposite, energy storage, electrocaloric cooling iii

Resumé:

Les polymères semi-cristallins à base de fluorure de vinylidène (VDF) possèdent différents polymorphes polaires et apolaires dans leur structure à l'état solide, qui sont responsables pour une large gamme de comportements ferroélectriques et ferroélectriques relaxeurs.

Cette thèse aborde l'ingénierie de la structure à l'état solide de ces polymères à base de VDF, plus précisément des terpolymères à base de VDF comprenant du trifluroethylene (TrFE), chlorofluoroéthylène (CFE) comme P(VDF-TrFE-CFE) et/ou du chlorotrifluoroéthylène (CTFE), à la fois sous leur forme pure, en mélange avec du poly méthacrylate de méthyle (PMMA) ou dans des systèmes nanocomposites organiques/inorganiques, contenant des nanoparticules d'oxyde de pérovskite à base de titanate de baryum.

Nous avons montré que la nature des structures à l'état solide dans ces polymères à base de VDF est fortement dépendante des paramètres qui gouvernent le processus de solidification, tels que la température et durée de dépôt ou la densité d'enchevêtrement des polymères en solution. En controllant ces paramètres, une large variété de structures à l'état solide, ainsi que les different comportements ferroélectriques ou relaxeurs correspondants, a été obtenue dans des polymères à base de VDF avec ~ 70/30 rapport VDF/TrFE. La formation de ces structures à l'état solide est également affectées par la densité d'enchevêtrement des chaînes en solution, qui influence la distribution de termonomère CFE ou CTFE, ainsi que la distribution des conformations trans-gauche et all-trans, dans la structure à l'état solide, conduisant à différentes propriétés ferroélectriques et relaxeurs. Il a également été montré que les conformations tg induites dans la structure ne sont pas stables lors du stockage à température ambiante et qu'elles évoluent vers des conformations tout-trans au bout d'un certain temps. Cette instabilité est minimisée en augmentant la température de coulée à 105°C pour les solutions concentrées.

Nous avons également montré que dans les systèmes polymères à base de mélanges PMMA: P(VDF-TrFE-CFE), la présence du PMMA, arrange les chaînes VDF dans des structures plus polaires sous l'influence d'interactions dipôle-dipôle. Le PMMA amorphe est bien miscible avec les polymères fluorés dans la phase amorphe, ainsi le mélange entraine une diminution de la cristallinité globale et un amincissement des lamelles cristallines des polymères fluorés.

Cependant, cette diminution de la cristallinité globale s'accompagne d'une augmentation de l'extension latérale a-b, qui contient la direction polaire. D'autant plus, nous avons montré que le PMMA peut augmenter de façon significative les champs de claquages, ce qui est hautement souhaitable pour les applications dans le stockage d'énergie à base de dielectrique. iv Afin de produire des nano-composites, nous avons également synthétisé des nanoparticules ferroélectriques de titanate de baryum de taille 100-200 nm, via deux méthodes différentes : hydrothermales et polyol. Pour manipuler le comportement ferroélectrique de ces nanoparticules diverses quantités d'ions Zr +4 ont été introduites dans la structure à travers la synthèse de titanate de baryum et de zirconium (BaZrxTi(1-x)O3) avec trois teneurs différentes en Zr (x = 0,12, 0,2, et 0,3). La transition de phase pseudo-cubique à cubique, qui est observé à 115 ° C pour les nanoparticules de titanate de baryum, est déplacé vers des températures plus basses lorsque les ions Zr +4 sont ajoutés à la structure. De larges transitions de phase polaire à apolaire sont observées à 70-80°C pour les nanoparticules avec 12% molaire de Zr par rapport à Ti, tandis que des transitions de phase polaire à apolaire encore plus larges sont détectées à 20-30°C pour les nanoparticules avec 20% molaire de Zr. Nous avons que le choix de systèmes nanocomposites organiques/inorganiques comprenant des polymères à base de VDF et des nanoparticules à base de titanate de baryum avec des transitions de phase polaires à apolaires proches les unes des autres peut entrainer une amélioration de la polarisabilité.

Enfin, la capacité de stockage d'énergie et l'effet électrocalorique ont été étudiés sur les systèmes conçus pendant cette thèse, et une amélioration de ces propriétés a été détectée.

. Brief history of ferroelectricity

The history of ferroelectricity begins with the creation of "Sel polychreste"(which is also known as "Rochelle salt") in the middle of the 17 th century in La Rochelle, a small but important city in the south-west coast of France. In the beginning, this salt was made by Elie Seignette for medical reasons, however, after a few years, the unique crystal structure and dielectric properties of this salt drew the attention of crystallographers and physicists. In 1824, David Brewster, while studying the phenomenon of pyroelectricity on various substances, found the abnormally high pyroelectric response in Rochelle salt. After about 60 years, by discovering the concept of "Piezoelectricity", Rochelle salt showed itself as a material with high piezoelectric properties [1].

Finally, in 1920, almost 250 years after the creation of the Rochelle salt, Joseph Valasek, a student at the University of Minnesota in Minneapolis discovered the specific dielectric properties in this material which is called "Ferroelectricity" [1].

Theory of ferroelectricity

By a simple definition, ferroelectricity is a spontaneous electrical polarization in a dielectric material which can be reversed upon application of an electric field [2]. When an electric field applies to the crystalline dielectric materials with non-centrosymmetric crystal structure (polar dielectrics), electrical charges can be stored on the surface (charge displacement, D), calculated by eq.1:

D=ԑ0E+P (eq.1)
Where ԑ0 is the vacuum permittivity (8.854*10 -12 F* m -1 ), E is the applied external electric field and P is the overall polarization which is the density of charges created on the surface and is defined as eq.2:

P= ԑ0(ԑr-1)E (eq.2)

Where ԑr is the relative permittivity of the material (ratio relative to the vacuum permittivity).

Thus, the net polarization of these materials is a variable with an external electric field. It should be noted that these crystalline polar dielectrics exhibit a high relative permittivity and perform spontaneous and reversible electrical polarization, even in the absence of an electric field. Such materials are called ferroelectrics.

From a structural point of view, poled ferroelectric materials are crystalline materials, composed of regions with similar orientations of dipoles which are called ferroelectric domains. These domains are separated by domain boundaries or walls (see fig. 1.1) [3]. When an electric field is applied to a ferroelectric material, the domains grow due to the orientation of the dipoles in the same direction as the electric field. This growth of the domains is accompanied by movement of the domain walls [3]. Fig. 1.2 schematically displays the polarization and a movement of these domain walls upon application of a sinusoidal electric field. After withdrawal of the electric field, polarized dipoles can return to their randomly oriented state only with difficulty which is mostly due to steric hindrance and the energy barrier created by the dipoles in the well packed crystal structure of these materials. As such, they remain partially polarized in the absence of an electric field leading to a hysteresis behavior. This remaining polarization is called remnant polarization (Pr). Thus, a reverse electric field is required to randomly orient the dipoles again which is called the coercive field (Ec). It is noteworthy that the ferroelectric structure of a material is highly dependent on temperature. Increasing the temperature expands the unit cell. This expansion increases the mobility of the dipoles and decreases the energy barrier of their switching. Above certain temperatures, the domains easily break up to single randomly oriented dipoles. This state is called the "paraelectric" state. The temperature at which the ferroelectric structure turns to a paraelectric, is referred to as "Curie temperature" or "polar-to-apolar phase transition temperature" (TC). Above the Curie temperature, the relative permittivity sharply decreases based on the Cure-Weiss law (eq.3):

𝜀 = 𝜀 0 + 𝐶 𝑇-𝑇 0 (eq.3)
Where C is the Curie constant (a material-dependent property), T0 is the Curie-Weiss temperature which is a few degrees below the Curie temperature. Thus, the permittivity of a ferroelectric materials is maximum in vicinity of the Curie temperature and decreases above this temperature. Fig. 1.3 depicts the structural transition and relative permittivity behavior of barium titanate, a well-known ferroelectric material, as a function of temperature.

According to this graph, a ferroelectric material depending on its structure, may have several Curie temperatures. Moreover, the macroscopic polarity of ferroelectric materials changes at the Curie temperature (from high polarity below Tc to very low polarity above Tc). The structure of the domains (including size and density) which have influence on the energy barrier for dipole switching not only changes the coercive field but also determines the Curie temperature of the material [6]. Hence, modifying the structure of ferroelectric materials can change their Curie temperature.

Relaxor ferroelectric behavior

Introduction of defects in the crystal structure of inorganic ferroelectric materials by chemical substitution of ions with different charges and valences can lead to distortion in the lattice orders. This distortion breaks up the macro-sized ferroelectric domains into nano-sized domains, called polar nanodomains (or nano polar regions (NPRs)) [4,6,7]. However, in organic ferroelectric materials, pinning the polar chains can increase the density of the unit cells, because nonpolar conformation act as defects in the structure. Similar to inorganic materials, these defects can break up the domains into nanodomains.

These polar entities behave independently; each region possesses the same relaxation time.

By increasing the number of defects, more nanodomains form in the structure leading to a relaxational behavior in the whole structure under an application of electric field. This defective crystalline structure which has polarity close to paraelectric state, is called relaxor ferroelectric (RFE) phase [4] As shown in Fig. 1.4, unlike the ferroelectric behavior, no polarization-electric field hysteresis is observed after withdrawal of the electric field (having low Ec and Pr) in relaxor ferroelectrics which is directly related to the polarization of these independent nanodomains (less energy barrier exists in the structure) [3,6,8].

In addition to possessing a narrow hysteresis loop and low Ec and Pr, a broad frequencydependent peak of the dielectric constant at relatively low temperatures, as opposed to the very sharp peak obtained at high temperatures for ferroelectric materials, is the most important characteristic of RFE materials. It is thus used as universal indicator of relaxor behavior [4,7,9]. Moreover, in RFE materials, the phase transition is very diffuse and the relative permittivity does not follow the Curie-Weiss law [10]. Therefore, the polarization can remain high even after Tc which is interesting for many applications [6].

Fig. 1.4. Hysteresis loop of typical relaxor materials [4].

Ferroelectric and relaxor ferroelectric materials

As noticed earlier, dielectric materials are materials with net dipoles which can be reoriented by application of an external electric field. This polarization behavior varies in different dielectric materials. Ferroelectric materials are polar dielectric materials which exhibit a pronounced hysteresis polarization-electric field (P-E) curve with high remnant polarization (polarization at zero field) large dipolar domains cannot completely relax when the electric field is removed.

Upon increasing the number of independent dipoles and decreasing the strength of the dipole-dipole and dipole matrix interactions, narrow hysteresis loops (single or double loop, see fig. 1.5) appear on the P-E curve [11]. Such materials are called relaxor ferroelectric materials (RFEs). In paraelectric and dipolar glass dielectric materials, single dipoles are randomly oriented and macroscopic polarization is zero in the absence of an external electric field.

Fig. 1.5. Dipole and domain structure of electroactive materials with their D-E loops [11].

Ferroelectric and relaxor ferroelectric materials are generally classified into two categories:

Inorganic and organic materials.

Inorganic materials

Inorganic or ceramic ferroelectric materials are crystalline materials with noncentrosymmetrical crystal structures which have spontaneous polarization that can be reversed by application of an electric field. Fig. 1.6 shows the four types of known ferroelectric ceramics: (1) the tungsten-bronze group, (2) the oxygen octahedral group, (3) the pyrochlore group, and (4) the bismuth layer-structure group. The oxygen octahedral group (ABO3 perovskite type) is the most studied category among these groups. PbTiO3, BaTiO3, Ba(ZrxTi1-

x)O3, (Pb1-xLax)(Zr1-yTiy)O3, and Pb(Mg1/3Nb2/3)O3 are some well-known examples from this group which constitute the highest fraction of the industrial ferroelectric ceramics in the world by now [10,12].

Fig. 1.6. Categories of inorganic ferroelectrics materials [12].

Fig. 1.7 depicts the typical ABO3 unit-cell structure of perovskite oxides. At T<Tc, the cation which is placed in the center of the octahedral cage (B), is slightly shifted away along the zaxis forming a tetragonal shape which can be displaced by an external electric field. Therefore, the ferroelectric domains which are composed of these polar unit cells can be polarized along this electric field. After withdrawal of the electric field, these domains partially stay in the polarized state (Pr) and lead to a hysteresis behavior. temperature. Non-cubic structure is responsible for ferroelectric properties [13].

However, at T>TC, the tetragonal crystalline phase converts to a cubic phase which is a centrosymmetric phase and results into a paraelectric state. In most of the simple perovskite oxides containing tetragonal structure, such as BaTiO3, this transition is sharp and is a firstorder transition. Substituting the ions in A or B site of the structure with ions with different

A B O

valencies leads to some defects in the structure and distorts the ordering of the crystals. Thus, the tetragonal structure turns to less polar rhombohedral/orthorhombic structures. Moreover, large domains break up into nano-sized domains which assists the dipoles to more independently switch along the electric field direction. These independent nano-sized domains also decrease the Tc of the material [9]. For instance, by 20% substitution of Ti 4+ in the B site with a larger cation like Zr 4+ , Tc of BaTiO3 which is around 120°C can be decreased to around room temperature [14]. In addition to this modification of the Curie temperature, by this ion substitution, the relative permittivity may be broadened (due to second-order phase transition) and become highly dependent on frequency [10]. All of these observations guide us to the fact that such substitutions convert ferroelectric properties in perovskite oxide to relaxor ferroelectric properties. Among all the ferroelectric materials, the inorganic ceramics are ferroelectric materials which were commercially used in industry first. However, the costly production, the low mechanical flexibility and low breakdown strength of these materials hampered their progress in industry.

Today, these materials are still utilized in form of bulk and thin films in various applications such as energy storage capacitors but their future likely lies in nanocomposite systems which minimize the disadvantages of bulky ceramics while improving the dielectric properties of polymers. Indeed, nowadays, micron and nano-sized ferroelectric ceramics are more desired than bulk ceramics in the field of ferroelectric materials. To succeed, it is important to gain control of the quality of the crystals, which are highly affected by morphology and size of the particles [15]. In nano-sized ceramics, the ferroelectric tetragonal structure, due to presence of an internal stress produced by particle interfaces, partially or fully converts to a nonferroelectric cubic structure, which significantly decreases the dielectric and ferroelectric properties of material [16,17]. Hence, producing nano-sized ferroelectric ceramics is challenging.

As mentioned above, perovskite oxide materials due to their significant ferroelectric properties are the most-used ferroelectric materials. Among all of the perovskite oxide materials, barium titanate (BT) due to its low toxicity is the dominant material in industry. In order to synthesize tetragonal BT particle, which is a promising material for nanocomposites, conventional solid-state methods are the most-used techniques in industry. However, very high calcination temperature (>1000°C) and long production times as well as lack of ability to produce nano-sized particles without control of morphology are considerable drawbacks of this approach. Recently, various other pathways were explored to produce tetragonal nanosized BT using soft chemistry techniques [18][19][20]. By this techniques, the BT nanoparticles are produced via hydrolyzing metal precursors in a liquid medium. Thereby co-precipitation, hydrothermal and more recently polyol methods were used for the synthesis of BT nanoparticles.

In the hydrothermal method, barium and titanium precursors are hydrolyzed in a basic aqueous medium at high temperature (~220°C) and pressure, while in the polyol technique, barium and titanium precursors are hydrolyzed in a polyol medium such as diethylene glycol (DEG) instead of a aqueous medium, as it has higher boiling point (~230°C) than water and application of high pressure is not needed [21]. Moreover, due to the chelating effect of polyol (e.g. DEG) during reaction, obtained particles possess more homogeneous sizes than particles obtained in the hydrothermal method which is more favorable in nanocomposites.

Organic materials

The high costs and the difficult processability of inorganic materials as well as their rigidity and low breakdown strength, seriously limit their usages in industrial applications. Recently, ferroelectric organic materials with easier processability into thin, light and flexible films attracted attention. In the beginning, due to the electrically insulator behavior of the polymers, they were known as electrically passive materials. However, with the growth of polymer science, other special properties such as piezoelectricity were discovered in some specific polymers which resulted in these polymers to be used as an active element in electrical circuits [22]. By discovering piezoelectricity and a few years later, pyroelectricity in polymers, the possibility of ferroelectricity was explored [23]. P(VDF) was found as first ferroelectric polymer. A few years later, other polymers based on P(VDF) were designed with various ferroelectric behaviors desirable for different applications [23]. The structure and properties of these polymers will be discussed in detail in section 1.3.

Applications of ferroelectric materials

After discovering ferroelectricity in 1920, this property was studied academically and theoretically for about 2 decades. Thereafter, by finding and developing this property in bulk ceramics, practical interests of this property slowly was raised [24]. Many ferroelectric ceramics were found suitable mostly for actuators, piezoelectric transducers as well as pyroelectric detectors. Developing thin-film ferroelectric ceramics in 1984, enhanced their breakdown strength and broadened the potential applications of these materials [24,25].

Recently, by the development of over 700 ferroelectric materials, organic and inorganic ones, many applications have been practically achieved. Non-volatile memories, sensors, electrooptic modulators, electron emitters, energy harvesting devices, acoustics, information storage devices, magnetoelectrics as well as electrocaloric devices are among these new applications [24,[26][27][28][29]. Focus here is on two important applications of such materials: energy storage capacitors and electrocaloric refrigeration.

Energy storage capacitors

In a simple definition, energy storage properties cover the ability of a material to store charges and electrical energies. Because of the high demand of small and low-cost electronic devices with high energy and power capacitive storage ability, the development of high power and energy density materials have become very interesting for the researchers [30].

Dielectric materials with high polarizability and consequently high charge and discharge energy densities, have great potential to display high energy storage ability. Generally, the energy density term is described as (eq.4):

Ue=∫ 𝐸𝑑𝐷 𝐷 0 (eq.4)
Where E is the applied electric field to the material and D is the charge displacement in the material. Fig. 1.8 illustrates the energy density value of a material in a D-E graph (blue area). material with highlighted energy density (blue area) and energy loss (yellow area) of material [START_REF] Zhu | Novel ferroelectric polymers for high energy density and low loss dielectrics[END_REF].

According to eq.4, a high E (which highly depends on breakdown strength) and D (which depends on polarization) is required to enhance the energy density of a dielectric material.

Dielectric non-ferroelectric materials exhibit a linear charge displacement as a function of electric field without any hysteresis behavior which mostly relies on electronic and atomic polarization in the material. These polarizations arise from the movement of the electron clouds and deviation of the skeletal atoms from the equilibrium position, upon applying an external electric field. These polarizations exist in most of polymers, independent whether they are polar or non-polar. The value of these polarizations is very small, thus, leading to a very small energy density [START_REF] Zhu | Novel ferroelectric polymers for high energy density and low loss dielectrics[END_REF]. However, ferroelectric materials which possess high polarization mostly from orientational polarization, exhibit a higher polarization versus an external electric field. Due to their large domains and high energy barrier for dipole orientation, the polarization behaves nonlinearly and some part of the electrical energy losses are in form of heat (yellow area in fig. 1.8) and decrease the efficiency of their energy storage ability. As mentioned before, the relaxor ferroelectric materials possess a more narrow hysteresis loop than ferroelectrics. This narrow hysteresis behavior can minimize the loss (with their low remnant polarization) while the polarization still remains high, compared to non-ferroelectric materials (Fig. 1.4). Another important factor when considering energy storage properties is the breakdown strength. Bulk crystals have the lowest breakdown strength among dielectric materials while thin film dielectric polymers exhibit remarkably high resistance to the electric field [START_REF] Zhu | Novel ferroelectric polymers for high energy density and low loss dielectrics[END_REF].

Therefore, relaxor ferroelectric materials exhibit high energy densities as well as a low energy loss. They thus are promising materials for high energy storage capacitors.

Electrocaloric refrigeration

Over the past decades, vapor-compression cooling technology has become the most-used cooling technology but it has a moderate efficiency and, more importantly, uses hazardous refrigerants. Hence, this technology has begun to be replaced with a more efficient alternative cooling technology, called solid-state cooling. Because of the large size of compressors used in the vapor-compression technologies, solid-state cooling is desirable for on-chip cooling and temperature regulation for small electronic devices and sensors [START_REF] Moya | Caloric materials near ferroic phase transitions[END_REF]. By increasing the demand of small and light electronic devices in modern life, demands for small cooling technologies is also raised. Recently, electrocaloric cooling (EC) based on the electrocaloric effect (ECE) in materials, as a solid-state cooling has thus attracted attention in the cooling field [8,[START_REF] Mischenko | Giant electrocaloric effect in thin-film PbZr0[END_REF].

In the 1930s, the EC started to become an interesting cooling technology [START_REF] Kobeko | Dielektrische eigenschaften der seignettesalzkristalle[END_REF], but it was of low efficiency due to the very small temperature changes the available materials displayed (i.e. 2.5°C at 750V on Pb0.99Nb0.02(Zr0.75Sn0.20Ti0.05)0.98O3 oxide, which was reported as the highest EC response in 1960 [START_REF] Thacher | Electrocaloric effects in some ferroelectric and antiferroelectric Pb (Zr, Ti) O3 compounds[END_REF]). This limited its commercially use in industry. It took almost half of a century that a giant ECE of ΔT=12°C, observed in Pb(Zr0.95Ti0.05)O3 thin films was achieved [START_REF] Mischenko | Giant electrocaloric effect in thin-film PbZr0[END_REF]. Finally, in 2008, a giant electrocaloric effect in organic P(VDF) copolymer was reported by Zhang [8] which opened an interesting platform for electrocaloric applications of polymeric materials. In this pioneering work, the temperature change of ΔT=12°C was comparable with the value obtained with inorganic PZT at Tc of 80°C in the P(VDF-TrFE) ferroelectric polymer. Thereafter, many reports were published on the electrocaloric behavior of VDF-based polymers.

The electrocaloric effect (ECE) is a reversible temperature change (dT) of a polar ferroelectric material upon application of an electric field (dE), under adiabatic conditions [13,[START_REF] Mischenko | Giant electrocaloric effect in thin-film PbZr0[END_REF][START_REF] Lines | Principles and applications of ferroelectrics and related materials[END_REF][START_REF] Lang | Cryogenic refrigeration utilizing the electrocaloric effect in pyroelectric lithium sulfate monohydrate[END_REF].

For better understanding this phenomenon and its mechanism, fig.1.9. schematically shows an electrocaloric cycle in comparison with a vapor-compression cycle (pistons). Application of an external electric field to the ferroelectric materials leads to the switching of dipoles and an increase in the macroscopic polarization which consequently decreases the entropy. Due to the adiabatic condition, the initial temperature abruptly increases in the material. This step is equal to compressing refrigerants in vapor-compression cooling which decreases the entropy and increases the temperature of this refrigerant. In second step, in order to extract this produced heat, the heated material is attached to a heat sink to induce heat transfer to this sink. Then by withdrawal of the applied electric field and reorienting the dipoles, similarly to expanding the refrigerant, the entropy increases again which causes a decrease of temperature at adiabatic conditions. This temperature change which causes a heat absorption from the environment, known as electrocaloric ΔT, and the phenomenon in which the ferroelectric material shows this temperature change is called the electrocaloric effect. In a thermodynamic point of view, this behavior can be described by the relationship between polarization (P), external electric field (E) and temperature. Changes in the entropy (S) which is induced by the applied electric field (E) can be expressed as a function of polarization (P)

and temperature (T) using the Maxwell equations (eq.5) [START_REF] Thacher | Electrocaloric effects in some ferroelectric and antiferroelectric Pb (Zr, Ti) O3 compounds[END_REF]:

( 𝜕𝑃 𝜕𝑇 ) 𝐸 = ( 𝜕𝑆 𝜕𝐸
) 𝑇 (eq.5)

𝐶 𝐸 𝑑𝑇 = 𝑇𝑑𝑆 (eq.6)

where CE is the specific heat capacity, with C=d*Cp where d is the density and Cp is the volume heat capacity. Accordingly the following expression (eq.7) can be obtained:

𝑑𝑇 = 𝑇 𝑑𝐶 𝑃 ( 𝜕𝑃 𝜕𝑇 ) 𝐸 𝑑𝐸 (eq.7)
Then ∆𝑇 𝐸𝐶 under ΔE=E2-E1 can be calculated by integration of the pyroelectric coefficient

( 𝜕𝑃 𝜕𝑇 ) 𝐸
over the electric field (eq.8):

∆𝑇 𝐸𝐶 = ∫ 𝑇 𝑑𝐶 𝑃 ( 𝜕𝑃 𝜕𝑇 ) 𝐸 𝑑𝐸 𝐸 2 𝐸 1
(eq.8)

These phenomenological expressions identify the required terms for achieving a large electrocaloric cooling. Since ΔTEC is proportional to the pyroelectric coefficient (

𝜕𝑃 𝜕𝑇

)E , the maximum ECE will be found at maximum entropy and polarization change of the material versus temperature. This large polarization change can take place in the vicinity of TC in which the ordered domains (ferroelectric state) converts to disordered independent dipoles (paraelectric state) [START_REF] Lines | Principles and applications of ferroelectrics and related materials[END_REF][START_REF] Sanlialp | Direct measurement of electrocaloric effect in lead-free Ba (SnxTi1-x) O3 ceramics[END_REF]. As mentioned earlier, ferroelectric materials mostly possess firstorder phase transitions which are very sharp and abrupt transition. This transition exhibits a large pyroelectric coefficient and consequently a high electrocaloric response (ΔTEC). Although the ΔTEC is high for ferroelectric materials, the functional temperature area (at vicinity of Tc) is very narrow and usually takes place at high temperatures (due to high phase transition temperatures in ferroelectric materials). Since coolers and refrigerators usually work in an environment with a wide temperature range close to room temperature, this narrow regime and the high temperature of ferroelectric materials limit their practical applications. Hence, shifting the phase transition to a temperature range of interest (room temperature for domestic refrigerators and air conditioners) and also widening this range, is a main challenge to exploit the ECE in ferroelectric materials. As discussed earlier, relaxor ferroelectric materials possess very diffuse and broad phase transition closer to room temperature than ferroelectric materials. Thus, they can be potentially a good choice. Indeed, unlike ferroelectric materials, which show an EC effect in a narrow range near the Tc at high temperatures, relaxor ferroelectrics can exhibit a broader temperature range for ECE at lower temperatures.

According to the eq.8, the other feature which affects the ΔTEC, is the dielectric strength of the material. Since ΔTEC is proportional to E, increasing E would increase ΔTEC. Therefore, materials are desirable that can withstand the high electric fields [3].

A low electrical conductivity of the material is another required feature for EC materials. Any increase in conductivity, either electronic or ionic, causes a leakage current under an applied field, which results in Joule heating in the system. Thus, such a system under an alternating field, becomes warmer instead of colder. This reverse effect considerably decreases the ECE [START_REF] Valant | Electrocaloric materials for future solid-state refrigeration technologies[END_REF].

VDF-based polymers

Poly (vinylidene fluoride)-based polymers, which nowadays are well-known ferroelectrics, are semicrystalline polymers with vinylidene fluoride (VDF) as repeat unit (see Fig. 1.10). The piezoelectric response of PVDF was first reported in 1969 [START_REF] Kawai | The piezoelectricity of poly (vinylidene fluoride)[END_REF]. Two years later, in 1971, Bergman et al [23], reported pyroelectricity in PVDF. Ferroelectricity of PVDF was discussed in 1974 by Tamura et al [START_REF] Tamura | Piezoelectricity in uniaxially stretched poly (vinylidene fluoride)[END_REF]. In the last three decades, ferroelectric PVDF homopolymer and its binary and ternary polymers have attracted increasing attention [23,27,[START_REF] Tamura | Piezoelectricity in uniaxially stretched poly (vinylidene fluoride)[END_REF], due to their high dielectric constant (comparing to the other polymers), ferroelectric and relaxor ferroelectric behavior and their high breakdown strength. Ferroelectric and relaxor ferroelectric properties of those polymers mostly (not fully) originate from their crystalline regions [START_REF] Furukawa | Ferroelectric properties of vinylidene fluoride copolymers. Phase Transitions[END_REF].

Crystal structure and morphology

As mentioned before, ferroelectric materials have a polar structure. But more importantly, in addition to their polarity, the ability of this polar structure to reorient upon an external electric field plays a key role to act as ferroelectric material. In that regard, it is important to first deeply understand the crystal structure of VDF-based polymers.

VDF-based semicrystalline polymers are composed of various crystalline phases with at least four known different polymorphs, called α, β, γ and δ phase [22,[START_REF] Lovinger | Electron diffraction investigation of a high-temperature form of poly (vinylidene fluoride)[END_REF][START_REF] Lovinger | Crystallization and morphology of melt-solidified poly (vinylidene fluoride)[END_REF][START_REF] Davis | Electric-field-induced phase changes in poly (vinylidene fluoride)[END_REF]. These polymorphs vary with respect to chain packing and chain conformations: all-trans, tg + tg -and tttg + tttg - Among these four crystal phases, the α phase is a centrosymmetrical nonpolar phase while the β phase has the highest polarity due to its non-centrosymmetrical unit cell. In the α phase, which is thermodynamically the most stable phase, two chains are arranged with face-to-face tg + tg -conformation. Hence, the fluorine atoms cancel out the dipole moments of each other and decrease the overall polarity of the unit cell. On the other hand, in the β phase, chains have all-trans conformation which unlike the tg + tg -conformation, is thermodynamically unstable within a single macromolecule. In this phase, all the fluorine atoms arrange in parallel all-trans conformation and consequently increase the overall polarity of the crystalline phase which has a net dipole moment per unit cell of 8*10 -30 C m [START_REF] Su | Density functional theory and molecular dynamics studies of the energetics and kinetics of electroactive polymers: PVDF and P (VDF-TrFE)[END_REF][START_REF] Martins | Electroactive phases of poly (vinylidene fluoride): Determination, processing and applications[END_REF]. Since the ferroelectric properties of the PVDF-based polymers arises from the polar crystal phases, the α phase is the least preferable phase among all of the crystal phases while the β phase has proven to provide the best ferroelectric properties [START_REF] Furukawa | Ferroelectric properties of vinylidene fluoride copolymers. Phase Transitions[END_REF]. The γ phase is the other polar phase which possesses a dipole moment that is smaller than the one of the β phase. The unit cell of this phase is formed by packing of two chains in tttg + tttg -conformation [START_REF] Lovinger | Annealing of poly (vinylidene fluoride) and formation of a fifth phase[END_REF]. The third and last discovered polar phase, called the δ phase, is a polar version of the α phase. The chains in this phase are packed in a tg + tg -conformation in a parallel form in which the fluorine atoms not only do not cancel each other out but also increase the dipole moment of the unit cell [START_REF] Bachmann | The crystal structure of phase IV of poly (vinylidene fluoride)[END_REF].

The chemical structure, solidification conditions as well as the influence of other organic or inorganic materials on the VDF-based polymers play key roles on the overall crystallinity as well as the nature of the crystalline phase and their content. In section 1.5, we will focus on the influence of these factors on crystalline phase formation as well the resulting ferroelectric and relaxor ferroelectric properties of them.

Ferroelectric and relaxor ferroelectric behavior in VDF-based polymers

The D-E hysteresis loop in VDF-based polymers which arises from the polar crystalline phases is a strong evidence on the existence of ferroelectricity in these polymers. Fig. 1.12 depicts the D-E hysteresis loop of the P(VDF) homopolymer. As described in section 1.3.1, polar phases, specifically the β phase, are responsible for P(VDF)'s ferroelectric behavior. The polarity of these phases are switchable upon the direction of applied electric field. This dipole switching happens at an electric field that is higher than the coercive field (Ec) so that the electric field can overcome the energy barriers of the switching the dipoles in the crystal. Thus, the energy barrier and ability of the dipoles to switch in the packed crystals creates the D-E hysteresis loop. A P(VDF) with 50% crystallinity, crystallized in the β phase, can provide a remnant polarization around 90 mC/m 2 [START_REF] Gregorio | Effect of crystallization temperature on the crystalline phase content and morphology of poly (vinylidene fluoride)[END_REF]. The β phase is a well-packed crystalline phase in which the polymer chains are in all-trans conformations. The energy barrier for dipole switching is high in these crystalline phases.

Pinning the polymer chains, either chemically by electron irradiation or physically by introduction of bulky monomers into the structure can expand the crystal phases. This leads to a decrease in the energy barrier of dipole switching which can decrease the coercive field and remnant polarization in the material. Decreasing the hysteresis loop, originating from this expansion, converts the ferroelectric behavior to relaxor ferroelectric [START_REF] Zhang | Giant electrostriction and relaxor ferroelectric behavior in electron-irradiated poly (vinylidene fluoride-trifluoroethylene) copolymer[END_REF]. By this expansion in ferroelectric phases, some gauche conformations gradually appear in the all-trans conformation which are stable due to the steric hindrance in the packed lattice (gauche conformation is thermodynamically more stable in single chain). This gauche defect which forms tg + tg -and tttg + tttg -conformations in the structure can build the non-polar α or the polar γ and δ phase in the structure [START_REF] Su | Density functional theory and molecular dynamics studies of the energetics and kinetics of electroactive polymers: PVDF and P (VDF-TrFE)[END_REF].

P(VDF) homopolymer

Poly vinylidene fluoride (P(VDF)) is a semi-crystalline ferroelectric polymer, synthesized via the polymerization of vinylidene fluoride (VDF) leading to a maximum of 50-70% overall crystallinity. This polymer has a glass and melting temperature in the range of -40 to -30°C and 155 to 192°C, respectively [START_REF] Ameduri | From vinylidene fluoride (VDF) to the applications of VDF-containing polymers and copolymers: recent developments and future trends[END_REF].

In the absence of an electric field, P(VDF) crystallizes from the melt into a non-polar phase (α phase) [START_REF] Gregorio | Effect of crystallization temperature on the crystalline phase content and morphology of poly (vinylidene fluoride)[END_REF]. The β phase, which is responsible for the ferroelectricity in P(VDF), is more stable in the crystalline form (due to steric hindrance) and can be induced from the α phase P(VDF) film through a poling process, i.e., by heating, stretching and/or applying large electric fields [START_REF] Valant | Electrocaloric materials for future solid-state refrigeration technologies[END_REF][START_REF] Davis | Electric-field-induced phase changes in poly (vinylidene fluoride)[END_REF][START_REF] Lando | Nuclear magnetic resonance and x-ray determination of the structure of poly (vinylidene fluoride)[END_REF][START_REF] Prest | The morphology and thermal response of high-temperaturecrystallized poly (vinylidene fluoride)[END_REF]. However, these processes just partially modify the phases. Recently, it has been reported that introducing of binary and ternary monomer in the VDF backbone could overcome this issue. The most studied co-and ter-monomers are trifluoroethylene (TrFE), chlorotrifluoroethylene (CTFE), chlorofluoroethylene (CFE) and etc [START_REF] Valant | Electrocaloric materials for future solid-state refrigeration technologies[END_REF]. The properties of these PVDF-based co-and ter-polymers will be discussed in section 1.3.4.

Binary and ternary VDF-based polymers

In 1968, Lando and Doll suggested that direct crystallization of P(VDF) into the β phase will be facilitated by introduction of trifluoroethylene (TrFE) as comonomer [START_REF] Furukawa | Ferroelectric properties of vinylidene fluoride copolymers. Phase Transitions[END_REF] This suggestion was a starting point towards binary and ternary VDF-based polymers and their influence on the crystalline structure and ferroelectric properties of P(VDF). The major advantages of incorporating such co-and ter-monomers include more polar phase formation, easier poling and achieving relaxor ferroelectric behavior.

P(VDF-TrFE

) is the most studied copolymer of P(VDF) which is produced via polymerization of VDF and -trifluoroethylene (TrFE) (Fig. 1.13). This polymer has a simple chemical structure but features a rather complicated crystal structure and morphology [27]. In this copolymer, the β phase is more stable than the α phase. This phase formation is due to steric hindrance of TrFE monomer in the structure. Because of the TrFE content in the polymer unit, the βcrystalline phase can be obtained directly after solidification (whether from the melt or solution) [START_REF] Zhu | Novel ferroelectric polymers for high energy density and low loss dielectrics[END_REF][START_REF] Martins | Electroactive phases of poly (vinylidene fluoride): determination, processing and applications[END_REF][START_REF] Yang | Novel polymer ferroelectric behavior via crystal isomorphism and the nanoconfinement effect[END_REF]. The Curie temperature of P(VDF-TrFE) varies from 55 to 128 °C depending on VDF content (from 55 to 88 mol%) [START_REF] Lovinger | On the Curie temperature of poly (vinylidene fluoride)[END_REF][START_REF] Furukawa | Ferroelectric behavior in the copolymer of vinylidenefluoride and trifluoroethylene[END_REF]. P(VDF-TrFE) (75/25 mol%) shows a relatively high remnant polarization of 7.5μC/cm in comparison to P(VDF) homopolymer. The room temperature dielectric constant of P(VDF-TrFE) is higher value (~18) than that of PVDF (~6-12) with a dielectric loss <0.1 at 1 kHz frequency [START_REF] Yagi | Transition behavior and dielectric properties in trifluoroethylene and vinylidene fluoride copolymers[END_REF]. relaxor rather than ferroelectric behavior due to their bulky termonomers. These terpolymers possess various crystalline phases which will be discussed in detail in section 1.5.2.

The phase transitions in these terpolymers are diffuse and take place at relatively low temperatures depending on their termonomer nature and content. These properties make these materials very interesting for many applications such as energy storage applications and electrocaloric cooling.

Engineering the solid-state structure of VDF-based polymers

As discussed earlier, the solid-state structure of VDF-based polymers is mainly responsible for the ferroelectric and relaxor ferroelectric properties of these polymers. Hence, engineering the solid-state structure can optimize certain properties (the Curie temperature, polarization, coercive field, switching speed, etc.).

Various factors such as chemical structure, solidification conditions and addition of other organic or inorganic components, morphology of dried film (thickness, surface area etc.) [START_REF] Zhong | Phenomenological study of the size effect on phase transitions in ferroelectric particles[END_REF][START_REF] Bune | Two-dimensional ferroelectric films[END_REF], play key roles in determining the overall crystallinity and crystalline phase formation of the polymer.

Effect of chemical structure

In order to explain the effect of the chemical structure on the solid-state structure of the VDFbased polymers, materials with different chemical structures are compared here. Unlike the P(VDF-TrFE-CFE) terpolymers which possess a double hysteresis loop, P(VDF-TrFE-CTFE) terpolymers where the CTFE termonomer is larger than CFE in term of size, features a single hysteresis loop. The interchain distance in this terpolymer is higher and the energy barrier for dipole switching is smaller. Thus, a narrower hysteresis loop can be achieved compared to P(VDF-TrFE-CFE).

Engineering the solidification

VDF-based polymers as a semicrystalline polymer can be crystalized either from the melt or solution. During this crystallization, different crystalline phases can be induced. This is highly dependent on the solidification conditions. For instance, in case of crystallization from solution, parameters such as casting temperature, nature of solvent, solution concentration and molecular weight of the polymer can lead to polymer films comprised of different polymorphs [START_REF] Gregorio | Effect of crystallization temperature on the crystalline phase content and morphology of poly (vinylidene fluoride)[END_REF][START_REF] Ma | β-Phase of poly (vinylidene fluoride) formation in poly (vinylidene fluoride)/poly (methyl methacrylate) blend from solutions[END_REF]. Hence, in addition to crystalline phase modification after solidification, such as stretching or poling, controlling the solidification of the polymer in order to achieve the interested polymorph, can be a useful way to manipulate the solid-state structure.

VDF-based polymers typically crystallize into the non-polar α phase which is the least desirable phase with respect to ferroelectric properties. The polar phases like β phase, due to its high ferroelectric properties, have thus been always in the center of attention of research on these materials, and engineering the formation of these phases was found to be an important challenge. Song et al [START_REF] Song | Formation of β-phase microcrystals from the melt of PVF 2-PMMA blends induced by quenching[END_REF] obtained predominantly the β phase during crystallization of P(VDF) from the melt with a high cooling rate. The γ phase is another polar phase that is responsible for the ferroelectric properties in this polymer. It is achieved during crystallization from DMF, DMA and DMSO solutions with slow heating [START_REF] Wendorff | Concentration fluctuations in poly (vinylidene fluoride)-poly (methyl methacrylate) mixtures[END_REF]. Again, in case of crystallization from the melt, Osaki and Kotaka [START_REF] Osaki | Electrical properties of form III poly (vinylidene fluoride)[END_REF] studied the variation in the γ phase content in P(VDF)

with respect to the melt crystallization temperature. They reported that the γ phase formation starts at 160°C and reaches its maximum value at 170°C.

In 1994, Gregorio et al [START_REF] Gregorio | Effect of crystallization temperature on the crystalline phase content and morphology of poly (vinylidene fluoride)[END_REF] investigated the influence of casting temperature on the crystallization of P(VDF). They reported that P(VDF) predominantly forms the β phase during crystallization from solution at temperatures below 70°C, while crystallization at a temperature between 70 and 110°C leads to a mixture of α and β phase. Above 110°C, a predominant α phase can be obtained. Fig 1 .16 summarizes the studied solidification conditions including post-solidification treatments in order to obtain the various polymorphs in solid P(VDF) [START_REF] Lovinger | Developments in crystalline polymers[END_REF]. Although an almost complete study has been carried out on the influence of solidification conditions on P(VDF), there is not any detailed research on other VDF-based polymers such as CFE and CTFE terpolymers.

Blending with other polymers

In addition to solidification conditions, it is observed that the presence of other organic polymers blended with VDF-based polymers can also have some influence on the crystalline phase formation in these polymers. PMMA is one of these polymers which can rearrange the VDF-based chains during crystallization and modify the solid-state structure of the polymer.

In the early 1960's, it was discovered that P(VDF) is compatible with poly(methyl methacrylate) (P(MMA)) when they are at a molten state [START_REF] Schmitt | Transparent molding compositions comprising methacrylate polymers blended with a minor amount of poly (vinylidene fluoride)[END_REF][START_REF] Noland | Compatible high polymers: Poly (vinylidene fluoride) blends with homopolymers of methyl and ethyl methacrylate[END_REF] below LCST of ~350°C [START_REF] Bernstein | LCST behavior in polymer blends[END_REF] (critical temperature which blends above this temperature are immiscible) with interaction parameter (χ12) value of -0.295 at 160°C between P(VDF) and P(MMA) [START_REF] Nishi | Melting point depression and kinetic effects of cooling on crystallization in poly (vinylidene fluoride)-poly (methyl methacrylate) mixtures[END_REF]. Noland et al [START_REF] Noland | Compatible high polymers: Poly (vinylidene fluoride) blends with homopolymers of methyl and ethyl methacrylate[END_REF] investigated this miscibility, by monitoring the glass transition temperature (Tg) of the blends in P(MMA)-rich systems. P(MMA) is an amorphous polymer with a relatively high Tg compared to P(VDF). In this work, P(VDF)/P(MMA) blends revealed a single Tg which indicates miscibility between these two polymers [START_REF] Noland | Compatible high polymers: Poly (vinylidene fluoride) blends with homopolymers of methyl and ethyl methacrylate[END_REF][START_REF] Paul | Polymer blends containing poly (vinylidene fluoride). Part IV: Thermodynamic interpretations[END_REF]. The dynamic mechanical behavior of these blends and the dielectric relaxation also showed a single transition supporting this view [START_REF] Paul | Properties of compatible blends of poly (vinylidene fluoride) and poly (methyl methacrylate)[END_REF][START_REF] Hahn | Dielectric relaxation of the crystal-amorphous interphase in poly (vinylidene fluoride) and its blends with poly (methyl methacrylate)[END_REF]. In addition, a depression in melting point in different blend compositions suggests molecular intermixing of the P(VDF) and the P(MMA) [START_REF] Nishi | Melting point depression and kinetic effects of cooling on crystallization in poly (vinylidene fluoride)-poly (methyl methacrylate) mixtures[END_REF].

Miscibility of P(VDF) in P(MMA) arises from some physical and chemical interactions between the P(VDF) and the P(MMA) which can be characterized by FT-IR and NMR techniques [START_REF] Coleman | A fourier transform infrared study of polymer blends. I. Poly (vinylidene fluoride)-poly (methyl methacrylate) system[END_REF][START_REF] Roerdink | Computerized infra-red study of the interaction of poly (vinylidene fluoride) with stereoregular poly (methyl methacrylate)[END_REF][START_REF] Douglass | Compatibility in PVF2/PMMA and PVF2/PEMA blends as studied by pulsed NMR[END_REF]. By looking at the FT-IR spectrum of P(VDF)/P(MMA) blend, it can be seen that the vibration of C=O bonds in the P(MMA) is shifted towards lower wavenumbers in presence of the P(VDF) which is similar to a hydrogen bonding of C=O [START_REF] Roerdink | Computerized infra-red study of the interaction of poly (vinylidene fluoride) with stereoregular poly (methyl methacrylate)[END_REF]. On the other hand, no shifts are observed for C-H in the P(VDF). As such, no hydrogen bonding exists between the C-H of the P(VDF) and the C=O of the P(MMA) [START_REF] Roerdink | Computerized infra-red study of the interaction of poly (vinylidene fluoride) with stereoregular poly (methyl methacrylate)[END_REF]. It can be concluded that the C=O bonds are just polarized and weakened by the dipoles of the P(VDF) [START_REF] Paul | Polymer blends containing poly (vinylidene fluoride). Part IV: Thermodynamic interpretations[END_REF][START_REF] Roerdink | Computerized infra-red study of the interaction of poly (vinylidene fluoride) with stereoregular poly (methyl methacrylate)[END_REF]. Stronger interaction of the P(VDF) segments with P(MMA) in form of isotactic than syndiotactic arrangements in addition highlight dipole-dipole interaction [START_REF] Roerdink | Influence of tacticity of poly (methyl methacrylate) on the compatibility with poly (vinylidene fluoride)[END_REF]. In isotactic P(MMA), all the methyl methacrylate groups are placed at one side of the chain and have more dipole-dipole interaction with P(VDF) segments.

Fig. 1.17. Interaction of various polymers with P(VDF) [START_REF] Paul | Polymer blends containing poly (vinylidene fluoride). Part IV: Thermodynamic interpretations[END_REF].

Due to these strong dipole-dipole interactions between P(VDF) and P(MMA) (Fig. 1.17) [START_REF] Paul | Polymer blends containing poly (vinylidene fluoride). Part IV: Thermodynamic interpretations[END_REF], P(MMA) being an amorphous polymer, it was concluded that P(MMA) is miscible with the semicrystalline P(VDF) in the molten state as well as the amorphous regions in P(VDF). By cooling the blend system from the melt or evaporating a solvent, crystallization of VDF chains sets in. During crystallization from the melt, P(VDF) segments have to migrate from the amorphous regions and to pack and consequently arranged into a crystalline phase. [START_REF] Nishi | Melting point depression and kinetic effects of cooling on crystallization in poly (vinylidene fluoride)-poly (methyl methacrylate) mixtures[END_REF]. Thus, crystallization is affected by two parameters: cooling rate and blend composition [START_REF] Nishi | Melting point depression and kinetic effects of cooling on crystallization in poly (vinylidene fluoride)-poly (methyl methacrylate) mixtures[END_REF]. Faster solidification can more vitrify the blend hindering the diffusion of VDF chains, therewith forming less and smaller crystallites. This phenomenon can also occur at elevated temperatures in case of crystallization from solution. On the other hand, increasing the P(MMA) content in the system also hampers the mobility of the VDF chains and decreases the radial growth rate of the crystallites as much as two orders of magnitude (Fig. 1.18) [START_REF] Wang | Spherulitic crystallization in compatible blends of poly (vinylidene fluoride) and poly (methyl methacrylate)[END_REF]. P(VDF)/P(MMA) blend systems with P(MMA) weight fraction of less than 30wt.%, comprises a crystalline phase of pure P(VDF) and a blend amorphous phase of both P(VDF) and P(MMA) [START_REF] Gregorio | Effect of PMMA addition on the solution crystallization of the alpha and beta phases of poly (vinylidene fluoride)(PVDF)[END_REF]. In addition to this two phases, a crystalline-amorphous interphase containing pure P(VDF) may also exist which results in partial phase separation [START_REF] Hahn | Evidence for a crystal-amorphous interphase in PVDF and PVDF/PMMA blends[END_REF]. Increasing the P(MMA) content to more than 30wt.%, prevents nucleation and leads the polymer to solidify into a nearly fully amorphous state. However, some nucleation agents, like PTFE, can assist in crystallizing P(VDF) with higher weight fractions of P(MMA), like 60wt.% [START_REF] Schneider | Impact of nucleating agents of PVDF on the crystallization of PVDF/PMMA blends[END_REF].

As discussed earlier, P(VDF) can be crystallized in various polymorphs which depends on chemical structure and solidification parameters. Presence of P(MMA) beside P(VDF) due to the strong dipole-dipole interactions of P(VDF)/P(MMA) affect crystalline phase formation.

Gregorio et al [START_REF] Gregorio | Effect of PMMA addition on the solution crystallization of the alpha and beta phases of poly (vinylidene fluoride)(PVDF)[END_REF] reported that P(MMA) with 10 and 15wt.% P(VDF) favors the formation of the β phase. In 2006, Zhang et al [START_REF] Zhang | Microstructure and electromechanical responses in semicrystalline ferroelectric relaxor polymer blends[END_REF] reported that not only P(VDF) homopolymer but also the VDF-based polymers like P(VDF-TrFE-CFE) terpolymer are miscible with P(MMA) in the amorphous phase, while P(MMA) is excluded from P(VDF-TrFE-CFE) crystals. As a consequence, P(MMA) gradually decreases the overall crystallinity of VDF-based polymers. Since low-and high-field dielectric properties of the system arises from the crystalline phase of VDF-based polymer, lower dielectric constant and smaller charge densities were obtained from such blends. On the other hand, dipole-dipole interactions of VDF and P(MMA) in the amorphous regions modified the Young's modulus and increase the breakdown strength of the system [START_REF] Ma | β-Phase of poly (vinylidene fluoride) formation in poly (vinylidene fluoride)/poly (methyl methacrylate) blend from solutions[END_REF].

Moreover, P(MMA), by decreasing the dielectric constant at low fields, can delay the saturation polarization as function of electric field and consequently make the D-E hysteresis loop more linear which is highly favorable for energy storage applications.

Inorganic/organic nanocomposites

Similar to the influence of other polymers on the solid-state structure formation of VDF-based polymers, inorganic nanoparticles by blending with these polymers, can have an effect.

Physical and chemical properties of these inorganic nanoparticles such as shape, size, surface chemistry as well as their ferroelectric properties can thus be used to engineer the solid-state structure and consequently tune the ferroelectric properties of VDF-based polymers [START_REF] Lee | Ceramic particle size dependence of dielectric and piezoelectric properties of piezoelectric ceramic-polymer composites[END_REF][START_REF] Uchino | Dependence of the crystal structure on particle size in barium titanate[END_REF].

In addition to modifying the solid-state structure of the polymer, in order to control the ferroelectric properties, inorganic ferroelectric particles with very high dielectric constants can synergistically increase the dielectric constant of the composite [START_REF] Nalwa | Ferroelectric polymers[END_REF]. However, due to the large difference between dielectric constant of particles and the VDF-based polymer (medium), this dielectric enhancement is followed by a decrease in the breakdown strength of the system. This weakening in dielectric strength is even more pronounced when the particles are not properly dispersed into the polymer matrix. Dispersability of the particles into the polymer matrix has a direct relation with their wettability in the matrix. When the surface energy of the nanoparticle and the matrix is as close as possible, the nanoparticle can easily wet and be dispersed in the matrix. Choosing particles with similar surface energy to the matrix highly limits however the design of a required system. Thus, modifying the surface of the nanoparticles with functional groups having surface energy close to the matrix can solve this issue.

Furthermore, due to the much higher elastic modulus of ceramics than polymers, the mechanical flexibility of the composites may also decrease [25]. Using nano-sized particles instead of macro-sized particles which have relatively larger surface areas, decreases the optimum particle content in the composite and improves the breakdown strength issue.

Recently, perovskite oxide (ABO3) nanoparticles with high dielectric constant such as lead and barium titanate (BaTiO3 and PbTiO3) and their derivatives demonstrated some desirable properties like high dielectric nanocomposite for various applications such as energy storage capacitors and electrocaloric devices [START_REF] Dang | Novel ferroelectric polymer composites with high dielectric constants[END_REF].

Thesis outline

Having highlighted well-known ferroelectric materials and mechanisms of ferroelectricity in these materials which arise from their solid-state structure, the importance of engineering the solid-state structure in both inorganic perovskite oxides and P(VDF) polymers for applications such as energy storage applications and electrocaloric cooling were reviewed in this introductive chapter.

As mentioned in this chapter, most of the existing research is focused on the influence of chemical structure of the materials on their solid-state structure, while the conditions of solidification, interaction of polymers and inorganic substances are less understood. Hence, in this thesis entitled "Engineering the solid-state structure of VDF-based systems for efficient energy storage and refrigeration", we focus on engineering the solid-state structure of VDF-based polymers with different chemical structures through investigation of the role of solidification conditions and interactions with other polymers or inorganic nanoparticles with these VDF-based polymers. In that regard;

In chapter 2 we will focus on the role of solidification parameters such as casting temperature, casting time and solution concentration on the solid-state structure formation of P(VDF-TrFE-CFE) compared to P(VDF-TrFE-CTFE) and P(VDF-TrFE). In the second part of the chapter, we will investigate the effect of blending PMMA with P(VDF-TrFE-CFE) with different PMMA content and solidification conditions.

In chapter 3, the synthesis of inorganic nano-sized barium titanate and its derivative barium zirconium titanate will be investigated to modify the ferroelectric properties of such nanoparticles and to manipulate the ferroelectric properties of the VDF-based terpolymers.

Finally, in chapter 4 and 5, we will use our understanding of the solid-state structure formation to tune the properties of VDF-based polymer systems for energy storage applications and electrocaloric cooling.

Motivation and context

Since the dielectric and ferroelectric properties of VDF-based polymers are structuredependent properties [1,2] and arising from the polymer chain arrangements in the crystalline and amorphous phases of a solidified polymer, understanding the solidification mechanisms is important in engineering the solid-state structure of these materials to tune their functional properties.

To achieve this, we need to understand which parameters affect the solidification of VDFbased polymers considering the following aspects:

Chemical structure of the polymer: Size and electronegativity of the atoms in each monomer determine the molecule-molecule interaction. Hence, they affect the local arrangement of the polymer chains [3]. VDF-based polymers are semicrystalline and can be crystallized in 5 different polymorphs (α, β, γ, δ, and ε), in which polymer chains can adopt three different conformations; all-trans -, tg + tg -and tttg + tttg - [4,5]. The stability of these chain arrangements directly relate to the size of the atoms (due to steric hindrance), the polarity of the monomer (due to dipole-dipole interactions) and in general the chemical structure of the polymer chains [6].

To give an example, the P(VDF) homopolymer is polymerized from VDF which contains fluorine, carbon and hydrogen atoms. The electronegativity of the fluorine and the consequent dipole that arises in the C-F bond affects the arrangement of polymer chains during solidification, and in turn defines the solid-state structure of the polymer. Crystalline phases generally adopt the α phase arrangement (the most thermodynamically stable polymorph) which is based on tg+tg-chain conformations. This polymorph, due to its centrosymmetrical unit cell is not polar, thus, does not display ferroelectric behavior.

Introducing TrFE with an extra fluorine compared to VDF, as a comonomer into polymer backbone, forces the chains to crystallize into the β phase with all-trans conformations which is the most polar polymorph and shows good ferroelectric properties (the third fluorine atom in TrFE leads to steric hindrance with neighboring fluorine atoms in case of tg+tgconformations, thus, limiting their formation). Indeed, termonomers such as CFE and CTFE (due to the presence of bulky chlorine atoms) can also limit chain packing, thus, increasing the intermolecular distance which in turn decreases the steric hindrance in the molecule as well as friction of dipole rotation. Decreasing the friction assists chains to solidify with a higher fraction of tg+tg-conformations. As a consequence, VDF-based terpolymers containing bulky termonomer typically exhibit a relaxor behavior arising from tg+tg-stabilization [7].

Although the chemical structure is an important factor in determining the solid-state structure, it is not the only factor. VDF-based polymers, for instance P(VDF) homopolymers, with a certain chemical structure, depending on its solidification protocol, cannot always be solidified in one specific structure. Therefore, in addition to chemical structure, the solidification process, which can be either from solution or melt, and deposition parameters used, can result in the polymer to adopt various solid-state structures. Some of these parameters are discussed below.

Entanglement density:

The mobility of the polymer chains during the solidification is one of the factors which drastically influences the final solid-state structure. In case of solidification from solution, low concentrated polymer solutions lead to a low entanglement density resulting in a higher chain mobility in solution. This freedom enhances the crystallization during solidification which directly affects the solid-state structure.

Casting temperature: Similar to entanglement density, higher casting temperatures lead to higher chain mobility, which can affect the structural formation. In addition to the chain mobility, the stability of each polymorph is also highly dependent on temperature, hence, considering casting temperature, polymers can be solidified directly in different polymorphs.

Casting time:

It is important to note that the solid-state structure is not only governed by thermodynamics but also kinetics [8]. A short casting time can hinder crystallization and vitrify the polymer at the beginning of the process which can affect the final solid-state structure.

Blending with other polymers: As mentioned above, intermolecular interactions during solidification are playing a key role in organizing and arranging polymer chains during solidification. Therefore, any dipole-dipole interaction of C-F bonds in a VDF-based polymer with another electropositive bond, like a C-H bond in PMMA can affect the VDF chain arrangement. Due to this interaction, PMMA is, for instance, miscible with VDF-based polymers in the amorphous phase.

As also already mentioned, the solid-state structure and which polymorph are most prominent also govern the ferroelectric properties of this class of polymers. Thus, gaining the understanding to induce specific polar polymorphs and to control the crystalline quality, quantity, size and direction, can directly enhance the ferroelectric or relaxor ferroelectric behavior of these materials. In literature, most of the documented reports on the structural properties of VDF-based polymers were focusing on the chemical structure and how this influences solid-state structure [9][10][11][12], while details of the solidification process often were neglected despite having a remarkable effect on the formed structure.

In this chapter, some of these factors such as chemical structure, entanglement density, casting temperature, solidification time and blending with PMMA, and their influence on solid-state structure of VDF-based polymers will be discussed in detail.

Experimental parameters

All polymers were dissolved in cyclopentanone at various concentrations (0.1, 0.5, 1, 2.5, 5, 10 and 15 vol.%) for viscosity measurements. For all structural characterization, solutions with targeted concentrations were drop cast on clean glass substrates at the desired temperatures in a vacuum oven. The dried films were gently peeled off from the substrates. Finally, the free standing films were characterized by DSC, FT-IR spectroscopy and WAXS.

Low temperature casting at 25°C and 40°C required 2 days to remove all residual solvent in vacuum, while at 55°C and 90°C, it was sufficient to dry films at ambient pressure, however, in order to remove all residual solvent (especially in case of samples with short casting times) the cast films were kept under vacuum at room temperatures for at least 1 day (the fraction of remained solvent was monitored by FTIR spectroscopy).

Other casting techniques such as blade coating and wire-bar coating were also used to investigate the influence of casting method on the crystallinity and phase formation of the polymers. In case of blade coating, 10vol.% solutions were used, combined with a 10µm blade height and a 10 mm/s speed to obtain a 1-µm-thick film. For wire-bar coating, a small bar (number 2) and 5 mm/s speed was used to coat 10vol.% solutions to obtain films of similar thickness. For blade and wire-bar coating of diluted solutions (0.5vol.%), due to the very low viscosity of the solution (~2mPa.s), 8 layers were coated to obtain the same dried thickness as for films produced from more concentrated solutions (1 µm).

For dielectric and polarization measurements, metal-insulator-metal capacitors were fabricated on glass substrates. First, the glass substrates were washed with isopropanol and water and acetone, respectively, in three steps via sonication for 15 min each. Silver was evaporated on the clean glass as a bottom electrode by thermal evaporation. Then, 10 vol.% solution was blade coated on the bottom electrode in order to get 10 µm-thick films. In case of dilute solutions (0.5vol.%), the bottom electrodes were dipped into a polymer solution bath and dried at target temperatures (to obtain a film of 10µm thickness and to avoid possible crack formation). Just like the bottom electrode, silver was evaporated on such films by thermal evaporation. The temperature was kept below the phase transitions of interest during thermal evaporation in order to avoid any structural change. Even though the chemical composition of CFE7 is kept constant, the formation of some ferroelectric (FE) and defective ferroelectric phases (DFE) were also reported in literature [13,14]. Hence, in order to define a phase formation mechanism for CFE7 and consequently investigate the influence of solidification parameters on the phase formation, the solidification of this polymer at various casting parameters was monitored using different techniques. Thereby, the entanglement density, casting temperature, casting time and casting method were some of the parameters investigated.

Neat P(VDF-

Entanglement density

Polymers with high molecular weight are typically entangled in concentrated solutions. Chain entanglement restricts chain mobility and hinders the organization of molecules during the solidification process. This can directly affect the solid-state structure of the dried polymer.

However, when diluting the solution with a good solvent (solvent with strong polymer-solvent intermolecular interactions) leads to disentanglement and, in turn, isolated polymer coils.

These isolated polymer coils have more space to move and rearrange in solution. Therefore, chain mobility and polymer-polymer intermolecular interactions are different in an entangled vs, non-or less-entangled situation [1,2].

Measuring the viscosity of polymer solutions provides answers about the chain mobility and intermolecular interactions. More specifically, the viscosity derives from 3 intermolecular forces: solvent-solvent, polymer-solvent and polymer-polymer intermolecular forces which makes this property highly dependent on molecular weight (Mw), concentration (C) and the entanglement density of the dissolved polymer [3].

At very low polymer concentrations with no entanglements or a low-entanglement density which is also known as the "dilute regime", polymer coils are separated and have only minor hydrodynamic interactions with each other if any at all. In this regime, the main intermolecular interactions are polymer-solvent interactions with a small influence of solvent-solvent forces. Thus, the viscosity shows a Newtonian behavior at any shear rates. In contrast to the dilute regime, in highly concentrated solutions with high entanglement densities which is known as "concentrated regime" (Fig. 2.1.c), lack of free space in solution forces polymer coils to overlap and entangle. Contrary to the dilute regime, the main force in the concentrated regime is arising from polymer-polymer interactions. Accordingly, a non-Newtonian viscoelastic behavior is observed at high shear rates. This results in a high viscosity change as function of solution concentration.

In between the dilute and concentrated regimes, there is a critical overlap concentration (C * )

where the polymer coils are starting to touch each other. The viscosity at this concentration, which is located in the regime called "semi-dilute regime" (Fig. 2.1.b), starts to rapidly change with concentration. The concentration range of this regime is highly dependent on the thermodynamic quality of the solvent. Solvents with strong interactions with the polymer broadens the regime while in poor solvents the entanglements occur in a narrower concentration regime.

In this work, in order to identify C * for CFE7 and determine the three viscosity regimes, the viscosity of polymer solutions was measured at 20°C using 5 different concentrations (0. shear thinning, may be originate from polymer-polymer interactions as well as chain entanglements in solutions. It should also be mentioned that diluted solutions at low shear rates featured a very noisy response: those results are not reported in the graph.

In order to distinguish the different viscosity regimes for CFE7, the reduced viscosity of the solutions is plotted as function of their concentrations in a logarithmic graph (Fig. 2.3). The reduced viscosity is calculated from the ratio of the relative viscosity increment (ηi) to the weight fraction of the polymer in the solution. The relative viscosity increment was calculated using the following equation: ηi=(η-ηs)/ηs, where ηs is the viscosity of the solvent. Different regimes were identified corresponding to the three previously mentioned regimes: dilute, semi-dilute and concentrated. In the dilute regime (below 2vol.%, red area), the reduced viscosity slightly increases with concentration while in the semi-dilute regime (from 3 to 5vol.%, blue area), the viscosity is starting to notably increase with concentration. Finally, in the concentrated regime (above 5vol.%, green area), the viscosity is highly dependent on the solution concentration. In addition to the observed non-Newtonian behavior in the highly concentrated solutions (Fig. 2.2), the sharp viscosity increase for concentrations >5 vol.% are clear indications for the onset of occurrence of entanglements at these CFE7 concentrations (molecular weight of 334 kg/mol). From the above, it is concluded that CFE7 with a molecular weight of 334 kg/mol in solutions of less than 3 vol.% in the dilute regime. The critical overlap concentration is found to occur between 3 and 5vol.%; the concentrated regime develops from 5vol.% on.

In order to investigate the effect of the entanglement density on the solid-state structure formation of CFE7, 0.5 and 10 vol.% solutions were, thus, chosen as solutions with low and high entanglement densities, respectively, and cast at 40°C and dried over 2 days (as described in the experimental section). 40°C was chosen as casting temperature to minimize chain mobility arising from thermal activation.

As already mentioned, the solid-state structure of VDF-based polymers depends on how the polymer chains arrange in the amorphous or crystalline phases. To identify specific chain conformations, FTIR spectroscopy was used to measure CFE7 films cast from 0.5, 5 and 10vol.% solutions. As discussed in chapter 1, VDF-based polymers can adopt 3 major chain conformations; tttt (all-trans), tttg + tttg -and tg + tg -(trans-gauche). These conformations and their corresponding infrared absorptions are well documented in literature [4][5][6][7][8]. Fig. 2.4

shows the obtained FTIR spectroscopy data. Generally, in VDF-based polymers, trans-gauche (tg) conformations are identified by their characteristic absorption bands at 608, 764, 795, 855, and 976 cm -1 , while the all-trans (tttt) conformation features characteristic absorption peaks at 840 (short range all-trans) and 1279 cm -1 (long range all-trans) [8]. Determination of the tttg chain conformations which possess both short range ttt and tg arrangements can be challenging. The small peak at 505 cm -1 and the shoulder at 1239 cm -1 are reported as the corresponding peaks for this conformation [9]. Note, though, that because of the presence of tg sequences also in tttg conformations, there are common peaks in FTIR spectra for tg and represent characteristic peaks for long and short-range all-trans conformations, respectively, while tg and t3g identify trans-gauche and trans-trans-trans-gauche arrangements, respectively.

According to the intensity of the detected peaks (Table 2.1), all the chain conformations (alltrans, tttg and tg) in cast CFE7 films are very similar at room temperature and no considerable difference is observed in the spectra of films cast from dilute, semi-dilute and concentrated solutions. However, the thermal evolution of the structure which is recorded by FTIR spectroscopy from low to high temperatures and is displayed in Fig. 2.5, shows differences in the conformational changes. In this graph, data is shown for films heated from 25°C to 85°C, comparing structures cast from 0.5vol% (Fig. 2.5.a) and 10vol.%. (Fig. 2.5.b) solutions. At room temperature, both CFE7 films contain long and short range all-trans conformations with associated peaks at 1288 and 843 cm -1 , respectively, with no trace of tg conformations (608 cm -1 ), while at 85°C, the all-trans conformation peak (1288 cm -1 ) disappears while the tg peak at 608 cm -1 increases. Fig. 2.6 shows the relevant regimes facilitating to see this evolution.

1600 1400 1200 1000 800 600 400 As illustrated in Fig. 2.6, upon heating, all-trans conformations gradually convert to the tg conformation. Above 60°C no all-trans peak was detected at 1288 cm -1 for film cast from 0.5 vol.% solutions (Fig. 2.6.a) while for films cast from 10 vol.%, this peak disappears at 65°C (Fig. 2.6.c). We conclude that the all-trans conformations in films cast from concentrated solutions show more resistance to conformational changes during heating, which indicates a denser all-trans chain arrangement in such highly entangled films. In order to better visualize these differences, the relative intensities of the vibration of the all-trans and tg conformations during heating are given in Fig. 2.7. The tm>4 (all-trans) relative intensities (black line) were obtained from the ratio of the tm>4 peak (at 1288cm -1 ) with respect to the baseline, while the tg (trans-gauche) relative intensities were obtained from the 608cm -1 vibration. In more entangled films, the development of the relative intensities for the all-trans and the tg peaks cross and shift towards higher temperatures, supporting the view that more stable all-trans conformations are present in high entangled CFE7. In addition to FTIR spectroscopy, which provides details on chain conformation as well as their thermal evolution in CFE7 films, DSC was employed to identify the phase behavior of the cast films. Since the films are cast from solution, heating from -50°C (far below Tg of CFE7) to 200°C

t
(far above the melting point of CFE7) can provide information of relevant phase transitions in this temperature regime. Note also that upon melting, the thermal and processing history of the polymer is erased. Hence, we used the first heating thermograms and data obtained from second heating are not reported here. Generally, we first cooled the polymer films to -50°C followed by an isothermal step at -50°C for 5 min to stabilize the heat flow. Thereafter, the samples were heated from -50°C to 200°C with 10°C/min.

58°C

leads to an endothermic transition that can be observed in DSC. For CFE7, cast from solutions of different concentrations this polar-to-apolar transition is found to occur around 60°C. As for the polar-to-apolar phase transition, the endotherms broaden similar to the melting endotherm but in contrast to the latter, the polar-to-apolar transition shifts to higher temperatures. Indeed, the single endotherm recorded at 60°C for films cast from 0.5 vol.% solutions evolves into a double feature at ~50°C and 70°C in both films cast from 5 and 10 vol.% solutions. This shows that the entanglement density not only can broaden the lamellae thickness distribution but also can force the polymer to adopt various phases with different thermodynamic stabilities, while less-entangled solutions crystallize into mostly one crystalline form.

Among all the possible crystalline forms that VDF-based polymers can adopt, the polymorph based on well-packed all -trans conformations displays the densest unit cell. This prevents chains to easily rotate, and limits changes in conformation through heating. Accordingly, this phase is the thermodynamically most stable compared to the other phases. The ferroelectric behavior of this dense polymorph originates from the strong cohesive forces between dipoles that form because of the trans conformation. This is referred to as the FE phase. Introducing gauche conformations, increases the interchain distance, which facilitates the chain rotation.

As consequence, the polar-apolar phase transition temperature shifts to lower temperatures.

These arrangements, depending on their gauche conformation content, do not display a pronounced ferroelectric behavior. Accordingly, such structures are referred to as defective ferroelectric (DFE) phase (with low gauche content) or relaxor ferroelectric (RFE) phase (with high gauche content). As mentioned earlier, all-trans conformations can change to tg conformations by heating to temperatures above the polar-apolar phase transition temperature. The stable polymorph above this phase transition, which only contains tg conformations, is paraelectric (PE). All these polymorphs can be identified in DSC. Fig. 2.8

shows that films cast from highly entangled solutions (10vol.%) display a broader endotherm for the polar-apolar transition, which suggests that such material is composed of at least two phases, DFE and FE with different all-trans/tg conformation ratios. On the other hand, films cast from dilute solutions (0.5 vol.%) feature notably sharper phase transition peaks, indicating that the DFE conformation is dominant.

In order to identify the longer scale coherence of the chain packing as well as the general thermal structural development, wide-angle X-ray scattering (WAXS) was performed on films cast at three temperatures: 25°C, 55°C and 90°C and measured during heating. As already mentioned, FE, DFE and RFE unit cells have orthorhombic symmetry while the PE unit cell possesses hexagonal symmetry, thus each structure can be identified by their specific WAXS profiles. As documented in literature, the PE phase is associated with a characteristic Bragg peak positioned at 2θ=~18° corresponding to an inter-planar distance dPE=d200=d110=~0.493 nm. Typically, a relatively low full-width-at-half-maximum (FWHM) is found at 25°C [10]. In contrast, the orthorhombic FE, DFE or RFE arrangement is associated with a broader Bragg diffraction (due to vicinity of two diffraction lines originating from ( 200) and (110) planes) located at higher 2θs in the following order: FE>DFE>RFE, and smaller inter-planar distances with: FE<DFE<RFE [10].

The Bragg peaks near the (200/110) reflection recorded in CFE7 films in WAXS during heating are presented in Fig. 2.9. To characterize each crystalline polymorph, first, all identified peaks were deconvoluted as indicated by different colors in Fig. 2.9. The sharp diffractions correspond to crystalline peaks (highlighted in color) while the broad, weak halo is assigned to the amorphous material fraction (grey area). This allows to extract structural parameters such as inter-planar distance, crystalline phase content and a-b plane lateral extension (see Annex).

The overall degree of crystallinity (χtotal) of the various CFE7 structures was found to be 12.5 % at room temperature for films cast from 0.5 vol.% solutions and 8 % in films produced from concentrated solutions (10 vol.%) and cast at 40°C. The crystallinity increases during heating

and reaches 20% at 90°C. This means that the lack of chain mobility hinders the crystallization more in samples made from concentrated solutions. Yet, during heating the chains can rearrange leading to crystallization. Similar to the results obtained from DSC and FTIR spectroscopy, the two main arrangements in CFE7 at 25°C were found to be: DFE (orange peak) and FE (blue peaks). By heating the samples to 55°C (which is around the polar-apolar phase transition temperature) the DFE phase converts to a PE phase (green peak) in both samples while the FE phase remains stable.

This suggests that the low temperature shoulder of the polar-apolar phase transition endotherm observed for CFE7 cast from 10vol.% solution in DSC, located at 50°C, can be associated with a DFE-to-PE phase transition, while the second feature is the Curie temperature, i.e. the FE-to-PE phase transition. Comparing the samples with high and low entanglement density, that is samples cast from concentrated vs. diluted solutions, consistent with DSC and FTIR spectroscopy, the latter shows a more intense DFE Bragg peak than films cast from 10 vol.% solutions. The phase transition also starts earlier while a larger PE peak evolves at 55°C.

In addition to this qualitative interpretation, the structural parameters extracted from the Bragg peaks for the various crystalline phases are listed in table 2.2 and 2.3 as well as Fig. 2.10.

All phases are defined based on their calculated inter-planar distance (Fig. 2.10.a&d). For both entanglement densities, at room temperature, the films are richer in the FE phase than DFE.

The decrease in DFE content at 55°C leading to the formation of an PE arrangement, along with the constant FE phase content, supports the view that that the FE-PE transition occurs above 55°C, while the DFE-PE phase transition happens below 55°C. From the above, it is clear that the entanglement density affects the solid-state structure of CFE7. CFE7 cast at 40°C from concentrated solutions crystallizes into two phases, FE and DFE, while films cast from diluted solutions, crystallize into the DFE phase with less gauche defects than the DFE phase in highly entangled film. In general, according to our FTIR spectroscopy data, the total amount of all-trans and tg conformations for both scenarios are similar at room temperature but the difference is how the distribution of these conformations evolve in the different phases.

According to literature, a free VDF-TrFE chain is more stable in a tg conformation while this conformation is less favorable in a packed crystal (due to steric hindrance resulting from the presence of a third fluorine in TrFE in vicinity of neighboring fluorine atoms). As already discussed, CFE7 cast from solutions in the dilute regime, the polymer chains are almost

Casting temperature

The casting temperature is another highly relevant parameter that can affect chain mobility and, consequently, chain arrangement in a cast polymer film. Therefore, investigating the role of this parameter allows us to obtain clearer picture of the solidification and solid-state structure formation of VDF-based polymers.

We start our discussions with CFE7 film cast from 0.5 and 10vol.% solutions using the following deposition temperatures: 25°C, 40°C, 55°C, 90°C (i.e. between Tg and the onset of Tm). The casting method used is described in the experimental section on page 4. In order to minimize confusion between casting temperature and measuring temperatures, in the text, we assigned specific names to the samples which are listed in table 2.4. 12. shows the DSC thermograms of cast films upon heating with 10 °C/min. In order to see the casting temperature effect with different entanglement densities, both, material produced from diluted and concentrated solutions (Fig. 2.12.a: 0.5 vol %, and Fig2.12.b: 10 vol %), were used here.

As observed in Fig. 2.12, increasing the casting temperature sharpens the melting endotherm and shifts it towards higher temperatures (from 123°C in 0.5&10-CFE7-25 to 131°C in 0.5&10-CFE7-90 samples) in both scenarios. Evidently, by solidification of the polymer at higher temperatures, not only are crystallites obtained with a narrower lamellae thickness distribution (leading to a sharper melting endotherm) but also the lamellae thickness is increased (deduced from the increase in melting temperature). The variation in structure induced by different casting temperatures goes beyond the thickness of the formed crystallites; it also affects the local chain arrangement leading to differences in the temperature where the polar-apolar phase transition is recorded. Using a casting temperature of 25°C, two overlapping transitions are observed at 45°C and 65°C, respectively, which can be assigned to the DFE and FE to PE transition, respectively. Using a casting temperature of 40°C, the FE transition observed for samples cast from 0.5 vol.% solutions is slightly shifted to higher temperatures (67°C) and its relative intensity of the DFE transition (i.e. its enthalpy) is increased in contrast to samples cast from concentrated solutions. This observation implies that we obtain slightly higher fractions and denser FE arrangements by casting films at 40°C from diluted solutions. When casting at 55°C, the observed endotherms can be clearly distinguished and they shift towards opposite directions.

Indeed, FE arrangements with a Curie temperature of 71°C and a DFE phase with a Curie temperature of 41°C are generated using this casting temperature. The intensity of the FE endotherm is higher in films cast from 0.5 vol.% solutions compared to those cast from 10vol.% systems. This means that more FE phase is formed from 0.5vol.% solutions. In addition to FE and DFE phase formation when using this casting temperature, one broad phase also appears around 19°C, mostly in films cast from 10 vol.% solutions. We tentatively attribute this observation to a diffuse RFE arrangement. Finally note that in case samples were cast at 90°C, the DFE endotherm is significantly weakened and the FE feature disappears while the RFE endotherm develops more independent from the solution concentration used.

Exo up

Since DSC only provides insights on transition temperatures, we went on and identified the actual local chain arrangement in these phases using FTIR spectroscopy (Fig. 2.13). Due to the similarity of conformation induced when casting from concentrated and diluted solutions, we focus here on results obtain on systems cast from 10 vol.% solutions. All the data reported here were obtained at room temperature. In order to better observe the chain conformation evolution in such films, we concentrated on two wavenumber regimes (see Fig. In order to identify the various phases in our CFE7 samples, WAXS was used. To conclude, we monitored the solid-state structure formation of the CFE7 terpolymer dropcast at various temperatures. Among all the samples, 10-CFE7-25 and 10-CFE7-40, feature similar FE and DFE content with slight different FE/DFE phase ratios. 10-CFE7-55, due to its higher casting temperature, displays more gauche defects, accordingly, more DFE than FE is obtained, while also same RFE is observed. Increasing the casting temperature for 10-CFE7-90 introduces even more gauche defects and leads to the RFE to be the dominant phase. This chain arrangement and phase formation at different casting temperatures is related to the thermodynamic stability of each arrangement which is also responsible for the differences in Curie temperatures for each polar arrangement. Fig. 2.15 schematically illustrates the free energy of the ferroelectric, paraelectric and melt phases in P(VDF-TrFE)(75/25) copolymer structures as function of temperature and as documented in literature [11]. μF, μP and μM represent the free energy of each phase. The intersections of the free energy curves are representing the Curie temperature for these phases. As shown in this figure, at temperatures below the Curie temperature, FE is more stable and has a lower free energy than PE as well as the molten state. At temperatures above the Curie temperature, PE is more stable than FE and the molten state. At temperatures above the melting temperature, the molten state is stabilized [11]. Casting and annealing CFE7 at certain temperatures can decrease the free energies of associated phases and consequently stabilizes the formed solid-state structure. Casting at low temperatures (25°C and 40°C), favors denser crystalline phases with high alltrans chain conformation contents, i.e. DFE and FE. Casting at these temperatures increases the all-trans content in the DFE phase, which leads to its densification and a shift of the DFEto-PE transition temperature (TDFE-PE) to higher temperatures. This is illustrated by the downward shift of the FE/DFE/PE free energy line. This phenomenon is more pronounced in films cast at 40°C compared to those made at 25°C, due to the higher chain mobility afforded by the slight temperature increase. As a consequence, the DFE-PE phase transition merges with the FE-PE phase transition, as is clearly visible in the DSC thermograms presented in Fig. 2.12.

While casting at 55°C also leads to phases with all-trans chain conformations, more gauche defects are included in the chain arrangements. Therefore, the all-trans conformation content of the DFE phase decreases, leading to a shift to lower temperatures of its Curie point. This is well illustrated by the downward shift of the FE/PE free energy line. The stabilization of the gauche defects also increases the DFE and RFE contents to the detriment of the FE phase, as observed in WAXS. The increased chain mobility contributes to an increase of the degree crystallinity and to the formation of larger crystalline lamellae, as indicated by the large lateral coherence found in WAXS (Fig. 2.14.i).

Finally, casting at temperatures where in the paraelectric phase (90°C), bulky -CFE comonomers can be included inside crystalline regions, leading to a pronounced lamellar thickening of the RFE crystals, in agreement with the observed increase in Tm in DSC. This suggests that the polymer chains are pinned in tg + tg -conformations inside the crystalline lamellar arrangement that remains stable when cooling to room temperature, thus strongly decreasing the all-trans content even at ambient conditions. This results in a shift of the FE-PE phase transition to lower temperatures. The downward shift of the PE free energy line when casting at 90°C reflects both the lamellar thickening (increase of Tm) and the expansion of the crystalline unit cell (decrease of Tc). Therefore, we can conclude that the casting temperature, by stabilizing different phases, can considerably affect the solid-state structure of CFE7. In order to test this hypothesis on other VDF-based polymers, it will be necessary to study more VDF-based polymers with various stable crystalline phases and different phase transition temperatures which will be also discussed in this chapter.

Casting time

Kinetics of solidification, as well as its thermodynamics, plays an important role on the solidstate structure formation of VDF-based polymers. Thus, in order to study the kinetics of the solidification, CFE7 films were drop cast at 90°C using various casting times ranging from 1 min to 3hrs. A casting temperature of 90°C was chosen as it allows fast solidification. solutions. It can be seen from the data presented in Fig. 2.18.a that crystallization of dilute CFE7 solutions starts with DFE phase formation which stays almost unchanged for around 15min. After 3hrs, this DFE phase is converted to a RFE phase (note: that after 1 min, solidification of 0.5vol.% CFE7 content, the solvent was not fully evaporated, hence, most of CFE7 solidification occurred afterwards, i.e. at room temperature). When using concentrated solutions (Fig. 2.18.b), CFE7 crystallizes at first into a FE and a DFE phase with a broad lamellar thickness distribution which then progressively converts to a DFE phase after 15 min. Similar to structures produced from dilute solutions, the film evolves into a RFE phase with a narrow thickness distribution after 3hrs. Information on the chain arrangement of cast samples was obtained by FTIR spectroscopy. Fig. 2.19 displays the FTIR spectra of such films at room temperature. The observation of a 1288 cm -1 peak, in agreement with DSC, indicates a conformational all-trans to tg evolution when moving from short to long casting times at 90°C when using concentrated solutions while this evolution was not that well-pronounced when processing from dilute systems. In order to further scrutinize this evolution of the phase structure, WAXS was performed on CFE7 film cast at 90°C for 15 min and 3hrs using 10vol.% solutions. Extracted parameters are plotted in Fig. In the previous part a free energy diagram was proposed for CFE7 to explain its thermodynamic phase formation probability at different temperatures. Concentrating here on casting time, we observe that CFE7 can be crystallized into a FE phase even at 90°C when using very short casting times despite of the casting temperature used, where usually the PE phase should be stabilized. According to the free energy diagram (Fig. It is worthwhile to emphasize that this structural evolution can be readily monitored for CFE7 cast from concentrated solutions; it is not very pronounced in structures cast from dilute systems. This can be attributed to the faster crystallization of the latter systems. As explained earlier, in case of concentrated solutions, there are entanglements between the CFE7 chains in solution and high polymer-polymer interactions hindering crystallization. However, in low concentrated solutions, free polymer coils can crystallize separately leading to thick and well crystallized lamellas in the gauche formation. That is why the less entangled CFE7 forms a defective phase earlier than systems produced from concentrated systems.

Casting method

There are many different methods to cast a polymer from solution which affects the various solidification parameters like solvent evaporation rate. Among all of these techniques, three techniques are frequently used in laboratories: blade coating, drop casting and wire-bar coating. Each method needs a specific viscosity in order to coat a proper film. To investigate the phase behavior of CFE7 when cast with these methods, DSC was performed on cast films using 2 concentrations: 0.5 and 10vol.% and a casting temperature of 55°C. In all scenarios, FE, DFE and RFE phases are forming (Fig. 2.22). According to the DSC graph, no considerable differences is observed between the methods, we thus conclude that capacitors can be fabricated from any method and be compared. Fig. 2.24 displays the DSC results of this series of polymers, obtained when heating them from 0°C to 180°C at 10°C/min. Among all of these polymers, CFE7 shows the lowest melting point (122°C) suggesting the presence of the thinnest crystallites likely because of the presence of 7 mol% CFE7 leading to defects in its structure. Indeed, increasing the CFE fraction would prevent the crystallites to grow in the z direction and lead to thinner lamellae (compared to polymers with lower CFE content). In CTFE4, the lamellae thickness is increased as deduced from the fact that the melting peak is shifted to 145°C which is very close to the melting point of TrFE30 (148°C). As mentioned earlier, the VDF/TrFE ratio of these two polymers are very similar. Moreover, the amount of CTFE (4 mol%) is small; hence it is not surprising that their melting temperatures are very close. Even though the melting peak of these two polymers are similar, the onset of melting and their peak widths are remarkably different. We, thus, conclude that the CTFE termonomer can also affect the crystallite thickness distribution. In TrFE55, the melting temperature is higher than what is observed for the other polymers (157°C), which indicates the presence of thicker crystallites. In addition to considerable differences in overall crystallinity and crystallite thickness, relevant phase transition temperatures and the fractions of specific crystalline phases in a given structure also vary in these polymers. We observe, for instance, that CFE7 and CTFE4

have almost similar polar-to-apolar phase transitions (around 60°C). TrFE55 has a very diffuse transition covering a range from 38°C to 50°C, while TrFE30 shows a pronounced transition at 100°C. Since CFE7 and CTFE4 have very similar VDF-TrFE content, this difference may be explained by the termonomer type used and its content. As reported earlier, the bulky termonomer often is excluded from the crystalline phase when casting film at temperatures below their polar-to-apolar phase transition. Therefore, at this temperature, the induced phases in these polymers can be similar. In TrFE55, since the dominant monomer is TrFE, the We also evaluated the effect of entanglement density, casting temperature and casting time. Temperature (°C)

Drop cast @80°C
CFE7, casting at a temperature above the polar-to-apolar phase transition temperature of CTFE4, can stabilize the PE phase and increase gauche conformation. This causes a shift in both, the DFE and FE phase transition temperatures to lower temperatures (from 70°C to 54°C). The melting temperature is also sifted to higher temperatures (147°C to 150°C) (due to the lamellar thickening). However, the polymer chains are not as pinned in tg+tgconformations as in CFE7, in which a RFE phase is formed as the dominant phase. As mentioned earlier, the bulky terpolymer in CFE or CTFE can be accommodated in the crystalline PE phase thus keeping the interchain distance as well as a tg conformation even at low temperatures. In CTFE4, due to the low CTFE content, the tg conformation can turn to a more stable all-trans conformation at low temperatures leading to a higher fraction of DFE phase compared to RFE.

In TrFE55, by increasing the casting temperature, the stability of the formed DFE phase increases and the phase transition peak shifts to higher temperatures. Since the VDF/TrFE ratio of this copolymer is different to the other polymers studied here, the proposed free energy diagram is dedicated by the high TrFE content. The regions of phase stability are, thus, quite different in polymers with other VDF/TrFE ratios.

Finally, in TrFE30 polymer, the same trend as CFE7 and CTFE4 polymer is observed when casting them at high temperatures, leading to the stabilization of gauche defects and consequently to a less dense FE phase. However, the induced tg content is not as high as in other terpolymers, which is due to lack of bulky termonomer resulting in thicker lamellas in this copolymer.

In order to investigate the kinetics of solidification in the same series of polymers, they were drop cast at temperatures above their polar-to-apolar phase transition temperatures for short and long periods of time (5 min and 3 hrs) using 0.5 and 10 vol.% solutions. DSC was then performed on such cast films. Fig. 2.26 displays the DSC thermograms of such cast films upon a heating at 10°C/min. trans conformations (FE phase) when using short casting times. This structure formation is less pronounced during solidification from diluted solutions. Similar to CFE7, structure formation continues at room temperature leading to all-trans conformations. Since crystallization from dilute solutions is faster (due to a higher chain mobility), more tg conformations are formed in such structures.

In TrFE55 (Fig. 2.26.c), due to its different VDF-TrFE ratio, regardless of the entanglement density, cast films show a similar structure formation after 5 min casting. This structure minimally changes when using longer casting times. Indeed, just a small trace of FE phase is observed in films cast from 0.5vol.% solutions after 3 hrs casting.

In TrFE30 (Fig. 2.26.d), similar to CFE7 and CTFE4, a less defective phase is obtained when using shorter casting times which means that, at room temperature, a higher fraction of all-trans conformations can be induced. This all-trans stabilization at shorter casting times is similar in films cast from dilute or concentrated solutions.

Temperature (°C)

Cast for 5min

Exo up d

To summarize this section, VDF-based polymers with similar VDF/TrFE ratio (~70/30), can be crystallized into phases with large fractions of all-trans conformation by casting at lower temperatures than the polar-to-apolar phase transition temperature, and high tg conformation content when casting at higher temperatures than the polar-to-apolar phase transition temperature. We emphasize that the content of induced all-trans or tg conformations contents still depends on chemical structure (termonomer type and content) of the polymer. Therefore, the free energy diagram of these polymers will have similar types of regions with different stabilities.

Solid-state structure-induced dielectric and ferroelectric properties changes

As discussed in the literature review section, polar polymorphs are responsible for the dielectric and ferroelectric behavior of VDF-based polymers. After investigating the effect of solidification parameters on the solid-state phase formation, in this section we elucidate how structure affects the low-field dielectric (with relative permittivity measurement) and highfield dielectric (with polarization vs electric field measurement) properties of VDF-based polymers.

Relative permittivity measurements were performed on capacitors produced with CFE7 cast at three casting temperatures: 40°C, 55°C and 90°C, using 0.5 and 10vol.% solutions and bottom and top metallic electrodes (the device fabrication procedure is described in the experimental section, page 5). Fig. 2.27 shows the relative permittivity values as a function of temperature heating from -30°C to 90°C (two cycles) at 1kHz. Heating and cooling rates were 10°C/min. As mentioned in chapter 1, according to the Curie-Weiss law, the permittivity peaks are associated with the polar-to-apolar phase transition in the material. The low temperature peak in the dielectric loss (tanδ) graph corresponds to the Tg of the film. For devices cast at 40°C, a high temperature phase transition is observed at 70°C for structures produced with 10vol.%, and at 65°C for films cast from diluted systems (0.5 vol.%).

We tentatively assign this feature to the FE to PE phase transition. This transition is also observed in the form of a small peak in the dielectric loss data at the same temperature. The Tg peak is located at -15°C for films cast from concentrated solutions while it is slightly shifted to (-12°C) for structures produced from dilute systems. In devices cast at 55°C, in agreement with our DSC data, the phase transition is shifted to lower temperatures (50°C when using 0.5vol.% solutions and 55°C for films cast from concentrated systems) in the first heating cycle. Finally, in the devices cast at 90°C, a very diffuse phase transition is identified at 26°C which is associated with the RFE to PE phase transition. No considerable change in Tg observed when comparing both films.

In order to investigate the effect of casting temperature on the dielectric properties of CFE7 films, the permittivity of films cast at 40°C, 55°C and 90°C from 10 vol.% solutions was measured during two heating cycles (Fig. 2.28). Reassuringly, all the phases identified in DSC and FTIR spectroscopy are also observed in the permittivity measurements. By increasing the casting temperature, the permittivity values are decreased from 60 (90°C-cast films), 55 (55°C-cast films) to 45 (40°C-cast films). This decrease in permittivity values may be related to the crystallinity and chain conformation of these phases; indeed, the overall crystallinity is higher in samples cast at 90°C compared to films produced at lower temperatures. Moreover, during the first heating cycle, the dielectric loss (tanδ) peak arising at Tg, is increasing and shifting to higher temperatures when increasing the casting temperature. This suggests that when casting CFE7 at high temperatures, the bulky termonomer is progressively excluded from the amorphous phase increasing the chain mobility in the amorphous phase.

In order to study the ferroelectric properties of the various CFE7 phases, the polarization and current density of CFE7 capacitors were measured as function of applied electric field (Fig. 2.29). For this, various phase morphologies with different polarities were induced using different casting temperatures. The large hysteresis loop capacitors were produced with CFE7 cast at 40°C, with a high remnant polarization (Pr) and coercive field (Ec) at room temperature, which reflects the FE behavior in such structures in good agreement with DSC, WAXS and FTIR spectroscopy. The current for this samples is however noisy and leaky, possibly due to some remaining solvent in the film (even when drying under vacuum for 2 days). The switching current at a voltage of > 0 V suggests the presence of some DFE phase in addition to the FE phase. Casting at 55°C leads to two identifiable switching current peaks and a narrower hysteresis loop. The switching current peak at zero electric field indicates a large interchain distance in the material. This could explain that the required electric field energy for chain rotation and relaxation is lowered [12]. We, thus, conclude that the CFE7 gradually evolves from a FE to a RFE phase when casting at 55°C. For capacitors cast at 90°C, two well separated peaks are observed when measuring the current density as function of field. The large peak at negative electric fields indicates the presence of unstable polarized nanodomains arising from the RFE phase. The second peak at positive field is associated with the rotation of some chain segments comprising some CFE termonomer, thus, requiring more energy to be polarized [12].

To conclude, ferroelectric and relaxor ferroelectric behavior is observed in polarization measurements in agreement with our data from DSC, FTIR spectroscopy and WAXS. We observe that for films cast at 40°C, the solid-state structure is minimally changed after 1 week, while structures cast at 55°C and 90°C are notably modified. These changes are more pronounced in CFE7 films cast at 55°C. In these films, the FE-PE transition is relatively unchanged after 1 week, however the broad RFE-PE phase transition at 19°C disappears and the DFE-PE phase transition increases significantly in intensity. We conclude that the main modification originates from transforming the RFE to the DFE phase, i.e. the tg conformation content is decreased. In films cast at 90°C, similarly to those cast at 55°C, the RFE phase slightly evolves into the DFE phase which is due to tg conformations evolving into tt arrangements. Compared to films cast at 55°C, this modification is less pronounced and the structure seems more stable.

As mentioned earlier, casting of CFE7 at temperatures close to room temperature is thermodynamically stabilizing all-trans conformations which can form better FE/DFE phases (depending on the all-trans content). During aging of CFE7 films at room temperature (temperature below the polar-apolar phase transition temperature) all-trans conformation 82 can be induced. This all-trans stabilization is more pronounced when casting at 55°C than 90°C. Thus, by increasing the casting temperature, the formed tg chain conformation is more stable at room temperature. That means that casting at temperatures above the polar-apolar phase transition temperature, crystallizes CFE7 structures with a high interchain distance. This high interchain distance not only facilitates tg chain arrangements but also accommodates bulky CFE termonomer in the crystalline structure. This accommodation physically pins the chains and stabilizes a high interchain distance even at lower temperatures. Hence, high temperature casting (such as casting at 90°C) despite having predominant all-trans conformations at room temperature (all-trans conformations are more stable at temperatures below FE-PE phase transition), assists tg formation at room temperature.

However, when casting at 55°C, the interchain distance is not sufficiently high to accommodate the CFE termonomers and the induced tg conformations are less stable at room temperature.

In fig. 2.31, the stability of the formed phases when cast at 55°C and 90°C using 0.5 and 10 vol.% solutions was monitored by DSC (10°C/min, first heating) investigating various times after casting (2hrs to 1 week). Since the lamellar thickness seems not to be influenced by room temperature aging, only the polar-apolar phase transition peaks are reported here. Fig. 2.31. DSC heating thermograms taken at 10°C/min of CFE7 films cast at 55°C and 90°C using 0.5vol.% and 10 vol.% solution and aged at room temperature for the times indicated in the graphs.

For films cast at 55°C from 10 vol.% solutions, one broad RFE-PE phase transition peak is observed at 19°C. This transition is shifted to higher temperatures in films cast from 0.5vol.% solutions. CFE7 films cast from 10vol.% systems seem more stable over time than those cast from dilute ones. Our observations thus support the view that remaining tg conformations and consequent phases with more gauche defect at room temperature are due to chain pinning by the bulky CFE termonomer. Indeed, in structures cast from concentrated solutions, due to the decreased chain mobility, the CFE termonomers are included in crystallites while for films cast from 0.5vol.% solutions, CFE7 can be crystallized with a relatively little CFE in the crystalline lamellae leading to a higher stability of films cast from 10 vol.% solutions over time.

Similarly, in films cast at 90°C, more RFE phase is formed using 10 vol.% solutions compared to films cast from 0.5vol.% solutions. However, as mentioned earlier, structures cast at 90°C are even more stable than films cast at 55°C. We, thus, conclude that increasing the casting temperature increases the stability of the tg conformations and consequently phases with high tg content such as found in the RFE phase, and achieved in CFE7 cast at 105°C for 3hrs using 10 vol.% solutions. Fig. 2.32 shows the DSC first heating thermograms of CFE7 films cast at 105°C using 10 vol.% and 0.5 vol.% solutions for 3hrs and stored at room temperature for 2hrs to 1 week from casting. As expected, in the films cat from 10 vol.% solutions, more RFE phase is formed, and the structure is more stable compared to casting films at 90°C. However, even at this temperature, a small trace of DFE transition is detected when aged for 1 week. For the films cast from 0.5vol.% solution, the solid-state structure is still unstable and DFE phase is formed even after 2 hrs. Hence, due to the longer evaporations in diluted solution, solid-state structures are formed similar to films cast at 90°C which is unstable at room temperature.

To identify the different phases formed over time by storing at room temperature, WAXS was performed at room temperature on CFE7 films cast at 55°C, 90°C and 105°C from concentrated solutions when aged for 1 week at room temperature. The inter-planar distance, crystalline phase content and lateral extension of the various phases were then extracted from the (110)/(200) Bragg peak and are plotted in Fig. 2.33.

For films cast at 55°C, comparable inter-planar distances and almost similar unit cells are observed with aging, which suggests that the nature of the crystalline arrangement is not changed with aging (just a slight densification of the DFE phase is observed). However, we find that the RFE phase evolves into the DFE/FE phases with aging in agreement with the DSC data. In addition to the reduction of the RFE phase content, the lateral coherence of RFE also decreases with aging. For films cast at 90°C, consistent with our DSC data, a DFE phase is detected for aging of 1 week which we assign to a tt stabilization in the structure. In contrast, for films cast at 105°C, despite the fact that we record a small DFE transition in DSC, no DFE phase is detected in WAXS. Since this phase transition is considerably less pronounced than the one found for the RFE phase in other DSC scan, the associated peak might be covered in WAXS by the amorphous halo. The only other difference that we observe is a decrease in lateral coherence of the RFE phase when aged for 1 week. In order to investigate the influence of structure destabilization on the dielectric properties of CFE7 films, we measured their relative permittivity and dielectric loss data (fig. 2.34) as function of temperature upon heating with 10°C/min. In accord to our WAXS data, films stored for 1 week, when heated till their casting temperatures, rearrange leading to a similar structure as "as-cast" structures without aging. Consistent with DSC and WAXS, films cast at 55°C, the more defective crystalline phase (DFE) is transformed to a less defective crystalline phase and the polar-to-apolar phase transition temperatures is shifted to higher temperature. Films cast at 90°C, show more stable phases.

This stability becomes more obvious in films cast at 105°C, in which the permittivity is just slightly changed after 1 week storing. Note: we also can identify the Tg our dielectric loss data near -10 °C for all films, although the crystalline phase is evolving during aging, the amorphous phases is unaffected. We observe that ԑ" is reduced more than ԑ' for all CFE7 films independent of casting temperature (see fig. 2.35). This suggests that the imaginary part of permittivity is less stable than the real part. i.e., the crystal-amorphous interface may play a role, rather than the bulk.

For films cast at 105°C, the imaginary and real permittivity similarly decreases with time, but the properties are notably more stable than for films cast at lower temperatures.

In most reports on CFE7 to date, in order to increase the overall crystallinity, films usually were annealed at 105°C after casting. In fig. 2.36, we have compared the real and imaginary part of films cast at 105°Cwith films cast at 90 °C and subsequently annealed at 105 °C for 3h. In contrast to films cast at 105°C and measured as is, films cast at 105°C but annealed were found to be as unstable as films cast at 90°C. Since the stability of CFE7's crystalline phases originates from chain pinning by the CFE termonomer, during casting at 105°C the inter-chain distance is sufficient to accommodate CFE in the crystalline bulk as well as at interfaces while casting at 90°C, this accommodation is less pronounced. During annealing, the interchaindistance increases although the crystallites formed already during casting. Therefore, the crystallites cannot accommodate the bulky termonomer as easily as during casting. In turn, the solid-state structure is not as stable as in films cast at that temperature.

In conclusion, the solid-state structure formation of CFE7 depends on its casting temperature.

It is not stable over time at room temperature. While all-trans conformations are more likely to be present at room temperature, the formed tg arrangements progressively evolve into tt sequences in the structure. This transition is more pronounced for films cast at lower temperatures (except for temperatures below 40°C where they already solidified into a structure with a high content of all-trans chain conformations). Since the stability of defective structures is arising from chain pinning by bulky CFE termonomers, casting at high temperature can introduce more CFE into the crystalline regions and more stabilize the formed defective phase. This stability can be achieved during casting followed by annealing. and the C-H (in P(VDF)), the VDF chain arrangement can be influenced by the methacrylate group in PMMA [16]. With isotactic and atactic PMMA, in which the methacrylate groups are fully or partially oriented to one side of the PMMA chains, the induced chain orientation in VDF is more pronounced [15]. Indeed, is documented in literature that the methacrylate groups in PMMA can rearrange C-H bonds to one side of the chain and consequently induce all-trans conformations in VDF-based polymers during solidification. As discussed earlier, the all-trans conformations are responsible for the ferroelectric properties in VDF-based polymers. Thus, blending with PMMA can provide a tool to manipulate the structure of VDFbased materials in order to tune their properties.

In the following sections, an atactic PMMA with 80 kg/mol molecular weight was employed to investigate the influence of PMMA on the solid-state structure formation of CFE7. Thereby, the impact of PMMA content as well as of specific solidification parameters such as entanglement density of blend solution and casting temperature were studied.

PMMA content

PMMA is an amorphous polymer which is known to be well-miscible with VDF-based polymers. VDF-based polymers are semicrystalline polymers; i.e. they can partially crystallize during solidification. Hence, in blends, PMMA acts as a crystallization barrier for VDF-based polymer. Indeed, via the PMMA content in a PMMA:VDF-based polymer blend the overall crystallinity of the VDF-based component can be adjusted (independent whether the blend was produced from solution or melt) [18]. Increasing the PMMA content in such a system can decrease the overall crystallinity of the VDF-based polymer until a critical content is reached.

Then, the blend solidifies only into an amorphous phase containing a mixture of PMMA and VDF-based polymer. This critical PMMA content is generally found around 50 wt.% [19]. Since the dielectric and ferroelectric properties of VDF-based polymers arise mostly from the crystalline part of the material, we chose various PMMA contents below 50 wt.% (1.5, 6 and 12wt.%) in order to investigate the effect of PMMA on the solid-state structure formation of CFE7.

To prepare PMMA:CFE7 films with various PMMA contents, blend solutions were drop cast at 90°C for 3 hrs using 0.5 vol.% (PMMA+CFE7) solid content. In previous sections, FTIR spectroscopy was employed to detect different chain conformations in CFE7; here, in addition to CFE7 chain conformations, we also monitor any molecular interactions between PMMA and CFE7 [5]. A small displacement of the infrared peak at 1275cm -1 in blends compared to neat PMMA suggest that the C=O bond in PMMA is affected by the presence of CFE7, likely due to a dipole-dipole interaction with the C-H moieties in CFE7 [5,20]. Due to this molecular interaction, the chain conformation of CFE7 can be affected leading to different amorphous and probably crystalline phases than found in neat CFE7. Unfortunately, the all-trans CFE7 conformations (at 1280 cm -1 ) is covered by the -C-O-C-vibration peak of PMMA, hence, any possible evolution of all-trans chain arrangement cannot precisely be followed. However, since no FTIR peaks from PMMA structure are placed at 806cm -1 , the tg conformation of the VDF chains can be detected in the spectra. We observe that increasing the PMMA content, decreases the content of tg conformations in CFE7 chains (fig. 2.38.d), resulting in the CFE7 chains to adopt arrangements with more tt sequences [16,21]. This means that the dipole-dipole interaction can rearrange the VDF chains in a more developed trans form. However, in addition to this tt stabilization upon addition of PMMA, the intensity of the tg conformation peak is still considerable, even at 12wt.% PMMA, which may be due to the high casting temperature (90°). 2.39. We find that adding PMMA to the system shifts the melting point of CFE7 to lower temperatures. This means that the lamella thickness distribution of CFE7 is broadened and thinner crystallites are formed in CFE7 in the presence of PMMA. Since the PMMA are believed to be in the amorphous phase of the blend, they prevent the CFE7 crystallites to grow along the chain direction. In addition to the lamellae thickness, the overall crystallinity (derived from the enthalpy of melting) of the system is also slightly decreased when adding PMMA. This overall crystallinity reduction results from PMMA to kinetically and thermodynamically hinder the other component's crystallization [22,23]. This is supported by the fact that the crystallization temperature (TC) upon cooling is clearly shifted to lower temperatures (from 103°C for neat CFE7 to 96°C and 82°C in 6 and 12wt.% blends, respectively) (Fig. 2.39.c). Clearly, during solidification from the melt, similar to solidification from solution, well-mixed PMMA in VDFbased polymers can hinder the crystallization of the system.

Other observations can be made. The broad RFE phase transition (at 19°C) in neat CFE7 shifts to higher temperatures (32°C) upon addition of 1.5 wt.% PMMA while it is 19°C in blends of 6 and 12wt.% PMMA content. However, a very broad peak is detected at 70°C in blends with 6 and 12wt.% PMMA, which is associated with the FE to PE phase transition. In general, the stabilization of the all-trans conformation is more pronounced in blends comprising 1.5wt.% PMMA, which contrasts the tg destabilization observed in FTIR spectroscopy. In order to elucidate the evolution of the local chain arrangement over time by DSC and FTIR spectroscopy, DSC measurements were performed on the blend films after 1 day of casting and, consequently, the crystalline phases can be deduced by the fact that the DFE phase content increases with increasing PMMA content. This can be deduced from the change in DFE and FE phase transition peak intensity at 50°C. This DFE stabilization supports the view that PMMA stabilizes phases with more all-trans conformations. However, this stabilization of all-trans conformations is slow. Due to the presence of VDF-TrFE sequences in CFE7, the all-trans conformation is more stable than the tg arrangements at room temperature; however, casting the polymer at a temperature above the polar-apolar phase transition temperature (which is 70°) crystallizes the CFE7 in a phase with more tg conformations. The bulky CFE termonomer pins the chains and, in turn, the tg arrangements at room temperature -at least for a given time. This induced tg conformation and, consequently, the DFE/RFE phases are found to be less stable in PMMA blends. This means that PMMA not immediately but gradually assists the VDF chains to rearrange and form phases with more all-trans conformations. for 3hrs using solutions with 0.5vol% (PMMA+CFE7) total content, kept at room temperature for 1 week (heating rate: 10°C/min).

In order to identify the different phases in all blends, WAXS was performed on blend films, as-cast and after 1-month storage at room temperature. consistent with our FTIR spectroscopy and DSC data, the FE phase appears in blends over time (not in neat CFE7). In as-cast samples, 83% RFE phase with 17% DFE phase is obtained for neat CFE7 and the blends with 1.5wt.% PMMA (with increased lateral coherence) while addition of 6wt.% PMMA reduces this RFE content to 55% and increases the DFE content to 45%. In 12wt.% PMMA blend samples, the amount of the slightly denser RFE phase increases while the DFE phase content is reduced. We, thus, conclude that 1.5wt.% PMMA just increase the lateral coherence of the DFE phase, while addition of 6 wt.% PMMA also stabilizes the DFE phase. An increase in PMMA content to 12wt.%, further reduces the d-spacing rather than stabilizing the DFE phase. As interestingly, in addition to decreasing the crystalline lamellas thicknesses, which is observed in DSC through a reduction of the melting point, the addition of PMMA leads to an increase in the lateral extension of the crystallites along the ab plane. This development is maximized in blends of 6wt.% PMMA. At higher PMMA content, a further decrease in the lateral extension and the overall crystallinity is observed (Fig. 2.41). The same trend is observed in samples stored for 1 month, but additionally some FE phase is observed to appear. This FE phase is more strongly produced at increasing PMMA content until at 12wt.% PMMA, the FE content is found to be higher than that of the DFE phase.

From the above, we infer that by incorporating PMMA into the amorphous part of the system, the overall crystallinity of the blends is reduced and leads to thinner CFE7 lamellae while increasing their lateral extension (see Fig. the neighboring VDF chains resulting progressively into phases with more all-trans confirmations, especially in samples cast at high temperatures. This all-trans rearrangements as well as the increasing in the lateral extensions improves the macroscopic polarization.

Intriguingly, in blends with 1.5 wt.% PMMA, the crystallinity of the system changes while alltrans conformations are stabilized. At higher PMMA contents, however, the overall crystallinity considerably decreases.

Solidification parameters

In chapter 2.3, it was shown that not only the chemical structure of VDF-based polymers but also the selected solidification protocol can noticeably influence the solid-state structure formation of these materials. In this section, we discuss the solidification of the PMMA:CFE7 blends cast at different temperatures from dilute vs. concentrated solutions. We investigate the influence of these solidification parameters on solid-state structure formation. More specifically, 0.5 and 10 vol.% (CFE7+PMMA total solid content) blend solutions with various PMMA fractions (1.5, 6 and 12wt.% PMMA content with respect to overall polymer content)

were drop cast at 3 temperatures: 40, 55 and 90°C. In order to monitor the phase behavior of such cast films, DSC was performed. For films cast at 40°C using 0.5 vol.% solutions, we find a similar crystalline phase structure independent of the PMMA content in the blend although a slight decrease in the enthalpy of fusion is recorded for blends due to the reduction of overall crystallinity. However, in films cast from 10 vol.% solutions, the influence of PMMA is more pronounced. A significant broadening of the lamellae thickness distribution is observed when increasing the PMMA content. Due to the high chain entanglement density in these solutions, the chain mobility is low, thus, PMMA chains cannot easily migrate during solidification. Furthermore, the low casting temperature (40°C) does not provide sufficiently high thermal energy for notable chain movement. Thus, the crystallization in the blends is hindered and a very broad thickness distribution is found, deduced from the broadening of the melting peaks. The decrease in the high temperature phase transition observed in blends with 12 wt.% PMMA, which corresponds to a less defected FE phase may be due to the overall crystallinity reduction as well as a decrease in the lamellae thickness. Generally, the phase found when using this casting temperature (40°), independent of PMMA content, is FE.

For films cast at 55°C, the reduction in both overall crystallinity and lamellar thickness upon addition of PMMA is less pronounced, likely due to the higher thermal energy and consequent higher mobility of the polymer chains at these conditions. However, the crystalline phase formation of CFE7 is highly dependent on PMMA content. Increasing the PMMA content tends to stabilize all-trans-rich arrangements up to a content of 12wt.% PMMA, there, we obtain a higher FE phase content than DFE. This behavior, which is more pronounced in films cast from dilute solutions, is attributed to the dipole-dipole interactions between PMMA and VDF chains. These stabilize more all-trans conformations and, in turn, the FE phase.

For films cast at 90°C, due to the high thermal energy that is supplied, the overall crystallinity reduction in films cast from concentrated solutions is even less than for films cast at 55°C. Since CFE7 tends to be crystallized in PE arrangements at casting temperatures above the polar-apolar phase transition (70°C), the structure formation remains unchanged in the presence of PMMA. In contrast, in films cast from dilute solutions, addition of only 1.5 wt.% PMMA shifts the broad polar-apolar phase transition at 19°C to 30°C. This shift does not occur for blends with 6 wt.% and in binaries comprising 12 wt.% PMMA. For these systems, a broad FE peak is observed at 70°C. Consistent with the inter-planar distance extracted from WAXS, blends with 6 wt.% PMMA feature a higher fraction of RFE phase while in blends with 1.5wt.% and 12 wt.% PMMA, a higher number of all-trans arrangements are stabilized. This stabilization is less pronounced for films cast at 55°C. We, thus, suggest that the tg stabilization at high temperatures decreases the stabilization of all-trans conformations induced by the addition of PMMA.

Note that the tt stabilization is a slow process. We, thus, performed DSC on the cast films after one-day aging at room temperature. Fig. 2.45 shows the corresponding DSC heating thermograms measured at 10°C/min. For samples cast at 55°C, the FE phase remains unchanged during aging while a stronger DFE transition peak appears at 40°C in all samples. This DFE transition is more pronounced in blend systems than in neat CFE7. Comparing the films cast from dilute vs. concentrated solutions,

Exo up T m T FE-PE

a)

DFE stabilization is found to be more noticeable in films cast from 10vol.% solutions. When we compare fig. 2.44 and fig. 2.45, we see two competing trends for films cast at 55°C using 0.5 vol.% solutions: ageing stabilizes the DFE over the FE, while adding PMMA stabilizes the FE over the DFE. In samples cast at 90°C, the PMMA seems to delay the destabilization of the various phases.

We conclude that PMMA has different effects on CFE7 structure formation when varying casting temperature and solution concentration. At 40°C, the crystallization from highly concentrated solutions is more hindered while at 55°C, the phase formation is more strongly modified upon addition of PMMA: higher fractions of FE phase (with more all-trans chain conformations) are stabilized, especially when using dilute solutions. At casting temperatures of 90°C, crystallization is less hindered and the FE stabilization is less pronounced. However, interestingly, blends with 1.5 wt.% PMMA showed a high content of all-trans arrangements.

Impact of PMMA blending on dielectric and ferroelectric properties

In the previous section, we observed that in the PMMA:VDF-based polymer blends, addition of PMMA, not only hinders the crystallization of the VDF-based polymer component, but also affects the arrangement of the VDF chains in the solid state, hence, the structure formation of VDF-based polymers can be controlled via the PMMA content. This manipulation in structure allows to tune the dielectric and ferroelectric properties of these systems.

In order to investigate the low-and high-field dielectric properties of PMMA:CFE7 blends, the relative permittivity as well as the polarization were measured. For this, blend films with various PMMA content ranging from 0 to 12wt.% PMMA, were cast at two casting temperatures: 55°C and 90°C using 10 vol.% solutions. in the solid-state structure of such films. Except for films with 1.5wt.% PMMA content, which has a slightly stronger DFE-PE transition peak than that found in neat CFE7 films, increasing the PMMA content in blends, decreases the content of the DFE arrangements, while the content of the RFE arrangements remains unchanged. We note that in addition to the relative permittivity data, the DFE-PE phase transitions are also observed in form of a small peak in tanδ. Similar to the relative permittivity results, the recorded DFE-PE transition peaks deduced from the tanδ are also decreased when increasing the PMMA content (except for the films with 1.5wt.% PMMA, which shows a stronger DFE-PE transition peaks compared to the other films).

As mentioned above, after heating to 55°C, the solid-state structure is rearranged and forms a structure as found in "as-cast" films. This evolution of structure can be monitored in the second heating cycle. Indeed, in the second heating cycle, just one diffuse polar-to-apolar phase transition is recorded at 39°C for neat CFE7 films as well as films with 1.5wt.% PMMA, which covers other possible phase transition features. This broad polar-to-apolar phase transition, corresponding to the DFE-PE transition, disappears upon increasing the PMMA content and a RFE-PE transition feature with lower intensities is observed in films with 6 and 12wt.% PMMA. The possible FE-PE transitions is still not detectable due to the broad DFE-PE and RFE-PE transitions. These relative permittivity results are consistent with the DSC data:

all films have chain arrangements with more tg conformations that act as defects after casting at 55°C.

Since the formed structure is unstable at room temperature, a different arrangement is detected in the first heating cycle compared to that in the second heating cycle. This supports the view that during this aging at room temperature, tg conformations are destabilized and less defective arrangements are formed. Among all blends, samples with 1.5 wt.% PMMA are more modified upon aging while blends with more PMMA content show a more stable solidstate structure. In addition to evolution of the chain arrangement upon addition of PMMA, in as-cast films and samples after aging at room temperature, the Tg peak recorded in the tan δ graph is also shifted to higher temperatures upon addition of PMMA, which agrees with the picture that CFE7 and PMMA are miscible in the amorphous state.

For the samples cast at 90°, the DFE-PE and RFE-PE transition features are observed at 38°C and 25°C, respectively, in the first heating cycle. Since the intensity of the RFE-PE transition peak is higher than that in the blends cast at 55°C and the DFE-PE phase transition occurs at a lower temperature independent of the PMMA content, we conclude that more tg conformations are stabilized in CFE7 structures cast at 90°C. This tg formation is less pronounced in blends than in neat CFE7. The RFE-PE transition peak is weaker than the DFE-PE transition peak in blend films compared to neat CFE7. This structure formation likely arises from the stabilization of all-trans conformations in VDF chains in presence of PMMA. This alltrans stabilization is more pronounced in films with 1.5wt.% PMMA content. Fig. 2.47 displays the polarization of the blend films cast at 55°C (Fig. 2.47) and 90°C (Fig. 2.47) as a function of electric field. In order to form the structure as found in "as-cast" samples, the films cast at 55°C need to be heated to 55°C while films cast at 90°C require heating to 90°C.

The polarization after such treatment leads to a polarization as found in as-cast films. For samples cast at 55°C, consistent with the relative permittivity data, except for films with 1.5wt.% PMMA, the polarization of the samples decreases with increasing PMMA content. This is due to the decrease in overall crystallinities in such blends. In films with 1.5wt.% PMMA, this overall crystallinity is minimally changed while the all-trans conformations content is increased. Thus, a higher polarization is observed compared to neat CFE7.

Increasing the PMMA content not only decreases the polarization of the system but also decreases the Pr and weakens the hysteresis loop which contrasts the stabilization of the alltrans conformations found in CFE7 in the presence of PMMA. Indeed, PMMA is a dipolar glass which can be linearly polarized using an electric field (without any hysteresis loop). Addition of this dipolar glass to the system decreases the thickness of the crystalline lamellae as well as the overall crystallinity of the VDF-based component. This structural modification decreases the low-and high-field dielectric properties and delays the saturation of the polarization as a function of electric field. This delay slightly changes the non-linear ferroelectric polarization behavior to a more linear polarization. After a heating and rearranging the structure to the one found in as-cast films, the hysteresis loops are even narrower. As discussed earlier, heating samples to 55°C temporarily induces more tg conformations accordingly, the hysteresis loops become more narrow. In samples cast at 90°C, as expected from the DSC thermograms and the relative permittivity data, a higher amount of tg conformations decreases the hysteresis loop. The difference between the polarizations are less pronounced than in samples cast at 55°C. This is due to the higher overall crystallinities in these samples. Thus, with this casting temperature, the solidstate structure of samples is less affected by addition of PMMA compared to the structures formed during casting at 55°C. Similar to samples cast at 55°C, as-cast films possess more tg conformations than structures formed after one-day aging at room temperature. Hence, the all-trans conformations are stabilized during this type of aging. The effect of the PMMA on this stabilization is not clear from the polarization of the systems.

Conclusions

In this chapter, the solid-state structure formation of amending VDF-based polymers was investigated using materials with different chemical structures as well as the solidification conditions both in neat polymer and in blends with PMMA. It was shown that different and thickness distribution were achieved using dilute solutions. In concentrated solutions, a higher entanglement density prevents the CFE termonomers to be randomly distributed in the structure, stabilizing the RFE (with a high content of tg conformations) as well as FE arrangements (with a high content of all-trans conformations) more than DFE structures. In dilute solutions, a lower entanglement density assists the CFE termonomer to be randomly distributed within the structure and induces a higher content of DFE phase (with a mixture of tg and all-trans conformations). It is also observed that the overall crystallinity and the chain arrangements in CFE7, as well as the stability of these arrangements, are highly dependent on casting temperature. A free energy diagram as a function of temperature was proposed for CFE7 allowing to predict its solid-state structure formation via the thermodynamic stability of the formed crystalline structures as function of casting temperature. Based on this diagram, the crystalline phases with more all-trans conformations, leading to stabilization of DFE and FE phases, are favored when films are cast at temperatures below the DFE-PE phase transition. Casting the CFE7 at temperatures between the DFE and FE phase transition, stabilizes the FE and DFE phase (not as cast but 2 hrs after aging at room temperature). A solid-state structure with high tg+tg-chain conformation content is achieved after casting at temperatures above the FE-PE phase transition temperature. These solid-state structures were found to be more stable than those obtained using lower casting temperatures.

Increasing the casting temperature, increases the interchain distances in the structure. It, thus, can partially accommodate the CFE termonomer into the crystallites interfaces. This accommodation pins the chains and freezes the solid-state structure, formed at 90°C. By increasing the casting temperature to the temperature close to the melting temperature (i.e.

105°C for CFE7), this accommodation and pining is more pronounced leading to more stable defective crystalline phases. It was also observed that the chain pinning is more pronounced during casting than annealing of dried films. Therefore, casting the solutions directly at the onset of the melting temperature achieves more stable RFE phase with more gauche defects compared to films cast at lower temperatures and annealed at the onset of melting.

In the casting time study, it is observed that the less defective phase achieves using short casting times even at high casting temperatures. Long enough casting times (i.e.3hrs for the CFE7 solutions) is required to pin the chains by the bulky CFE termonomers in order to arrange and freeze more tg conformations leading to form more defective structures even at room temperature.

It is shown that increasing the bulky CFE termonomer content in the chemical composition of the terpolymer displays more defective crystalline structure with more tg chain conformations. It is also shown that independent from the termonomer content, crystalline phase formations follow almost similar trend (achieving more defective phases via casting at higher temperatures) in the VDF-based polymer structures with same VDF-TrFE ratio (70/30 in this work). The P(VDF-TrFE) (45/55) copolymer shows opposite trend due to its different crystallization behavior originating from its different chemical structure.

By investigating the polarization of the structures, it is shown that increasing the tg conformation content in the structure decreases the hysteresis loop and provides more relaxor ferroelectric properties. For instance, a VDF-based terpolymer with 7 mol.% CFE content has a stable RFE phase with phase transition to PE at 20°C when it is cast at high temperatures. This relaxor behavior decreases the energy loss during application of the electric field which provides more efficiency to the designed capacitors suitable for the energy storage applications.

In the CFE7:PMMA blends, PMMA is known to be miscible with CFE7 in the amorphous state and hinders the crystallization of CFE7 specifically along the z-direction. Therefore, thinner but bigger lamellas are obtained upon addition of PMMA content in the systems. In addition to modifying the morphology of the crystallites, atactic and isotactic PMMA chains, due to their dipole-dipole interactions with VDF chains, arranges the all-trans conformations in the VDF chains. These all-trans arrangements are more pronounced when cast at 55°C where the FE phase is more favored in the CFE7 structure. It is also shown that entanglement density of the blend solutions has significant influence on the solid-state structure of systems. In the 10vol.% blend solutions, the systems are solidified with less crystallinity and broader lamellae thickness distributions comparing to the films dried using 0.5 vol.% solutions. That means that PMMA hinders the crystallization more in the concentrated solutions. On the other hand, in dilute solutions, PMMA has more influence on the all-trans stabilization than hindering the crystallization. Since the overall crystallinity decreases with the PMMA content, blend films with 1.5wt.% PMMA in which the overall crystallinity is minimally decreased, higher polarization is detected as a function of electric fields. In addition to the all-trans stabilization in the blends, saturation of the polarizations is also hindered which changes the shape of the hysteresis loop leading to a more linear and hysterises free polarization. This polarization behavior can increase the energy density of the system which make them the good choices for energy storage applications.

Motivation and context

Inorganic ferroelectric ceramics became attractive in the 1940s via the discovery of the ferroelectricity in barium titanate (BaTiO3) capacitors with very high dielectric constants [1].

This breakthrough introduced these ceramics as promising systems for ferroelectric-related applications such as piezoelectric transducers, energy storage capacitors, medical diagnostic transducers, electro-optic light valves, ferroelectric thin-film memories, and solid-state cooling devices. [2]. Nowadays, barium titanate perovskite oxide, possessing an ABO3 structure (Fig. 3.1), and its derivatives, due to their high dielectric constant (ԑr > 1000) and high polarizability, combined with their low toxicity, remain dominant materials for these applications. This directly affects the ferroelectric properties of the material. The size and the charge disorder introduced by the ion substitutions lead to a breaking-up of the ferroelectric domains into short-range ordered polar nanodomains and, thus, induce a more relaxor behavior [4,5].

As a result, the polar-to-apolar phase transition is shifted towards lower temperatures.

Accordingly, modifying the chemical structure of barium titanate is a well-known path to modify its ferroelectric properties [6,7]. For instance, Smolensky et al. [8] introduced Sr +4 ions into the Ti position in the barium titanate lattice (BaSrxTi(1-x)O3) to change the ferroelectric properties to typical relaxor ferroelectric properties using 30 mol% Sr (x=0.3) with respect to Ti. At 0.175<x<0.25 content, both ferroelectric and relaxor ferroelectric features are recorded

In literature, there are various ways to synthesize perovskite oxide nanoparticles. These are generally classified into two methods: solid-state and soft routes. The solid-state method is a based on the synthesis of oxides based on mixing and grinding oxides and carbonates in the solid-state at high temperatures in order to synthesize more complex micron-sized oxide particles or bulk. This method is appropriate for the synthesis of many ceramic oxides such as perovskites at high temperatures and pressures. This very high energy and time consumption, as well as the micron size of the products, limit utilization of this technique for the synthesis of nano-sized inorganic particles [15].

The soft-route method is another class of techniques based on dissolving ions in a solution in order to synthesize ceramics into fibers, coatings, micron-and nano-sized particles, etc. The ability to control this as well as the energy and time consumption, renders this method more favorable, especially for the synthesis of the nano-sized particles. Co-precipitation, sol-gel, hydrothermal, and more recently, polyol methods are known as common soft-route methods.

In this chapter, we use hydrothermal and polyol methods in order to synthesize nano-sized ferroelectric barium titanate oxide (BaTiO3) and its derivatives modified using Zr +4 ions in order to enhance the dielectric and ferroelectric properties of VDF-based nanocomposites.

The hydrothermal method is a soft-route method based on hydrolyzing metal precursors in a basic aqueous solution under high pressures. Thereby, the concentration of metal precursors and hydrolysis agent play an important role in the synthesis of BT through a hydrothermal procedure. Accordingly, the influence of the concentration of hydrolysis agent, the reaction time as well as the annealing process on the synthesis of barium titanate batches is investigated in this chapter.

The polyol method, similarly to the hydrothermal method, is a soft-route method based on the forced hydrolysis of metal precursors in a polyol medium such as diethylene glycol (DEG) [17]. In the hydrothermal method, the hydrolysis of precursors is carried out in the presence of water at high pressures (due to the low boiling point of water (100°C) at ambient pressure). In contrast, in the polyol method, due to the high boiling point of the polyol medium, reactions can be carried out at ambient pressure. In addition to this milder pressure, the polyol medium with its -OH units can participate in the reaction and may act as a barrier between the growing particles and isolates them. This isolation is known as chelating effect.

The chelating behavior provides control of the morphology and size of the particles, and results in producing more homogeneous particles.

During the synthesis of barium titanate using a soft-route method, barium carbonate (BaCO3) with a free energy that is lower than in barium titanate (BT) can be produced as a secondary phase in the particles. The amount of this side-product that is synthesized due to the presence of CO2 can be minimized if the reaction is carried out under neutral gas atmosphere such as argon or nitrogen. Furthermore, annealing the produced nanoparticles at very high temperatures not only can decompose and eliminate this barium carbonate from the product but also results in the crystallization of the remaining amorphous phases leading to highly crystalline single-phase BT nanoparticles. Similar to the hydrothermal method, the influence of the concentration of the hydrolysis agent as well as the annealing temperature was investigated in this chapter on the synthesis of BT nanoparticles via the polyol method.

As mentioned in chapter 1, the unit cell of crystalline BT can take on in 4 different symmetries: cubic, tetragonal, rhombohedral, and orthorhombic. All of these symmetries are polar and provide ferroelectric properties except the cubic symmetry, which is a non-polar centrosymmetrical structure. All of these polymorphs can undergo the polar-to-apolar phase transitions (tetragonal/rhombohedral/orthorhombic to cubic phase transition) at various temperatures. Since the rhombohedral and orthorhombic structures in BT are not stable at room temperature (Fig. 3.2), the tetragonal polymorph is responsible for the ferroelectric properties in this material. Accordingly, an ideal ferroelectric BT nanoparticle is a pure and well-crystalline nano-sized (i.e., 100-200nm) particle comprising mostly the tetragonal structure at room temperature. However, due to the presence of higher internal stresses in the lattice in nanoparticles compared to the bulk, the tetragonal structure is less stable in nano-sized particles compared to the bulk; hence, producing nano-sized tetragonal BT has been found challenging in form of nano-sized particles.

Although the breakdown strength in the nanocomposites is high compared to that in bulk ceramics, aggregation of nanoparticles can lead to regions of low resistance to electrons resulting in a reduced dielectric strength of the nanocomposite. Moreover, aggregation of particles decreases the interfacial area between particles and the polymer, which results in reducing the interfacial polarization in the nanocomposite. Therefore, in order to reduce this local charge accumulation in nanocomposites and increase the interfacial polarization of the system, the surface of the nanoparticles is in this thesis chemically modified using different organic functional groups. This surface modification manipulates the surface energy and facilitates the dispersion of the nanoparticles into the polymer matrix, thus limiting possible agglomerations.

More specifically, in this part of the thesis, we present the synthesis and surface modification of nano-sized barium titanate and barium zirconium titanate (BaZrxTi(1-x)O3) having different Zr content (x=0.12, 0.2 & 0.3) to produce nanoparticles with different ferroelectric and relaxor ferroelectric properties. These nanoparticles, with their various polar-to-apolar phase transitions (depending on their chemical structure), allow manipulation of the ferroelectric properties of VDF-based nanocomposites rendering them suitable for applications, such as electrocaloric solid-state cooling. In the next part, the synthesized nanoparticles were blended with P(VDF-TrFE-CFE) and P(VDF-TrFE-CTFE) polymers in order to tune the solid-state structure of the polymers and enhance the polarization of the composite hybrid systems.

Materials and methods

Barium titanate (BT) nanoparticles were produced via a hydrothermal method using a protocol adapted from Huarui Xu's [18]. Some changes were applied to the parameters in order to tune the required properties, such as the particle size and the crystalline phase structure. For the synthesis of BT via the hydrothermal method, different barium and titanium precursors have been reported in literature [18]. In this work, barium and titanium chloride (BaCl2 & TiCl4) were utilized as the metal precursors. Sodium hydroxide (NaOH) with different concentrations was used as the hydrolyzing agent. First, 0.0156 mol barium chloride and 0.0156 mol titanium chloride were dissolved in 25 ml deionized water at room temperature and stirred for 1 hr. During this step, titanium chloride, as a very unstable precursor, starts to be hydrolyzed in the presence of water and produces titanium dioxide (TiO2) and hydrogen chloride (HCl) (see scheme 3.1). Since the hydrolysis of the precursors only occurs at a basic solution, the production of HCl in solution, which decreases the pH of the system, can stop the hydrolysis of the TiCl4 and prevent the BaCl2 from hydrolyzing [17]. To increase the pH, NaOH as a hydrolyzing agent was added to the solution, and then the solution was transferred to an autoclave. In this step, BaCl2 in the presence of water and NaOH at high temperatures (240°C) gradually diffuses into the formed TiO2 lattice and produces a BT perovskite oxide with an ABO3 unit cell structure (see scheme 3.2). This structure consists of a corner-sharing network of oxygen octahedra, creating an octahedral cage (B-site) with interstices (A-sites), as shown in fig. 3.1 thereby, the Ba 2+ ions occupy the A-site while the Ti 4+ ions occupy the Bsite. This synthesized BT can likely comprise -OH groups inside and on the surface of the lattice. In order to eliminate all of these -OH groups from the nanoparticles, the assynthesized powders were annealed at different temperatures for 8hrs. The reaction was carried out in a ~50ml Teflon-lined autoclave placed in a high-temperature oven. After the synthesis, the produced nanoparticles were washed three times using water and ethanol, respectively, and were dried at 90°C overnight. In the next step, the dried nanoparticles were annealed at 650°C and 800°C in a high-temperature furnace.

BT nanoparticles were also produced via a polyol method adapted from Ulises Acevedo-Salas' work [19]. Some changes were applied in the parameters in order to synthesis BT nanoparticles with desired size and structure. For this purpose, barium acetate (Ba(C2H3O2)2)

as the barium precursor and titanium isopropoxide (Ti(OCH(CH3)2)4) as the titanium precursor were used with sodium hydroxide (NaOH) as the hydrolysis agent in diethylene glycol (DEG)

as the media (Scheme 3.3). The reaction was carried out under heat reflux at the boiling point of diethylene glycol (DEG) (220°C). After 5hrs, the precipitated paste (without washing) was dried at 90°C overnight under vacuum and then annealed in a high-temperature furnace.

After identifying the desired synthesis parameters, barium zirconium titanate (BaZrxTi1-xO3) with three different Zr content (x=0.12, 0.2 & 0.3) was synthesized via the polyol method using obtained experimental parameters from the synthesis of BT. Stoichiometric amounts of Zr precursor were used in order to achieve the target compositions for BZT.

In the next step, the surface of the synthesized BT and BZT nanoparticles were modified via a sol-gel method using functional groups with various hydrophilicities, such as hydroxyl groups (-OH), (3,3,3-trifluoropropyl) thriethoxysilanes with a -CF3 terminal group, and perfluorododecyl-1H,1H,2H,2H-triethoxysilanes with a highly fluorinated -C6F13 terminal group. First, the surface of nanoparticles was modified using hydrogen peroxide at high temperatures in order to introduce the -OH groups. Then, various silane-based coupling agents containing different functional groups were used to modify the surface of the nanoparticles.

In order to prepare the nanocomposites, the nanoparticles were sonicated in 5vol.% (with respect to the solvent) polymer solutions for 40 min. Metal (silver)-insulator (nanocomposite)-metal (silver) capacitors were fabricated via drop-casting the well-dispersed nanocomposite dispersions on a bottom electrode at 90°C for 3hrs and evaporating a top electrode onto the cast films using a shadow mask to achieve 2 mm 2 active surface areas. The thickness of the nanocomposite films was chosen to be from 10 to 20µm for the different analyses. All the fabricated devices were annealed at 105°C for 30min before use.

Barium titanate (BaTiO 3 ) perovskite oxide nanoparticles 3.3.1. Barium titanate synthesized via a hydrothermal method

In this part, nano-sized barium titanate particles were synthesized via a hydrothermal method. As mentioned before, there are various parameters that affect the chemical and crystalline structure as well as the morphology of the produced BT using this method. These parameters include the concentration of the metal precursor and the hydrolyzing agent, reaction time and temperature, as well as the annealing temperature and time. In this work, the influence of the concentration of hydrolysis agent, the reaction time as well as the annealing temperature was investigated. Table 3.1 shows all the batches with the detailed parameters in addition to the size of the synthesized nanoparticles from each batch (last column). In order to characterize the chemical composition of the synthesized nanoparticles, FTIR spectroscopy was performed. Fig. 3.5 displays these FTIR spectra in comparison with commercial 200-nm BT nanoparticles. One main peak is observed at 495 cm -1 in the commercial BT spectra (black spectra) as well as the synthesized BT produced from run 3 (green spectra), while is slightly shifted towards lower wavenumbers in the synthesized BT from Run 1 (red spectra) and Run 2 (blue spectra). This peak, which likely originates from the stretching vibration of the Ti-O bond, is usually accompanied by a bending vibration of the Ti-O unit in 400 cm -1, mostly in the tetragonal BT. This bending vibration feature is beyond the range of the FTIR spectrometer used. The Ti-O vibration band on the BT nanoparticles produced from run 1 and 2 have slightly different absorption bands compared to the commercial BT, while the nanoparticles produced from run 3 possess a Ti-O band close to a tetragonal BT structure. Moreover, the nanoparticles produced from Run1 also display a broad -OH band in its spectra, which is likely due to the presence of some residual -OH units in the nanoparticles. These -OH units can be either inside the lattice or on the surface of the particles. Note that this sample was not annealed at high temperatures, while the other In order to find the proper annealing temperatures, thermogeravimetric analysis (TGA) was first performed on the samples produced from run 1 (non-annealed nanoparticle) and then performed on the annealed nanoparticles produced from run2. Fig. 3.6 displays the mass loss of these nanoparticles (synthesized from Run1 and 2) upon heating to 900°C. Consistent with FTIR spectroscopy, the samples synthesized from run 1 show some degradable species which were degraded at 100°C to 300°C. According to FTIR spectroscopy, -OH functional groups may be attributed to this degradable species. Considering the degradation temperatures, the presented -OH functional groups are divided into two groups: surface -OH and lattice -OH groups. The surface -OH is eliminated at 100°C to 200°C, while the lattice -OH is degraded around 300°C [20].

In addition to the presence of -OH groups, another degradable species is detected at 500°C to 600°C in the produced nanoparticles from run 1, which are likely due to the presence of some produced barium carbonates during the reaction. According to the TGA data of the nanoparticles produced from run1, at 800°C, all the degradable components are eliminated from the particles, and a pure product is achieved. This temperature thus is an appropriate temperature for annealing of the nanoparticles to remove all impurities. Supporting this view, no mass loss is observed in the samples synthesized from Run2, which were annealed at 800°C. These results exhibit the effect of annealing on the purity of the product. In order to identify the crystalline structure as well as the crystallinity of the produced nanoparticles, X-ray diffraction (XRD) was performed on the nanoparticles synthesized in different batches with different parameters. Compared to the commercial BT, nanoparticles produced from run 3 and 4 show comparable lattice dimensions (similar to a cubic perovskite oxide structure), while as we expected from the FTIR spectroscopy, the XRD pattern of the nanoparticles produced from run 2 shows a pattern similar to a crystalline structure of TiO2. Thus, the reaction is not completed in these nanoparticles, which is due to the low concentration of the hydrolyzing agent (NaOH) in the reaction-increasing the concentration of the hydrolyzing agent results in the synthesis of BT instead of TiO2.

(3 1 0) (3 0 0) (2 2 0) (2 1 1)
(2 1 0)

(2 0 0) (1 1 1)
(1 1 0)

Run4

Run3

Run2

Counts

2 (°)

Commercial

(1 0 0) a) b)
123 Since a clear and pure tetragonal phase is not detected via XRD, DSC was employed in order to detect any possible polar-to-apolar phase transitions. Fig. 3.8 shows the DSC thermograms of the synthesized BT nanoparticles measured with 10°C/min heating. A small endothermic feature is observed at 115°C, corresponding to the tetragonal to cubic phase transition in BT, in the commercial and synthesized BT from run 3 and 4. Therefore, a polar tetragonal phase is also present beside the possible non-polar cubic phase in our BT structures. Note that based on the XRD patterns, this polar tetragonal phase is not pure, and the structure is more like a pseudo-cubic structure. In conclusion, the nanoparticles produced from run 1 (without annealing) contained some -OH impurities, which were detectable by SEM, FTIR spectroscopy, and TGA. Although

annealing the nanoparticles at 800°C eliminated these impurities, the produced nanoparticles were based on TiO2 (from run 2). The shortage of hydrolysis agent in the reaction did lead to a low pH, and, consequently, the reaction was terminated after TiO2 formation. The nanoparticles produced from run 3 provided a well-crystalline pseudo-cubic BT nanoparticle with 50 to 200 nm particle size without any secondary phases. Increasing the reaction time in run 4 increased the size of the particles while the crystalline phase remained pseudo-cubic.

The polar-to-apolar phase transition was found to occur at ~115°C, which is very close to a tetragonal to cubic phase transition in the commercial BT.

Barium titanate synthesized via a polyol method

Similar to the hydrothermal method, nano-sized BT particles were synthesized via a polyol method using different reaction parameters. The influence of the concentration of the hydrolyzing agent and annealing temperature on the chemical and crystalline structure, as well as the morphology of the produced BT particles, was investigated in this part. Table3.2

shows the parameters in detail in these batches. In order to investigate the morphology of the synthesized BT nanoparticles, fig. 3.9 shows SEM graphs of the particles synthesized from the different runs listed in Table 3.2. SEM of the particles produced from run 1 show very small 10-50nm particles size, which are highly stacked together, while the produced particles from runs 2 and 3 were produced with a homogeneous size around 100 to 200nm. Non-homogeneous, larger particles (100 nm to 1 micron-sized) were produced in run 4. Increasing the concentration of hydrolysis agent in run 4, compared to the other runs, kinetically increases the rate of the particle growth, which results in the production of non-homogeneous micron-sized particles. ). The synthesized nanoparticles from run 1 show a complex pattern which means the material is still full of ions and organic matter before annealing. The obtained peaks in these samples with their corresponding bands are listed in table 3.3. According to these absorbed bands, the majority of the characteristic peaks correspond to the presence of diethylene glycol (due to absorbed -OH and -CH bonds). The lack of asymmetric stretching vibrations of the carboxylic acid in these FTIR spectra (~1700 cm -1 ) which present in the acetate groups of barium acetate, indicates that the barium acetate is fully dissolved and washed off. The remaining diethylene glycol is due to the possible bonding of -OH in DEG with the Ti-O lattice. As mentioned in the introduction section, one of the advantages of the polyol method over the other methods is the chelating effect of polyol (DEG in this work) during the reaction, which separates the growing crystalline particles as the DEG molecules "attach" to the particles separating them. This chelating effect results in the production of more homogeneous particles. In the produced particles from runs 2 and 3, which are the annealed particles, the remaining DEG is degraded, and only stretching vibrations of the Ti-O and Ti-O-Ti bonds are observed at 495 cm -1 which is comparable to that in the commercial BT. In the produced samples from run 2, in addition to the Ti-O unit peak, another small feature is observed at 860 cm -1 which corresponds to the C-O bond in barium carbonate. Hence, in samples produced from run 2, which was annealed at 650°C, barium carbonate may be produced as a secondary phase. after 650°C and 800°C annealing measured at room temperature. The FTIR spectra of the commercial BT is used for comparison.

Table 3. FTIR absorption bands with their corresponded functional groups deduced from FTIR spectra of the produced BT from run1. In order to determine an appropriate annealing temperature to obtain pure BT without any secondary phase, TGA was performed on non-annealed samples (from run1). Wavenumber (cm -1 )

Absorbtion (cm

Commercial BT Run2 Run3
127 for run 2. Still 0.5 wt.% mass loss was observed when heating to 800°C (Fig. 3.11.b). According to the FTIR spectra of the produced nanoparticles from run 2, this weight loss is due to the release of CO2 through the decomposition of barium carbonate at high temperatures.

Therefore, the barium carbonate can be produced as a side-product even under an argon atmosphere. This side-product can be eliminated during annealing to temperatures above 800°C. This formation of the pure phase is observed in the TGA of such nanoparticles produced from Run 3, which were annealed at 800°C. from the hydrothermal method, no peak splitting is observed for the (200) feature, and the produced BT nanoparticles from run 3 and 4 are possibly pseudo-cubic. In conclusion, very homogeneous, single-phase, highly crystalline BT particles with a particle size of 100-200nm were synthesized via a polyol method at 220°C in DEG followed by annealing at 800°C. Barium carbonate was produced as a side product which is eliminated through annealing at 800°C. Increasing the concentration of hydrolysis agents also increased the particle size.

Barium zirconium titanate (BaZr x Ti (1-x) O 3 ) perovskite oxide nanoparticles

Barium zirconium titanate (BaZrxTi(1-x)O3), similar to barium titanate, is a perovskite oxide (ABO3) compound that is produced via substitution of titanium ions (Ti 4+ ) with zirconium ions (Zr 4+ ) in the B-site of the barium titanate structure. Partially adding Zr 4+ cations with bigger ion radii compared to that of the Ti 4+ cation can change the tetragonal symmetry of BT to a less stable rhombohedral/orthorhombic symmetry in barium zirconium titanate (BZT), which evolve into non-polar cubic structures at lower temperatures compared to polar-to-apolar phase transitions in conventional BT structures (at ~115°C). More specifically, Zr 4+ cations as defects can break the micron-sized ferroelectric domains into crystalline BT structures of nano-sized relaxor ferroelectric domains. This leads to structures with relaxor ferroelectric properties. These relaxor properties fully depend on the amount of Zr 4+ cation in the structure. As reported in literature, only 20 mol% of Zr in B-site can cause a shift in the polarto-apolar phase transitions to room temperature [24], which is desirable for some applications such as electrocaloric cooling applications. In this subchapter, a polyol method was used to synthesize BZT nanoparticles with various Zr content in order to obtain relaxor ferroelectric properties with various phase transitions. 3) to obtain the desired composition in the produced compounds. Except for the concentration of zirconium precursor, all the utilized parameters were adopted from the successful synthesis of BT (from run 3) discussed in the previous part. These parameters are listed in table 3.4. Since the reaction parameters were adopted from the synthesis of BT from run 3, all the BZT particles, similar to BT, have a homogeneous ~ 50 to 200 nm size. The morphology of the synthesized particles is also spherical. Since Ti-O and Zr-O bonds have comparable stretching vibrations at 495 cm -1 , similar peaks are recorded in FTIR spectra for BT and BZT samples. The only difference is a slight shift towards higher wavenumbers, which is likely due to the slight difference in the unit cell volume of BZT structures compared to BT.

Synthesis of

Since FTIR spectroscopy is not a sufficient characterization to distinguish BZT and BT structures, energy dispersive X-Ray analysis (EDX) was carried out on the nanoparticles, Fig. 3.15 shows EDX features corresponding to the elements present in the compounds as a function of applied energy measured on BZT20. This analysis supports the view that zirconium is present beside barium and titanium. Hence, zirconium which is a less reactive ion compared to a titanium ion is oxidized and participates in the BT lattice. This spectrum was repeated 4 times at different locations in the samples. The amounts of the detected elements were estimated using the intensity of the recorded peaks. These are listed in table 3.5. The estimated values are comparable with the amounts of elements expected to be present in BZT20. Table 5. Estimated amounts of detected elements from EDX features using 5 different spectra.

In order to quantitatively identify the crystalline structure of the synthesized nanoparticles, Note that distinguishing an orthorhombic/rhombohedral symmetry from a cubic symmetry in nano-sized BZT particles are more difficult than distinguishing a tetragonal symmetry from cubic symmetry in nano-sized BT particle. Furthermore, due to the type of polar-to-apolar phase transitions which is second-order in BZT ceramics, no clear phase transition is also detectable by DSC. Therefore, additional characterization techniques are required to identify the presence of any possible polar non-cubic phases in the synthesized BZT nanoparticle structures, which would lead to a relaxor behavior. Dielectric spectroscopy is a suitable technique, in which even very broad pseudo-cubic to cubic phase transitions are detectable.

However, since this methodology is generally performed on bulk samples, using it to measure nanopowders was found challenging. To overcome this issue, the synthesized nanoparticles were sintered produce BZT nanoceramics with comparable grain sizes.

Effect of Zr content on structure and dielectric properties of nanoparticles in spark plasma sintered nanoceramics

Spark plasma sintering (SPS) is a sintering technique in which DC or AC currents pass through a graphite die containing powder in order to sinter them without any grain growth to occur (see Fig. 3.17). Due to Joule heating produced from this current, uniformly distributed internal heat is generated inside the material that, along with applied pressure, can densify the powders. In contrast to conventional sintering techniques in which the thermal energy is provided using an external source, in SPS, the generated heat is internal. This internally generated heat facilitates heating and cooling and, in turn, accelerates the sintering process.

This acceleration prevents any grain growth during densification. In order to identify the required parameters for sintering the BZT nanoparticles without any grain growth, the synthesized BZT20 nanoparticles were densified in several SPS runs where different sintering temperatures were used. The utilized parameters are listed in table 3.6.

Since we used carbon dye as the pressing arm in the device and carbon is more reactive than metal, the metal oxide nanoceramics can be reduced in contact with these carbon dyes during sintering at high temperatures. In this reduction process, the lost oxygen from BZT at high temperatures oxidize the carbon dye and release CO and CO2 gas as a product. Hence, in order to reoxidize the nanoceramics, the sintered nanoceramics were heated at 800°C for 3 hours in a furnace under an air atmosphere. Theoretical density of BaZr0.2Ti0.8O3 =5.96 g/cm 3 Fig. 3.18 shows a piece of the densified BZT20 nanoceramics sintered in several runs before and after the reoxidation step. The density of the produced nanoceramics was calculated using the following equation: ρ=m/V, where m is the mass of the densified nanoceramic and V is the volume of that piece. Then, the calculated densities were divided by the theoretical density of BZT20 in order to calculate the relative densities. The calculated values are listed in table3.6. Fig. 3.19 shows SEM cross-sectional micrographs of the BZT20 nanoceramics densified in several runs listed in Table 3.6. Nanoceramics with ~1-2 µm grain sizes were produced after sintering through run 1 and 2. Hence, heating to 1200°C and 1100°C during the sintering leads to coarsening of the grains. In contrast, using sintering temperature of 1000°C in run 3, this coarsening is limited, and nanoceramics with comparable grain sizes to the initial nanoparticles (~200-300 nm) are produced. Clearly, at temperatures above 1000°C, the high thermal energy that is generated results in more diffusion of the neighboring BZT particles, leads to the movement of grain boundaries and, consequently, increases the grain size. While at 1000°C temperature, this diffusion only produces "neck" like paths between the particles, and no grain growth occurs during the sintering. We find that the nanoceramics' lattice constants are comparable to that of the initial powders' unit cell, except for the nanoceramics sintered in run 2, which has slightly higher lattice constants. Thus, independent of grain growth during densification, the material's structure is minimally affected by the process. Consistent with the observed grain sizes in SEM, the Bragg peaks' width is also decreased when the sintering temperature increased. The FWHMs of the (110) Bragg peaks are plotted for all the produced nanoceramics (Fig. 20.c). As we already know, the broadening of a crystalline peak indicates a decrease in crystallite size. Hence, by decreasing sintering temperature, the grain size and the crystallite sizes are reduced. In conclusion, SPS using sintering temperatures of 1000°C densified BZT20 nanoparticles to produce nanoceramics with ~88% relative density compared to the theoretical density of BaZr0.2Ti0.8O3, and comparable grain sizes with respect to the initial nanoparticles. According to the lattice parameters deduced from XRD, the crystalline structure of the sintered nanoceramics is also identical to the nanoparticles. Therefore, 1000°C temperature was used for the sintering of BZT nanoparticles with 12 and 30% Zr content in the B-site of structure to produce nanoceramics (Table 3.7). After the structural characterization of the sintered nanoceramics compared with the original nanoparticles, we went on to investigate the dielectric properties of these nanoceramics to identify cubic and non-cubic crystalline phases in the nanoparticles. Indeed, non-cubic symmetries can evolve into cubic phases at the Curie temperature. Based on the Curie-Weiss law, the dielectric constant is maximum at this phase transition temperature. In fig. As expected, BZT12 nanoceramics show a broad polar-to-apolar phase transition around 70-80°C which is close to the reported phase transition of this material in literature [3,10,25].

This polar-to-apolar phase transition is highly dependent on the frequency, which is an essential characteristic of relaxor ferroelectric materials. In BZT20, the relative permittivity and the polar-to-apolar phase transition temperature are reduced when the Zr content is increased. These reductions result from the modification in the crystalline structure (indicated by XRD) induced by the Zr +4 ions in the lattice. This broader phase transition also shows relaxor ferroelectric behavior, as reported in literature [26][27][28]. Increasing the Zr content to 30% further, reduces the relative permittivity, and no phase transition is observed in the range of -40°C to 120°C. This constant and lower relative permittivity induced by the addition of 30 mol% Zr in the structure indicates a very unstable relaxor phase that evolves into a cubic phase below this temperature range.

In conclusion, Zr +4 cations, due to their bigger ionic radius as a defective element, can distort the tetragonal phase of barium titanate when added to the structure. This distortion leads to an orthorhombic/rhombohedral phase which is less stable compared to the tetragonal one.

This defective element can break up the microscopic domains in the system and change the ferroelectric to a relaxor behavior, deduced from the very diffuse phase transition at lower temperatures than the phase transition of barium titanate.

Surface modification of nanoparticles

Surface modification with different functional groups via sol-gel

To stabilize and prevent the nanoparticles from aggregation in the polymer matrix before and during solidification, the surfaces of the BT nanoparticles, synthesized via a polyol method, were functionalized using several functional groups. Such functionalization can modify the surface free energy of the particles' surface, which consequently affects their wettability. The wettability of the nanoparticles is improved if this surface energy is close to the surface energy of the matrix. Improving the wettability prevents nanoparticles from aggregation and assists uniform distribution into the matrix. Such distribution not only optimizes the loading of nanoparticles but also increases the particle-polymer interfacial area.

To find suitable functional groups, three different organic groups with different hydrophilicity were attached to the nanoparticles' surface. These functional groups with their coupling agents are listed in table 3.8. Perfluorododecyl-1H,1H,2H,2H-triethoxysilane (PFO) as a heavily fluorinated substance leads to a hydrophobic surface on the particles while -OH as a hydrophilic group increases the hydrophilicity of the particles' surface. (3 3 3-trifluoropropyl) triethoxysilane (TFPTES) with terminal -CF3 groups also leads to hydrophilicity between the -OH and highly fluorinated groups on the particle's surface. Materials with various hydrophilicities can, thus, result in various surface energies, which change the nanoparticlepolymer interactions in a nanocomposite. To measure the amounts of the attached functional groups on the particles' surface, TGA was performed on the non-modified and modified nanoparticles. For this, particles were heated to 600°C (10°C/min), so that all the organic groups are degraded and removed from the particles. During the degradation process, mass losses provides the amounts of attached organic functional groups on the particles' surface. In fig. 3.25, the TGA of the modified nanoparticles is compared with that of non-modified BT nanoparticles. Since non-modified BT nanoparticles are not degradable in this temperature range, no degradation is observed during the heating. In contrast, for modified particles, as expected from FTIR spectroscopy, a small mass is observed during heating (mostly at ~350°C). This mass loss is ~1% for OHmodified particles, while ~5 and 8% losses are detected for PFO and TFPTES-modified particles, respectively. Such low a mass loss for OH-modified particles is likely due to the relatively small mass of -OH functional groups around the nanoparticles. This mass loss is increases for the other modified particles due to their larger molecular size compared to -OH groups. However, despite the bulkier molecular structure of PFO compared to TFPTES, the mass loss is almost 3% less. It seems that the high steric hindrance of PFO is detrimental to the yield of surface modification; thus, fewer functional groups are attached to the particles' surface.

Fig. 3.25. TGA thermograms of surface-modified BT nanoparticles heated to 600°C.

Impact of functional groups on surface energy of nanoparticles

After surface characterization of the modified nanoparticles, the influence of these various functional groups on polymer-nanoparticles interactions was evaluated. Typically, the dispersion of nanoparticles into polymers due to different densities and surface energies of the components are thermodynamically unstable. Modifying the surface energy of dispersed components (here: BT nanoparticles) to achieve a surface energy close to the surface energy of the matrix (here: VDF-based polymers) increases nanoparticle wettability leading to improved their physical stabilization in the matrix. The different surface energies of the components can destabilize the nanoparticles in the matrix and keep the produced air bubbles during sonication around the particles, which consequently decreases the dielectric strength of the system. Thus, a functional group with a surface energy close to the VDF-based terpolymer used here ( i.e.,CTFE4 and CFE7) is required. To calculate the surface energy of the modified nanoparticles, silicon wafers were functionalized with similar functional groups. For the -OH groups, an oxygen plasma was employed to introduce -OH groups onto the wafers' surface. Other functional groups were attached via a sol-gel method using silane coupling agents similar to the ones used in the nanoparticle surface modification procedure (without stirring).

According to the young's equation (eq.3.1), the surface free energy of a solid (γS) can be calculated by knowing the surface tension of a liquid (γL), the contact angle (θ) and the interfacial tension (γSL) between the liquid and the solid [29]:

γS= γSL+ γL.cos θ (eq.3.1)
To calculate the surface free energy (γS), the interfacial tension (γSL) must be determined first.

Based on the Fowkes' method, the interfacial tension can be calculated from the solid and liquid surface tensions (γS and γL) plus their disperse (γ d ) and polar (γ P ) parts (eq.3.

2) [30]:

γSL= γS+ γL-2((γs d . γL d ) 1/2 +(γs P . γL P ) 1/2 ) (eq.3.2)

At least two liquids with known polar and disperse parts are required to calculate the surface tension of the polar and disperse parts of the solid. The surface free energy is the sum of these two surface tensions. For this purpose, three liquids with known surface tensions were employed to estimate the surface energy of the modified wafers. Water with a high polar part, diiodomethane with a low polar part, and diethylene glycol are the liquids used for this work. Fig. 3.26 shows the droplets of these liquids on the surface of the modified wafers. The surface free energy of the modified wafers and polymer matrices were calculated with eq.3.2 and listed in tables 3.9 and 3.10. Table 3.9. The droplets' contact angle on the modified Si-wafers and calculated surface free energy of the modified Si-wafers with their polar and disperse parts.

Water Diethylene glycol Diiodomethane

Si

Impact of functional groups on particle dispersion and stability in suspension

To evaluate the influence of these functional groups on the physical stability of the nanoparticles when dispersed in VDF-based terpolymer solutions, destabilization of the modified and non-modified nanoparticles was investigated during 12 hrs after dispersion. nanoparticles precipitate after a while, similar to non-modified nanoparticles. TFPTES and PFO modified particles, due to their more comparable surface energies with the polymer, provide very stable dispersions even after 12 hrs. Comparing these two functional groups, TFPTES, due to its closer surface energy to the terpolymer ( ~32 mJ/m 2 ) shows better stability compared to the dispersed PFO-modified nanoparticles.

Impact of functional groups on dielectric properties of nanocomposites

Despite all the advantages that inorganic nanoparticles may add to ferroelectric polymers, charge accumulation around the nanoparticles can notably decrease the dielectric strength of the nanocomposites. This charge accumulation occurs due to a considerable difference between the dielectric constants of VDF-based terpolymers (~60) and BT nanoparticles (~7000). This charge accumulation depends on the morphology, size, and interfacial particlematrix interactions. Surface modification of the nanoparticles, by improving the nanoparticles' wettability, enhances the polymer-nanoparticle interactions and decreases the charge accumulation probability around the nanoparticles. In fig. 3.29, the breakdown strength of nanocomposites made with modified nanoparticles is shown. These breakdown strengths were calculated using the Weibull probability method (explained in the annex) from several devices. 15vol.% particle content with respect to the polymer was used for this study.

As expected, the neat polymer has the highest breakdown strength (i.e., 335 MV/m), while nanocomposites with non-modified nanoparticles lead to the lowest breakdown strength (149 MV/m) in the device. Consistent with the obtained surface energies, TFPTES-modified BT nanoparticles with a strength of 244 MV/m have the least detrimental impact on the breakdown strength of the nanocomposites for the same particle content. BT as a ferroelectric material, when dispersed into VDF-based polymer, can enhance the lowand high-field dielectric properties of such systems. However, the question is whether that BT is the best choice to enhance these properties?

BT has a ferroelectric to paraelectric phase transition around 120°C where its relative permittivity and, consequently, polarization is maximum. If this polar-to-apolar phase transition is shifted to lower temperatures close to the matrix's (the VDF-based polymer) polar-to-apolar phase transition, both the inorganic and organic components of the nanocomposites display their maximum permittivity and polarization in the same regime.

Accordingly, the polarizability of the nanocomposites will be increased compared to the values found for BT nanocomposites. As discussed in chapter 2, CFE7, solidified at 90°C, has a relaxor ferroelectric to paraelectric phase transition in the vicinity of room temperature (20°C), while BZT20 has a polar-to-apolar phase transition at similar temperatures. Thus, their nanocomposites are expected to display an improved performance in terms of polarizability compared to nanocomposites composed of BT and CFE7. Fig. 3.31 shows the relative permittivity of the nanocomposites composed of CFE7 and BT and BZT20 nanoparticles as a function of temperature upon heating with a rate of 10°C/min. As expected, the relative permittivity of BZT20 nanocomposites is slightly higher than the one for BT nanocomposites, which is likely due to the very similar phase transition regime of BZT20 and CFE7. This high relative permittivity leads to and increased polarizability of the system. Fig. 3.32 displays the polarization of these nanocomposites measured at 800 kV/cm and 100

Hz at room temperature. Consistent with the relative permittivity data, the polarization of BZT20 nanocomposites is higher than the one found for BT nanocomposites, which supports this view that BZT20 with a polar-to-apolar phase transition close to the polar-to-apolar phase transition of CFE7 can more strongly enhance the polarizability of the nanocomposite.

. Solidification behavior

In this section, the influence of addition of BaZrxTi(1-x)O3 nanoparticles with x=0.12 as a relaxor ferroelectric system with polar-to apolar phase transitions around 70°C on the solid-state structure formation and, consequently, dielectric and ferroelectric properties of P(VDF-TrFE-CTFE)(61/35/4) (CTFE4) as a relaxor ferroelectric polymer with a polar-to-apolar phase transition around 60°C was investigated using different nanoparticle content.

From the knowledge gained in chapter 2, CTFE4, due to its chemical structure, can be solidified in solid-state structures composed of defective all-trans arrangements (having tg conformations as a defect). This solid-state structure evolves into a non-polar α phase having only tg arrangements at 50-60°C. Due to the ferroelectric properties of such films, the defective phase below the polar-to-apolar phase transition is called DFE, which evolves into the PE phase at the polar-to-apolar phase transition temperatures.

BZT nanoparticles 100-200nm in size were blended with CTFE4 to manipulate the latter's solid-state structure and to decrease its overall crystallinity. Since the dielectric and ferroelectric properties of such terpolymers originate from the crystalline phase and their All the melting enthalpies were deduced from the DSC heating thermograms of CTFE4 films and nanocomposites to visualize the changes in the overall crystallinities. These values are plotted in fig. 3.33.b. According to the enthalpies, the overall crystallinities linearly decrease when the BZT particle content is increased. For 15vol.% nanocomposites with respect to the polymer, the overall crystallinity of CTFE4 is decreased. In addition, the melting points are shifted to lower temperatures (by ~5°C). This shift means that the thickness of the polymer's crystalline lamellae is decreased. Likely, the nanoparticles, are introduced into in the amorphous phase of the polymer, which hinders its growth in the z-direction and results in the formation of crystallites with thinner lamellae.

According to the polar-to-apolar phase transition of the samples, introducing the nanoparticles, independent of the particle content, slightly shifts the phase transitions to lower temperatures (from 55°C in the CTFE4 films to 50°C in the nanocomposites). Therefore, This dielectric enhancement is also apparent at high electric fields. Fig. 3.35 displays the polarization of BZT12 nanocomposites with 5 and 15 vol.% particle content at 60 MV/m. At this electric field, the maximum polarization is increased when the BZT12 particle content is increased while the remnant polarization and coercive field are reduced. This reduction is due to the stabilization of more tg conformations which is inferred from the shift in phase transition toward lower temperatures. In conclusion, use of BZT12 particles with comparable polar-to-apolar phase transitions with CTFE4 films enhances the dielectric properties of the system, whether at low and high electric fields. In addition to this enhancement, due to the high dielectric constant of BZT12 (especially at its polar-to-apolar phase transitions), more tg conformations are stabilized in CTFE4. This stabilization slightly shifts the polar-to-apolar phase transitions toward lower temperatures.

Decreasing the lamella thickness by the presence of the nanoparticles could be the reason for this observation.

Conclusions

In this chapter, crystalline barium titanate (BaTiO3) nanoparticles as a ferroelectric system with tetragonal to cubic phase transitions at 120°C was synthesized in different sizes (from BT. Moreover, annealing the produced particles at 800°C temperature eliminates the barium carbonate (BaCO3), formed as a side product even under an argon environment. Pure and well crystalline non-cubic BT was achieved using both techniques. Non-cubic to cubic phase transitions at 115°C were observed in DSC heating thermograms for synthesized BT using both methods, which is very close to tetragonal to cubic phase transitions in the ferroelectric BT.

100nm
To increase the compatibility between the nano-sized BT with VDF-based terpolymers such as P(VDF-TrFE-CFE) and P(VDF-TrFE-CTFE) used as a matrix in the nanocomposites, the surface energy of BT nanoparticles was modified by using different functional groups. For this, three different functional groups with various hydrophilicity were tested. It was observed that the -OH groups as highly hydrophilic groups increased the surface energy of the nanoparticles, while -CF3 and -PFO groups decreased the surface energies. This increase in the surface energy that arises from the polar part makes a considerable difference between the surface energy of particle and matrix, which is detrimental for the compatibility of these two components in nanocomposite systems. In contrast, PFO, a very hydrophobic group, notably decreased the surface energy of the nanoparticles, which in turn decreased the compatibility of the components. -CF3 functional groups with surface energy close to CFE7 and CTFE4 terpolymers showed the best compatibility with such terpolymers. This compatibility increased the particles' physical stabilization in nanocomposite suspensions and enhanced the dielectric strength of nanocomposite films.

In the next part, nano-sized barium zirconium titanate (BaZrxTi(1-x)O3) was synthesized with three compositions (x=0.12, 0.2, 0.3) via a polyol method using parameters that were identified from the synthesis of 100-200nm-sized BT particles. The composition and structure of the obtained particles were characterized using EDX and XRD analyses. It was observed that Zr +4 cations as larger ions than Ti +4 in the B-site of ABO3 perovskite structures not only increased the unit cell size of the crystallites but also changed the nature of the non-cubic (likely tetragonal) structure of BT into a less stable, non-cubic (likely orthorhombic/rhombohedral) structure. This phase modification changed BT's ferroelectric behavior into a relaxor ferroelectric one, which showed broader polar-to-apolar phase transitions at lower temperatures compared to the tetragonal to cubic phase transitions in BT. The dielectric properties of the synthesized BZTs confirmed this phase modification.

Finally, BZT/CTFE4 nanocomposites with different particle contents were prepared using BZT with 12 mol% Zr with polar-to-apolar phase transitions around 70°C and CTFE4 with polar-toapolar phase transitions around 50°C. BZT12 nanoparticles, hindered the crystallization of the polymer and decreased the thickness of the latter's lamellae. This decrease in the overall crystallinity was estimated from the deduced melting enthalpies measured in DSC. Moreover, the BZT nanoparticles influenced the solid-state structure formation of the CTFE4 stabilizing

Motivation and context

Nowadays, fossil fuels are the most widely used energy sources in the world, however due to their limited renewability and, as importantly, their hazardous influence on the environment, renewable and environmentally friendly alternatives have attracted great attention.

However, the low efficiency as well as the intermittence of many these platforms limits their industrial usage. Designing an efficient and low-cost capacitor with the ability of storing energy could help to overcome these limitations. Electrolytic systems, polymers, ceramics and other materials (e.g. mica) with a high dielectric constant, high breakdown strength and low dielectric losses are well-known classes of materials which are used for designing such capacitors. Among these capacitors, polymer-based devices due to their high voltage endurance, good mechanical robustness and their ability to be shaped in any configuration, are highly promising [1,2].

As mentioned in chapter 1, the energy density (Ue) of dielectric materials corresponds to the integral Ue=∫ 𝐸𝑑𝐷 (shaded area in fig. 4.1.a), where E is the applied electric field and D is the electric displacement or charge density [2]. Hence, in addition to good endurance of the material at high electric fields, a high D is also required to achieve a high energy density material. In linear dielectric materials, the energy density is proportional to the square of E:

Ue=1/2Kԑ0E 2 , where K represents the dielectric constant and ԑ0 is the vacuum permittivity (8.85*10 -12 F/m). Thus, the charged and discharged energies are linear. In ferroelectric and relaxor ferroelectric materials, due to changes in molecular conformation between the nonpolar and polar structure, K is not constant and changes with E. Hence, the discharged energy is not linear. A high K in low fields results in the polymer to adopt an early D-saturation while a low K at high fields leads to a delayed D-saturation. Late or early saturation of D is also a key factor to achieve a high energy density (Fig. 4.1.b) [3]. Despite the high dielectric strength of polymer films, their low dielectric constant and low D seriously limit their usage. For instance, biaxially oriented polypropylenes (BOPP) based capacitors, which are the main commercial polymer-based capacitors, despite of their high dielectric strength (~640 MV/m), display a very low dielectric constant (~2.2). This low dielectric constant consequently decreases the energy density of the capacitor to just ~1-2 J/cm 3 [4,5]. Therefore, in order to achieve a higher energy storage performance with these capacitors, size and numbers of applied capacitors should be increased which is to the contrary to the goal to decrease size and weight of electronic devices. Hence, it is essential to increase the dielectric constant while maintaining a high dielectric strength.

Recently, PVDF-based ferroelectric polymers, due to their relatively high dielectric constant (>50) [6], and charge density as well as high breakdown strength (300-400 MV/m), have been introduced as an excellent choice for polymer-based capacitors [6]. However, due to the ferroelectric behavior of VDF-based polymers, part of the stored energy is released in form of heat. This heat loss is generated through the resistance of polarized domains to reorient.

Therefore, in contrast to linear dielectric materials, as shown in fig. 4.2, in ferroelectric materials, the number of recovered charges is always less than the values found for the charged state. The efficiency of the discharging process is defined as ; η=

𝑈 (𝑟𝑒𝑙𝑒𝑎𝑠𝑒𝑑) 𝑈 (𝑟𝑒𝑙𝑒𝑎𝑠𝑒𝑑)+𝑈 (𝑙𝑜𝑠𝑡)
, As a consequence, relaxor ferroelectric properties can be used to produce capacitors with higher efficiencies [7]. As studied in chapter 2, the solid-state structure, which determines the dielectric and ferroelectric properties of VDF-based polymers, can be controlled via their chemical structure as well as their solidification. Indeed, the solidification process possesses a remarkable influence on the thermodynamic stability of the formed solid-state structure at a given temperature. Moreover, the solidification at temperatures higher than the phase transition temperature stabilizes in a more pronounced fashion tg conformations while casting at lower temperatures leads to polymer chain arrangements in all-trans conformations. Due to presence of large termonomer units which contain bulky chlorine atom, the formed structures are not stable upon a heating and can temporarily or permanently (depending on the temperature) evolve into more defective structures by stabilizing tg conformations. For instance, the most stable solid-state structure of P(VDF-TrFE-CFE)(64/29/7) (CFE7) terpolymer when casting at 90°C contains a high tg content . The tg conformations increase the interchain distances and reduce the barrier for chain rotation. Consequently more relaxor ferroelectric properties are induced rather than a ferroelectric behavior. This relaxor ferroelectric behavior, which can lead to high polarization but with a narrower hysteresis loop than what is found for ferroelectrics, is required for a material of choice for energy storage applications.

Therefore, decreasing the all-trans conformation content and narrowing the area of the hysteresis loop while keeping a high polarization (electric displacement), is a promising way to achieve high energy density capacitors with high efficiency.

In the first part of this chapter, the influence of solid-state structure on the energy density and efficiency of 3 different VDF-based polymers is investigated. Achieving reproducibly a high energy density with high efficiency is the main goal. For this, relevant energy-structure correlations are investigated on polymers with a most pronounced relaxor ferroelectric behavior (i.e., the material with the most tg conformations). We will discuss that stable solidstate structures with a high fraction of tg conformations are achieved by choosing suitable solidification parameters using the understanding gained in chapter 2.

Materials and methods

Metal-insulator-metal capacitor devices were fabricated in which the different VDF-based polymers were cast as the insulator. Gold of 100nm thickness was then thermally evaporated on glass substrates. The glass substrates were previously cleaned with water, isopropanol and acetone, respectively, in three steps, 15min each, via sonication. Solutions (5vol.%) were blade cast on the bottom electrodes at 105°C and dried for 3hrs in order to cast 5µm-thick films. Gold top electrodes were thermally evaporated through a home-made shadow mask with circles of 6 mm 2 surface area. Hence, several capacitors with 6mm 2 active area were fabricated on one glass substrate (Fig. 4.3). Some devices were prepared with 2-layered active films. For this, solutions (5vol.%) were blade cast using the same protocol as above but adjusting it to obtain 2.5µm-thick films.

Thereby, the film thickness is controlled by the blade height. After 3hrs drying of the first layer, the second layer was cast on top of the first layer. Finally, by evaporating the top electrode, similar capacitors as the one-layer devices were fabricated. Comparing CFE7 and CTFE8, we find that although CTFE8 possesses more tttg + tttg - conformations than CFE7, and CFE7 has more tg + tg -arrangements that are less polar than tttg ones, the polarization of CFE7 is higher than CTFE8. This difference in polarization, thus, likely originates from differences in the overall crystallinity of these two terpolymers, since CFE7 is clearly more crystalline than CTFE8. Although the maximum polarization is notably different in the three polymers, the discharged energy density values are similar (as alluded to above). Therefore, in addition to the maximum polarization of the materials which has a remarkable effect on energy density, the saturation of the polarization is also effective in influencing the charged and discharged energy densities of these materials. Indeed, the saturation strongly depends on the polarizability of the chains in a given polymer, which is dictated by its solid-state structure. possesses a more polar structure, the higher Pmax in CFE7 seems to originate from a higher overall crystallinity. conformations. This enhances the ferroelectric properties of CFE7. On the other hand, PMMA has a detrimental effect on the crystallinity of CFE7, decreasing the overall crystallinity of the blend. Since the dielectric and ferroelectric properties of the blends predominantly arise from the crystalline regions, these properties are suppressed when increasing the PMMA content.

Therefore, at first glance, CFE7:PMMA blends, due to their more ferroelectric behavior rather than a relaxor ferroelectric one, leading to a larger hysteresis loop and causing a decrease in efficiency, and polarization due to the decrease in overall crystallinity (resulting in a decrease in energy density), would not seem to be a very suitable material for energy storage applications. However, the discharged energy density of the blends is similar to neat CFE7 and their efficiencies are higher than that of neat CFE7 at low electric fields.

The observed effect cannot be attributed to the dielectric constant. Fig. 4.8 shows the dielectric constant of the blends and neat CFE7 at various frequencies at room temperature.

Increasing the PMMA content decreases the dielectric constant, especially at low frequencies.

This likely is due to the decrease in overall crystallinities when introducing PMMA. As a consequence of this dielectric constant depression, the maximum polarization of the blends is decreased with increasing PMMA content. Fig. 4.9 displays the measured maximum polarizations as a function of electric field for the blends and neat CFE7. is also important to note that PMMA as a polar amorphous material not only decreases the maximum polarization but can also delay the polarization saturation of the system resulting in the polarization hysteresis loop to become more linear with electric field. This more linear polarization behavior increases the energy density compensating, in turn, the decrease of energy density arising from the maximum polarization depression when blending CFE7 with PMMA. This more linear hysteresis loop not only increases the energy density of the system but also makes the hysteresis loop narrower which results in a decrease in heat loss and an increase in device efficiency. This would explain why the efficiency of the blends was higher than neat CFE7 at lower electric fields.

In addition to limiting the saturation in polarization, PMMA as an amorphous polymer increases the amorphous phase content in blends, rendering the system to display a higher breakdown strength than that of the more crystalline, neat CFE7. Since the electric field is another effective parameter determining the energy density that a given system can reach, increasing the breakdown strength will achieve higher energy densities in a material. The breakdown strength of the CFE7:PMMA blends with various PMMA contents were, thus, analyzed using 6 devices. The breakdown strengths were calculated using a two-parameter Weibull distribution analysis [8]. According to β and λ obtained for the neat CFE7 devices, we deduce a breakdown strength of 393 MV/m while the breakdown distribution is large (β=2.64). These calculated breakdown strengths and shape factors are worse for blends comprising 1.5 and 3 wt.% PMMA (Eb=330 MV/m and β=2.4 for 1.5wt.% PMMA, and Eb=313 MV/m and β=2.57 for 3wt.% PMMA), while more reliable and higher breakdown strengths were found for blends with higher PMMA content (6 and 12wt.%; Eb=472 MV/m and β=3.5 for 6wt.% PMMA and Eb=515 MV/m and β=4.31 for 12wt.% PMMA). Since blends with a low PMMA content possess an overall crystallinity that is comparable to the one of neat CFE7 (which displays a more pronounced ferroelectric behavior), we infer from our data that addition of PMMA not only improves breakdown strength but it also can act as structural defect, which decreases the breakdown strength. The latter is especially true for blends of a low PMMA content, while blends with 6 and 12wt.% PMMA, as expected, feature a more reliable and higher breakdown strength, arising from a lower crystallinity of the samples.

Breakdown strength and energy density of the blends at their Eb are plotted in fig. 4.12 as a function of PMMA content. Considering the observed similar energy densities of blends vs. neat CFE7 at various electric fields (Fig. 4.7), we infer that higher breakdown strengths dictate higher energy densities in blends with 6 and 12 wt.% PMMA over other samples. Therefore, the breakdown strength is a key advantage of blends over neat CFE7. Finally, note that In fig. 4.12.b, the calculated efficiencies at specific Eb are plotted as a function of PMMA content.

Independent of PMMA content, the efficiency is constant for all samples (except for blends with 1.5wt.% PMMA, which features a more crystalline ferroelectric behavior). Although the breakdown strength and, consequently, the energy density is increased with PMMA content, the obtained values are not remarkably higher than for neat CFE7. Clearly, not only the solid-state structure of a polymer system can determine the dielectric strength but also the morphology of the polymer film (thickness, size, shape, number of layers and etc.) affects dielectric strength. In order to explore pathways to further improve the breakdown strength, two-layered films with the same thickness as used for the previous set of one-layer samples (7 µm) were cast using a similar casting temperature (90°C). Fig. 4.13 shows the Weibull distribution probability for breakdowns of two-layered devices at room temperature. Compared to one-layer devices, the shape and scale factor of devices made with two blend layers are remarkably enhanced compared to a single layer while neat CFE7 shows comparable scale factor but higher shape factors. Unambiguously, two-layered devices drastically improve the breakdown strength and, as a consequence, higher energy density values are achieved. In two-layered devices, there is an interface between the two cast layers. This interface can stop any current flow when an electric field is applied, by deviation of possible pathways along these interfaces. By blocking these electron pathways throughout the films, higher electric fields can be applied. Interestingly, this improvement of the breakdown strength is more pronounced in blends than neat CFE7.

Conclusions

Using knowledge gained in chapter 2, in this chapter, the energy storage ability and efficiency of VDF-based polymers was investigated using different terpolymers with different phase morphologies. We found that VDF-based polymers with comparable VDF/TrFE ratios, independent of their termonomer type and content, display similar energy densities when an electric field is applied. However, the energy density efficiency is low for CTFE4 at lower electric fields which is attributed to the more ferroelectric behavior of this polymer at room temperature. In CTFE8 and CFE7, the polar-to-apolar phase transition temperatures are shifted to lower temperatures (around room temperature) and a more defective phase is formed. Accordingly, a narrower hysteresis loop and a higher efficiency is recorded for these polymers. Comparing these two terpolymers, CFE7, due to its high content of tg conformations, has a narrower hysteresis loop at low electric fields while CTFE8 possesses more tttg conformations which results in this polymer to have the largest hysteresis loop.

In CFE7:PMMA blends, PMMA hinders the saturation of the polarization due to its low K at high fields and the increased energy density. Moreover, the decrease of overall crystallinity decreases the maximum polarization that can be achieved which reduces the energy density as well. These two parameters results in the polymer blends to achieve energy densities close to those of neat CFE7, independent of PMMA content. In polymer blends, the maximum polarization decreases (detrimental for energy density) while the D-saturation is delayed Energy density (J/cm 3 )

1 layer 2 layers

Motivation and context

The electrocaloric effect (ECE) is a reversible temperature change (ΔT) of a polar material under the application of an electric field (ΔE), in adiabatic conditions, that shows great potential for solid-state cooling applications. The EC effect depends on the pyroelectric coefficient (p = dP/dT, where P is the polarization). Hence, by enhancing the pyroelectric coefficient, the ECE in ferroelectric materials can be improved [1].

As already discussed in detail in this thesis, ferroelectric materials display a spontaneous polarization at temperatures below the Curie temperature (TC). They feature a polar structure below Tc, while above Tc, they evolve into a non-polar centrosymmetrical arrangement.

Therefore, the pyroelectric coefficient of ferroelectric materials is maximized near Tc as a large change in P can be achieved when changing the temperature.

The intensity of P also depends on the order of the phase transition. The phase transition from ferroelectric to paraelectric (FE to PE) is a first order phase transition, with a high latent heat, while the relaxor ferroelectric to paraelectric (RFE to PE) is a second order transition [2].

More specifically, for the former, highly polar FE structures abruptly evolve into nonpolar PE structures over a narrow range of temperatures. In contrast, the transition of the less polar RFE structures is based on a progressively evolution into non-polar PE structures. That is the reason for the presence of a broader phase transition peak as measured by DSC for relaxor ferroelectric materials. The ECE is more pronounced for ferroelectric materials compared to relaxor ferroelectric materials but this high ECE is only maintained over a very narrow temperature range, while in relaxor ferroelectric materials, although the ECE is lower, it can be maintained over a broad temperature range, which facilitates its use in EC-based cooling applications. Since the ECE is maximum around the phase transition, it is necessary for room temperature cooling, to use materials with such a transition close to this temperature.

Relaxor ferroelectric polymers not only have broader polar-to-apolar phase transitions but also display these at lower temperatures compared to their ferroelectric counterparts. As a consequence, these materials have become prime candidates for practical ECE cooling systems in industry [3].

More specifically, relaxor ferroelectric polymers due to their high breakdown strength, flexibility, good processability as well as having polar-to-apolar phase transitions close to room temperature, were introduced as highly promising materials for cooling devices [2]. As shown in chapter 2, P(VDF-TrFE-CTFE)(61/35/4) (CTFE4) solidifies in a stable DFE phase when cast at the onset of its melting temperature. This DFE phase can evolve into a PE phase at 50°C to 60°C. As this system contains a high fraction of all-trans chain conformations, its macroscopic polarization is higher than in the RFE phase. This results in a higher pyroelectric coefficient and ECE response at around 50°C to 60°C compared to other terpolymers. We, thus, focus in this chapter on this polymer.

As second model system, we use BaZr0.12Ti0.88O3-based nanocomposites. The reason is that such BZT nanoparticles display a defective ferroelectric structure with a polar-to-apolar phase transition around 50°C to 70°C, which is close to the phase transition of CTFE4. In addition, surface modified BZT nanoparticles, due to some interfacial polarization, can lead to a dielectric enhancement in the system. This increase of polarization can, in turn, enhance the pyroelectric coefficient of the system around the phase transition. Since CTFE4 and BZT12 display polar-to-apolar phase transitions at similar temperatures, we explore in this chapter whether, in CTFE4/BZT12 nanocomposites, an enhancement in the pyroelectric coefficient is observed.

Direct measurement of the electrocaloric effect

In this chapter, the ECE of CTFE4 and CTFE4/BZT12 nanocomposites with different particle content is investigated via direct electrocaloric measurements using calorimetry [4]. More specifically, an in-house modified differential scanning calorimeter was employed to obtain the ECE in the chosen materials. The principle of this technique is to measure the heat released and absorbed by a material system while an external electric filed is applied. For this purpose, devices with specific geometries were fabricated with the materials of interest. In order to fabricate these devices, silver was thermally evaporated on a thin polyethylene naphthalate (PEN) substrate through a shadow mask to create circular bottom electrodes of 40 mm 2 surface area. The target polymer:nanocomposite was, then, blade cast on to these bottom electrodes to achieve 15 µm-thick films. As top electrode, silver was thermally evaporated on the cast films using a similar shadow mask as was employed for the fabrication of the bottom electrodes. These assemblies were subsequently placed inside the DSC Due to presence of the PEN substrate and probably air, which have low thermal conductivities, between device and heat flow detector in DSC chamber, we had some heat loss during the measurement. Nonetheless, in order to measure the heat release, an AC electric field was applied to the devices while inside the DSC chamber. Reassuringly, the released heat, which was deduced from the hysteresis loop of polarization vs. electric field was found to be comparable with the one measured by DSC (Fig. This structural change, which notably changes the entropy of the system, leads to a high pyroelectric coefficient and, thus, a high ΔTEC. Above the polar-to-apolar phase transition temperature (60°C), the PE phase which is considerably less polarizable than the FE phase, decreases the polarization, the pyroelectric coefficient and ΔTEC. By increasing the electric field, ΔTEC also increases. Indeed, at 100 and 118 MV/m, ΔTEC is more stable above 60°C (i.e.

above the polar-to-apolar phase transition temperature), which is due to the diffuse nature of phase transition in relaxor ferroelectric materials. This improvement in polarization, which mostly arises from interfacial polarization between polymer matrix and particles, can enhance the pyroelectric coefficient leading to a high ΔTEC.

Since the pyroelectric coefficient is maximum at temperatures around the polar-to-apolar phase transition temperature, the phase transitions in the VDF-based polymer matrix and the inorganic nanoparticles play key roles in determining the ΔTEC of nanocomposites. Choosing components with phase transitions close to each other, thus, not only has a beneficial effect on the pyroelectric coefficient but also can enhance ΔTEC.

As shown in chapter 3, BZT12 features a defective ferroelectric phase with a broad polar-toapolar transition at 70 to 80°C. CTFE4 was selected here as it displays a polar-to-apolar phase transition at 50 to 60°C. Therefore, CTFE4/BZT12 nanocomposites may lead to a materials combination with enhanced ECE. Increasing the particle content further, enhances the ΔTEC of the system. Since the ECE originates from the achievable polarization in a given system as a function of temperature, the polarization of nanocomposites increases with increasing particle loading (Fig. 5.4.b). We suggest that the observed enhancement in polarization arises from an increase in dielectric constant in BZT12/CTFE4 nanocomposites.

Note that despite of the better ECE at 63 MV/m of the nanocomposites with 15 vol.% particle content compared to those comprising 5 vol.% BZT12, due to the decreased breakdown strength with increasing nanoparticles content, applying higher electric fields to this nanocomposite was found quite challenging. 16vol.% (particle content) nanocomposites consistently broke after 70MV/m at the temperatures used. However, nanocomposites with 5 vol.% BZT12, due to their low particle content have a higher dielectric strength (close to neat CTFE4). Fig. 5.5.a shows the dependency of ΔTEC with temperature of 5 vol.% nanocomposites at high electric fields (80 and 100 MV/m). Data obtained for neat CTFE4 are also shown. At 80MV/m, ~13% enhancement in ΔTEC is recorded for the 5 vol.% nanocomposite at 55 and 60°C (from 5.5°C in neat CTFE4 to 6.2°C at nanocomposite system). This enhancement increased to ~25% at 100 MV/m (from 6.4°C at CTFE4 to 8.0°C in the nanocomposite).

Conclusions

In this chapter, the electrocaloric response of P(VDF-TrFE-CTFE)(61/35/4) cast at 90°C (above phase transition temperature) was analyzed at different temperatures (below and above the polar-to-apolar phase transition temperature) at different electric fields (from 60 to 110MV/m) using an in-house modified differential scanning calorimetry (DSC) technique.

Since the polarization of a material depends on its relative permittivity (which is maximum at phase transition temperature), and the electrocaloric effect (ECE) is proportional to the integral of the pyroelectric coefficient (∫ 𝑑𝑃/𝑑𝑇), the polarization of CTFE and, consequently, ΔTEC that can be achieved, is maximum around 55°C (i.e. around the polar-to-apolar phase transition). This maximum polarization and ΔTEC decreased above the DFE-to-PE and RFE-to-PE phase transitions.

In order to manipulate the electrocaloric effect of a polymer, we tuned the polarization of a given system by adding BZT12 nanoparticles with a phase transition around 70°C. Similarly to the neat polymer, the electrocaloric response of such BaZr0.12Ti0.88O3/P(VDF-TrFE-CTFE)(61/35/4) nanocomposites with 5 and 15vol.% particle content, cast at 90°C (above phase transition temperature) were analyzed at different temperatures (below and above their polar-to-apolar phase transition) and at different electric fields (from 60 to 100MV/m) using an in-house modified differential scanning calorimeter technique. Since 15vol.% nanocomposite has low dielectric strength, the electrocaloric effect was recorded only at In this thesis high-and low-field dielectric properties of polymer-based systems, and the relations with these materials' solid-state structure, have been investigated. The aim was to deliver a better understanding of the origin and mechanisms of the ferroelectric (FE) and relaxor ferroelectric (RFE) behavior in these materials, in order to enhance their performance for ferroelectric-related applications such as energy storage and solid-state electrocaloric refrigeration. VDF-based polymers and nanocomposites with perovskite oxide particles were, thereby, used as model systems.

In chapter 2, we focused on organic VDF-based polymers and characterized their solid-state structure and its relation with the dielectric properties of these polymers. The first challenge was to identify relevant parameters which can be exploited to influence the solid-state structure of VDF-based polymers. The second challenge was to understand the role of these parameters on the structure formation of these polymers. Finally, identifying the relation between structure and low and high field dielectric properties was the last objective of this chapter. We revealed that the chemical structure, as well as the solidification process, are critical in terms of engineering the solid-state structure of the VDF-based polymers. In case of the solidification from solution, solution concentration and chain entanglement in solution, solidification temperature and solidification time were found to be important parameters affecting the solidification process and, consequently, the solid-state structure. It was shown that the entanglement density, which limits the chain mobility in solution during solidification, can change the distribution of chain conformations (all-trans and trans-gauche arrangements) in the resulting structures, and in turn lead to different phases responsible for either FE or RFE behavior in the polymer. A higher entanglement density leads to stabilization of tg+tg-conformations and, consequently, the RFE phase. In solutions with a lower entanglement density, due to the higher chain mobility, one phase only is stabilized that we defined as a defective ferroelectric (DFE) phases. This phase contains a mixture of all-trans and tg conformations. It was also observed that the entanglement density can modify the crystallite size distribution in the polymers' solid-state structure, with a broader crystallite size distribution obtained when solidifying the material from concentrated solutions with a high entanglement density.

Since the stability of the chain conformations and crystalline phases in VDF-based polymers is highly dependent on temperature, different phases were found in structures cast with P(VDF-TrFE-CFE)(64/29/7) (i.e. CFE7) at varying casting temperatures. We demonstrated that solidification at temperatures in which all-trans conformations forms, stabilizes the FE phase in the polymer, while casting at the temperatures in which tg conformations are more stable, leads to the RFE phase to be stabilized. This behavior depends on the chemical structure of the polymer as well. The temperature stability of the various polymorphs directly arises from

In chapter 4, the energy storage ability of P(VDF-TrFE-CFE) and P(VDF-TrFE-CTFE) terpolymers, neat and blended with PMMA , was investigated. Using the understanding gained in chapter 2, the neat polymers were cast at high temperatures in order to stabilize the RFE phase. By stabilizing the RFE phase and decreasing the hysteresis loop area, the energy storage efficiency was successfully increased. In PMMA:terpolymer blends, we found that PMMA limits the saturation polarization of the system, resulting in a more linear and less hysteretic P-E response. This compensates for the decrease of overall crystallinity upon addition of PMMA, which decreases the maximum polarization. Thus, comparable energy storage performances are achieved for neat and blend systems at the same applied electric fields. It was also observed that PMMA improves the dielectric strength resulting in higher energy densities due to the ability of the blends to sustain higher applied electric fields. Producingbi-layer structures further enhances the breakdown strength of such polymer-based systems.

This enhancement is attributed to the presence of interfaces between the two layers which can hinder formation of electrical pathways that lead to shorts.

Barium titanate nanoparticles

The commercial 200 nm-sized barium titanate nanoparticles were purchased from abcr co.

(Germany). The particle size was checked via SEM before use.

Coupling agents for surface modification

The synthesized barium titanate and barium zirconium titanate were surface modified via a sol-gel method (discussed in chapter 3). The process involves converting monomers into a colloidal solution (sol) that acts as the precursor for an integrated network (or gel) of either discrete particle. The precursor sol can be either deposited on a substrate to form a film or

Techniques

Rheometry

To measure the polymer solutions' viscosity, an Anton Paar MCR302 rheometer was used with a CP50-1°geometry (cone 50 mm-1°/plate) (Fig. 7.6). Different shear rates and shear stresses were applied to measure the shear viscosity and the viscoelastic behavior of the solutions. I am really indebted to one of my wonderful friends. A person who was also an advisor to me as a post-doc, Florian LE GOUPIL. Not only for his scientific helps but I could count on him for everything since my arrival in Bordeaux. His brilliant, skillful supervision enriched this study higher than my expectation. His unassuming approach to research and science is a source of inspiration. This approach is reflected by his simple but clear writing style, which is something I hope to carry forward throughout my career. Thanks for the full encouragement and enthusiasm that you gave me during the work (especially during writing the thesis). I would not finish it without these motivating guidance.

I am fortunate to have been a part of the organic electronic group in LCPO. Very special thanks to all the teammates for creating a wonderful environment in the lab. I will never forget the fun moments with you including all the parties and pethanque games. Thanks for accepting me in your group. I also have special thanks to my officemates for the lively physics conversations which was fun even in the very serious situations.

I would like to thank ElorPrintec platform for their facilities that made my work possible. I would also like to thank Eric Lebraud for XRD analyses and U-chan Chung seu and Catherine Elissalde for SPS which are both performed in the ICMC lab. Many thanks to Ahmed Bentaleb for WAXS analyses which is performed in the CRPP lab in bordeaux.

  Mots clés: polymères à base de VDF, ferroélectricité, nanocomposite, stockage d'énergie, .......................................................................................................................... i Resumé: ........................................................................................................................

Fig. 1 . 1 .

 11 Fig.1.1. Domain and domain walls in a microscopic view of ferroelectric crystal[4].

Fig. 1 . 3

 13 Fig.1.3 Relative permittivity of barium titanate as a ferroelectric material with different phase transitions [5].

Fig. 1 . 7 .

 17 Fig.1.7. Cubic and non-cubic (tetragonal) ABO3 unit cell of perovskites above and below Curie

Fig. 1 . 8 .

 18 Fig.1.8. Polarization-electric field hysteresis loop of dielectric (linear) and ferroelectric (nonlinear)

Fig. 1 . 9 .

 19 Fig.1.9. Schematic view of an electrocaloric cooling cycle in comparison with conventional vaporcompression cooling cycle (pistons)[38].

Fig. 1 . 10 .

 110 Fig.1.10. Repeat unit of vinylidene fluoride in PVDF homopolymer.

(Fig. 1 .

 1 11)[22]. The chain conformation and arrangement of the fluorine as a high electronegative atom leading to a high dipole moment (3.34 * 10 -30 C*m -1 ) along the backbone, can determine the polarity of the crystal phase.

Fig. 1 .

 1 Fig.1.11. PVDF unit cells of crystals; A) α B) δ C) γ and D) β viewed along the c-axis and crystalline chain conformations; E) tg + tg -, F) tttg + tttg -and G) All-trans. F,C and H are represented by red, cyan and blue, respectively. Image is taken from[31].

Fig. 1 .

 1 Fig.1.12. Uniaxially-drawn PVDF of 1 µm in thickness at 1 Hz and 20°C.

Fig. 1 .

 1 Fig.1.13. Chemical structure of P(VDF-TrFE) copolymer.

Fig. 1 .

 1 Fig.1.14. Chemical structure of P(VDF-TrFE-CTFE) (left) and P(VDF-TrFE-CFE) (right) terpolymers.

Fig. 1 .

 1 fig.1.15)[START_REF] Zhang | Giant electrostriction and relaxor ferroelectric behavior in electron-irradiated poly (vinylidene fluoride-trifluoroethylene) copolymer[END_REF][START_REF] Jeong | Electro-optical response of the ferroelectric relaxor poly (vinylidene fluoride-trifluoroethylene-chlorofluoroethylene) terpolymer[END_REF][START_REF] Bao | Phase transitions and ferroelectric relaxor behavior in P (VDF-TrFE-CFE) terpolymers[END_REF][START_REF] Klein | Influence of Crystallization Conditions on the Microstructure and Electromechanical Properties of Poly (vinylidene fluoride-trifluoroethylenechlorofluoroethylene) Terpolymers[END_REF]. This effect, indeed, depends on the termonomer content. The physically pinned bulky-CFE can show resistant to rotation in the low electric fields but increasing the field can also orient the termonomers and giving rise to a double hysteresis loop behavior (see fig.1.15)[11].

Fig. 1 . 15 .

 115 Fig.1.15. Nanodomain formation in P(VDF-TrFE)-based terpolymers and e-beam-irradiated P(VDF-TrFE). Room-temperature bipolar D-E loops at 10 Hz (triangular wave function) are shown for (a) P(VDF-TrFE-CFE) 59.2/33.6/7.2, (b) P(VDF-TrFE-CTFE) 62.2/30.2/7.6, and (c) e-beamed P(VDF-TrFE) 50/50 (60 Mrad at 70°C)[11].

Fig. 1 .

 1 Fig.1.16. Solidification of P(VDF) into different crystalline phases and evolution of these phases depending on conditions selected [73].

Fig. 1 .

 1 Fig.1.18. Diameters of spherulites of P(VDF) crystals versus temperature in blend with P(MMA) at various temperature[85].

Fig. 1 .

 1 19 shows the degree of crystallinity of the β phase as a function of casting temperature in blends with different PMMA contents.

Fig. 1 . 19 .

 119 Fig.1.19. The degree of crystallinity of β phase as a function of crystallization temperature for blends with different compositions derived from FTIR spectroscopy [86].

Fig. 2 .

 2 Fig.2.1.a schematically illustrates dissolved polymer chains in a dilute regime in a good and poor solvent.

Fig. 2 . 1 .

 21 Fig.2.1. Polymer chains in a dilute, semi dilute and concentrated regimes in thermodynamically good and poor solvents.

Fig. 2 . 2 .Fig. 2 . 3 .

 2223 Fig.2.2. Viscosity of CFE7 solutions of various polymer concentrations as function of shear rate at 20°C.

Fig. 2 . 4 .

 24 Fig.2.4. FTIR spectra of CFE7 films cast at 40°C from 0.5, 5 and 10vol.% solutions. tm>4 and tm>3

Fig. 2 . 5 FTIR

 25 Fig.2.5 FTIR spectra of CFE7 films cast at 40°C from a)0.5 and b)10vol.% solutions, measured during heating from 25°C to 85°C. tm>4 and tg represent characteristic peaks for long-range alltrans and trans-gauche conformations, respectively.

Fig. 2 .

 2 Fig.2.6. a) All-trans peak of CFE7 films cast from 0.5vol.%, b) and 10vol.% solutions and c) tg peak of the same film cast from 0.5vol.% and d) 10vol.% solutions.

Fig. 2 . 7 .

 27 Fig.2.7. The development of all-trans and tg FT-IR relative intensities with temperature.

Fig. 2 . 8 .

 28 Fig.2.8. shows the DSC data obtained on films cast from solutions of three different concentrations below, near and above C * , at 40°C. Independent on casting concentration, the melting point (Tm) is found around 125°C. However, depending on casting parameters, the transition from a polar to a nonpolar α phase, the Curie temperature for ferroelectrics is affected by the solution concentration used to cast the CFE7 films. Below the polar-to-apolar phase transition temperature, according to our FTIR data, CFE7 can contain different fractions of all-trans and tttg conformations, which are β and γ polymorphs. Above the phase transition temperature, these all-trans phases evolve into the α phase, which is formed from tg chain conformations. This polar-to-apolar phase transition depends on the thermodynamic and kinetic stability of the formed phases occurring at a temperature between Tg and Tm, and

Fig. 2 . 8 .

 28 Fig.2.8. DSC heating thermographs of CFE7 films cast at 40°C from 0.5, 5 and 10vol.% solutions. Dashed lines are drawn as guides for the eye.

Fig. 2 . 9 .Table2. 3 .

 293 Fig.2.9. WAXS patterns of CFE7 cast from 10 vol.% (a to c) and 0.5vol.% (d to f) solutions at 40°C at the vicinity of the (200/110) Bragg peak during heating and measured at 25, 55 and 90°C.

Fig. 2 .Fig. 2 . 10 .

 2210 Fig.2.10. Calculated inter-planar distance, phase content and lateral a-b plane extension deduced from the WAXS (200/110) Bragg peak of CFE7 films cast at 40°C from 10vol.% solutions.

  is no hydrodynamic force to generate interconnectivity between the chains during solidification. Therefore, each molecule may crystallize individually, leading to packing defects and increasing the interchain distance. The increased chain mobility assists formation of thicker crystallites at the expense of more gauche defects, while at highly entangled systems, high polymer-polymer intermolecular interaction hinder crystallization. The lower chain mobility limits the distribution of CFE bulky termonomer in the crystallites leading to less distribution of gauche defects. Hence, in the films cast from the concentrated solutions, we obtain various crystallites with different solid-state structures (see fig.2.11.a), while in the films cast from dilute solutions we have more homogeneous crystallites with comparable tg defects (see fig.2.11.b). This interpretation would explain all the observed differences in the polar-apolar phase transition observed in DSC, FTIR spectroscopy and WAXS and the thermal evolution of the structure.

Fig. 2 .

 2 Fig.2.11. Suggested solid-structure formation for CFE7 terpolymers cast from a) concentrated solutions and b) diluted solutions at 40°C.

  Fig.2.12. shows the DSC thermograms of cast films upon heating with 10 °C/min. In order to

Fig. 2 .

 2 Fig.2.12. Heating DSC thermograms obtained with CFE7 film cast at various temperatures using a)0.5 vol % and b)10vol.% solutions.

Fig. 2 .Fig. 2 . 14 .

 2214 Fig.2.13. c).

Fig. 2 . 14 summarizes 90 (

 21490 the extracted parameters such as inter-planar distance (d(200)/d(110)), crystalline phase content and a-b plane coherence deduced from the (200)/(110) Bragg peak measured for 10-CFE7-25 (Fig.2.14.a), 10-CFE7-40 (Fig.2.14.b), 10-CFE7-55 (Fig.2.14.c) and 10-CFE7-Fig.2.14.d) upon heating from 25°C to 90°C.In 10-CFE7-25 and 10-CFE7-40, in agreement with DSC and FTIR spectroscopy, CFE7 crystallizes into FE and DFE arrangements, with a higher FE content in 10-CFE7-40. Regarding the thermal evolution of these structures, at 55°C, DFE and FE is converting to PE while some FE phase remains (crystalline a-b plane coherence and content of FE fraction decrease). The presence of more FE phase in 10-CFE7-40 when compared to 10-CFE7-25 at 55°C is in accord with the observed slight shift of Curie temperature in DSC and the higher all-trans conformation content in FTIR observed in these systems. This evolution continues until all of the FE phase transforms into the PE phase.For 10-CFE7-55, all three arrangements, i.e. DFE, FE and RFE, are formed. Compared to the lower temperature cast samples, the FE fraction is considerably reduced while the RFE phase starts to appear. At 55°C, similar to 10-CFE7-25 and 10-CFE7-40, the less stable RFE and DFE evolve into two PE phases with different but close inter-planar distances while the FE fraction remains unchanged. Finally, for 10-CFE7-90, the RFE phase is the dominant phase. This is also observed in DSC.

Fig. 2 .

 2 Fig.2.15. Schematic illustration of the free energy of P(VDF-TrFE)(75/25) copolymer versus temperature[11].

Fig. 2 .

 2 Fig.2.16. Schematic free energy diagram for CFE7 as function of temperature.

Fig. 2 .

 2 Fig.2.17 displays the effect of different casting temperatures on the free energy diagram of CFE7. Casting at a given temperature increases the region of phases stability of the various phases found at that temperature, which is expressed by a downward shift of the corresponding free energy line.

Fig. 2 .

 2 Fig.2.17. CFE7 free energy diagram for structures cast at 40°C (a), 55°C (b) and 90°C (c). The free energy after casting are shown by dotted line with different colors.

Fig. 2 .

 2 Fig.2.18 displays the corresponding DSC thermograms for CFE7 cast at 90°C and dried at that temperature for 1, 5, 10, 15 min and 3hrs using 0.5 vol.% (Fig.2.18.a) and 10vol.% (Fig.2.18.b)

Fig. 2 .

 2 Fig.2.18. DSC first heating thermograms of CFE7 cast at 90°C from a) 0.5 and b) 10vol.% solutions using different casting times.

Fig. 2 .

 2 Fig.2.20. Measured inter-planar distance, phase content and lateral a-b plane extension as extracted from the WAXS (200/110) Bragg peak of CFE7 cast at 90°C for a)15 min and b)3hrs. Dotted lines are plotted as guide to the eye to separate the parameters at different temperatures.

2

 2

  .21), by casting CFE7 at 90°C, the PE free energy line starts to shift downward, however due to the very short casting time, this stabilization cannot occur, leading only to a slight change in free energy (Fig.2.21.a). However, if the film is kept at room temperature, crystallization may continue at temperature. As a consequence, the final structure of CFE7 films cast at 90°C for 1 min is identical to CFE7 structures cast at 25°C and 40°C. Hence, as is shown in fig.2.21.a, in this phase formation competition, shifting the FE/DFE/PE free energy at room temperature can overcome the slight PE stabilization at 90°C and render the FE/DFE to be the dominant phases in such CFE architectures. Prolonging the casting time at high temperature (90°C), the PE phase stabilization becomes more pronounced (Fig.2.21.b), until after 15 min, only an intense DFE-PE transition peak is recorded in DSC. Casting for 3 hrs at 90°C (Fig.2.21.c) leads to a large shift in the PE free energy line and in turn, to the RFE phase to be stabilized. This observation indicates that the CFE termonomer cannot be accommodated within the crystalline regions at short casting times and requires room temperature aging to induce the FE phase.

Fig. 2 .

 2 Fig.2.21. Free energy diagram of CFE7 films cast at 90°C for a) 1min, b) 15min and c) 3hrs. The free energy reduction through high temperature casting and room temperature aging are shown with red and green dotted lines, respectively.

Fig. 2 .•

 2 Fig.2.22. DSC heating thermograms of CFE7 films cast using different techniques at 55°C and dried at that temperature for 1 day

Fig. 2 .

 2 Fig.2.23. VDF-based polymers used as testbeds for free energy vs. temperature diagrams, given their chemical structures as well as their Curie temperatures.

Fig. 2 .

 2 Fig.2.24. DSC heating thermograms of VDF-based polymer films cast at the temperatures given in table 2.5 (25°C for CFE7, 40°C for CTFE4 and TrFE55 and 80°C for TrFE30).

  phase formation is occurring in TrFE-rich crystals, which is very different from the other polymers studied here. More specifically, since TrFE comprises three fluorine atoms, a less stable, diffuse DFE phase is formed. In contrast, for TrFE30 with just 30 mol% TrFE content, due to its specific VDF/TrFE ratio and more favorable all-trans conformation in this structure, and dense FE phase is observed with a polar-to-apolar phase transition temperature identified in DSC at ~100°C. Reassuringly, all the polymers except TrFE55 (due to its different TrFE content), are crystallized in FE and DFE phases when cast below their transition temperatures.

Fig. 2 .

 2 Fig.2.25 displays the DSC thermograms of these polymers cast at different temperatures using 0.5vol.% and 10vol.% solutions (the molecular weight of the polymers and consequently their C * in cyclopentanone are similar). Two temperatures below and above the recorded polar-toapolar phase transition temperatures were used as casting temperatures (see Fig.2.24).

Fig. 2 .

 2 Fig.2.25. DSC heating thermograms of a) CFE7, b) CTFE4, c) TrFE55 and d) TrFE30 cast at mentioned temperatures from 0.5 and 10 vol.% solutions.

Fig. 2 .

 2 Fig.2.26. DSC heating thermograms of a) CFE7, b) CTFE4, c) TrFE55 and d) TrFE30 films cast at mentioned temperatures using 0.5 and 10 vol.% solutions and casting times of 5 min and 3 hrs.

Fig. 2 .

 2 Fig.2.27. Relative permittivity and dielectric loss (tanδ) of CFE7 films cast using 0.5 and 10 vol.% solutions and casting temperatures of a) 40°C, b) 55°C and c) 90°C measured as function of temperature at 1kHz.

Fig. 2 . 28 .

 228 Fig.2.28. Relative permittivity and dielectric loss (tanδ) of CFE7 films cast from 10 vol.% solutions as function of temperature (first heating at 1kHz).

Fig. 2 . 29 .

 229 Fig.2.29. Polarization and current density of CFE7 films cast at different temperatures measured as function of applied electric field at room temperature at 100Hz frequency.

Fig. 2 .

 2 Fig.2.30. DSC heating thermograms (10°C/min) of CFE7 films cast at a)40°C, b)55°C, c)90°C, as cast and after 1-week storage at ambient conditions.

Fig. 2 .

 2 Fig.2.32. DSC first heating thermograms for CFE7 films cast from a) 0.5 vol.% and b) 10 vol.% solutions at 105°C for 3hrs when aged from 2 hrs to 1 week at room temperature.

Fig. 2 . 33 .

 233 Fig.2.33. Calculated inter-planar distance, lateral extension and phase content extracted from WAXS data taken for CFE7 films cast at 55°C, 90°C and 105°C, as cast and after 1-week aging.

Fig. 2 . 34 .

 234 Fig.2.34. Relative permittivity of CFE7 films cast at 55°C, 90°C and 105°C, as cast and after 1-week aging, measured as function of temperature.

Fig. 2 .

 2 Fig.2.[START_REF] Thacher | Electrocaloric effects in some ferroelectric and antiferroelectric Pb (Zr, Ti) O3 compounds[END_REF] shows the real (ԑ') and imaginary (ԑ") part of the relative permittivity measured for CFE7 films at room temperature as a function of time (at 1 kHz). For this measurement, CFE7 films were heated to their casting temperatures and then cooled back to room temperature using a cooling rate of 10°C/min. The real and imaginary part of permittivity were recorded every 30 minutes for a day.

Fig. 2 . 35 .

 235 Fig.2.35. Relative permittivity (real and imaginary part) of CFE7 films cast at 55°C, 90°C and 105°C measured as function of time after heating.

Fig. 2 . 36 .

 236 Fig.2.36. Relative permittivity (real and imaginary part) of CFE7 films cast at 90°C and 105°C and measured as function of time after heating

1 .

 1 Miscibility and solidification of terpolymer:PMMA blendsPoly (methyl methacrylate) is an amorphous polymer with methyl methacrylate as repeat unit. The methacrylate group in these unit can be positioned along the chains at various positions which categorize PMMA into 3 different isomers: isotactic, syndiotactic and atactic PMMA[15].In an isotactic PMMA, all the methacrylate groups are located on one side of the molecular backbone while in syndiotactic PMMA, methacrylate groups are alternatively positioned along the chains. When the methacrylate groups are randomly oriented along the chains, PMMA is atactic (Fig.2.37). The most common PMMA is polymerized by a free-radical mechanism; they thus are atactic. The specific chain arrangement can affect the physical properties of PMMA such as the glass transition temperature (Tg), solubility in solvents and miscibility with other polymers[15]. Furthermore, in blends, the position of the methacrylate groups can modify the arrangement of neighboring chains and induce a specific chain orientation of the other blend component whether in their amorphous or crystalline state (due to intermolecular interactions between the methacrylate group and specific moieties in the other polymer). For instance, in PMMA:PVDF blends, due to the dipole-dipole interactions between the C=O unit (in PMMA)

Fig. 2 .

 2 Fig.2.37 Repeat unit of PMMA and schematic views of PMMA with different tacticity: isotactic, syndiotactic and atactic. Shaded circles represents methacrylate groups[17].

Fig. 2 .

 2 38.a shows the FTIR spectra of these blend films at room temperature and fig.2.38.b shows just the FTIR vibration corresponding to the C=O bond in PMMA at 1725 cm -1 . Moreover, in order to see the influence of PMMA on the VDF-chain arrangement, fig.2.38.c and d show the FTIR peak corresponding to the C-H bond in VDF-chains with all-trans and tg conformations.

Fig. 2 .

 2 Fig.2.38 a) FTIR spectra of CFE7 blends with different PMMA content as measured at room temperature. B) C=O vibration peak c) C-H vibration peak with all-trans conformations. c) C-H vibration peak with tg conformations.

Fig. 2 .

 2 Fig.2.39.DSC thermograms of CFE7/PMMA blends with various PMMA content, cast at 90°C for 3hrs using solutions with 0.5vol% PMMA+CFE7 total content: upon a) heating and b) cooling employing 10°C/min.

(Fig. 2 .

 2 [START_REF] Valant | Electrocaloric materials for future solid-state refrigeration technologies[END_REF], 10°C/min). The role of PMMA on modifying the chain conformations of the CFE7

Fig. 2 .

 2 Fig.2.40. DSC heating thermograms of CFE7:PMMA blends with various PMMA content cast at 90°C

Fig. 2 . 41 .Fig. 2 .

 2412 Fig.2.41. The calculated overall crystallinity (extracted from WAXS) of CFE7:PMMA blend films with different PMMA content cast at 90°C for 3hrs.

  2.43.b). Moreover, addition of PMMA rearranges

Fig. 2 . 43 .

 243 Fig.2.43. Suggested solid-state structures for a) neat CFE7 and b)CFE7:PMMA systems cast at 90°C.

Fig. 2 .

 2 Fig.2.44. DSC heating thermograms of as cast CFE7/PMMA blends with various PMMA content cast at a) 40°C, b) 55°C and c) 90°C using 0.5 vol.% (dotted line) and 10vol.% (solid line) solutions upon heating with 10°C/min.

Fig. 2 .

 2 Fig.2.45. DSC heating thermograms for CFE7:PMMA blends with various PMMA content cast at a) 55°C and b) 90°C using 0.5 vol.% (dotted line) and 10vol.% (solid line) solutions and measured after 1 week aging at room temperature (10°C/min rate).

Fig. 2 .Fig. 2 . 46 .

 2246 Fig.2.46. Relative permittivity and tanδ of CFE7:PMMA blend films with various PMMA content cast at a) 55°C and b) 90°C and measured during two heating cycles (10°C/min) in the temperature range of -30°C to 55°C for films cast at 55°C and in the range of -30°C to 90°C for films cast at 90°C.

Fig. 2 . 47 .

 247 Fig.2.47. Polarization and current density of CFE7:PMMA blend films with various PMMA content cast at 55°C as a function of electric field a) before and b) after heating to 55°C.

Fig. 2 . 48 .

 248 Fig.2.48. Polarization and current density of CFE7:PMMA blend films with various PMMA content cast at 90°C as a function of electric field a) before and b) after heating to 90°C.

Fig. 3 . 1 .

 31 Fig.3.1. Archetypal ABO3 perovskite structure, featuring a 3D network of B-O6 octahedra and A ions (red circles). B ions are given as black circles, and oxygen ions as blue circles[3].

Fig. 3 .

 3 3 shows the phase stability diagram of barium titanate (Fig.3.3.a) and titanium dioxide (Fig.3.3.b) at 90°C[16].

Fig. 3 .

 3 Fig.3.3. a) Yield diagram of BaTiO3 synthesized using the hydrothermal method from metal precursors at 90°C. b) Phase stability diagram for titanium-H2O systems at 90°C[16].

Scheme 3 . 1 .

 31 Scheme 3.1. Oxidation of Ti precursor in aqueous solutions (Hydrolysis).

3 .

 3 Scheme 3.3. Hydrolysis of Ba and Ti precursors in the presence of NaOH in DEG at 220°C (close to the boiling point of DEG).

Fig. 3 .Fig. 3 . 4 ,

 334 Fig.3.4 shows the morphology of the BT nanoparticles synthesized using a hydrothermal method from the listed batches. From run 1, particles were produced with a diameter of 50 to 100 nm, covered by some residual organic compounds. Produced particles from run 2 and 3 were of 50 to 200 nm diameters, whereby run 3 produced more spherical particles compared to those from run 2. Differently from the others, in run 4, the produced particles were non-homogeneous micron-sized particles, whereby Increasing the reaction time to 24h in run 4, resulted in an increase in the particle size and produced micron-sized particles, which are not desirable for our objectives, while 12 hrs reaction time, independent of the concentration of precursor and hydrolyzing agent, or the annealing, produced 50nm to 200nm-sized particles.

Fig. 3 . 5 .

 35 Fig.3.5. FTIR spectra of the synthesized BT via a hydrothermal method using different parameters listed in table.3.1. and measured at room temperature. The FTIR spectra of the commercial BT are given for comparison.

Fig. 3 . 6 .

 36 Fig.3.6. TGA thermograms of annealed and non-annealed synthesized BT via hydrothermal method measured from room to 900°C.

Fig. 3 .

 3 7 displays the XRD spectra of the synthesized BT nanoparticles in comparison with commercial tetragonal BT nanoparticles. All samples display a highly crystalline structure without any amorphous phase. The commercial BT exhibits a perovskite oxide structure with the labeled (hkls) in the corresponding X-Ray patterns. (hkl) or miller indices define the orientation of a crystal plane by considering how the plane intersects the main crystallographic axis. For instance, (100) shows a plane that intersects the x-axis in parallel with the y and z-axis. Dimension of the lattice were accordingly estimated via miller indices.

Fig. 3 .

 3 Fig.3.7 a) XRD patterns of the synthesized BT via a hydrothermal method using several batches measured at room temperature. b) Comparison of the (200) hkl peak of the synthesized BT using 12 hrs and 24 hrs reaction times. The XRD pattern of the commercial nano-sized BT is also shown.

Fig. 3 . 8 .

 38 Fig.3.8. DSC thermograms of the BT nanoparticles synthesized via a hydrothermal method produced from runs 3 and 4. The DSC thermogram of the commercial BT is used for comparison.

Fig. 3 .

 3 Fig.3.9 SEM micrographs of BT nanoparticles synthesized via a polyol method from different runs listed in Table 3.2.

Fig. 3 .

 3 Fig.3.10 shows the FTIR spectra of the synthesized particles from run 1 (Fig.3.10.a), run 2 and 3 (Fig.3.10.b). The synthesized nanoparticles from run 1 show a complex pattern which means

Fig. 3 .

 3 Fig.3.10. FTIR spectra of the synthesized BT particles via a polyol method a) before annealing and b)

Fig. 3 .

 3 11 shows the weight loss of such samples upon heating with 10°C/min. By heating the non-annealed nanoparticles, almost 30 wt.% of organic residuals are degraded in a temperature range between room temperature to 550°C. Accordingly, 650°C annealing temperature was chosen

Fig. 3 .

 3 Fig.3.11. TGA thermograms of the synthesized BT via a polyol method before and after annealing at 650°C and 800°C.

Fig. 3 .

 3 Fig.3.12 shows the XRD patterns of the nanoparticles produced from the different runs measured at room temperature. Consistent with FTIR spectroscopy and TGA, small amounts of barium carbonate are observed in the XRD patterns of the samples produced from run 2 (with 650°C annealing temperature), while the particles that were annealed at 800°C, possess a pure and highly crystalline perovskite structure of BT. Similar to the produced nanoparticles

Fig. 3 .

 3 Fig.3.12. XRD patterns of the BT nanoparticles synthesized via a polyol method from different runs measured at room temperatures

  Similar to the synthesis of BT via a polyol method discussed in the previous part, in order to synthesize BZT via this method, barium acetate (Ba(C2H3O2)2) and titanium isopropoxide (Ti(OCH(CH3)2)4) were used as barium and titanium precursors, respectively, along with zirconium isopropoxide (Zr(OCH(CH3)2)4) as zirconium precursor. According to the reported phases for BZT with various Zr content in literature (which is shown in fig 3.2 in the introduction part of the chapter), BZT with 12 mol%, 20 mol%, and 30mol% Zr content in Bsite was targeted in this work. Depending on the Zr content in the target compounds, stoichiometric amounts of zirconium isopropoxide were used ([Zr]/[Ti]+[Zr]=0.12, 0.2, and 0.

Fig. 3 .

 3 Fig.3.13 shows SEM micrographs of the synthesized particles having different Zr content.

Fig. 3 .Fig. 3 .Fig. 3 .

 333 Fig.3.13. SEM micrographs of synthesized BZT nanoparticles with various Zr contents.

Fig. 3 .

 3 Fig.3.15. EDX pattern of synthesized BZT20 measured at room temperature using 10 eV energy.

Fig. 3 .

 3 Fig.3.16.a shows the XRD patterns of the synthesized BZT nanoparticles measured at room temperature. Since BT and BZT, both are perovskite oxides, comparable Bragg peaks are recorded for both BT and BZT nanoparticles. Doping larger Zr +4 cations (compared to Ti +4 cation) into the crystalline structure of BT, the tetragonal symmetry of the structures slightly change to larger orthorhombic/rhombohedral symmetries which, indeed, have bigger lattice constants (a,b,c). This increase in the lattice constants and, consequently, volume of the unit cells is observed via a shift of the Bragg peaks towards lower angles in the XRD patterns of the synthesized BZT nanoparticles upon increasing of the Zr content. These observations thus support the view that Zr is doped into the structure. Fig.3.16.b shows a linear evolution for the calculated lattice constants upon increasing the Zr content. These calculated lattice constants are in a good agreement with reported similar BZT nanoparticles lattice constants in literature[3].

Fig. 3 .

 3 Fig.3.16. a) XRD patterns of the synthesized BZT nanoparticles with various Zr content measured at room temperature. b) (100) Bragg peaks of all BZT samples. c) Lattice constants were deduced from the XRD patterns of BZT nanoparticles measured at room temperature.

Fig. 3 .

 3 Fig.3.17 Schematic view of a spark plasma sintering device.

Fig. 3 . 18 .

 318 Fig.3.18. Picture of 200 nm-sized BZT20 nanoparticles sintered via SPS a) before and b) after reoxidation.

Fig. 3 .Fig. 3 .

 33 Fig.3.19. SEM micrographs of sintered 200nm-sized BZT20 using different parameters listed in table 3.5.

Fig. 3 .

 3 Fig.3.20 a) XRD patterns of the sintered BZT20 nanoparticles measured at room temperature. b) Lattice parameters of the sintered BZT20 nanoparticles deduced from XRD. c)FWHM deduced from the (110) XRD peak of sintered BZT20 nanoparticles.

Fig. 3 .

 3 Fig.3.21 shows SEM micrographs of the sintered nanoceramics produced with BZT12 and BZT30 nanoparticles at 1000°C. Like BZT20, the grains in the BZT12 and 30 nanoceramics do not grow during the sintering, and ~100-200 nm-sized nanoceramics are produced.

Fig. 3 .Fig. 3 .

 33 Fig.3.21. SEM micrographs of the sintered 200 nm-sized BZT12 and 30 via SPS at 1000°C.

  3.23 the relative permittivity of the sintered nanoceramics are shown measured at 100 and 1000 Hz as function of temperature using a heating rate of 10°C/min.

Fig. 3 .

 3 Fig.3.23 Relative permittivity of BZT nanoceramics as a function of temperature measured at 100 and 1000 Hz.

  ,1H,2H,2Htriethoxysilane Since the functional groups are attached to the BT nanoparticles' surface, FTIR spectroscopy is a suitable analysis method to characterize these modifications. In fig.3.24, FTIR spectra of surface-modified nanoparticles are compared with non-modified BT nanoparticles. For OHmodified nanoparticles, a broad peak is detected at ~3400 cm -1 , corresponding to the stretching vibrations of -OH units present on the surface. Particles modified with (3 3 3trifluoropropyl) trimethoxysilane, which introduces -CF3 units on the particles' surface (in addition to -OH peak), displays peaks at ~1050 and 1250 cm -1 in the FTIR spectra, corresponding to Si-O and C-F bonds, respectively. Since (3 3 3-trifluoropropyl) trimethoxysilane is a silane coupling agent, one layer of silicone oxide forms around the particles, onto which the -C2H4CF3 groups are bonded. The recorded peaks confirm the introduction of this silane coupling agent on the particles' surface. Since (3 3 3-trifluoropropyl) trimethoxysilane and perfluorododecyl-1H,1H,2H,2H-triethoxysilane have comparable chemical bonds in their structures, the recorded peaks in the FTIR spectra of the modified particles from these two coupling agents are similar. The only difference is the higher contents of -CF bonds in PFO, which is observed as a more intense peak in the FTIR spectra of the PFO-modified nanoparticles.

Fig. 3 .

 3 Fig.3.24. FTIR spectra of a)-OH b)-CF3 and c)-PFO surface-modified nanoparticles measured at room temperature.

Fig. 3 .

 3 Fig.3.26. Water, diiodomethane and diethylene glycol droplets on the modified Si-wafer captured at 25°C.

Fig. 3 . 27 .

 327 Fig.3.27. Polar and disperse parts of the surface energy calculated for modified Si-wafers at 25°C. The calculated polar and disperse parts of CFE7 and CTFE4 are also shown for comparison.

Fig. 3 .

 3 Fig.3.28 shows the destabilization index of the dispersed non-modified and modified nanoparticles in 5vol.% CTFE4 solutions (polymer content with respect to the solvent). The nanoparticle content is 15vol.% with respect to the polymer. The destabilization index is a quantified term based on the intensities of transmission signals detected and calculated by a turbiscan machine. Increasing the destabilization index indicates the precipitation of the dispersed components in a given medium. The figure's inset shows the stability of the dispersions during the first 1 hr, while the main graph shows the stabilities of the nanoparticles over 12hrs.

Fig. 3 . 28 .

 328 Fig.3.28. Destabilization index of 15vol.% non-modified and surface-modified BT nanoparticles dispersed in P(VDF-TrFE-CFE) terpolymer solutions measured at room temperature.

Fig. 3 .

 3 Fig.3.29. Weibull probability for the breakdown of 15vol.% particle in CFE7 nanocomposites composed of various modified nanoparticles.

Fig. 3 .

 3 Fig.3.30 displays the polarization at their breakdown fields of CFE7 films and their nanocomposites comprising -OH and -CF3 modified nanoparticles. BT as an inorganic ferroelectric material increases the polarizability of CFE7, leading to an increase in the maximum polarization of the system to 11 µC/cm 2 . However, due to their very different dielectric constants compared to CFE7, these inorganic nanoparticles have a detrimental effect on the dielectric strength of the system, which can be partially compensated by the surface modification of the nanoparticles. As expected from the surface energies provided from the functional groups and CFE7, nanocomposites with -CF3 modified nanoparticles show higher maximum polarization due to their high breakdown strengths compared to nanocomposites comprising -OH modified nanoparticles. Furthermore, the remnant polarization of the nanocomposites is also increased due to BT's ferroelectric behavior, which slightly decreases the relaxor behavior of the system. The ferroelectric behavior of the nanocomposites comprising different nanoparticles and terpolymers is discussed in detail in the next subchapters.

Fig. 3 . 31 .

 331 Fig.3.31. Relative permittivity of CFE7 films compared to BT and BZT20 nanocomposites with 15vol.% particle contents as a function of temperature.

Fig. 3 .

 3 Fig.3.32. Polarization of CFE7 films compared to BT and BZT20 based nanocomposites with 15vol.% particle contents measured at 100 Hz and room temperature.

Fig. 3 .

 3 Fig.3.33 displays the DSC thermograms of CTFE4 films blended with BZT12 nanoparticles varying the particle content from 5 to 15 vol.% and using a heating rate of 10°C/min. Since DSC measures only the phase behavior of the medium (not the nanoparticles), the weight of nanoparticles was subtracted from the overall weight of nanocomposites. In order to welldisperse the particles into the polymer solutions, 5 vol.% solutions were chosen. 0.5vol.% polymer solutions as very dilute solutions have very low viscosities (close to the one of the solvent), facilitating the nanoparticles' precipitation in the solutions. On the other hand, 10vol.% polymer solutions as very concentrated solutions have a high viscosity, in which the nanoparticles cannot be easily dispersed.

Fig. 3 .

 3 Fig.3.33. a) DSC heating thermograms of the BZT12/CTFE4 nanocomposites with different particle content (10°C/min). b) Normalized melting enthalpies.

Fig. 3 .

 3 Fig.3.34 shows the relative permittivity and tanδ of BZT12/CTFE4 nanocomposites having different particle content measured at 10°C/min. The relative permittivity of the nanocomposites is increased with increasing particle content. According to the polar-toapolar phase transition feature, which also appears in tanδ, consistent with DSC, the polarto-apolar phase transitions are slightly shifted toward lower temperatures. This shift is not very clear in the graph due to the broad phase transition peak in the relative permittivity.

Fig. 3 . 34 .

 334 Fig.3.34. Relative permittivity and tanδ of BZT12/CTFE4 nanocomposites as a function of temperature (heating rate: 10°C/min).

Fig. 3 . 35 .

 335 Fig.3.35. Polarization of BZT12/CTFE4 nanocomposites having different nanoparticle content at 100 Hz.

  to 1 μm) via two different methods: hydrothermal and polyol routes. It was observed that the concentration of the hydrolyzing agent (NaOH in this work) and the reaction time play a vital role in determining the chemical composition and morphology of the synthesized

Fig. 4 .

 4 Fig.4.1. (a) Electric displacement and discharged energy density (shaded area) with electric field. (b) Effect of the slope of the D-E curve on discharged energy density. (I) shows the early D-saturation and consequently lower energy density while (II) shows delayed D-saturation and higher energy density[3].

Fig. 4 . 2 .

 42 Fig.4.2. Polarization-electric field dependency of a relaxor ferroelectric terpolymer with the discharging process and energy loss highlighted in blue and red, respectively.

Fig. 4 . 3 .

 43 Fig.4.3. Photograph of prepared devices (left) with top and side schematic view (right).

Fig. 4 .

 4 Fig.4.4.a shows the discharged energy density of the polymers investigated as function ofapplied electric field at room temperature. In Fig.4.4.b, the discharge efficiency is given as function of applied electric field at room temperature for each polymer. All polymers display similar energy densities, while their discharge efficiencies differ significantly. CFE7 shows an almost stable 70% efficiency at different electric fields while CTFE8 displays a low efficiency at low electric fields, which improves at higher fields until fields above ~300 MV/m are reached. There, the efficiency becomes higher (72%) than the one of CFE7. Note: The efficiency of CTFE4 remains far below that of the other polymers for the whole electric field range (from 41% at 190 MV/m to 55% at 440 MV/m).

Fig. 4 .

 4 Fig.4.4. a) Discharged energy density and b) efficiency of neat CTFE4, CTFE8 and CFE7 at different electric fields at 100Hz.

Fig. 4 .

 4 Fig.4.5. a) Maximum polarization of the CTFE4, CTFE8 and CFE7 at different electric fields and 100Hz. b) DSC heating thermograms of cast films measured at 10°C/min. c) FTIR spectra of cast films obtained at room temperature.

  The electric displacement versus electric field of fabricated capacitors are shown in fig.4.6. As expected for CTFE4, large hysteresis loops are obtained when measuring at various electric the bulky CTFE termonomer content is low in this terpolymer, the content of alltrans conformations is still considerable. Thus, less RFE phase is formed, as can be deduced from DSC (Fig.4.5.b). Clearly, the high content of all-trans conformations not only increases the maximum polarization that is achievable, but also increases the remnant polarization and hysteresis, which is shown in fig.4.6.a. By increasing the termonomer content, the less polar tttg+tttg-conformation is stabilized over the all-trans conformation (Fig.4.5.c), thus, decreasing the hysteresis (Fig.4.6.b). Although the maximum polarization is lower in CTFE8 compared to CTFE4, the potential decrease in energy storage density can be compensated by the lower Pr resulting from tttg+tttg-stabilization. As shown in fig.4.6.c, in CFE7, when compared to CTFE8, the hysteresis loop becomes narrower at lower electric fields. This narrower hysteresis loop likely arises from the presence of less polar tg+tg-conformations in CFE7 compared to CTFE8 with more tttg+tttg-arrangements. Note also that although CTFE8

Fig. 4 . 6 .

 46 Fig.4.6. Electric displacement-electric field hysteresis loops of CTFE4, CTFE8 and CFE7 obtained at different applied electric fields and 100Hz.

Fig. 4 . 8 .Fig. 4 . 9 .Fig. 4 .

 48494 Fig.4.8. Dielectric constant of CFE7 blended with different amounts of PMMA measured as a function of frequency at room temperature.

Fig. 4 .

 4 Fig.4.11 shows the Weibull distribution probability for a breakdown to occur at room temperature. Thereby, the slope represents the shape factor (β), and the values at zero give the scale factor (λ). λ is the natural logarithm of the breakdown strength (LnEb), and β provides information about the reliability of the obtained breakdown values. A higher β indicates a more reliable data set (narrower distribution of data).

Fig. 4 .

 4 Fig.4.11 Two-parameter Weibull distribution analysis of the breakdown strength (Eb) of CFE7:PMMA blends with the shape parameter β indicated in the graph.

Fig. 4 .

 4 Fig.4.12 Calculated discharged energy density and efficiency of the blends at their breakdown strength.

Fig. 4 .

 4 Fig.4.13 Two-parameter Weibull distribution analysis of the breakdown strength (Eb) with the shape parameter β deduced for two-layered blends of different PMMA content.

Fig. 4 .Fig. 4 .

 44 Fig.4.14.a shows the calculated energy densities compared to one-layer blend systems of various PMMA contents. Fig.4.14.b shows the efficiency of the corresponding blend devices in one-and two-layered blend-and neat CFE7 devices. No considerable change in two-layered devices compared to one-layer devices is found. This suggests that the efficiency of blends is stable independent of PMMA content.

  chamber and the bottom and top electrodes were connected to a high voltage amplifier via thin wires. The polarization of the device was recorded as well as the heat release and absorbance during application of an electric field. Fabricated device picture and schematic view of the system used are shown in fig 5.1.

Fig. 5 .

 5 Fig.5.1. a) Picture and schematic view of an exemplary fabricated device used for direct ECE measurements via DSC. b) Installed device inside the DSC chamber.

Fig. 5 . 2 . 5 . 3 . Influence of polymer and nanocomposite solid-state structure on electrocaloric performance 5 . 3 . 1 .

 5253531 Fig.5.2. DSC thermogram of fabricated device versus time during application of an electric field

Fig. 5 .

 5 Fig.5.3. a) Electrocaloric temperature change and b) maximum polarization of CTFE4 cast at 90°C and annealed at 105°C for 1hr as a function of temperature measured at different electric fields.

  Ceramic ferroelectric materials, due to their high polarizability and consequent high pyroelectric coefficient, can display a high EC response. However, their low breakdown strength and the often poor mechanical robustness of these materials limit their usage in cooling applications. Nanocomposites based on VDF-based polymers and ferroelectric inorganic nanoparticles, due to the inorganic component, the composite can display a high polarization while the breakdown strength remains high (due to VDF-based polymer matrix).

Fig. 5 .

 5 Fig. 5.4.a shows the measured ΔTEC of nanocomposites with two particle contents as a function of temperature measured at 63MV/m. As can be deduced from fig.5.4, adding the ferroelectric nanoparticles to CTFE4, increases the polarization of the system (Fig.5.4.b), and

Fig. 5 .

 5 Fig.5.4. a) Electrocaloric temperature change and b) maximum polarization of BZT12/CTFE4 nanocomposites with 5 and 15 vol.% particle content, cast at 90°C and annealed at 105°C for 1hr measured as a function of temperature at 63 MV/m. For comparison, data for neat CTFE4 are also shown.

Fig. 5 . 5 .

 55 Fig.5.5. Electrocaloric temperature change of BZT12/CTFE4 nanocomposites with 5 vol.% particle content cast at 90°C and annealed at 105°C for 1hr measured as a function of temperature at a) 80 MV/m and b) 100 MV/m. For comparison, data for neat CTFE4 are also shown.

Finally, in chapter 5 ,

 5 the electrocaloric performance of neat P(VDF-TrFE-CTFE) and BZT/ P(VDF-TrFE-CTFE) nanocomposites was investigated with view on electrocaloric refrigeration applications. Again, films were cast at high temperatures (90°C) in order to induce the RFE phase. Following the study of the crystalline structure of BZT performed in chapter 3, a relaxor inorganic BZT with a phase transition temperature close to that of the terpolymer matrix was chosen. It is observed that nanocomposites with BZT nanoparticles with a 12 mol% Zr content in P(VDF-TrFE-CTFE)(61/35/4) displayed a higher ΔTECE than the neat terpolymer. This electrocaloric enhancement is attributed to an increase of the polarizability of the nanocomposite, afforded by the proximity in polar-to-apolar phase transition temperatures of the nanoparticles and the matrix (~ 60°C).In conclusion, control over the solidification conditions can be effectively exploited to manipulate the solid-state structure of VDF-based polymers and, consequently, their highand low-field dielectric properties. By engineering their solid-state structure, we were able to design architectures with enhanced energy storage abilities as well as an increased electrocaloric performance. In this thesis, we mainly focused on terpolymers with a VDF/TrFE ratio of ~ 70/30 solidified from solution. There exists a plethora of other VDF-based polymers with various monomer compositions which can show different dielectric properties interesting for many applications (e.g., copolymers with a VDF/TrFE ratio of ~ 45/55, which show a drastically different phase behavior than the polymers studied in this work).Solidification of these polymers from the melt provides other ways to control structure formation. Cooling rate and molecular weight of the polymer are additional, useful parameters to tune structure. Moreover, during solidification from solution, parameters such as type of solvent and molecular weight of the polymer can lead to further options to target specific phase formation. In the blend systems, we used atactic P(MMA). Further work may concentrate on the employment of isotactic P(MMA) which would allow to control the crystallization sequence of the components in the blend. In the nanocomposites, we mainly focused on barium titanate and barium zirconium titanate particles to achieve more relaxor ferroelectric properties. The morphology of the nanoparticles such as size and shape are known to give different crystalline phases providing further means to induce desired properties. Accordingly, a plethora of options still seems unexplored for the design of ferroelectric systems targeted toward specific applications. VDF-TrFE-CFE) (64/29/7) terpolymer was purchased from Arkema-Piezotech (France). Repeat units of this terpolymer are shown in fig.7.1. The molecular weight of the polymer was measured ~300 kg/mol. According to the required concentrations, the polymer powders were added to Cyclopentanone and left overnight under stirring at room temperature to prepare the polymer solutions. Then, clear solutions were used to cast the films.

Fig. 7 . 1 .

 71 Fig.7.1. Repeat units of P(VDF-TrFE-CFE) terpolymer.

Fig. 7 . 2 .

 72 Fig.7.2. Repeat units of P(VDF-TrFE-CTFE) terpolymer.

  The P(VDF-TrFE)(70/30) and (45/55) copolymers were purchased from Arkema-Piezotech (France). Repeat units of this copolymer are shown in fig.7.3. The molecular weight of the polymer was measured ~300 kg/mol. Similar to the other polymers, according to the required 204 concentrations, the polymer powders were added to cyclopentanone and left overnight under stirring at room temperature. Then, clear solutions were used to cast the films.

Fig. 7 . 3 .

 73 Fig.7.3. Repeating units of P(VDF-TrFE) copolymer

Fig. 7 . 4 .

 74 Fig.7.4. Chemical structure of a) titanium isopropoxide, b) zirconium isopropoxide, and c) barium acetate

  In this thesis, the silane coupling agents, carrying different functional groups (Fig.7.5) as precursors, were polymerized on the surface of -OH modified nanoparticles. This polymerization was associated with forming a 3-dimensional network of siloxane (Si-O-Si) bonds around the nanoparticles which contains the mentioned functional groups in fig.7.5.

Fig. 7 .

 7 Fig.7.5. a) (3 3 3-trifluoropropyl) triethoxysilane and b) perfluorododecyl-1H,1H,2H,2H-triethoxysilane as silane coupling agents selected for the surface modification of the nanoparticles.

Fig. 7 . 6 .

 76 Fig.7.6. Schematic view of the used cone and plate rheometer.

  This study was carried under the financial assistance of MARBLE International Associated (Chair program) within the framework of IdEx Bordeaux program (reference ANR-10-IDEX-03-02) and would not have been possible without the support of countless people.First of all, I would like to express my deepest gratitude to my dear supervisors, Professors Georges HADZIOUANNOU and Natalie STINGLIN, for their marvelous patience, guidance, and support. I have benefited greatly from your wealth of knowledge and meticulous editing. I am extremely grateful that you took me on as a student and continued to have faith in me over the years. I will not forget our endless meetings and advices that gave me a correct vision in research. I would also thank my committee members, professors Paul W.M BLOM, G.H.(Gerwin) GELINCK, Mario MAGLIONE, Graziella GOGLIO, Katja LOOS and Kamal ASADI for their efforts and contributions to this work.

TrFE-CFE) (64/29/7) terpolymer 2.3.1. Solidification behavior and crystalline phase formation

	P(VDF-TrFE-CFE) (64/29/7) as a VDF-based terpolymer containing 64 mol% VDF, 29mol% TrFE
	and 7mol% CFE, called CFE7 in this work, usually crystallizes in a relaxor ferroelectric phase
	(RFE) with a phase transition to the paraelectric phase (PE) that is close to room temperature.

Absorbance (a.u.) Wavenumber (cm -1 )

	0.5 vol%
	5 vol%
	10 vol%	t m>3	3 g
	t m>4
		tg	tg
			49

Table 2 .

 2 1. Intensity of FTIR peaks (relative to the baseline) corresponding to all-trans, tttg, and tg chain conformations in CFE7 cast from 0.5, 5 and 10vol.% solutions at 40°C.

	Conformation	All-trans	trans-gauche
	Peak	1284 cm -1 1239 cm -1 768 cm -1 608 cm -1
	0.5 vol.% CFE7	0.017	0.005	0.003	0.004
	5 vol.% CFE7	0.017	0.005	0.004	0.004
	10 vol.% CFE7	0.015	0.007	0.005	0.006

2.3.3. Time stability of induced solid-state structure

  In order to investigate the time stability of the induced CFE7 solid-state structures solidified at various casting temperatures, films were stored at room temperature for a week. In fig.2.30, DSC heating thermograms taken at 10°C/min of 1 week-stored films and, for comparison, as-cast films, are shown. The films were cast at different temperatures; 40°C,

	55°C and 90°C.

Absorbance (a.u.) Wavenumber (cm -1 )

  In order to observe the structure formation of CFE7:PMMA blends, we performed DSC on blend films drop cast at 90°C for 3hrs using 0.5vol.% solutions. The heating (fig.2.39.a) and cooling (fig.2.39.b) thermograms measured at 10°C/ min are shown in fig.

						0 wt.%
						1.5 wt.%
						6 wt.%
						12 wt.%
						100 wt.%
							C=O	t m>4		tg
					2000	1750	1500	1250	1000	750	500
		0 wt.% 1.5 wt.%			b		0 wt.% 1.5 wt.%			c		0 wt.% 1.5 wt.%	tg	d
	Absorbance (a.u.)	6 wt.% 12 wt.% 100 wt.%				Absorbance (a.u.)	6 wt.% 12 wt.% 100 wt.%			Absorbance (a.u.)	6 wt.% 12 wt.% 100 wt.%
		1800	1750	1700	1650	1320	1280	1240	1200		660	640	620	600	580	560
		Wavenumber (cm -1 )		Wavenumber (cm -1 )				Wavenumber (cm -1 )

Table 3 .

 3

1. Hydrothermal reaction parameters for BaTiO3 synthesis in several batches.

Table 3 .

 3 2. Polyol reaction parameters for the BaTiO3 synthesis through several batches.

	Run N°	[Ti] (mol/l)	[Ba]/[Ti]	[NaOH] (mol/l)	Reaction temperature (°C)	Reaction time (hr)	Annealing temperature (°C)	Annealing time (hr)	Particle diameter (nm)
	1	0.8	1.6	1	220	5	-	-	~20-50
	2	0.8	1.6	1	220	5	650	8	~50-200
	3	0.8	1.6	1	220	5	800	8	~50-200
	4	0.8	1.6	2	220	5	800	8	~500-1000

Table 3 .

 3

	4. Polyol reaction parameters for the synthesis of BaZrxTi(1-x)O3 with different Zr content
				(x=0.12, 0.2, 0.3)
	Sample Composition	[Ti] (mol/l)	[Zr]/[Ti] [Ba]/([Ti]+[Zr])	[NaOH] (mol/l)	Annealing temperature (°C)	Annealing time (hr)	Particle diameter (nm)
	BZT12 BaZr0.12Ti0.88O3	0.8	0.136	1	1	800	8	~50-200
	BZT20 BaZr0.2Ti0.8O3	0.8	0.25	1	1	800	8	~50-200
	BZT30 BaZr0.3Ti0.7O3	0.8	0.428	1	1	800	8	~50-200

Table . 3

 . .6 Used parameters for spark plasma sintering of synthesized BZT20 nanoparticles.

Table . 3

 . .7 Used parameters for spark plasma sintering of synthesized BZT12 and 30.

Table 3 .

 3 8. Functional groups used to surface modify BT nanoparticles.

Table 3 .

 3 10. The droplets' contact angle on CTFE4 and CFE7 terpolymer films and calculated surface free energy of these terpolymer films with their polar and disperse parts.

	Matrix	ΘWater (°) ΘDEG (°) ΘDIM (°) γs d (mJ/m 2 ) γs p (mJ/m 2 ) γs=γs d + γs p
	P(VDF-TrFE-CFE)	85±1	63±1	64±1	26.09 ±0.71 4.92 ±0.36 31.01 ±1.06
	P(VDF-TrFE-CTFE) 84±2	64±1	62±2	25.27±1.34	4.75 ±1.94 30.02 ±2.42
	To compare the calculated surface energies, all of the values, including polar and disperse
	parts, are plotted in fig.3.27.

. Energy density in VDF-based terpolymer blends 4.3.1. P(VDF-TrFE-CFE) and P(VDF-TrFE-CTFE) neat terpolymers

  In this part, the influence of the solid-state structure of VDF based polymers on energy density and efficiency is investigated. For this purpose, we selected three VDF-based polymers of similar VDF/TrFE ratio. The selected materials are listed in table1. Due to the importance of structural stability, all the capacitors were fabricated at 90°C. As revealed in chapter 2, decreasing the entanglement density in polymer solutions (i.e. using dilute solutions) can improve crystallinity and stabilize tg conformations in the materials' solid-state, which can be desirable in terms of increasing the energy density and efficiency of capacitors. However, a low entanglement density requires solutions of very low concentration, which result in the very thin dried films (100nm). Due to the higher chance of electrical shorts at high electric field and the importance of the dielectric strength to energy storage, it was used for film preparation solutions of 5vol.% polymer content.

	4.3
	Bottom electrode	Top electrode	Polymer film
		Top view
		Side view

Table 1 .

 1 VDF-based terpolymers used in this chapter.

	Sample	Polymer	Composition
	CFE7	P(VDF-TrFE-CFE)	(64/29/7)
	CTFE4	P(VDF-TrFE-CTFE)	(61/35/4)
	CTFE8	P(VDF-TrFE-CTFE)	(61/31/8)

Fig.7.8. Weibull probability plot obtained by linearized cumulative distribution function formula.

63MV/m. Due to the increase of maximum polarization by addition of the particles, higher ΔTEC values were achieved at temperatures around the polar-to-apolar phase transition of the nanocomposites (55-60°C).

Acknowledgement:

Chapter 2

Engineering the solid-state structure of VDF-based terpolymers and blends processed from solution Chapter 3

Interfacial engineering of VDF-based terpolymer nanocomposites towards improved high-field dielectric properties

in the material [9]. A similar study was carried out with Zr +4 ions, where relaxor ferroelectric properties are induced in BaZrxTi(1-x)O3 ceramics using 25 mol% Zr with respect to Ti (x=0.25) [10,11]. Fig. 3.2 displays the phase diagram of these modified barium titanate ceramics as a function of the substituted ion content and temperature. Recently, ferroelectric hybrid organic/inorganic composites ferroelectric polymer and ferroelectric inorganic particles have become interesting to design systems combining the high dielectric constant of the inorganic component with the high breakdown strength of the polymer matrix [12][13][14]. In addition to the high dielectric constant and high polarizability of the particles, the interfacial polarization in polymer-particle interfaces can increase the overall polarization of the system. Thus, increasing the surface area of the particles can enhance the polarizability of the system. Using nano-sized particles instead of micron-sized particles not only provides a higher interfacial surface area but also minimizes the particle content loading while minimally affecting the mechanical properties of the polymer matrix.

a larger fraction of tg arrangements. This behavior was comparable to the role of CFE termonomers on terpolymers, pinning the chains, increasing the interchain distance, and stabilizing tg conformations. The dielectric properties of the nanocomposites showed an enhancement in the dielectric constant of the systems when increasing the BZT particle content. Considering the decrease in polymer crystallinity by the particles, this dielectric improvement likely arises from the relaxor ferroelectric properties of the particles, which also increases the macroscopic polarization of the system. This enhancement in low and high field dielectric properties may improve the terpolymer's efficiency for many applications, which require high polarization and relaxor behavior, such as energy storage applications.

Chapter 4

Impact of solid-state structure and interfacial engineering on energy storage properties

To conclude, CTFE polymers, due to their more polar (all-trans and tttg) conformations in their chains (even at 8 mol.% CTFE), possess a more pronounced hysteresis before their saturation polarization is reached. Since the all-trans conformation content in CTFE4 is higher than in CTFE8, this hysteresis loop is larger in CTFE4 than in CTFE8. As the applied electric field increases, this hysteresis loops become narrower and the efficiency increases. However, in CFE7, due to the more pronounced relaxor structure of this polymer (more tg conformations), the hysteresis loop is narrower than in CTFE terpolymers even at low electric fields. Indeed, this hysteresis loop is stable at any electric field. Therefore, the efficiency is essentially stable independent of electric field.

P(VDF-TrFE-CFE): PMMA blends

In this section, the energy storage ability and efficiency of P(VDF-TrFE-CFE) polymers blended with PMMA is investigated using various PMMA contents. Since P(VDF-TrFE-CFE) with 7 mol.% of CFE termonomer display a more stable device efficiency upon application of an electric field, this polymer was chosen here. Fig. 4.7 shows the calculated discharged energy densities and efficiencies of CFE7/PMMA blend films of various PMMA content at room temperature as a function of electric field. The energy density of these blends, similarly to neat CFE7, increases with electric field, however, the overall efficiencies of the blends contrasts with that of neat CFE7, following a different behavior. In the blends, the efficiencies decrease with electric field, while in the neat CFE7, the efficiency increases with electric field. In order to rationalize our observations, we consider that PMMA, due to dipole-dipole interactions with VDF repeat units in CFE7, induces, on the one hand, all-trans chain (beneficial for energy density) when compared to neat CFE7. Due to this delayed D-saturation, the device efficiency of the blends is higher than for neat CFE7 especially at low electric fields.

Decreasing the overall crystallinity does not have a detrimental effect on energy density values. Moreover, it increases breakdown strength of the system. 7µm-thick neat CFE7 films feature a breakdown strength of 400 MV/m, which improves to 515 MV/m for blends with 12 wt.% PMMA.

It was also shown that using two-layered systems instead of one-layer films can further increase the breakdown strength, which we attribute to the presence of more interfaces in the system. These interfaces can deviate electron pathways that may form during application of an electric field. This prevents formation of possible shorts. As a consequence, the breakdown strength is increased to 600 MV/m for blends with 12wt.% PMMA.

Chapter 5

Impact of solid-state structure and interfacial engineering on the electrocaloric solid-state cooling performance

Chapter 6

General conclusions and future work this chemical differences. It is shown that polymers with similar VDF/TrFE ratio (~ 70/30 in this work) follow a similar trend while polymers with other compositions may behave differently (such as P(VDF-TrFE)(45/55)). We also observed that the inclusion of bulky CFE or CTFE termonomers can pin the polymer chains into tg conformations and keep the interchain distances high, even at room temperature, provided that the polymer films are cast at high temperatures. Thus, the effect of the casting temperature on copolymers comprising only VDF and TrFE, is less pronounced than that on terpolymers. Decreasing the casting temperature lowers the pinning of the CFE units and decreases the stability of the formed tg conformations.

The presence of bulky termonomers also highlights the role of casting time. We found that longer casting times at high temperatures results in higher interchain distances and favors the integration of CFE units in the crystalline moieties in the polymer, thereby stabilizing the tg conformations. In contrast, short casting periods decreases the influence of the CFE moieties on the structure formation, leading to the stabilization of all-trans conformations. When annealing such films, the induced thermal energy provides the required chain mobility for the CFE termonomers to be integrated in the crystalline lamellae stabilizing the tg conformations.

It thus was concluded that the combination of chemical structure and the solidification parameters, determines the solid-state structure of VDF-based polymer.

In the second part of chapter 2, we focused on blending VDF-based polymers with PMMA, which is a non-ferroelectric amorphous polymer miscible with VDF-based polymers. This allows to engineer the solid-state structure of such blends. It was found that the PMMA-VDF interactions, arising from dipole-dipole interactions between the C-H bond in PMMA with C-F bonds in VDF-based polymers, stabilizes the all-trans conformations leading to more polar phases. However, this stabilization by PMMA is highly dependent on the solidification parameters used such as casting temperature. In addition to phase stabilization, PMMA hinders crystallization and decreases the overall crystallinity of the blends. The combination of the stabilization of more polar phases and the decrease in overall crystallinity highly modifies the dielectric properties of such structures.

In chapter 3, the dielectric properties of inorganic perovskite oxides and the influence of their chemical structures were investigated using barium titanate (BT) as model system because it is a well-known non-toxic perovskite oxide ceramic. In order to change the chemical structure of this ceramic, different contents of Zr +4 cation were introduced into the BT structure.

Hydrothermal and polyol methods were employed in order to synthesize this BT in form of nanoparticles. During the synthesis, the influence on the chemical and crystalline structure as well as the morphology of the synthesized ceramic were elucidated as function of reaction parameters such as concentration of hydrolyzing agent, reaction time and annealing temperature. 100-200 nm sized pure and highly crystalline non-cubic BT nano-particles were produced using these methods. It was shown that increasing the hydrolyzing agent concentration not only increases the size of the particles to micron-size scale but also changes the chemical structure of the product. Since the polyol method requires milder conditions than the hydrothermal method, the former method was employed to synthesize barium zirconium titanate (BZT) particles with 12, 20 and 30 mol% of Zr content. We found that the introduction of Zr +4 cation on the B-site of the tetragonal phase of the BT perovskite stabilizes the non-cubic rhombohedral phase, and shifts the phase transition temperature to lower temperatures. A polar-to-apolar phase transition close to room temperature was achieved in BZT particles with 20 mol% Zr content; no phase transition was observed in BZT with 30 mol% Zr content. Due to the lower permittivity of the rhombohedral phase, the dielectric constant of the BZT ceramic decreased when increasing the Zr content. In order to measure the dielectric properties of the system, the synthesized nanoparticles were spark plasma sintered and compacted into nano-ceramic pellets with the same grain size as the prepared nanopowders.

Although inorganic ceramics have remarkably high dielectric constants when compared to VDF-based polymers, their low dielectric strengths limit their application. We thus used these inorganic nanoparticles to enhance the dielectric properties of VDF-based polymers studied in chapter 2. The as-synthesized nanoparticles were surface modified with various functional groups of different hydrophilicity. Via this surface modification, the surface energy of the nanoparticles was modified to limit agglomeration and, thus, to increase the dielectric strength of the nanocomposites. It was concluded that -CF3 functional groups, is the most compatible functional group for these terpolymers. Finally, low and high field dielectric properties of such nanocomposites using P(VDF-TrFE-CFE) and P(VDF-TrFE-CTFE) terpolymers as the matrix material and BZT nanoparticles with 12 and 20 mol% Zr content were investigated. For comparison, nanocomposites with BT nanoparticles were produced as well.

We demonstrate that in such nano composites the polar-to-apolar phase transitions of the organic and inorganic component are comparable leading to an enhancement in dielectric constant and, in turn, in the polarizability of the nanocomposite.

Finally, having identified pathways to engineer the solid-state structure and dielectric properties of organic and hybrid organic/inorganic systems, the performance of these systems was investigated in chapter 4 and 5 for energy storage and electrocaloric refrigeration applications.

To indicate the critical overlap concentrations in the polymer solutions, reduced viscosity was plotted as a function of concentration. The reduced viscosity is equal to the relative viscosity increment ratio (ηi) to the mass concentration of the species of interest (ηi/C). The relative viscosity increment was calculated from the following eq.7.1: ηi=(η-ηs)/ηs (eq.7.1)

Where η is the obtained viscosity for the polymer solutions, and ηs is the obtained viscosity of cyclopentanone, used as the solvent.

Drop casting

Drop casting is one of the easiest and, in the meantime, common casting techniques in laboratories that are used for the polymer solutions having low viscosities. In this technique, polymer solutions are spread homogeneously on a substrate and let in an oven or on a hot plate for the solvent to be evaporated and films prepared. In this thesis, polymer solutions were spread on a glass substrate and left in the oven at different temperatures for different times. The drying was performed under ambient pressure. The cast films, then, were peeled off for further characterizations. These samples were used for DSC, FTIR, and WAXS characterizations.

To fabricate the metal/polymer film/metal capacitors, the polymer solutions were drop cast on the glass substrates which bottom electrodes were already evaporated on them. After casting the films, top electrodes were evaporated on the polymer films. The thickness of the films was set by the amount of the polymer solutions cast on the substrates.

Blade coating

In addition to the drop casting, blade coating is also one of the most common techniques for coating uniform thin films, ranging from a few microns to a few millimeters-thickness. The instrument is simple, consisting of a hot plate containing some holes with vacuum depression, which keeps the substrates attached to the plate during the casting. After deposition of the solutions on the substrates, a blade with a chosen rate and height spread the solutions on the substrate (fig. 7.7). By varying these two parameters and/or the polymer solution concentrations, it is possible to modulate the prepared films' thickness. 10µm-thick polymer films were prepared on the bottom electrodes using this technique to fabricate metal/insulator/metal capacitors for the energy storage study (discussed in chapter 4). For the electrocaloric study (discussed in chapter 5), 15µm-thick films were blade coated on Polyethylene naphthalate sheets as substrates.

Profilometer

The thickness of the thin films was measured employing a profilometer Bruker Dektak XT-A.

A tip is let to scan the surface across two scratches previously created on the films. The difference between the height of the films and the height of the substrates, measured through the scratches, gives the films' thickness. Several measurements are performed in order to decrease the errors in the evaluations.

Differential scanning calorimetry (DSC)

The thermal analysis experiments and electrocaloric measurement were carried out using a Mettler Toledo DSC under the continuous flux of azote. For the thermal analyses, aluminum pans were filled with tens of milligrams of samples, while an empty pan was used as the reference.

Concerning the cast polymer films, samples were cooled down to -50°C and kept for 5 min at that temperature. The first heating cycle was performed from -50°C to 200 °C at a rate of 10 °C/min to record the thermal history of the films. Afterward, a cooling cycle was performed on the samples with the same 10 °C/min to record the polymer's crystallization from the melt.

Like the first heating cycle, another heating cycle was performed with the same 10°C/min.

The recorded data was exported by STARe software.

Blade Substrate 208

Fourier transform-infrared spectroscopy (FTIR)

Bruker Vertex 70 FT-IR spectrometer in Attenuated Total Reflectance (ATR) mode at 16 cm -1 resolutions was used to perform FTIR spectroscopy on free-standing polymer films and synthesized nanoparticles from 400 to 4000 cm -1 wavenumber. Temperature-dependence FTIR was performed using a connected heating plate in the same device with heating up to 210°C.

Wide-angle X-ray scattering (WAXS)

Wide-angle X-ray measurements were performed at the Centre de Recherche Paul Pascal (CRPP, Pessac, France) on a Rigaku Nanoviewer (XRF microsource generator, MicroMax 007HF), with a rotating anode coupled to a confocal Max-Flux® Osmic mirror (Applied Rigaku Technologies, Austin, USA) producing a beam with a wavelength of 1,5418A° or energy of 8KeV. The free-standing films were put in glass capillaries exposed to the X-Ray beam in a high-temperature oven (ambient to 250°C). The detector was placed at a distance of 135mm, providing access to a 2θ angle in the 2,8° -50° range. The images were collected on a MAR345 image plate detector (MARResearch, Norderstedt, Germany), and spectra were integrated with the FIT2D software.

WAXS patterns were plotted and fitted using the Fityk 0.9.8 software [1] in a wide 2θ range of 5° to 50°. Using this software, we decomposed WAXS spectra into crystalline and amorphous contributions. The crystalline peaks were associated with the sharp peaks, located at the Bragg angle 2hkl, while amorphous signals were fitted with broader peaks. The inter-planar distances (dhkl) were deduced from the Bragg angles 2 hkl using the Bragg law: 2dhkl*sinθhkl=λ (eq.7.

2)

If a Bragg peak is associated with only one (hkl) family of planes, its corrected full width at half maximum (FWHM) hkl gives the lateral extension, Dhkl, perpendicular to (hkl) planes, using the Scherrer formula [8]: Dhkl = 0.9 * 𝜆 hkl * coshkl (eq.7.

3)

With the hkl full width corrected for the instrumental resolution. A mass index of crystallinity, χc of films was calculated using the following equation: Where Ac and Aa are the areas under the deconvoluted crystalline peaks and amorphous halos, respectively. In the orthorhombic pseudo-hexagonal structure of the VDF-based polymers, the two prominent Bragg peaks at 2200 and 2110, are too close to be resolved in this experiment; thus, the obtained single Bragg peak is the juxtaposition of the ( 200) and

(110) Bragg peaks. For each of them, the relationship between the Bragg angle 2200

(respectively 2110)and the inter-planar distance d200 (respectively d110) is giving by the Bragg law. The measured inter-planar distance d200/d110 is somewhere between d200 and d110 (weighted average).

X-ray powder diffraction (XRD)

The crystalline phase of the synthesized nanoparticles was identified using a PANalitycal X'pert PRO MPD diffractometer with a Cu Kα1 radiation source (λ=1.54184 Å). The data were collected over an angular range of 2θ=10-80°with a step size of 0.0167°. The powders were ground and sifted with a 40μm sieve for better homogeneity of the particle size. The samples were carefully prepared on stainless steel sample holder using a razor blade to prevent preferential orientations. Profile matching was performed using the FullProf Suite using the known space groups of BaTiO3 [2].

Spark plasma sintering (SPS)

To prepare the nanoceramics, 2g of the synthesized nanoparticles were consolidated using spark plasma sintering (515 S model, Fuji Electronic Industrial Co., Ltd., Kanagawa, Japan) Alumina was added on the top and bottom. SPS runs were carried out under temperaturecontrolled mode using an optical pyrometer (IR-AHS2, Chino, Tokyo, Japan) focused on the die's external surface. After the sintering, the die was cooled down to room temperature with a cooling rate of 50°C/min. After the sintering, the bulk samples were polished on their top and bottom surfaces using progressively finer abrasive paper to finally obtain 1.5 mm thick disks. Then, the pellets were post annealed for 10 hrs at 800 •C under air to re-oxidize the samples and remove the carbon contamination from the graphite tools.

Relative permittivity measurement

The relative permittivity was measured with a Solartron (1296 dielectric interface, solartron analytical) and Impedance Analyzer (Solartron Analytical, SI 1260). To control the temperature, a Linkam LTS 350 temperature-controlled element was used. These tools were interfaced with a computer whose command can be sent by the software called "Smart".

The dielectric permittivity describes a medium's response to an applied electric field in the function of the frequency of the electric field. The dielectric permittivity is a dynamic complex parameter, 𝜀r (𝜔) defined as:

Where the real part is ε' (stored energy) and the imaginary part is ε" (dissipated energy) that vary in function of ω. The ratio between imaginary ε" and real part ε' is defined as the charge losses tan(δ) that occur if charges pass through the material or go out from the material. The magnitude of ε' depends on the degree of polarization that can occur. For air, ε'= 1. For inorganic ferroelectric materials such as BaTiO3, ε' = ~10 3 -10 4 . For polymeric ferroelectric such as P(VDF) ε' = ~15, which increases to ~50 in terpolymers. To increase ε, one should increase the density of electrical dipole moments.

Ferroelectric characterization

The polarization hysteresis loops of the metal/ferroelectric/metal capacitors were recorded at room temperature using the TF Analyzer 2000E of aixACCT Systems. A continuous sinusoidal wave with a 100 Hz frequency was used, and various electric fields from 100 to 2000 mC/m 2 were applied using a high voltage amplifier. Several devices were measured per sample, and a statistical study was performed to corroborate the reproducibility of the polarization loops. The polarization is expressed in mC/m 2 and is proportional to the number of charges accumulated at the electrode surfaces. The integration of the measured current in the function of the time and the electric field applied gives the dielectric polarization. The software aixPlorer30 already gave the I-E and the P-E curves.

Weibull probability used for breakdown strength

The Weibull distribution is a continuous probability distribution named after Swedish mathematician Waloddi Weibull. He originally proposed the distribution as a model for material breaking strength but recognized the distribution's potential in his paper published in 1951 [3], a statistical distribution function of broad Applicability. Today, it is commonly used to assess product reliability, analyze life data, and model failure times. In this thesis, we used this model to obtain the breakdown strength of the fabricated capacitors. ) 𝛽 (eq.7.8) ln(-ln(1-F(x)))=βlnx-βlnλ (eq.7.9) Eq.7.9 can be seen as a straight line in the standard form (Fig. 7.8). In this figure, β is the slope of the data, and λ is the exponential of E where the Y-axis is zero. The achieved scale factor (λ) from the graph was taken as the breakdown strength of the capacitors.

Last, but not least, my warm and heartfelt thanks go to my lovely wife, Shadi, for her endless support. You have always stood behind me, and this was no exception. I am also grateful for my parents whose constant love and support keep me motivated and confident. My accomplishments and success are because they believed in me. I love you all!