Table of contents

Thesis introduction

The energy demand of the world is increasing day by day. In particular, the world's electricity demand is steadily increasing from≈13 TW in 2012 to an expected 26 TW in 2050. [1] Currently, we are depending mainly on renewable energy sources such as sunlight, wind and tidal energies. They can be used repeatedly and are naturally replenishing sources of energy. The conversion efficiency of solar energy to useful electricity by solar panels and solar thermal plants is estimated as only 40%. More than 60% of global energy is wasted in the form of heat. [2] [3] The conversion of waste heat to electricity may play an important role in our current challenge to develop alternative energy technologies to reduce our dependence on fossil fuels and reduce greenhouse gas emissions. Highly efficient thermoelectric materials can be one solution for converting this huge amount of waste heat to electricity. Inorganic materials such as silicides, Heusler alloys and metallic oxides are widely used for high temperature (800K) thermoelectric applications. These materials provide excellent thermoelectric performances and are mainly composed of non-toxic elements. [4] Moreover, the situation below 500 K is different. The room temperature and commercially available thermoelectric generators are mainly composed of a special class of inorganic materials called chalcogenides. They are composed of highly toxic elements such as Lead (Pb). The high cost, toxicity and poor processability of these heavy elements limited their applications in various thermoelectric technologies [3] .

The ability to convert conjugated polymers from insulating to highly conducting form triggered their use in different electronic applications such as organic field-effect transistors [5] [6] (OFET), organic solar cells, organic light-emitting diodes [7] (OLED) and organic thermoelectrics [8] . Pioneering work by Heeger and coworkers on doped polyacetylenes stimulated the field of organic thermoelectrics (use of flexible polymers Thesis introduction 3 as thermoelectric materials). [9] Unlike inorganic materials, these organic materials are characterized by their easy processability. They are composed of abundant elements such as carbon, nitrogen, sulphur and oxygen and they possess lower thermal conductivity (0.37 W m -1 K -1 ) than inorganic materials. [10] . Moreover, they are flexible and they can be processed into large-area devices using suitable methods such as slotdie coating, inkjet printing, brush coating and roll-to-roll coating. Both, flexibility and easy processability of organic conducting polymers results in reduced production costs and lead to the production of cheap organic-based devices, which are now available on the market.

The field of organic thermoelectrics received considerable attention after the publication of the seminal work of Crispin [8] and coworkers on poly (ethylene dioxythiophene) (PEDOT). Crispin et. al [8] have shown that accurate control of the oxidation level combined with the low thermal conductivity of PEDOT can lead to high ZT values of  0.25 at room temperature. Such performance is good enough for thermoelectric device applications. This discovery triggered much interest, on TE polymers. Doped polythiophenes such as P3HT and PBTTT showed high thermoelectric (TE) performances due to their semi-crystalline or liquid crystalline nature [11] [12] . Moreover, studies by Chabinyc [13] and coworkers proposed that the thermoelectric potential (PF) of a doped conducting polymer can be improved by increasing the conductivities of conjugated polymer from the relation PF 1/2 . Following this work, most of the studies on thermoelectric polymers focused mainly on improving the electrical conductivity of the doped polymers. A similar approach was followed in the frame of this thesis. Our main focus was to improve the thermoelectric properties of conducting polymers by improving their crystallinity and by introducing long-range order and orientation. We used a method Thesis introduction 4 called high temperature rubbing, which allows to align a large variety of conjugated polymers [14] . Sequential doping of the oriented thin films prepared by this approach produced highly conducting anisotropic films with high degrees of optical, structural and thermoelectric anisotropies. These highly anisotropic conducting polymer thin films are ideal model systems to probe the TE properties both parallel and perpendicular to the chains. They provided experimental evidence of the anisotropic charge transport mechanisms. Moreover, combining polarized UV-Vis-NIR spectroscopy and transmission electron microscopy (TEM) provided a better understanding of the location of the dopant anions in these oriented polythiophene thin films. A clearer picture of dopant induced structural changes in the polymeric unit cells was obtained from electron diffraction measurements.

The first chapter (Chapter 1) in my thesis is mainly dealing with the state of the art and basic concepts, which are relevant to my topic of research. The first section of this chapter is mainly dealing with general aspects of conjugated polymers such as their bonding, energy levels, chemical structure, crystal structure and different orientation methods. A short discussion about crystal structure, experimental evidence for the semicrystalline nature of conventional regioregular P3HT and liquid crystalline PBTTT is included. The second part of the chapter is dealing with the basic principles and concepts of thermoelectricity. This is followed by a brief discussion about different doping methods used for doping conducting polymers. At the end of this chapter, we present the impact of doping on the crystal structure of PBTTT and P3HT. Finally, a summary of representative thermoelectric performances in p-doped polythiophenes PBTTT and P3HT is discussed.

In Chapter 2, we mainly investigate the impact of the side chain length of PBTTT (C8, C12, C14, C18) on the diffusion coefficients of the dopant anion 2,3,5,6-tetrafluoro-7,7,8,8-Thesis introduction 5 tetracyanoquinodimethane (F4TCNQ) into the side chain layers of PBTTT. Doping of polymer semiconductors such as PBTTT with acceptor molecules such as (F4TCNQ) is widely used to tune the charge transport and thermoelectric (TE) properties in thin films.

However, the mechanism of dopant insertion in the polymer matrix, the insertion kinetics and the ultimate doping levels were only marginally investigated. Electron diffraction measurements on these doped thin films showed that the Intercalation of F4TCNQ into the layers of side chains of PBTTT results in an expansion of the lattice along the side chains and the contraction along the -stacking direction for all polymers. The extent of lattice expansion decreases with increasing side chain length. UV-Vis-NIR spectroscopy demonstrates integer charge transfer for all investigated PBTTTs. The doping kinetics and final doping level depend on both the side chain length and packing.

Highly disordered n-octyl and crystalline n-octyldecyl side chain layers tend to hamper dopant diffusion in the side chain layers contrary to n-dodecyl side chains that can host the highest proportion of dopants. Consequently, the best TE properties are observed for C12-PBTTT films. Alignment of the polymers enhances significantly the TE performance by increasing the charge conductivity and the thermopower along the rubbing direction. Aligned films of C12-PBTTT show charge conductivities of 193 S/cm along the rubbing direction and power factors of approx. 100 µW.m -2 .K -1 .

Chapter 3 compares the TE properties of P3HT and PBTTT upon doping with a strong oxidizing agent, namely FeCl3. FeCl3 doping of aligned films results in charge conductivities up to 2×10 5 S/cm and metallic-like thermopowers similar to iodine-doped polyacetylene. The films are almost optically transparent and show strongly polarized near infra red polaronic bands (dichroic ratio > 10).

The comparative study of structure-property correlations in P3HT and C12-PBTTT identifies three conditions to obtain conductivities beyond 10 5 S/cm: Thesis introduction 6 i)

Achieve high in-plane orientation of conjugated polymers with a high persistence length ii) ensure uniform chain oxidation of the polymer backbones by regular intercalation of dopant molecules in the polymer structure without disrupting the alignment of -stacked layers and iii) maintain a percolating nano-morphology along the chain direction.

The highly anisotropic conducting polymer films are ideal model systems to investigate the correlations between thermopower S and charge conductivity . A scaling law S -1/4 prevails along the chain direction but a different S -ln relation is observed perpendicular to the chain direction, which was previously observed by Mateeva [15] et. al for stretch aligned polyaniline. This suggests two different charge transport mechanisms parallel and perpendicular to the polymer chains. The simultaneous increase of charge conductivity and thermopower along the chain direction results in a substantial improvement of thermoelectric power factors up to 2 mW . m -1. K -2 in C12-PBTTT Chapter 4 is a comparative study of the effect of F6TCNNQ and F4TCNQ doping on the crystal structure and thermoelectric properties of oriented C12-PBTTT films. F4TCNQ doping of organic semiconductors such as P3HT and PBTTT has been widely documented in the recent literature. This chapter focuses mainly on doping of C12-PBTTT using a new doping agent, namely F6TCNNQ. Unlike F4TCNQ, F6TCNNQ doping of polythiophenes has been marginally studied. This dopant is chosen because of its deeper LUMO energy level of -5.3 eV, versus -5.1 eV for F4TCNQ, that can enhance charge transfer between C12-PBTTT and F6TCNNQ. Oriented C12-PBTTT thin films are prepared by high temperature rubbing and sequentially doped with F6TCNNQ using a new doping method, called "incremental concentration doping" (ICD). The ICD method implies that a given sample is doped by dipping it successively in dopant solutions of increasing concentration. ICD Thesis introduction 7 of C12-PBTTT using F6TCNNQ leads to record electrical conductivities of 2430 S/cm and thermoelectric power factors (PF) 530±200 µWm -1 K -2 along the rubbing direction.

Moreover, this method leads to a higher electrical conductivity than the conventional doping method (direct doping). This study demonstrates that a progressive intercalation of dopants in PBTTT crystals using the ICD method helps preserve the high level of order initially present in the aligned C12-PBTTT films.

The diffusion coefficients of both F6TCNNQ and F4TCNQ anions have been determined using the doping kinetics followed by UV-Vis NIR spectroscopy. A lower diffusion coefficient is observed for F6TCNNQ anions, demonstrating the size dependence of diffusion coefficients into the polymer matrix. The longer and bulkier the anion, the lower its diffusion coefficient. Polarized UV-Vis-NIR spectroscopy on these doped thin films showed that F6TCNNQ -anions are better oriented in the layers of the alkyl side chains as compared to F4TCNQ. The estimation of charge carrier mobilities indicates that ICD leads to higher mobilities than DD. The lower degree of ordering of F4TCNQ -anions produced lower electrical conductivity than F6TCNNQ doped C12-PBTTT thin films.

Thesis introduction

Chapter 1. Fundamental concepts & state of the art

This chapter is dedicated to the fundamental concepts and literature survey, that are relevant to this PhD thesis. The chapter is divided in two sections. The first section mainly focuses on the general aspects of conjugated polymers such as their bonding, energy levels, chemical structure, crystal structure and different orientation methods. A short description of the crystal structure and experimental evidence for the semi-crystalline nature of conventional regioregular P3HT is briefly introduced. Similarly, the crystal structure and experimental evidence for the thin film microstructure of PBTTT is also included. Later, a brief description of different polymer orientation techniques such as rubbing, strain alignment, friction transfer and epitaxy are discussed.

The second part of this chapter is mainly dealing with the thermoelectric properties and doping of polythiophene thin films. A short introduction and discussion of thermoelectric parameters such as electrical conductivity, Seebeck coefficient and power factors are provided. It is followed by a brief summary of different chemical doping methods used for polythiophenes e.g. solution doping, sequential doping, vapor phase doping. The consequence on the resulting thermoelectric properties is discussed.

At the end of this chapter, we present the impact of doping on the crystal structure of PBTTT and P3HT. Finally, a summary of representative thermoelectric performances in p-doped polythiophenes PBTTT and P3HT is given.

1. General aspects of conjugated polymers 1.1. Energy level design of conjugated polymers.

Conjugated polymers are organic macromolecules; they are characterized by a conjugated backbone with alternating double and single bonds. They have been the focus of enormous attention in the industrial and academic sector for various optoelectronic applications. [1] The main interest of conjugated polymers stems from the π electron cloud. π electrons in the double bonds can be delocalized over several monomeric units in the polymer chains. Because of this electron delocalization from one monomer to another monomer, they have the ability to conduct electrical charges mainly along the conjugated backbone. Let us consider the simplest case of the ethylene molecule. The ground state electronic configuration of the carbon atom is 1S 2 2S 2 2P 2 . In the case of ethylene, the carbon atom is in sp 2 hybridization and hence one electron in the 2S orbital can be promoted to the vacant P orbital. Now the excited state configuration of the carbon atom become 1S 2

2S 1 2P 3 . The electrons in the 2S, 2Px and 2Py orbital undergo sp 2 hybridization, which gives rise to three hybridized orbitals. Two of these three orbitals overlap with 2 hydrogen Chapter 1: Fundamental concepts and state of the art 13 atoms, which results in two carbon-hydrogen sigma bonds and the third orbital can be used to form a sigma bond with the next -CH2 units, resulting in a C-C sigma bond in ethylene. The unused 2PZ orbitals undergo sidewise overlapping which produces the  bond in the ethylene molecule (see figure 1.1). The sidewise overlap of the 2PZ orbitals results in a restricted rotation along the bond axis and gives rise to a planarized and stiff polymeric backbone. [2] The electrons present in the overlapped Pz orbitals do not belong to one single bond but are delocalized over several atoms. The length scale of delocalization of these  electrons is the conjugation length. The presence of alternating double and single bonds in conjugated polymers is responsible for the semiconducting properties of conjugated polymers. The smallest identical repeating unit of a polymer is called a monomer.

Chapter 1: Fundamental concepts and state of the art

During polymerization reactions, the chemical coupling of the monomers leads to orbital interactions and hence it results in an energy level splitting of the π (bonding) and π* (antibonding) orbitals. The π (bonding) orbitals are known as the highest occupied molecular orbitals (HOMO) and π* orbitals are known as the lowest unoccupied molecular orbitals (LUMO). [3] The valence band is arising from the π orbitals while the π* orbitals form the conduction band (see figure 1.2). The energy gap between the HOMO and LUMO is the bandgap (Eg). The bandgap of a polymer depends on its chemical structure and on the extent of conjugation. The longer the conjugation length, the smaller the bandgap. The sigma bond between the monomer preserves the structure of the molecule. The electrons in the π orbitals are free to undergo electronic and optical excitation and the charges can be delocalized along the polymer chains. The conjugated polymers are semiconductors or insulators in their intrinsic form and they can be converted to a conductor via doping. In 1977, Heeger and coworkers Chapter 1: Fundamental concepts and state of the art 15 reported high conductivity in iodine-doped polyacetylene. [4][5] Although doped polyacetylene has limited commercial applications, its discovery led to major developments in the field of organic electronics. Soon after the discovery of doped polyacetylene, extensive research work was carried out on other conducting polymers such as polythiophenes, polyanilines and polypyrroles (see figure 1.3) for different industrial and scientific applications.

P3HT: The working horse among organic semiconductors

Soluble semi-conducting polymers attracted wide attention from the scientific community because of their excellent semiconducting properties and easy solution processability. This made them excellent candidates for flexible electronic devices prepared by roll-to-roll manufacture at low cost. Among semi-conducting polymers, regio-regular poly (3-hexylthiophene) (P3HT) is one of the most popular polymers as it combines unique crystalline and optoelectronic properties. The first synthesis of unsubstituted polythiophene was reported by Lin and coworkers in 1980. [6] This discovery was followed by highly conducting Iodine doped thin films of these unsubstituted thiophenes by Yamamoto et al. [7] Alkyl side chains were introduced on the polymeric backbone to help solubilize the polymer in common organic solvents. [7][8] Introduction of alkyl side chains on polythiophene gives rise to three main regio isomers that differ in the position of the coupling between two successive monomers. The term "regio regularity" indicates the control of the configuration of the hexyl side chains. Depending on the position of the hexyl side chain, different regioisomers can be found in P3HT. The HH (head to head) regioisomer is produced by the coupling of two monomeric units through 2,2 position.

The 2,5 coupling leads to the HT (head to tail) isomer and 5,5 coupling leads to the TT (tail to tail) isomer. The different monomers are represented in figure 1.4. Yamamoto

Chapter 1: Fundamental concepts and state of the art 16 and coworkers were the first to predict the existence of regioisomers in P3HT. The isomer with HT-HT coupling is called regioregular (rr) isomer. The HH and TT isomers are regioirregular isomers. In a fully regioregular P3HT chain, all the monomeric units are coupled in head to tail manner. [8] In regioirregular P3HT (HH and TT coupling), the steric repulsions between the adjacent alkyl side chains in the monomeric units can produce twists in the polymeric backbone. Such twisting reduces the conjugation length and packing efficiency and finally leads to poor optical and electrical properties. The rr-configuration produces thin films of P3HT with very high crystallinity. The high crystallinity is a key parameter for various optoelectronic and charge transport applications.

Chapter 1: Fundamental concepts and state of the art 17 rr-P3HT has a very high HOMO level (-5.1 eV) and possesses a bandgap of 2 eV. Highly crystalline P3HT was one of the first polymers to show high mobilities of 0.1 cm 2 /V.s in field-effect transistors [9] . The main advantage of P3HT is that it can be crystallized very easily from the solution to produce well-oriented crystals of large size. However, the grain boundaries are detrimental for charge transport and mobility. Depending on the processing method, different morphologies were obtained in P3HT. Crossland and coworkers were able to obtain a high FET mobility of 0.1 cm 2 /V.s in films with spherulitic

P3HT. [10] The high FET mobility was obtained due to the reduction in the grain boundaries as a result of precise control of the nucleation and growth of anisotropic spherulitic structures in P3HT. Most interesting optical and charge transport properties in P3HT derived from its high crystallinity.

Single crystals were also obtained from P3HT. Similar to inorganic semiconductors, P3HT

is also characterized by a crystal lattice (monoclinic form I) with well-defined unit cell parameters [11] . Unlike inorganic crystals, single crystals of P3HT show high anisotropy in optical and electrical properties. For example, charge transport along the chain direction (c axis) is higher than that along the π stacking (b axis) direction (see figure 1.5 A). The poorest charge transport is along the alky side chain direction (a-axis). P3HT slowly emerged as a benchmark semiconductor polymer due to its easy synthesis procedure, solution processability and highly crystalline microstructure. More detailed investigations of the microstructure and transistor performance of P3HT provided important guidelines for the design of new polymer semi-conductors with enhanced device performance. [9][12,13] [14-17]

Semi-crystalline P3HT-lamellar morphology

Many conjugated polymers are semi-crystalline in nature, which means that they consist of ordered polymer chains in the crystalline regions while the chain ends and chain folds are mainly present in the amorphous parts of the polymer. So-called tie chains connect adjacent crystals through the amorphous regions of the polymer. Chain folding is an inherent nature of semi-crystalline polymers. The lamellar morphology of the polymer is mainly produced due to the crystallization mechanism. During crystallization, some of the chains tend to align in a specific manner and create ordered regions in the form of lamellar crystals or fibrils. The chain ends and chains folds cannot be accommodated in the crystals and are rejected to the amorphous phase of the semi-crystalline polymer.

Hence, the lamellar structure implies a regular alternation of crystalline regions with ordered chains and amorphous regions with chain folds, chain ends, and chemical defects. The chain folding of poly(alkylthiophenes) was first revealed in the initial work of Mena -Osteritz, and coworkers. They were able to observe chain folding in regioregular poly(alkylthiophenes) prepared by epitaxial orientation on highly oriented pyrolytic graphite (HOPG) using scanning tunnelling microscopy (STM). [12,13] STM studies evidenced the existence of chain folds in regioregular polythiophenes and also evidenced the interdigitation of the alkyl side chains (see figure 1. 6 A and B).

Further evidence for the semi-crystalline structure in P3ATs was obtained by Brinkmann et al. using transmission electron microscopy (see figure 1.5 B). [14][15][16][17] . Transmission electron microscopic studies by Brinkmann and Rannou on epitaxially oriented P3HTs determined the molecular weight (Mn) for which the chain folding occurs. They proposed that chain folding occurs for Mn  7.3 kDa. The presence of chain folding at higher molecular weight leads to a saturation of the lamellar periodicity. Typically, a lamellar period of 28 nm is observed when Mn ≥ 18.8 kDa. Lower molecular weight fractions form a crystalline phase without chain folds i.e. extended chain crystals. As the molecular weight increases, the number of chain folds inside the amorphous zones increases and the polymer becomes semi-crystalline. Moreover, another interesting observation is the increase in lattice parameter (aP3HT) with the increase in molecular weight of the Chapter 1: Fundamental concepts and state of the art 19 polymer [18] . As the number of chain folds increases in the amorphous zones, chains reenter the crystalline part of the lamellae and exert a stress on the crystalline lattice i.e.

increases the separation along the alkyl side chains. [17,19,20] Apart from the molecular weight and regio regularity, the crystallinity and morphology of P3HT thin films can also be affected by processing conditions. The crystallization kinetics is an important factor that determines the crystallinity of a thin film. It has been observed that the polymer thin films deposited from higher boiling point solvents had higher crystallinity than the films deposited from low boiling solvents. For example, solution-processed thin films cast from chlorobenzene showed higher crystallinity than the films prepared from chloroform. This indicates that slower solvent evaporation is beneficial for higher crystallinity in solution-processed thin films. [21][22][23] Film deposition conditions can also affect the crystallinity of thin films. It has been reported that fast

Reproduced from (12&13)

Since our main focus of the thesis is to study the orientation and the anisotropic thermoelectric properties, the subsequent discussion is limited to the relevant crystal structures and orientation techniques of polymers.

Crystal structure of P3HT

As discussed earlier, regioregular poly (3-hexylthiophene) (rr-P3HT) is one of the most studied semi-crystalline polymers. Since it has high levels of crystallinity, the crystal structure of P3HT was an interest for crystallographers. Several studies helped to

A) B) C) D)
Chapter 1: Fundamental concepts and state of the art 21 elucidate the crystal structure of P3HT. In 1992, Prosa and coworkers reported a structural model for P3HT, which consists of well-defined "lamellae" made of π-stacked polymer backbones uniformly spaced by the layers of hexyl side chains. [24] Two polymorphs were identified: form II has interdigitated side -chains contrary to form I, that is dominant in thin films used in devices. An important structural refinement of form I P3HT was performed by Tashiro et. al [25] in 1997. The structural model was refined from X-ray data obtained on stretch-aligned P3HT from trichlorobenzene. [11] According to this model, P3HT crystallizes in a monoclinic unit cell (see figure 1.5 A) with lattice parameters a= 16.6 Å, b= 7.8 Å, and c = 7.8 Å,  = 86.5 °, a is the direction along the alkyl side chains (aP3HT), b is the direction along the π-π stacking (bP3HT) and c is the direction along the polymeric backbone (cP3HT). From the electron diffraction studies and crystallographic calculations, Kayunkid [11] et. al predicted that the unit cell of P3HT contains 4 thiophene monomers. One chain is shifted with respect to the second chain along the c axis (see figure 1.7 A). The shift along the backbone is produced as a result of crystallization of the hexyl side chains into an orthogonal sublattice. [11] Moreover, they were able to assign a new space group P21/C for the proposed structure. In the scope of this thesis, we used the P3HT crystal structure proposed by Kayunkid [11] and coworkers.

Liquid -crystalline polymers with enhanced structural order

Structure-property studies on P3HT provided important guidelines for the design of new polymer semiconductors with enhanced efficiency. The hexyl thiophene monomer lacks a centre of symmetry and this limits the various polymerization techniques to produce highly regioregular polymers, which is necessary to obtain highly crystalline films. [26,27] Also, the delocalized π-electron system of P3HT is susceptible to various oxidation reactions. Thirdly the high side-chain density of P3HT prevents the interdigitation and reduces the long-range order in the polymer. Because of these limitations, new polythiophenes, mainly thienothiophene-based polymers were designed.

In 2006, McCulloch and coworkers reported very high mobilities of 0.2 to 0.6 cm 2 /V.s for a new class of liquid crystalline polymers Poly(2,5-bis(3-alkylthiophene-2-yl)thieno [3,2-b]thiophene) (Cn-PBTTT). [28] It is a copolymer of thieno [3,2-b]thiophene and 4,4-dialkyl 2,2-bithiophene monomer units (see figure 1.8).In comparison to P3HT, the density of the alkyl side chains is large in P3HT with respect to PBTTT since the thienothiophene bears no side chains. In addition, the electron delocalization from the thienothiophene ring to the backbone of the polymer is less energetically favourable because of the high resonance stability of the thienothiophene unit. This reduced electron delocalization and the limited inductive effect of the alkyl side chain (+I effect) reduces the electron delocalization and lowers the HOMO of PBTTT. [28,29] The increase in ionization energy (IE) with respect to P3HT is consistent with a redshift of 20 nm in the absorption maxima of PBTTT (see figure 1.9). It has been found that the polymer backbone of PBTTT is more rigid than P3HT and it can go through a liquid crystal phase transition around 140 C. PBTTT can form ordered domains with planarized backbones over a distance of 100 nm. [28] in 2007, Hamadani and coworkers reported ). [30] Such high mobilities observed in PBTTT indicate that a high degree of structural order in conjugated polymers is essential for the application of polymers in organic electronics

Crystal structure of PBTTT

The thin film microstructure of PBTTT was investigated quite intensively because of its excellent performance in thin-film transistors. Simple models of crystalline structure of PBTTT were proposed by McCulloch and coworkers in 2006. [28] Synchrotron X-ray scattering studies on annealed C12-PBTTT gave evidence for the lamellar structure of this family of polymers.

X-ray scattering studies provided evidence for a highly ordered lamellar structure (see figure 1.10) characterized by numerous h00 reflections (out-of-plane) which are representative of the interlayer separation along the alkyl chains of PBTTT (aPBTTT). The obtained d100 spacing for the first order interchain separation along the alkyl chain is equal to 19.6 Å in C12-PBTTT. The intensity of the d100 reflection was enhanced by annealing at a higher temperature. However, there was no clear evidence for the interdigitation and tilting of side chains. It can be either interdigitated or closely packed without interdigitation and tilted with respect to the plane of the backbones. The (0k0) peaks in the in-plane scan is a fingerprint of the -stacking (3.72 Å). Moreover, the (00l) reflections are representative of the periodic order along the chain direction (c=13.5 Å).

These results proved the existence of highly ordered crystalline domains in the C12-PBTTT polymer. C12-PBTTT films showed mobilities of 0.2 to 0.6 cm 2 /V.s and mobilities of 0.7 cm 2 /V.s were obtained for the corresponding C14-PBTTT. These mobilities are about one Chapter 1: Fundamental concepts and state of the art 25 order of magnitude larger than that of P3HT because of the enhanced structural order.

Structural order in the polymer is significantly enhanced by annealing the liquid crystalline phase.

Reproduced from (28&31)

Chabinyc and coworkers proposed a more detailed investigation of the crystal structure of PBTTT in 2007. They investigated the substrate effect on crystal structure and charge mobilities of PBTTT. [31] It has been previously demonstrated that for P3HT, the crystalline domains in the thin films can nucleate from the film/substrate interface. These results proved the importance of the interfacial interactions on the thin film microstructure. [32] Chabinyc and coworkers observed high mobilities of (0.2 to 0.5 cm 2 /V.s) on silicon dioxide dielectrics, which were treated with alkyl trichlorosilane. However, poor mobilities (0.005 cm 2 /V. s) were observed on bare silicon dioxide. Alkyl side chains are interdigitated. Reproduced from (31) X-ray scattering studies of C12-PBTTT and C14-PBTTT polymers on trichlorosilane-coated silicon dioxide produced similar X-ray scattering patterns as observed by McCulloch and coworkers in 2006. [28] For both polymers, the lamellar stacking was observed along the (qz) direction while the peaks along the (qxy) direction contain the 003 and 010 peaks:

crystals were oriented edge-on the substrate. The interplanar spacing (d003) and (d010)

were nearly the same for both C12-PBTTT and C14-PBTTT, but a larger d100 spacing was Brédas and coworkers did a detailed investigation of the structure of PBTTT in 2012. The structural model of Cn-PBTTT (n=C12, C14, C16, C18) was obtained by a combination of experimental X-ray diffraction studies and computational modelling. [33] They proposed The side chains are tilted 26° to the conjugated backbone direction (see figure 1.12).

The alkyl side chain length is an important parameter, which can affect the structure and optoelectronic properties of a conjugated polymer. [34,35] Molecular models of the structure of C12-PBTTT and C14-PBTTT were proposed based on grazing incidence X-ray diffraction and molecular simulations. It has been observed that the size and shape of the simulated unit cell were slightly modified by increasing the length of the alkyl side chain from dodecyl (C12) to hexyl decyl (C16). The proposed packing determined by simulation was in a good agreement with the experimental data. The addition of two more methylene groups to C12-PBTTT resulted in an increase of the lamellar spacing by 2 Å. Also, the angles β (the angle between c and a) and ϒ (angle between a and b)

increase by 1° and 3° respectively. In the refined models, the alkyl side chains are interdigitated and form a very dense sublattice. The π-stacking distance remained constant in the range between 3.5 Å -3.6 Å. [33] The band structure measurement of these polymers confirmed that the bandgap is independent of the side-chain length. The measured bandgap is equal to 2.2 eV for all polymers, which indicates that they must have similar electronic properties.

Controlling orientation of conjugated polymers

Polymer orientation is an interesting technique because the resulting oriented films show improved optical and electronic properties over non-oriented polymers. During Chapter 1: Fundamental concepts and state of the art 29 polymer orientation, randomly oriented polymer chains can be aligned or crystallized in a specific manner leading to improved anisotropic physical properties. Different polymer orientation techniques were developed for industrial and scientific applications.

Depending on the method of orientation, different structures and morphologies can be produced, which affect the charge transport and optical properties. [10,14,36,37] Orientation techniques can be classified mainly in three families. The first family involves the use of shearing forces (Rubbing, Friction transfer) on the surface of the polymers. These forces can align the polymer chains in the direction of the applied forces. The second family of orientation methods involves the use of orienting substrates on top of which the polymers can nucleate and crystalize.

The third class of orientation methods rests on flow-orientation. The flow of a solution of polymers can lead to alignment of fibrillar nanocrystals of semiconducting polymers. [38] However, the efficiency of orientation depends on molecular weight and usually poor alignment is achieved for large Mw polymers.

Orientation by mechanical rubbing

Rubbing is an established method used for aligning polyimides layers and for the fabrication of liquid crystal displays. [39,40] This method was also used for aligning small molecules such as α-sexithiophenes. [41,42] The rubbing process involves the use of shearing forces on the surface of polymer films to align them in the direction of shearing.

Usually, the procedure (see figure 1.13) involves the use of a rotating cylinder covered with a microfiber tissue, which can be translated over the conjugated polymer on a substrate maintained at a suitable temperature (see figure 1.13). In 2002, Heil and coworkers reported an enhancement of mobility in regioregular P3HT which was aligned by mechanical rubbing at room temperature. [43] They were able to prepare highly From a structural point of view, more work on the orientation of P3HT was done by Brinkmann et. al in 2011. [43] Upon rubbing, the polymer chains were aligned in the direction of rubbing and crystals change the contact plane from edge-on to face-on with an increasing number of rubbing cycles (see figure 1.14). Polymer orientation was also studied as a function of molecular weight. A higher degree of chain orientation was observed for low molecular weight P3HT (≤ 7.9 kDa), while very poor orientation was observed for higher molecular weights (>50 kDa). The poor alignment of higher molecular weight fractions was observed due to the high degree of chain entanglement, which hindered the chain orientation. In addition to this, hole mobilities were also measured along the parallel and perpendicular directions to the rubbing for low Mw samples. It has been observed that the mobility was enhanced along the direction parallel to the rubbing and an anisotropy (µ///µ⊥) of 20 was obtained. In most of the studies using rubbing, alignment was performed exclusively at room temperature. [39,40] This rubbing technique was improved later in the Brinkmann group by introducing temperature as an additional parameter. Biniek [44] observed in C12-PBTTT by rubbing the polymer at a temperature in the range 75 °C ≤ TR ≤125 °C. [44] It has been demonstrated that temperature is an important parameter for polymer orientation. The applied temperature helps to melt the alkyl side chains which imparts some plasticity to the film that can be easily aligned. Highly oriented films of C12-PBTTT were obtained upon rubbing at 125 °C (see figure 1.15). Films show a high degree of anisotropy in optical and charge transport properties. The oriented films were birefringent in an optical microscope and electron diffraction patterns reveal a high degree of order along the chain direction. Because of this, the Chapter 1: Fundamental concepts and state of the art 33 mobilities were enhanced along the rubbing direction. High-resolution TEM revealed the nanomorphology of the oriented PBTTT films. Similar to P3HT, the presence of sulphur in the oriented backbone generates a specific contrast relative to the alkyl side chain in the high-resolution images of PBTTT (see figure 1.15D). [16,19,36,45] Reproduced from (36) The molecular weight of a conjugated polymer is also an important parameter that influences the degree of orientation. Hence, the correlation between molecular weight and degree of alignment is an important point to discuss in this chapter. In 2014, Brinkmann and coworkers studied the effect of molecular weight on the orientation of P3HT and PBTTT batches with different molecular weights. It was observed that the lower molecular weight P3HT (13 KDa) was aligned at room temperature while the higher molecular weight (43 kDa) showed an onset of alignments at 50 °C (see figure 1.16). The P3HT batch with a molecular weight of 53 kDa was poorly oriented at low temperature and showed a high degree of alignment at higher rubbing temperature. [36] The same trend was also observed in the case of PBTTT.

Chapter 1: Fundamental concepts and state of the art 34

High temperature rubbing was used by Hamidi-Sakr et al. to fabricate highly oriented and crystalline films of P3HT. They were able to obtain the semi-crystalline morphology in oriented films with a very high dichroic ratio > 25. [15] Mobility measurements were also carried out as a function of rubbing temperature on organic field-effect transistors (OFET). The mobilities along the rubbing direction were larger than those measured perpendicular to the rubbing. The measured mobilities did not improve with increasing TR because of the existence of amorphous interlamellar zones that slow down the hole migration between crystalline domains along the rubbing direction.

Later in 2017, Hamidi-Sakr et al. showed that doping of oriented P3HT with suitable electron acceptors such as 2,3,5,6-Tetrafluoro-7,7,8,8-tetracyanoquinodimethane (F4TCNQ) produces oriented thin films with highly anisotropic electrical conductivity. In addition, they observed that the measured conductivity was enhanced along the rubbing direction. Anisotropy in the 8-10 range was observed for the conductivity and a maximum conductivity of 22 S/cm for F4TCNQ doped P3HT along the parallel direction. [46] High temperature rubbing was shown to be an effective method for preparing highly oriented and crystalline polymer films with enhanced conductivity along the rubbing direction.

In the frame of my thesis, I used high-temperature rubbing to align different polythiophenes for thermoelectric applications, as discussed in chapter 2, 3 and 4.

Orientation by friction transfer

Smith and Wittmann pioneered this method of polymer orientation for oriented PTFE (Polytetrafluoroethylene) layers. [47,48] Friction transfer is a versatile method and it can be widely applied to align a large palette of polymers including P3HT, poly(9,9′-dioctylfluorene) (PFO) etc.

Reproduced from (49&50)

The polymer to be aligned is first compressed in the form of a cylindrical pellet under high pressure. This pellet is translated over a hot substrate at a constant speed and a thin oriented polymer film is deposited on the hot substrate (see figure 1.17). [49] Nagamatsu and coworkers applied this technique for the fabrication of oriented films of P3HT and poly(3-dodecylthiophene). [50] The oriented P3HT displayed an order parameter close to unity. The distribution of in-plane orientation of P3HT chain axis is of around 10 ° as evidenced by Grazing X-ray diffraction measurements. X-ray diffraction measurements indicate that the π stacking (b axis) is perpendicular to the substrate plane and the polymer crystals adopt a face-on orientation after alignment. Poly(3dodecyl thiophene) films oriented by friction transfer also showed high degree of Chapter 1: Fundamental concepts and state of the art 36 orientation along the drawing direction. Mobility measurements were carried out parallel and perpendicular to the direction of rubbing in OFET devices. The measured mobilities were highly anisotropic and anisotropy factors of µ///µ⊥ =8 were observed in these films.The main disadvantage of this method is that the substrate must be heated to very high temperature and the thickness of the deposited film is difficult to control.

Moreover, large amounts of polymers are necessary to prepare the pellets.

Orientation by strain alignment

Strain alignment is an alternating technique used for aligning semiconducting polymers.

In this method, the active polymer is deposited on a stretchable substrate such as Polydimethylsiloxane (PDMS). After deposition, the substrate can be stretched in a specific direction and hence the crystals and polymeric chains (backbones) can be aligned in the direction of stretching. (PTOPT) with a dichroic ratio of 2.4. [49,51] DeLongchamp and coworkers [52] applied strain Chapter 1: Fundamental concepts and state of the art 37 alignment on regioregular P3HT. The substrate layer was PDMS. High levels of orientation were observed in strained P3HT in which polymer backbones were aligned in the direction of stretching with a dichroic ratio of 4.8. The high degree of polymeric backbone alignment along the stretching direction resulted also in anisotropy of charge mobilities. The mobilities measured along the stretching direction were higher than those along the direction perpendicular to the rubbing (see figure 1.18). An anisotropy (µ///µ⊥) ≈ 9 was reported.

Orientation by using an orienting substrate (Epitaxy)

Epitaxy is defined as the natural or artificial growth of a crystal (guest) over a crystalline substrate (host) that determines their mutual orientations. [53] If the host and guest have a different chemical nature, then it is known as heteroepitaxy. If they have the same chemical nature, then it is known as homo epitaxy. Lattice matching is an important criterion for epitaxial growth. Epitaxy of semi-crystalline polymers on aromatic crystals produced thin films of high orientation with a regular orientation of crystalline domains. [49] The epitaxy of semi-crystalline polymers can be achieved on a large variety of substrates such as NaCl, KCl and aromatic crystals. [49,54,55] 10 to 15% of lattice mismatch is the upper limit for successful epitaxial growth. In the beginning, epitaxy was widely used for only inorganic materials but, the epitaxial orientation of polymers on organic or polymeric materials was widely exploited later in the 1980s following the pioneering work of Lotz and Wittmann [53,56] .

A systematic investigation of polymer epitaxy was performed by Lotz and Wittmann in the 1980s. [53,56] They studied epitaxial growth of polymers such as polyethylene, nparaffins and aliphatic polyesters on aromatic crystals such as benzoic acid. Later Brinkmann and Wittmann attempted epitaxial orientation of P3HT. P3HT was oriented Chapter 1: Fundamental concepts and state of the art 38 on the surface of aromatic crystals of 1,3,5-trichlorobenzene (TCB). TCB plays a two-fold role in the epitaxial crystallization of P3HT. Firstly, TCB plays the role of a solvent when molten and once it crystallizes, TCB acts as a substrate for the oriented crystallization.

1-dimensional epitaxy explains the alignment of P3HT on TCB. It has been found that C P3HT ≈ 2CTCB. After the orientation, the TCB layer was removed by vacuum evaporation that creates large areas of highly oriented P3HT crystals on a glass substrate. epitaxial orientation of P3HT on aromatic potassium 4-bromobenzoate (KBrBz). [55] The Chapter 1: Fundamental concepts and state of the art 39 procedure consists of three steps. The first step involves the preparation of large surface of KBrBz crystal. In the second step, the polymer is deposited on the top of the KBrBz crystals (see figure 1.19). Finally, isothermal annealing is applied to promote the epitaxial orientation of P3HT. The orienting substrate is removed by rinsing with ethanol.

Finally, highly oriented films of P3HT were obtained after epitaxy, as demonstrated from the electron diffraction patterns (figure 1.19)

Anisotropy of charge transport

Conducting polymers such as PBTTT and P3HT are characterized by well-defined unit cells. As described above, the charge transport in these unit cells is highly anisotropic.

For instance, charge transport along the chain direction (c axis) is higher than that along the π stacking direction (b axis). The poorest charge transport is along the alky side chain direction (a axis).

At a larger length scale, the polymers show a semi-crystalline morphology where crystalline and amorphous regions alternate. The crystalline region is composed of highly extended and ordered chains but the amorphous zones contain chain ends, chain folds and tie chains which connect successive crystalline regions. Taking this microstructure into account, Lan [57] and coworkers predicted the charge transport behavior of regioregular P3HT films by using theoretical calculations (see figure 1.20). [57] They proposed that within the ordered regions of P3HT, the hole mobility along the intrachain route ( along the polymer backbone) is significantly larger (1 cm 2 V -1 s -1 ) than the mobility along the π stacking direction (10 -2 cm 2 V -1 s -1 ). The chain ends and chain folds present in the amorphous zone reduce charge transport. The tie chains (bridging chain) help channel charges from one crystalline region to another. However, the presence of multiple crossing points between the chain ends in the amorphous zones further slows down charge mobility. Detailed investigations were carried out to determine the charge transport anisotropies in P3HT. For instance, Jimison [58] and coworkers showed that in directionally crystallized P3HT films (face-on oriented crystals), the charge transport is highly anisotropic with high mobilities measured along the chain direction in the crystalline regions (grains) because of the difference in charge transport along the chain direction (cP3HT) and along the direction of alkyl side chains (aP3HT). The mobilities along the chain were at least one order of magnitude higher than along the perpendicular direction. O'Connor [52] et. al were able to obtain enhanced mobilities along the chain direction in strain aligned P3HT

films with an anisotropy of 9. A similar situation was also seen by Hartmann [19] and coworkers in highly anisotropic thin films of P3HT prepared by rubbing at room temperature. The reported anisotropy in mobility was µ///µ⊥  22 for the best in-plane oriented samples of low molecular weight P3HT (6.4 kDa). But only an anisotropy of 2 was obtained for the high molecular weight P3HT due to highly interconnected crystalline domains, which prevent the alignment of crystalline domains upon rubbing at room temperature. [19] Anisotropic charge transport of PBTTT was also investigated [59] . Biniek [44] et. al applied high temperature rubbing to liquid crystalline C12-PBTTT and a detailed investigation was done on the anisotropy of mobilities. Highest mobilities of 4×10 -3 cm 2 V -1 s -1 were observed parallel to rubbing at a temperature of 100 °C, at which the films show high inplane alignment. Xue et. al applied strain alignment on liquid crystalline PBTTT and observed that the polymer chains were oriented in the direction of the applied strain and the resulting mobility 1.67 cm 2 V -1 s -1 was the highest reported in the case of oriented PBTTT. Similar results on mobility anisotropy were previously observed by Lee et. al in zone cast films of uniaxially oriented PBTTT [60] .

All these results on the polymer orientation prove that obtaining a high degree of inplane alignment is mandatory to reach higher charge mobilities in conjugated polymer thin films. Based on these results, we have decided to apply high-temperature rubbing to a series of polythiophenes including P3HT and PBTTT with different alkyl side chains for thermoelectric applications. The present project aims at developing new polymeric thermoelectric materials in oriented and highly crystalline films. Our strategy is to control the level of crystallization and the alignment of the conducting polymer before transforming the semi-conducting films into their doped conducting form by sequential doping. Both crystallinity and orientation of conducting polymers were recognized as key points to obtain more efficient thermoelectric films. In our approach, we will Chapter 1: Fundamental concepts and state of the art 42 decouple the doping of the polymer from the crystallization and take advantage of the very high level of order in undoped semiconductor polymers that can be obtained by high-temperature rubbing. Once oriented and crystallized, the semiconducting films will be doped by different methods. The thermoelectric properties of the semi-crystalline doped polymers can be controlled by precise tuning of the crystallinity and by changing the doping conditions.

Thermoelectricity

Thermoelectricity refers to the direct conversion of heat to electricity and vice versa.

Highly efficient thermoelectric materials are important for power generation applications because they have the ability to convert waste heat to electricity. At present, inorganic materials are the best performing TE materials. The thermoelectric material's efficiency at a given temperature T depends on its conductivity (σ), on its Seebeck coefficient (S) and thermal conductivity () via the figure-of-merit ZT=σS 2 /.

The numerator σS 2 corresponds to the power factor (PF) of the thermoelectric material.

To determine the usefulness of a material in a thermoelectric generator or a thermoelectric cooler, the power factor is calculated by its Seebeck coefficient (S) and its electrical conductivity (σ) under a given temperature difference. Materials with a high power factor are able to 'generate' more energy. Bismuth Telluride (Bi2Te3), Alloys of Bismuth-Antimony -Telluride and Bismuth chalcogenides are three important classes of inorganic TE materials with a ZT ≈1 at room temperature. The expensive raw materials and reduced abundance of the components lead the scientific community to propose to replace inorganic TE materials by organic conducting polymers. Recently, remarkable TE performances were reported in poly(3,4-ethylenedioxythiophene) polystyrene sulfonate (PEDOT:PSS) by Crispin and coworkers. [61] A ZT value of 0.25 was obtained at room temperature for this conducting polymer. This initial result obtained for a conducting polymer such as PEDOT:PSS motivated the scientific community to Chapter 1: Fundamental concepts and state of the art 43 investigate the TE performance of other conducting polymers, for instance, polythiophenes.

Basic principles of thermoelectricity

In the early 1800s, Thomas Seebeck observed that when the junctions of two dissimilar materials are kept at two different temperatures (T and T+T), a voltage difference is produced between them (V). [62] This Voltage difference is proportional to the applied temperature gradient(T).

S =

𝑉 𝑇

(1)

The ratio of the voltage gradient to the applied temperature gradient is called the Seebeck coefficient (S), which is an intrinsic property of a material. The Seebeck coefficient is usually small for metals (µV/K) and large (up to several mV/K) for insulators.

The sign of the Seebeck coefficient indicates the type of charge carriers. S>0 for holes (p-type semiconductors) and S<0 for electrons (n-type semiconductors). The physicist A.F loffe introduced the concept of figure of merit. The thermoelectric performance of a material is determined by the figure of merit ZT. The dimensionless ZT represents the measure of thermoelectric efficiency. [63] zT =

σ.S 2  T (2)
where σ is the charge conductivity, S is the Seebeck coefficient, T is the absolute temperature and  is the thermal conductivity. The thermal conductivity () of a conducting polymer is the sum of the electronic and lattice contribution i.e =e+L,

where e is the contribution from the charge carriers and L accounts for the contribution from the lattice vibrations.

Chapter 1: Fundamental concepts and state of the art 44

The inorganic Bi2Te3 shows ZT ≈1 at room temperature while the organic materials were characterized by very low ZT values usually below 0.1. [64] Due to the difficulties of thermal conductivity measurements in polymer films, the power factor (PF) is used to compare the TE performance of different polymers. [3] PF = S 2 σ (3)

The thermoelectric parameters such as S, σ and  are mutually dependent and this affects the optimization of TE performances. The optimization of one of these factors adversely affect the other ones and hence the optimization is impossible without compromising the other parameters.

σ = nqµ (4)
The electrical conductivity () depends on the mobility (µ) and the number of charge carriers (n) (see equation 4). n is determined by the amount of dopants and the number of free charge carriers. µ is determined by solid-state order (crystallinity, orientation).

Numerous studies demonstrate that the electrical conductivity and the mobility of charge carriers depend critically on the initial crystalline structure of the polymer. [36,46] The electrical conductivity increases upon doping but the Seebeck coefficient decreases.

The electrical conductivity can vary from 10 -7 S/cm for undoped to 10 5 S/cm for heavily doped polymers. [65] The Seebeck coefficient varies from 10 3 µV/K for undoped conjugated polymers to 10 µV/K for heavily doped semiconductors and S <10 µV/K for conductors.

A thermoelectric device is made of an array of individual TE legs connected in series (see figure 1.21) which is the smallest repeating unit of a thermoelectric device. [3][66] The two legs of a thermoelectric element must be made of dissimilar materials with different Seebeck coefficients to make sure that the resulting thermovoltage does not cancel, i.e. V OC = N(S 2 -S 1 )T (5) Where N is the number of legs and Voc is the open-circuit voltage and T is the temperature gradient. The maximum power (Pmax) output is obtained when the external load (Rload) is equal to the internal resistance (R internal), i.e. Rload =R internal, which is related to the Voc:

P max = V OC 2 4R int (6)

Thermoelectricity in conducting polymers

The ability to tune the conductivity of a polyacetylene film from insulating to metallic level was achieved by the joint work of Alan J. Heeger, Alan G. MacDiarmid and Hideki

Shirakawa. [67,68] Remarkable electrical conductivities of 10 5 S/cm and power factors of 1350 µWm -1 K -2 were obtained for Iodine doped polyacetylenes. [69,70] However, high instability of the iodine dopant in air and difficult processing of the oriented polyacetylene films prevented their use for thermoelectric applications. As a result, researchers turned their attention to air-stable and high-performance polymers for thermoelectric applications. In 2010, air-stable thin films of polyelectrolytes were developed by Mai et al. with high Seebeck coefficients, S  195 µV/K and very low thermal conductivity (). However, the PF was only of 0.84 µWm -1 K -2 due to their poor conductivity of 0.22 S cm -1 .

An important contribution to the field of organic thermoelectrics was achieved by Bubnova et al. in 2011. Crispin and coworkers managed to achieve a PF  324 µWm -1 K -2 on poly (3,4-ethylenedioxythiophene):tosylate (PEDOT:TOS) with a high ZT  0.25 at room temperature. [61] They showed that the thermoelectric properties of organic semiconductors can be enhanced by suitable processing and precise tuning of doping levels. Organic materials were envisioned to replace in part inorganic materials currently used for thermoelectric applications. As compared to inorganic materials, the main advantage of polymer semiconductors is due to their facile processability and low thermal conductivity. The work of Crispin et al. triggered a large interest for research on organic thermoelectric materials. The dominant strategy to improve TE properties was accordingly based on improving the electrical conductivity of doped polymers New polymers, new dopants and new doping strategies were introduced to boost the TE performances of conducting polymers. An important milestone in the field of organic TE was achieved by Yim [71] and coworkers in 2008. They studied the influence of the HOMO energy levels of various p-type polymers on doping strength. The high electrical conductivity of 10 -1 S/cm was obtained for P3HT with respect to other polymers due to a favourable energetical offset between the HOMO (-4.8 eV) of P3HT and LUMO of F4TCNQ. [71] Solution doping of P3HT mainly triggered by the work of Yim et al in 2008. [71] Thermoelectric properties of other polythiophenes such as PBTTTs were also exploited by Glaudell et al. in 2015. The studies evidenced that the deeper HOMO (-5.1 eV) level of PBTTT is also favourable for doping with conventional F4TCNQ and results in electrical conductivities of  3 S/cm and α  60 µV/K. [72] P3HT and PBTTT were two suitable model systems for chemical doping due to their high structural order. [14] Both the semi-crystalline structure of P3HT and the liquid crystalline nature of PBTTT are characterized by high degrees of structural order, which can lead to high mobilities and conductivities in thin films.

Doping of conjugated polymers a) Electronic Process involved upon P-type doping

Doping can be defined as the addition of impurities into an extremely pure semiconductor in order to enhance its electrical performance. In the case of an inorganic semiconductor, the highest occupied energy levels form a conduction band (CB) while the lowest occupied energy levels form the valence band (VB). The energy gap between the two bands is known as the bandgap (Eg). The electronic conductivity is possible only if an electron from the VB reaches the CB. This happens only if the two bands are close enough.

Unlike inorganic semiconductors, doping of an organic semiconductor is a complex mechanism in which a "charge transfer agent is used to generate, by oxidation or reduction, positive or negative charges in an intrinsically conducting polymer". [3] The doping of conducting polymers can be achieved by two methods.1) Redox doping, which involves the transfer of an electron between the conducting polymer and the dopant species to form a charge-transfer complex. 2) Acid-Base doping, which involves the Chapter 1: Fundamental concepts and state of the art 48 transfer of a cation or an anion to the polymer backbone (see figure 1.22). A common example for the acid doping is camphor sulphonic acid (CSA) and dodecyl benzene sulphonic acid (DBSA) doping of polyaniline. The protonation of the emeraldine form of polyaniline gives rise to highly conducting polyanilines with an electrical conductivity of  100 S/cm. [73] Acid doping is also applicable to polythiophenes such as PBTTT and P3HT.

Ethyl-benzene sulphonic acid (EBSA) and perflurooctyl-trichlorosilane (FTS) doping has been widely experienced with PBTTT. [74] In the case of P-type doping, an added dopant molecule can remove one electron from the highest occupied molecular orbital (HOMO) of the polymer which creates a hole in the VB and reduces the bandgap. This hole can be delocalized to the neighbouring polymeric chains and this migration of holes is at the origin of electrical conductivity. By adjusting the dopant concentration, the conductivity of a polymer can be tuned from 10 - 7 S/cm (for the undoped polymer) to 10 5 S/cm for highly doped polymers. [65] Hereafter I will mainly focus on P-type doping of conjugated polymers such as P3HT and PBTTT due to its relevance for my thesis. P3HT is a widely exploited polymer for solar cells and OFET applications because of its high mobility performance and solution processability. Due to the same reasons, P3HT is a well-studied system for thermoelectric applications. Doping the semi-crystalline structure of P3HT with highly stable and strong molecular dopants can produce highly conducting thin films. In the following, we will focus on the model system P3HT/F4TCNQ.

Chapter 1: Fundamental concepts and state of the art 50 Different models were proposed in order to explain the electron transfer between P3HT and F4TCNQ. Mainly, two contradictory models were proposed so far (see figure 1.23).

The first one implies integer charge transfer i.e. one electron from the polymer is transferred to the dopant molecule, which creates a hole in the HOMO of the polymer. [75] These holes in the polymer can either be trapped by the anion or move within the lattice of the polymer in the form of "free" charge carriers. Another model was proposed by Aziz et al. [76] They suggested that there is only a partial charge transfer between the polymer and the dopant and the electron transfer is taking place through the formation of a "supramolecular" Charge Transfer State (CTS). Both the HOMO and LUMO of the CTS are derived from the HOMO and LUMO of the neutral P3HT and F4TCNQ respectively. This model was questioned by the work of Pingel et al. [75] who followed the charge transfer in P3HT and F4TCNQ from a very low concentration (ppm) of dopant to very high concentration with the help of optical spectroscopy, Kelvin probe measurements and conductivity. They demonstrated that there is always a one-electron transfer between P3HT and F4TCNQ. Similarly, integer charge transfer was clearly demonstrated by Hamidi-Sakr [46] and coworkers in doped P3HT thin films using polarized UV-Vis-NIR spectroscopy.

Similar to P3HT, PBTTT is another important candidate for thermoelectric applications due to its high degree of structural order and high charge in mobility OFETs. [26,29,77,78] Evidence of charge transfer between PBTTT and F4TCNQ was obtained by the detailed investigation of Cochran [79] et al. in 2014. They concluded that due to a favourable energy offset between the HOMO of the PBTTT and the LUMO of the dopant, F4TCNQ can be used as an effective dopant for C14-PBTTT. [79] Later, in 2015, Glaudell [72] and coworkers reported promising TE performances in F4TCNQ doped PBTTT. More efforts to enhance the thermoelectric properties were done by Karpov [80] and coworkers in 2016. They have introduced a new dopant F6TCNNQ which had a slightly lower LUMO energy level than for conventional F4TCNQ. Molecular doping of P3HT with this strong electron-deficient molecular acceptor produced a high electrical conductivity of  7 S/cm, greater than conventional P3HT:F4TCNQ systems. [80] It was concluded that Chapter 1: Fundamental concepts and state of the art 52 a polymer: dopant system with more positive energy offset between the HOMO of the polymer and LUMO of the dopant can produce better charge transfer compared to a polymer dopant pair with similar energy levels.

FeCl3 is a strong Lewis acid, which is capable of accepting an electron pair from a Lewis base to form an acid-base adduct. FeCl3 has been used extensively to dope polyacetylene [81] leading to high electrical conductivities of 100 S/cm. FeCl3 + e - FeCl2 + Cl -(1)

FeCl3 + Cl - FeCl4 - (2)
In the proposed mechanism, there are two kinds of FeCl3 dopants which have their individual roles in polyacetylenes: one acts as an electron acceptor from polyacetylene and the other stabilizes the Cl -ion produced as a result of the reduction of FeCl3.

In the framework of my thesis, I will be mainly investigating the impact of F4TCNQ, F6TCNNQ and FeCl3 doping on the crystal structure and TE properties of oriented PBTTT and P3HT thin films prepared by high temperature rubbing. New dopants and new polymers were synthesized by utilizing the wide toolbox of organic chemistry which is represented above (see figure 1.24). A more detailed investigation of doping of polythiophenes such as PBTTT and P3HT is described in section 7.1 of chapter 1 and in the following chapters 2,3 and 4.

b) Charge carriers in conducting polymers

Polarons in a conducting polymer can be generated by several methods such as chemical doping or by electrochemical methods. As discussed in the previous charge transfer Chapter 1: Fundamental concepts and state of the art 53 models, the doping of polythiophenes such as PBTTT and P3HT can generate highly conducting radical cations, called polarons. Brédas and coworkers made a detailed investigation of the energetics of polarons and bipolarons upon suitable doping. [82] Experimental evidence for the formation of polaron and bipolaron in polymers such as P3HT was obtained by Christina [83] and coworkers using electron paramagnetic resonance spectroscopy (EPR) and Infrared spectroscopy (IR). The presence of a charge on a polymer chain modifies the bond lengths in the thiophene chain and produces a (see figure 25D). Vardeny and coworkers also observed these two polaronic transitions in Iodine doped P3HT (see figure 1.25 E). [84] This polaronic distortion in the polymer chains extends over several monomers depending on the structure of the polymer.

Further removal of an electron from this polaronic state generates a dication named bipolaron (see figure 1.25 C).

Delocalization of polarons is responsible for the electrical conductivity. In the case of amorphous materials, the polarons are highly localized, leading to very low charge carrier mobilities and conductivities even at high doping levels. For example, amorphous regiorandom P3HT displays poor conductivities of  0.01 S/cm upon vapor phase doping with F4TCNQ while the regioregular P3HT(RR  84%) showed enhanced conductivities up to 5 S/cm. [85] In the case of semi-crystalline polymers, extended intrachain order helps to delocalize these polarons to a larger number of monomers while the amorphous parts containing the chain ends and folds tend to localize polarons. More generally, polaronic transport in conducting polymers can be i) intrachain via delocalization along the polymer backbone. ii) via hopping between the polymer segments.

Factors influencing doping and thermoelectric properties of polythiophenes

The ability to dope polythiophenes with strong acceptors such as F4TCNQ stimulated the exploitation of thermoelectric properties of the PBTTT and P3HT family. One has to optimize several parameters in order to maximize TE performance for a dopant: polymer pair and hence it is a complex process. The easiest way to improve the TE performance of a conducting polymer is by improving its electrical conductivity, whereby one can xylene treated P3HT was attributed to the increased charge carrier mobility. [85][86] Detailed investigations on the crystallinity and electrical conductivity were reported by Scholes et al. in 2017 [87] . The properties of molecularly doped films of conjugated polymers were explored as a function of polymer crystallinity. [87] The electrical conductivity of sequentially doped P3HT-F4TCNQ was higher in films of P3HT with higher crystallinity. The higher electrical conductivities were attributed to the enhanced polaronic delocalization in the crystalline samples, which was evidenced by the UV-Vis-NIR spectroscopy. [87] b) Thermal conductivity

As discussed in the previous section (6.1), the ZT value of a conducting polymer is also depends on the thermal conductivity. This thermal conductivity is the sum of the electronic (e) and lattice (L) contributions. Few studies reported that the thermal conductivity of conducting polymers are in the range of 0.1 to 0.7 Wm -1 K -1 , similar to Chapter 1: Fundamental concepts and state of the art 56 amorphous solids. [88,89] Such low values indicate that the contribution from the electronic thermal conductivity is not significant. The electronic contribution of thermal conductivity (e) is estimated  0.002 Wm -1 K -1 based on the electron gas model and using the Wiedermann-Franz law, e = LT, where L is the Lorentz factor, which is equal to 2.4×10 -8 J 2 K -2 C -2 ,  is equal to 300 S/Cm for a conducting polymer and T is the temperature. [90] Thermal conductivity measurements on doped polyaniline further supported this low contribution from e. [87] Hence the main contribution to the thermal conductivity is coming from the lattice vibrations (L). This can be affected by doping, as it changes the number of charge carriers and the volume of the unit cell.

Conducting polymers can either be amorphous, semi-crystalline or liquid crystalline. No systematic studies have been reported so far to investigate the relationship between morphology, crystallinity and the thermal conductivity. Moreover, the room temperature thermal conductivity measurement demands high stability of doped semiconductor.

The thermal conductivity can also be anisotropic. The anisotropy in thermal conductivity is coming from the orientation of polymer chains. Some research work reported an anisotropy of 4 (///⊥ ) in oriented polyimide thin films [91] and 10 in polyethylene. [92] Some of our preliminary results on thermal conductivity measurements in oriented P3HT

further support this observation. Engineering of  is an alternative approach to tune the TE properties but the existing strategies were not reported in this introduction.

c) Size and electron affinity of dopant and IE of polymer

Size and electron affinity of dopants can have a direct impact on the polaronic delocalization and electrical conductivity. It has been already proved that the electron transfer between polythiophenes and dopant molecule is possible only if they have an energetically favourable offset between the HOMO of the polymer and LUMO of the Chapter 1: Fundamental concepts and state of the art 57 dopant. [71][79] Marder [93] and coworkers studied a series of dopants with different sizes and electron affinities combined with a number of polymers of different ionization energies. They investigated how these parameters influence the electrical conductivity and TE properties. [93] They have observed that, at lower dopant loading, the molybdenum based dopant Mo (tfd3) with high electron affinity and large size gave 800 times greater electrical conductivity than FeCl3 doping on the same polymer. The observed high electrical conductivity was due to high EA of the dopant and also to the shielding of the Coulombic interaction between the polaron and anion by the large-sized molybdenum dopant. Unfortunately, the doping efficiency of the molybdenum complex was saturated at low doping level but FeCl3 managed to effectively dope the polymers to reach higher doping level. Another example is the enhanced electrical conductivity observed by Karpov [80] and coworkers in F6TCNNQ doped P3HT due to a deeper LUMO of the dopant than F4TCNQ.

Electron transfer between the polymer and the dopant is possible only if they have an energetically favourable offset. For instance, the importance of IE energy of the semiconducting polymers on doping was proposed by Kiefer and coworkers in 2019. [94] They proposed that sometimes the dopants can accept two electrons instead of a single electron from the polymer backbone. This can be possible only when the IE energy of the neutral polymer is not only less than that of the EA of the dopant, but also less or similar to the EA of the dopant anion. In this case, the dopant can remove one electron to form a di-anion and the dopant anion is also able to remove one electron to become the di-anion. Hence the use of stronger doping agents with the ability to extract two electrons from the polymers of higher IE can produce highly conducting materials.

d) Polymer orientation techniques such as high temperature rubbing [46] , epitaxy [55] and tensile stretching [95] are some common methods used to produce highly oriented and crystalline polymers. Several studies showed that the electrical conductivities of Chapter 1: Fundamental concepts and state of the art 58 oriented polymers are highly anisotropic and high electrical conductivities and mobilities were measured mainly along the direction of the applied force. [46] For instance, Hamidi-Sakr and coworkers observed an enhanced electrical conductivity of  22 S/cm along the rubbing direction in oriented P3HT thin films. The film was aligned using high temperature rubbing. [46] . Hynynen and coworkers employed tensile drawing on P3HT to create free-standing films with a high degree of uniaxial alignment. Sequential doping of these aligned films with a molybdenum tris(dithiolene) complex leads to 5-fold enhancement of the power factors up to 16 µWm -1 K -2 and electrical conductivity of  13 S/cm along the drawing direction. [95] e) Doping methods and resulting thermoelectric properties of polythiophenes Three important doping procedures are mainly used for doping of polythiophenes. They are i) Solution doping ii) Sequential doping from solution and iii) Vapor phase doping.

i) Solution doping of polythiophenes

Solution doping refers to a doping method where both the polymer and the dopant are mixed together in a common organic solvent. The resulting solution is then used to cast thin films. (See figure 1.26 A). Original work by Aziz [76] et al. and Yim et al. reported that rather low conductivities in the range  0.1-1.0 S cm -1 are obtained in P3HT doped with F4TCNQ in chloroform. [71][72,96-98] Duong [97] et al. demonstrated that in solution mixed doping of P3HT-F4TCNQ, there exists two doping regimes, the weak and strong doping regimes. [97] At low doping regime (molar concentration of dopant  0.012), the dopant remains heavily solvated, does not donate the electron to the polymer and stays preferentially in the amorphous regions of P3HT(see figure 1.27 A). Within the strong doping regime ( 0.03 to 0.17), most of the F4TCNQ undergoes spontaneous charge transfer with the polymer and forms P3HT + / F4TCNQ -. This leads to high electrical conductivities in the strong doping regime. Duong [97] and coworkers hypothesized that Solution doping of PBTTT with F4TCNQ was first reported by Cochran and coworkers in 2014. [79] The main difficulties during the solution doping were the formation of aggregates, charge transfer salts and precipitation at room temperature. This mixed solution was maintained at 110 °C to 120 °C in order to avoid gelation prior to film casting, which leads to poor doping efficiencies and conductivities. Similarly to P3HT, Glaudell [72] and coworkers showed that solution doping of PBTTT always results in poor conductivities.

Figure 1.27. Illustration of doping dependence of electrical conductivity based on the location of dopant anion.

A) At lower doping concentration, the ionized F4TCNQ prefers to reside in the amorphous regions of P3HT. This results in poor conductivity. B) At higher doping concentration, the ionized F4TCNQ prefers to reside in the crystalline regions of P3HT. This improves electrical conductivity. Reproduced from (97) To conclude on solution doping, the main disadvantage is the aggregate formation due to the spontaneous charge transfer between the polymer and dopant in solution. As a consequence, the cast thin films show poorly ordered domains which prevent effective charge transport between ordered P3HT crystals and limits ultimately the conductivity. [96] Chapter 1: Fundamental concepts and state of the art 61 ii) Sequential doping of polythiophenes

In order to avoid the aggregation mechanism and poor conductivities in solution doping, a new doping procedure was introduced by Scholes [99] et. al in 2015 which is known as sequential doping. The polymer is first dissolved in a suitable solvent and spin-coated.

In the next step, the polymer film is doped by a solution of dopant in an orthogonal solvent (see figure 1.26 B). High electrical conductivity of  5.5 S/cm was obtained in sequentially doped F4TCNQ -P3HT. [99] During sequential doping, the introduction of F4TCNQ into P3HT preserves the initial structure and orientation of the pristine polymer crystallites and produces a film with lower surface roughness than the mixed solution doping. Later in 2016, the sequential doping procedure was studied in detail by Jacobs et al. [96] They demonstrated that sequential doping in P3HT-F4TCNQ occurs rapidly (<1 sec) and that the doping level of the film can be precisely controlled by varying the concentration of the doping solution. Atomic Force Microscopy (AFM) confirmed that sequential doping produces more uniform films as compared to mixed-solution doping.

Sequentially doped films contain a large number of tie chains that improve the electrical conductivity in these films.

The sequential doping protocol was successfully applied to other polymer/dopant systems such as P3HT: F4OCTCNQ, P3HT:Mo(tfd)3, P3HT:DDQ and PBTTT:F4TCNQ and demonstrated the wide applicability of this method [96] .

The sequential doping method was also applied to oriented thin films. [46] Highly oriented and crystalline polymer films of P3HT were prepared by high-temperature rubbing and subsequently doped using followed by sequential doping from Acetonitrile (ACN)-F4TCNQ.

The results on sequentially doped P3HT demonstrate that, in order to reach enhanced electrical conductivities and mobilities, it is important to preserve the initial thin film iii) Vapor phase doping.

An alternative doping method similar to sequential doping makes use of vapor phase deposition of the dopant on the polymer film. A precast polymer film is exposed to the vapors of the dopant for a certain exposure time which helps to control the amount of dopant incorporated in the film. [85] The advantage of this approach is to avoid any contact of the polymer with an orthogonal solvent. Some of the first examples of such doping were reported by Chabinyc [59] and coworkers for PBTTT and F4TCNQ. They compared the effect of solution and vapor phase doping with FnTCNQ (n=2 and 4) on the thin-film microstructure and resulting electrical conductivities of PBTTT. It was observed that the vapor phase doping results in 2 to 3 orders of magnitude higher conductivities than solution doping. Soft X-ray scattering studies on the vapor doped sample indicated that the higher conductivity was due to higher orientation correlation length (OCLlength scale of aligned backbones). Due to higher OCL, vapor doped F4TCNQ-P3HT films showed very high electrical conductivity of  670 S/cm and a Seebeck coefficient of 42±6 µV/K. This leads to a PF of  120 µWm -1 K -2 for PBTTT: F4TCNQ, one among the highest PF ever reported for P-type semiconducting polymers.

Various groups have shown that this doping method can be applied to numerous polymers and dopants with similar beneficial effects over solution doping. [85][100] For instance, Hynynen and coworkers showed that the vapor phase doping of P3HT using F4TCNQ lead to high electrical conductivity of 12.7 S/cm. [85] Chabinyc and coworkers further demonstrated that the vapor phase infiltration of F4TCNQ into the P3HT thin films does not strongly change the initial film morphology and lead to higher electrical conductivities than the solution doped P3HT films. [100] The reported results also indicate that the way the dopant molecules are introduced into the polymer film impacts very strongly the resulting properties of the doped layers. This is why we will investigate in details different doping protocols for oriented PBTTTs (see chapter 2,3 and 4).

Dopant induced structural changes in polythiophenes

Doping of P3HT and PBTTT with molecules such as F4TCNQ opens important questions on the exact location of dopants in the polymer matrix. Indeed, those polymers are Chapter 1: Fundamental concepts and state of the art 64 either semi-crystalline or liquid-crystalline in structure, which must impact the doping mechanism. In the case of P3HT, one must determine if the dopant prefers to reside in the crystalline domains or in the amorphous zones that contain tie chains, chain ends and chain folds. Several groups have tried to address these issues. An initial indication on the location of F4TCNQ -anion in solution doped P3HT was proposed by Salleo and coworkers [97] . They have proposed that the F4TCNQ -anion prefer to reside in the amorphous regions at low dopant concentration and move to the crystalline regions upon higher dopant loading which leads to high conductivities. This is reflected by the changes in the lattice parameters of P3HT thin film. The d100 and d020 spacing are changed to 18.5 Å and 3.59 Å respectively as evidenced by GIXD (see figure 1.29 A to C).

Chapter 1: Fundamental concepts and state of the art 65 Following this, Moulé [96] and coworkers reported that the location of the dopant anion is governed by the solvent in which the dopant is dissolved. They reported that sequential doping of P3HT with F4TCNQ in acetonitrile results in localization of the F4TCNQ -anion in the amorphous regions. Moreover, doping from chlorobenzene allows the infiltration of the dopant into the crystalline part (see figure 1.28 B). A more convincing picture of the dopant localization in the crystalline regions was obtained later in the study by Hamidi-Sakr [46] and coworkers on oriented P3HT. A combination of electron diffraction and polarized UV-Vis-NIR spectroscopy were applied to follow the doping mechanism in oriented P3HT films. Electron diffraction Chapter 1: Fundamental concepts and state of the art 66 measurements showed that the crystal lattice of P3HT was expanded in the direction of the alkyl side chain layers (d100) and compressed along the direction of π stacking (d020).

Polarized UV -Vis-NIR spectroscopy showed that the dopant F4TCNQ -lies in a plane perpendicular to the oriented P3HT backbone. The dopant anions are therefore located in the layers of alkyl side chains within the crystalline regions (see figure 1.30 E). This model was widely accepted and later other groups observed the same structural changes in doped P3HT unit cell irrespective of the doping method. [85][99] Doping PBTTT with F4TCNQ results in quite similar effects on the crystal structure as for P3HT, namely, an expansion of the lattice along the alkyl side chains and a contraction of the lattice along the π-stacking. An initial investigation of the dopant induced structural changes in solution doped PBTTT was done by Cochran et al. in 2014. [79] Based on GIXD measurements, they proposed that the dopant can be intercalated between conjugated backbones within the π stacks of the polymer. A clear structural model was not proposed due to relatively low intensity of diffraction peaks. Same structural changes in PBTTT irrespective of doping methods. Reproduced from (59).

Watanabe and coworkers [101] proposed that F4TCNQ -anion mainly reside inside the alkyl side chain layers in the crystalline regions of PBTTT. In vapor doped C14-PBTTT thin films, Chapter 1: Fundamental concepts and state of the art 67 the presence of the anions in the amorphous part was not evidenced experimentally. [97][46] Later, more details on dopant induced structural changes and the orientation of the F4TCNQ -anion in PBTTT matrix was evidenced by the work of Hamidi-Sakr et. al [46,102] and by the work of Chabinyc and coworkers (see figure 1.33 A and D). [59] Similar to P3HT, they have observed an expansion of the PBTTT crystal lattice along the alkyl chain direction (d100) and a slight compression along the π stacking direction

Moreover, the F4TCNQ -anions orient perpendicular to the PBTTT backbone similarly to P3HT.

It is worth to mention that the real reason for the observed decrease in the π-π stacking is still controversial. Hamidi-Sakr et.al [46] proposed that it might be due to the reorganisation of the crystal structure due to the incorporation of the dopant molecules into the side chain layers. However, Liu and coworkers suggested another mechanism based on DFT calculation. [103] They proposed that the polaron delocalization between the adjacent polymer chains and the resulting attractive forces between them is responsible for the reduced π-π stacking distance in doped polythiophenes. The incorporation of the dopants into the side chain region contributes only 0.1 Å to the observed reduction in π stacking. The main contribution is coming from the attractive forces between the charged polymer chains.

It is important to mention that structural characterization methods such as electron diffraction and grazing incidence X-ray diffraction are only able to reveal the structural changes, that occur in the crystalline part of the polymers upon doping. Hence, the presence of dopant anions in the amorphous phase of the polymer is not accessible via TEM. However, our polarized spectroscopic measurements on PBTTT and P3HT doped from F6TCNNQ and F4TCNQ (see chapter 2 and 4) do not show any reduction in the absorption of the amorphous part (⊥ direction) upon doping. Some of the recent publication revealed the presence of the dopants in the amorphous part of the polymer.

Chapter 1: Fundamental concepts and state of the art 68 Hynynen and coworkers proposed the intercalation of Mo (tfd-COCF3)3 dopant into the amorphous part of aligned P3HT. Wide-angle X-ray scattering (WAXS) does not show any evidence of incorporation of this dopant in the crystallites. [94] Yee and coworkers recently demonstrated the dopant induced ordering in amorphous regions of regiorandom P3HT using dopant such as F4TCNQ and FeCl3. This can be considerd as an evidence for the presence of dopant in the amorphous part. [104] In conclusion, doping P3HT and PBTTT with F4TCNQ results in the intercalation of dopant molecules in the crystal lattice of the polymers and more precisely inside the layers of alkyl side chains. The presence of the dopant anions in the amorphous part of the polymer is still a question of debate but seems to depend strongly on the type of dopant.

TE properties of a few representative using P3HT and PBTTT

Table 1.1 is a non-exhaustive collection of recent results on doped polythiophene thin films (P3HT and PBTTT) including hybrid materials with inorganic nanoparticles, graphene or carbon nanotubes. Mixing of Polymers with suitable fillers like carbon nanotubes, Tellurium nanoparticles, Bismuth Telluride [105] nanoparticles and graphene [106] was also tempted to improve thermoelectric properties of thin films. For instance, Müller and coworkers obtained high conductivities of 10 3 S/cm and a PF of 25 µWm -1 K -2 in composite films of P3HT and single walled carbon nanotube (SWCNT). [107] The strong π -π interactions between the carbon nanotubes or graphene and P3HT ensures uniform mixing. [106,[108][109][110][111][112] FeCl3 and different iron-based salts were used for higher thermoelectric output because of the strong oxidizing nature of the dopant for P3HT. Ferric salts of triflimide anions (Fe (TFSI)3) [113] and Iron(III) p-toluenesulfonate hexahydrate (Fe +3 -tos3 .6H2O) [114] were also used for P3HT doping.

Chapter 1: Fundamental concepts and state of the art 69 al [114] P3HT CNT 345 ± 88 97±11 325±101 Hong et. al [112]

P3HT

(rr) Fe(TFSi3) Sequential (Solution) 96.1 42.2 17.10 Qu et. al [113] P3HT FeCl3 Sequential Solution SWCNT 2760 ± 170 31.1±2.1 267±38 Hong et. al [115] P3HT SWCNT 501 ± 32 37.5±1.1 71.8±5.2 Lee et. al [108] P3HT FeCl3 254 37.2 35 Hong et. al [116] P3HT I2 Vapor 4.7×10 -1 386 7 Zhu et. al [117] P3HT MWCNT 1.3×10 -3 131 Du et. al [109] P3HT Te nano wires 285 Yang et. al [110] P3HT Graphen e 1.2 35.56 0.16 Du et. al [106] P3HT FeTf 87.5 44 20 Müller et. al [107] P3HT MWCNT 0.11 11.3 0.0012 Du et. al [111] P3HT FeCl3 SWCNT 10 3 48 95 Müller et. al [107] P3HT F4TCNQ Solution 0.63 Glaudell et. al [72] P3HT F4TCNQ Solution 0.1 Yim et. al [71] P3HT F4TCNQ Solution 0.1 Kiefer et. al [98] P3HT F4TCNQ Solution 1 Aziz et. al [76] P3HT F4TCNQ Solution 1.8 Duong et. al [97] P3HT F4TCNQ solution 8 Jacobs et. al [96] P3HT F6TCNNQ solution 7 Karpov et. al [80] PBTTT F4TCNQ Solution 2 Cochran et. al [79] Chapter 1: Fundamental concepts and state of the art 70

It is worth to mention that most of the iron-based dopants and a combination of them with carbon nanotubes showed satisfactory thermoelectric performances. The solution doping of P3HT and PBTTT using TCNQ derivatives such as F4TCNQ and F6TCNNQ

produced highly stable and conducting polythiophene thin films. [71][72] l [76][97] l [98][96] [58 [99]

Conclusion

This literature survey indicates that enhancing TE properties of conjugated polymers is possible by improving the electrical conductivity of the polymer. This can be PBTTT F4TCNQ Solution 3.5±0.05 60±9 1.3±5.2 Chabinyc et. al [59] P3HT F4TCNQ Sequential (Solution) 5.5 Scholes et. al [99] P3HT F4TCNQ Sequential (Solution) 3 Jacobs et. al [96] P3HT oriented F4TCNQ Sequential (Solution) 22

Hamidi-Sakr et. al [46 al [101] P3HT FTS Sequential (Vapor) 27±0.1 60±9 10±3 Glaudell et. al [72] P3HT FTS Sequential (Vapor) 50±20 Podzorov et. al [118] PBTTT FTS Sequential (Vapor) 604±0.7 19±3 22±7 Glaudell et. al [72] PBTTT FTS Sequential (Vapor) 1100 Podzorov et. al [118] PBTTT F4TCNQ Sequential (Vapor) 220±0.02 39±5 32±9 Chabinyc et. al [59] PBTTT F2TCNQ Sequential (Vapor) 36±3 140±20 70±20 Chabinyc et. al [59] Chapter 1: Fundamental concepts and state of the art 71 achieved by polymer orientations technique or by selecting different doping methods and processing of polymers.

An extensive literature survey on doping methods and resulting TE properties provided some important guidelines to enhance thermoelectric performances in polymer-based systems.

1. Casting a neat semi-crystalline film that forms locally π stacked domains with a long-range conjugated backbone can provide excellent microstructure for efficient charge transport [59] .

2. The molecular dopant must be introduced into the polymer microstructure [86,94,98,101] (Vapor or Sequential doping) without disturbing the local order while maintaining or enhancing the long-range order of conjugated backbones. Preserving the orientation, crystallinity and improving the length scale of backbone order are two key methods to enhance thermoelectric properties.

Highly crystalline thin films prepared by high temperature rubbing show enhanced electrical conductivity along the rubbing direction. [34,46,77] Preliminary studies on sequential doping of oriented P3HT films gave important informations on the delocalization of polarons and the location of dopant anions in these systems. Based on these preliminary results, in the framework of my thesis, I will be working on orientation of PBTTT thin films. Once the polymers are aligned by rubbing, different doping methods and different doping agents will be used and correlations between structure and TE properties will be established. enhance on/off ratio and charge mobilities in organic field-effect transistors (OFET). [1,2] Doped interface layers can ease charge injection at metallic electrodes. [3] More recently, Crispin and coworkers demonstrated remarkable thermoelectric properties of conducting polymers such as poly(3,4-ethylenedioxythiophene) -tosylate (PEDOT-Tos) by fine-tuning the doping level in the thin films. [4] Another challenging system, of high interest to the scientific community, is regioregular poly(3-hexylthiophene) (P3HT) doped with 2,3,5,6-tetrafluoro-7,7,8,8-tetracyanoquinodimethane (F4TCNQ). As shown in chapter 1, doping impacts strongly the resulting charge conductivity in thin films.

Blending P3HT with F4TCNQ in solution results in poor charge conductivities. [5][6][7][8][9] Conversely, sequential doping of crystalline P3HT films by spin-coating or by dipping into a solution of F4TCNQ in an orthogonal solvent results in improved conductivities approaching 10 S/cm. [4][5][6][7][8][9][10] As a consequence, the charge conductivity in sequentially doped P3HT can be up to ten times higher than that of doped materials prepared from blends in solution.

Different mechanisms of doping have been considered: formation of a charge-transfer complex with partial charge transfer in quaterthiophene versus integer charge transfer in P3HT/F4TCNQ. [11,12] The crystallinity/morphology of the pristine polymer films prior to doping is also a key parameter that controls the final conductivities. In F4TCNQ vaporphase doped films, conductivities of 12.7 S/cm are common for highly crystalline P3HT [13] versus 220 S/cm for annealed PBTTT films. [14] Given the high charge transport anisotropy Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 86 of P3HT, alignment of the polymer films prior to doping is one additional and very effective method to enhance charge conductivity in doped thin films. [15,16] F4TCNQdoping of thin films oriented by high temperature rubbing yields highly oriented conducting polymer films with enhanced conductivities and thermopowers along the rubbing direction. [17] Transmission electron microscopy and polarized UV-Vis-NIR spectroscopy analysis of highly oriented thin films of F4TCNQ-doped P3HT demonstrated that the F4TCNQ -anions are located in the layers of alkyl side chains within crystalline P3HT domains with the molecular long axis of F4TCNQ -oriented perpendicular to the polymer backbone. [17] This result raises new questions about the doping mechanism in the majority of semiconducting polymers bearing solubilizing alkyl side chains. In particular, to what extent do the packing (interdigitated or not) and the length of alkyl side chains impact the doping kinetics and ultimate doping level. This is of high importance for the design of new TE materials as the choice of alkyl side chains may help achieve well-defined doping levels and thus to tune the TE properties of doped polymer semiconductors.

The present chapter focuses on a series of poly(2,5-bis(3-alkylthiophen-2-yl)thieno[3,2b]thiophene) (PBTTT) polymers that differ by the length of their linear alkyl side chains, from n-octyl to n-octyldecyl (C8, C12, C14, C18) (see figure 2.1). PBTTT is a semi-conducting polymer that exhibits high charge carrier mobilities (up to 1 cm 2. V . s). It has a layered structure composed of highly π-stacked backbones separated by layers of alkyl side chains. [18,19] PBTTT can readily be oriented by both blade coating or high temperature rubbing. [20][21][22][23] PBTTT with dodecyl side chains (C12-PBTTT) showed a very high dichroic ratio of more than 10 and high anisotropy in charge mobilities in thin films rubbed at 125°C. [21] Since doping in F4TCNQ/Acetonitrile (ACN) solution preserves the alignment of the oriented PBTTT films, the combination of rubbing and solution doping is an effective Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 87 method to prepare highly oriented conducting polymer films. This strategy is extended here to a family of PBTTTs with different alkyl side chains from n-octyl (C8) to n-octyldecyl (C18). To that aim, the doping protocol is identical for all polymers i.e. using acetonitrile as a solvent and fixing the F4TCNQ concentration to 1 mg/ml. Acetonitrile has been chosen as it is an orthogonal solvent for all PBTTTs. As anticipated, all investigated PBTTTs Structure-property correlations are drawn as a function of doping time and alkyl side chain length.

Results and discussion

Fabrication of highly oriented and conducting PBTTT thin films

As demonstrated in our previous work [17] , combining high temperature rubbing and doping of the oriented films from a solution of F4TCNQ in an orthogonal solvent such as acetonitrile affords highly oriented films of conducting polymers. The structure and charge transport properties of which are readily controlled by adjusting the dopant concentration fixed to 1mg/ml in this study [17] . Hereafter, we show that the doping level can also be controlled via the doping time i.e. the contact time between the oriented polymer films and the solution of dopant. (a) is apparent as a shoulder around 500 nm. [21] All PBTTTs show a dominant absorption of the F4TCNQ -anion polarized perpendicular to the backbone of PBTTT indicating that Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 89 the long axis of the F4TCNQ -anion is oriented essentially perpendicular to the polymer backbone. [17] For all investigated side chain lengths (C8 to C18), the dopant molecules are hosted within the layers of alkyl side chains. The small remaining absorption of F4TCNQ - anions for POL//R is either due to the mis-oriented domains of PBTTT or to a possible small tilt of the long axis of F4TCNQ -anions away from the normal to the polymer backbone. Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 90

Having set the basis of the orientation and doping procedures and outcome, we will successively consider in the following, the kinetics of doping, the integer charge transfer and doping level reached at saturation, the structural changes induced by doping and finally the evolution of TE properties versus side chain length. For POL//R, the main absorption peaks after doping correspond to the polaronic P1 and P2 bands with a fraction of undoped polymer (N). For POL⊥R, the spectrum is quickly dominated by the F4TCNQ -anion peaks (413 nm, 770 nm and 875 nm) beside a contribution from amorphous PBTTT that varies little with doping time. For all polymers, the absorption bands of the polaron (P1 and P2) and of the F4TCNQ -anion increase with doping time and saturate for long doping times (>10 min). The saturation indicates that a maximum doping concentration is reached for all polymers. For a very short doping time of 5 s, the polaronic bands and anion bands are already intense for C12-PBTTT whereas for C18-PBTTT no doping is observed. In the same time interval, the absorption anion located in the layers of alkyl side chains. [24,25] When the polaron and the anion are Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 92 placed at infinite distance, the P1 band is most red-shifted. Upon increasing Coulombic interactions, the P1 band progressively shifts to higher energy. Our observation of an apparent larger P1 energy for C8-PBTTT is therefore consistent with a shorter polaronanion distance in C8-PBTTT as compared to the other PBTTTs with longer side chains. A more quantitative analysis of the doping kinetics is obtained by plotting the intensity of the F4TCNQ -anion band at 875 nm as a function of time (see figure 2.5). The absence Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 93 of UV-Vis signature of neutral F4TCNQ in the films [26] indicates that each neutral F4TCNQ diffused into the polymer film is ionized almost instantaneously upon redox reaction in the polymer film. Accordingly, the kinetics of the absorbance of F4TCNQ -anions should reflect the diffusion of the ionized dopant in the polymer matrix. For all polymers, the absorbance of the anion intensity scales like √𝑡 for short doping times and it tends to saturate at long doping times(see figure 2.5 A). Each polymer is characterized by a typical doping time . However, this time constant  increases with film thickness. This is related to the dependence between the characteristic doping time and the film thickness l. In a diffusion-limited regime: 𝑙 = √𝐷𝜏. The finite film thickness is responsible for the saturation of the doping kinetics. (see Accordingly, plotting the normalized F4TCNQ -absorption against t 1/2 /l helps extract the diffusion coefficients of F4TCNQ -in the polymer matrix, D, for each PBTTT (In particular, it is possible to verify that two films with different thickness fall on a single master curve (see figure 2.5 B). [27] The time-dependence of the band intensity of the absorption peak at 875 nm, A875(t), was accordingly adjusted using the phenomenological equation:

𝐴 875 (𝑡) = 𝐴 0 (1 -exp (- √𝐷 𝑙 (𝑡) 1 2 )) (1)
Where A0 is the final absorbance reached at saturation, l is the film thickness and D is the diffusion coefficient of F4TCNQ -in the polymer. As seen in figure 2.5 B, the kinetics of doping of all polymers are very well described by equation (1). The extracted diffusion coefficients are collected in table 1. C12-PBTTT and C14-PBTTT show the highest diffusion coefficients of approx. 9 . 10 -12 cm 2 /s which is close to that reported by Maliakal for the diffusion of iodine into P3HT films. [28] In strong contrast, both C8-PBTTT and C18-PBTTT show significantly smaller diffusion coefficients of approximately 0.6 . 10 -12 cm 2 /s.

Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT melting/crystallization peaks for C18-PBTTT are consistent with those observed for the corresponding poly(3-alkylthiophene). [29] Accordingly, only C18-PBTTT shows evidence for a well-defined crystalline packing of side chains, contrary to the other PBTTTs. Thus, if C18 side chains are effectively packed to form a tight polymethylene sublattice, the inclusion of dopant molecules will be more difficult leading to lower doping levels and slower diffusion kinetics.

The detrimental effect of crystalline side chains on dopant intercalation is further investigated using hot solution doping. For this experiment, an oriented PBTTT thin film is doped by dipping in a hot solution of F4TCNQ/Acetonitrile mixture at a temperature of 70 C. This temperature is slightly above the melting temperature of C18 side chains.

The polarized UV-Vis-NIR spectroscopy showed a higher oxidation level and a higher intensity of P1 and P2 bands in samples doped at 70 C (see figure 2.6 B). This further supports our previous argument that highly crystalline side chains can have a detrimental effect on dopant diffusion.

The intermediate cases of C12 and C14 side chains seem most adapted to host the dopant molecules as the layers of side chains are only imperfectly packed. In the crystal structure of C14-PBTTT proposed by Cho et. al the area per alkyl side chain (stem) is 24 Å 2 versus 18 Å 2 for a polyethylene crystal. [30][31][32] This implies that there is substantial space for hosting dopant molecules in the imperfectly packed layers of C12 and C14 side chains.

The reason why the doping kinetics is also slow for C18-PBTTT is not yet clear. However, one may suppose that the side chain layers of C8-PBTTT are too disordered so that the F4TCNQ molecules cannot be well intercalated between alkyl side chains. This would also be consistent with the UV-Vis-NIR results showing that the F4TCNQ -anions are not well oriented in the layers of side chains (anion peaks Visible for both // and ⊥ polarization).

Overall, the results on doping kinetics give strong evidence on the important role of side Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 97 chain packing and length on the ultimate doping levels and doping kinetics. This is further supported by the dependence of doping level on alkyl side chain length as seen hereafter.

Estimation of doping level and integer charge transfer.

The oriented character of the doped PBTTT films helps to establish the saturation level of doping since the contributions of polarons and F4TCNQ -are dominant for light polarization parallel and perpendicular to the rubbing direction, respectively. Therefore, the concentration of F4TCNQ -can be more readily quantified as compared to nonoriented films for which multi-peak fitting is necessary to disentangle polaronic and F4TCNQ -contributions. [7] Regarding F4TCNQ -, the extinction coefficient  of the 0-0 absorption at 880 nm is known (= 50000 mol -1 L cm -1 ). [33] We also know the lattice parameters of the doped phases for all PBTTTs at saturation. Hence, it is possible to extract the ratio of concentrations of both F4TCNQ -anion and thiophene cycles in the doped films at saturation.

The doping level i.e. the number of F4TCNQ -anion per thiophene cycle was calculated at the saturation of the doping in solution from the UV-Vis-NIR absorption spectra for POL ⊥ R. To that aim, we have first determined the absorbance of the anion and used the extinction coefficient of F4TCNQ -reported in the literature for the 0-0 component (= 50000 L mol -1 cm -1 ) to obtain the corresponding concentration inside the thin films following the equation CF4TCNQ-=Abs(0-0)/ t where t is the film thickness determined from the absorbance of the undoped film. The absorbance of the 0-0 component of F4TCNQ -overlaps partially with the P2 polaronic band as well as the tail of the absorption of the amorphous undoped fraction of the polymer. Rather than performing an arbitrary multipeak fitting using several Gaussians to account for the underlying absorbance of P2 and amorphous fraction [7] , we decided to use a linear baseline between 615 nm and The amount of generated anion depends on the possibility of charge transfer between the PBTTT backbone and the neutral F4TCNQ molecules that have diffused inside the side chain layers. Charge transfer between PBTTT and F4TCNQ depends on the energetic difference between the donor's HOMO and the acceptor's LUMO.

Cyclic-voltammetric measurements [34] indicate that all four PBTTTs have very close HOMO levels (see figure 2.7 and table 2.4) and therefore, the slight differences in HOMO positions cannot account on their own for the observed differences in doping levels for the PBTTTs. We observe a decrease in doping level from the C12 to C18-PBTTT suggesting that the length of side chain layers is one important parameter influencing the final Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 99 doping level. However, C8-PBTTT shows a doping level quite similar to C14-PBTTT which suggests that another parameter must influence the final doping level. As shown previously for the doping kinetics, the diffusion coefficient D of F4TCNQ -in the layers of side chains depends also on the crystallinity/packing of the side chains. As demonstrated by DSC, the layers of alkyl side chains are particularly well and tightly packed for C18-PBTTT, which explains the slow and difficult diffusion of F4TCNQ for this Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 100 polymer, hence the low doping level. In strong contrast, the absence of crystallization/melting observed for C8-PBTTT indicates that the disordered packing of C8 side chains is also detrimental for the diffusion and intercalation of dopant molecules.

Finally, it is also worth mentioning the potential role of the solvent used for the doping.

Fujimoto et al. demonstrated that doping levels are also determined by the choice of solvents e.g. acetonitrile versus fluorinated solvents. [35] Different solvents can show different miscibilities with the polymers. In the case of acetonitrile, the miscibility in the polymer should decrease from C8 to C18-PBTTT. However, C8-PBTTT should have the highest miscibility with acetonitrile and does not show the largest doping level. In addition, the diffusion kinetics is slow in C8-PBTTT. Hence, the differences seen in doping level and kinetics with side chain length cannot be explained based on different extents of acetonitrile diffusion in the four PBTTTs.

It is further important to determine if the doping mechanism is typical of integer charge transfer as observed for F4TCNQ-doped P3HT. [12] The orientation of the PBTTT films helps analyze quite precisely and independently the time dependence of polaronic and F4TCNQ -bands (see figures 2.3 and 2.4). Accordingly, it is relatively simple to identify correlations between the absorbances of the PBTTT polaron and the F4TCNQ anion as expected in case of integer charge transfer between PBTTT and F4TCNQ. Low dose electron diffraction (ED) was used to investigate the structural changes upon doping for the oriented PBTTT thin films. Figure 2.9 collects diffraction patterns before and after doping (saturation regime) for four PBTTTs with increasing side chain length.

All electron diffraction patterns are characterized by a set of equatorial (h00) (h=1,2,3,4)

and a meridional 003 reflection indicating that they are mostly made of face-on domains. [20] A minority of edge-on domains is also present as indicated by the weak equatorial 020 reflection. For all polymers, ED confirms the previous UV-Vis-NIR and POM observations: doping in F4TCNQ-acetonitrile preserves the alignment and the original orientation of PBTTT crystals observed in the rubbed films. In particular, the intensity of the meridional (003) reflection along the PBTTT backbone direction is preserved upon doping for all four polymers without important in-plane broadening. This is consistent with previous observations on P3HT oriented films. [17] Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT For C12-PBTTT and C14-PBTTT, the changes in unit cell parameters occur within a few seconds and saturation of the unit cell parameters is observed after 1 min of doping, in Importantly, for t≥1min, there is no coexistence of doped and undoped phase: only one 100 reflection corresponding to the doped phase is observed. This is similar to previous observations for F4TCNQ-doped P3HT and indicates that the remaining absorbance of neutral C8-PBTTT seen in the UV-Vis spectrum must correspond to non-doped chain segments enclosed inside the crystals of the doped phase and not to chains within undoped PBTTT crystals. [16] In this regard, the doping of PBTTT with F4TCNQ is clearly different from the co-crystallization of PBTTT and PCBM that shows the coexistence of the pure PBTTT and the co-crystal phases. [32] For C12-PBTTT, expansion of the interlayer spacing is from

Anisotropy of thermoelectric properties

Doping oriented thin films of P3HT with F4TCNQ induces a substantial increase in power factors because both charge conductivity and thermopower are enhanced in the alignment direction. [17] In order to evaluate the benefits of alignment on the TE efficiency of doped PBTTTs, we have measured the evolution in charge conductivity and for non-oriented vapor-phase doped PBTTT. [14,36,37] For C12-PBTTT, the conductivity values tend to saturate very rapidly since the doping kinetics is very fast. The fact that C12-PBTTT reaches the highest conductivity is also consistent with the fact that this polymer shows the highest doping level as determined from the absorbance of the F4TCNQ -anion (see table 2. The relatively low conductivity of C8-PBTTT may have different origins. First, this polymer shows the most blue-shifted polaronic bands, suggesting stronger polaron localization due to stronger Coulombic interactions between the polaron and the F4TCNQ - anions. [24,25] Second, C8-PBTTT has a relatively low molecular weight. It is known that

Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 109 molecular weight significantly impacts structure and morphology thus the charge mobility in polymer semiconductors. [38] As a matter of fact, C8-PBTTT is the only investigated polymer that shows clearly a semi-crystalline lamellar morphology with a 22 nm periodicity much alike P3HT (see figure 2.14). [17] Analyzing the ratio ///⊥ for the four polymers helps identify the impact of side chain length on the charge transport anisotropy (see figure 2.16 B). In the direction perpendicular to the rubbing, the highest values of ⊥ are close to 35 S/cm for C12-PBTTT which results in anisotropies ///⊥ in the range 5-6. For C8-PBTTT, the anisotropy is close to 2. It is much larger for C18-PBTTT (typically 16-18). The trend in the charge transport anisotropy ///⊥ versus length of side chain is clear (see figure 2.16 B): the longer the side chains, the larger the anisotropy ///⊥. This is consistent with the fact that the rubbed films consist mainly of aligned face-on oriented crystals (see ED patterns in figure 2.9). In other words, charge transport measured perpendicular to the rubbing should be dominated by transport along the insulating alkyl side chains, which should be lowest for C18-PBTTT The Seebeck coefficients of all samples exhibit an important anisotropy in oriented thin films. The thermopower parallel to the rubbing S// is always larger than perpendicular to the rubbing S⊥. Anisotropy of Seebeck coefficient in oriented conducting polymers has been reported previously for doped polyacetylene (PA) and polyaniline. [39][40][41] In Iodine-doped PA, the anisotropy S///S⊥ was of the order of two ). The anisotropy in Seebeck coefficient parallels that of the charge conductivity. However, as already noted for rubbed films of P3HT doped with F4TCNQ, the anisotropy of S is always lower than that of the charge conductivity . [17] The Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 112

dependence
The simultaneous increase of charge conductivity // and thermopower S// along the rubbing direction results in a substantial increase of power factors PF=S 2  along the rubbing direction. As seen in figure 2.13 C, the values of PF are substantially enhanced along the rubbing direction and the anisotropies can be very high. They reach 66 for C12-PBTTT and 320 for C18-PBTTT. The maximum PF for the C8-PBTTT, C14-PBTTT and the C18-PBTTT are 2.4 µW . m -1. K -2 , 32.6 µW . m -1. K -2 and 20.6 µW . m -1. K -2 , respectively. However, C12-PBTTT surpasses the three other polymers, with power factors of the order of 100 µW . m -2. K -1 that are comparable to those observed for vapor-phase doped PBTTT. [36,37] For C12-PBTTT and C18-PBTTT, the PF shows an optimum versus doping time and tends to decrease slightly at long doping times, when the doping level increases beyond a critical concentration. This trend is well known in many systems such as PEDOT-Tos. [4]

Conclusion

The influence of alkyl side chain length on the F4TCNQ-doping mechanism of a series of . [1] Among all conducting polymers, iodine-doped polyacetylene (PA) demonstrated some of the highest conductivities -up to 10 5 S/cm at room temperature for stretchaligned films as well as very high thermoelectric power factors of 1.3 mW . m -1. K -2 . [2-4][5] However, PA could only be processed in the form of free-standing films several tens of microns in thickness that are difficult to use for TE device fabrication. To alleviate this lack of processing of pure conjugated polymers, alkyl side chains were introduced to design new semi-conducting polymers that are soluble in organic solvents, and thus easily processable via spin-coating, inkjet printing and other solution-based coating methods. [6,7] With the further control of regioregularity, highly ordered polymer semiconductors (PSCs) with charge mobilities beyond 1 cm 2 /V . s ideally suited for OFETs were designed. [7] In the subsequent decades, research efforts focused on to the way to fine-tune the electronic properties of the semi-conducting polymers used in the fabrication of organic light-emitting diodes, field-effect transistors and organic solar cells. [8] Chemical engineering helped identify the required structures of conjugated polymers necessary to tune their bandgap and/or their absorption spectra to match the solar spectrum. [9] Recently, new interest in conducting polymers was triggered by the pioneering work of Crispin et al. on thermoelectric (TE) applications based on the conducting polymer poly(ethylenedioxythiophene) (PEDOT) doped with tosylate (Tos). [10,11] Conducting polymers such as PEDOT-Tos have a layered structure of positively charged Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 122 poly(ethylenedioxythiophene) chains alternating with layers of tosylate anions. [11] PEDOT-Tos shows remarkable thermoelectric properties due to the combination of high charge conductivity, reasonable thermopowers and low thermal conductivity. [10,11] Despite promising TE properties, the crystallization of PEDOT-Tos in thin films is difficult to control. So far, no simple method is available to either orient in-plane the PEDOT chains or to select the preferential contact plane of crystals on a given substrate.

For a large part the synthetic method is to blame: it requires the use of in-situ polymerization techniques with low control over the polymerization process and finally the structure of the polymer thin films. The only improvements in charge transport were obtained via control of secondary crystallization e.g. by using co-solvents such as ethylene glycol that helps extract the non-conducting polystyrene sulfonate from the bulk of the PEDOT-PSS films. [12,13] In strong contrast, advanced control of crystallization, crystal orientation and alignment have all been demonstrated for polymer semiconductors (PSCs) such as regio-regular poly(3-hexylthiophene) (P3HT) or poly(2,5-bis(3-dodecyl-2-thienyl)thieno[3,2b]thiophene) (C12-PBTTT) (see figure 3.1A). [14,15] Numerous methods including selfseeded growth, blade-coating, epitaxy or high temperature rubbing help control precisely the crystal dimensions and their orientation in thin films, resulting in diverse morphologies: single crystals, spherulites, aligned crystalline lamellae. [14][15][16][17][18][19][20] It is therefore natural to take profit of the high structural control gained on polymer semiconductors such as P3HT or PBTTT to further improve their charge transport and thermoelectric properties after doping. Doping of thin films of P3HT or PBTTT with fluoroalkylsilanes (FTS) or F4TCNQ has led to high conductivities typically of the order of a few hundred to one thousand S/cm. [21][22][23] Vapor-phase doping of P3HT and PBTTT with F4TCNQ leads to conductivities of 12.7 S/cm and 200 S/cm, respectively. [22,23] Controlled Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 123 doping of well-crystallized and oriented conjugated polymer can further lead to anisotropic conducting polymer films. [24] Sequential doping of rubbed P3HT films with F4TCNQ afforded aligned conducting polymers with conductivities of the order of 22 S/cm whereas vapor phase doping of highly crystalline P3HT afforded 12.7 S/cm conductivities. [22,24] In the second chapter of this thesis, the impact of the alkyl side chains on the efficiency of doping, its kinetics and the ultimate TE properties, was investigated for the family of PBTTT with alkyl side chains from C8 to C18. [25] However, the obtained TE properties remain below the best conductivities reported for doped polyacetylene (PA) or even PEDOT-Tos. [11] The main aim of this chapter is to enhance the electrical conductivity and thermoelectric performance of polythiophenes by taking advantage of their high degree of chain orientation and crystallinity that can be maintained upon doping with strong inorganic Lewis acid-based doping agents such as FeCl3. We will mainly investigate the following issues:

i) Is it possible to reach the conductivities of doped PA close to 10 5 S/cm for oriented films of P3HT or PBTTT (see figure 3.1 A) after doping?

ii) What limits the ultimate conductivities?

To answer these questions, we have investigated the possibility to reach high doping levels in films of P3HT and PBTTT oriented by high temperature rubbing. [14,24,26,27] The choice of the two polymers is guided by their structural differences in terms of nanomorphology in rubbed thin films: P3HT has a periodic semi-crystalline lamellar morphology whereas C12-PBTTT shows a liquid-crystalline-like structure. [14,15] The choice Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 124 of FeCl3 is motivated by the fact that it is a stronger oxidant than F4TCNQ. [28] Accordingly, this chapter is divided into three sections dealing with:

1) The structural and spectroscopic features induced by FeCl3-doping on aligned P3HT and C12-PBTTT films.

2) The dependence of charge transport and thermopower on [FeCl3] concentration.

3) The analysis of the S- correlations in the oriented films of P3HT and C12-PBTTT. As demonstrated previously for P3HT and F4TCNQ [24] , the combination of an orientation method such as mechanical rubbing and sequential doping in an orthogonal solvent (see shows that doping does not suppress the orientation of the films, even for high FeCl3 concentrations (10mM) (see figure 3.1 D).

Results and discussion

Fabrication of highly oriented and conducting thin films

Structure and nanomorphology of doped thin films by Transmission electron microscopy

Electron diffraction and high-resolution techniques were used to investigate the dopant induced structural changes in the unit cell of PBTTT and P3HT. Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 127

For both polymers, doping results in a small decrease of the π-stacking periodicity similar to that observed with F4TCNQ doping. [24,25] The main difference between C12-PBTTT and P3HT resides in the evolution of the side chain spacing. In C12-PBTTT, the maximum change in d100 is equivalent to 3 CH2-CH2 bonds i.e. 3.9 Å whereas in P3HT it is only 1.7

Å. In both cases, the increase of d100 is smaller than the 7.4 Å van der Waals diameter of FeCl4 -. Even more telling is the variation in unit cell volume ∆V in both cases. Assuming orthorhombic unit cells with two chains per unit cell for both P3HT (a=16.5 Å, b=7.8 Å and c=7.77 Å) and C12-PBTTT (a=19.7 Å, b=7.4 Å, c=13.5 Å), ∆V=20-30 Å 3 for P3HT versus ∆V=230-250 Å 3 for C12-PBTTT. Intercalation of small molecules such as phenyl-C71-butyric acid methyl ester (PC71BM) into the structure of PBTTT has been extensively analyzed and modelled. [29][30][31] It has been shown that PC71BM molecules intercalate in the structure of PBTTT and form bimolecular crystals of PC71BM and C14-PBTTT. The bimolecular crystal phase is characterized by an expanded unit cell parameter along the alkyl side chain direction.

Since the van der Waals diameter of FeCl4 -(7.4 Å) is below that of PC71BM, it is tempting to propose a similar scenario for FeCl3-doped C12-PBTTT. Comparing the value of ∆V with the molecular volumes of FeCl2 (42 Å 3 ) and FeCl4 -(69 Å 3 ) suggests that up to four dopant molecules i.e. 2 FeCl2 and 2 FeCl4 -can be hosted in the unit cell of C12-PBTTT (see Figure

3.3.D)

. This would imply a stoichiometry of one FeCl4 -plus one FeCl2 per four thiophene rings in a unit cell at saturation (corresponding thus to an atomic ratio S/Fe=2). Such a stoichiometry is further consistent with that observed in the heavily doped phase of PEDOT:Tos. [32,33] To confirm this stoichiometry, we further analyzed our samples using EDX -STEM. distribution of both S and Fe atoms in the doped layers, indicating that doping is uniform in the samples. Second, the EDX spectra help determine the atomic ratio of S and Fe elements. For C12-PBTTT, one obtains a value close of 1.9±0.2 whereas for P3HT it is approximately 2.5±0.2. The S/Fe ratio for C12-PBTTT is in full agreement with the value deduced from the unit cell parameter variation assuming it is related to the uptake of 2

FeCl4 -and 2 FeCl2.

Corresponding energy dispersive X-Ray (EDX) spectroscopy (EDX) 2D maps for S K B) and for Fe K C). D) Bright field STEM image of an oriented P3HT film doped with 5mM FeCl3/nitromethane. Corresponding energy dispersive X-Ray (EDX) spectroscopy (EDX) 2D maps for S K E) and for Fe K F). G) and H) represent the EDX

spectra of C12-PBTTT and P3HT film doped with 5mM FeCl3/nitromethane respectively.

For P3HT, the situation is different as the unit cell expansion is moderate and does not reflect the uptake of dopants. The absence of important unit cell variation in P3HT suggests that the alkyl side chain layers of P3HT can accommodate more readily the intercalation of FeCl4 -and FeCl2 molecules possibly because the side chains of P3HT are Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 130 not interdigitated. However, for both polymers, the incorporation of dopants in the layers of alkyl side chains requires a structural reorganization especially when side chains are interdigitated (see the schematic illustration in figure 3.3.D).

Lattice expansion along the side chains, partial removal of interdigitation as well as tilting of the backbones and side chains are the potential mechanisms by which the structure of the doped phase can adapt to a given concentration of dopants. It is worth to mention that both the alkyl side chain sublattices of P3HT and of C12-PBTTT are particularly loose and offer the possibility to harbor dopant molecules. Indeed, the areas per alkyl stem are 24.5 Å 2 for C14-PBTTT and 27 Å 2 for P3HT (form I). Both values are substantially higher than the 18.4 Å 2 for the classical polymethylene subcell, indicating that the layers of side chains can indeed host small molecules such as F4TCNQ, FeCl4 -or FeCl2. [24,25,34,35]

Nanomorphology of the thin films

Both C12-PBTTT and P3HT show coexistence of crystalline and amorphous domains but their nano-morphologies are substantially different as revealed by bright field and HR-TEM (see figure 3.5).

Firstly, bright field TEM shows that only the P3HT films have a characteristic periodic lamellar morphology whereas no such morphology is observed for rubbed C12-PBTTT.

This difference is further visible in the HR-TEM images of figure 3

Low dose HRTEM image of FeCl3-doped P3HT A) and C12-PBTTT B) films aligned by high temperature rubbing ([FeCl3] = 4 mM). The insets correspond to the Fast Fourier Transforms. Schematic illustration of the nanomorphology in the C) semi-crystalline lamellar morphology of P3HT and D) LC phase of PBTTT

These differences in nano-morphologies are related to different persistence lengths of the two polymer chains: 3 nm for P3HT versus 9 nm for C12-PBTTT. [36,37] Accordingly, in strong contrast to the periodic semi-crystalline lamellar structure of P3HT, C12-PBTTT displays rather a liquid-crystal-like morphology with extended planarized chain segments along the chain direction (see illustration in figure 3.5 D&E). As a result, the connectivity between ordered domains is far better along the chain direction in C12- PBTTT films (see illustration in figure 3.5 E). As seen hereafter, these different nanomorphologies impact the ultimate conductivities of doped oriented films. suggesting an oxidation level of the polymer backbones similar to that of PEDOT-PSS (30-40%). [10] As a result, the films doped with 10 mM FeCl3 are almost optically transparent with a light blue color similar to PEDOT:PSS (excess FeCl3 gives rise to a small band at 374 nm). [38] The presence of the P2 band indicates also that the regime of pure bipolaron formation is not attained for [FeCl3]= 10 mM (for a nondegenerate ground-state polymer, the P2 transition is forbidden for the bipolaron. [39] P1 and P2 are more redshifted in C12-PBTTT than in P3HT: max(P2) = 809 nm for P3HT and 854 nm for C12-PBTTT.

Although the position of the P1 peak cannot be observed directly and lies beyond 2500 nm, the slope of the absorbance curve at 2500 nm indicates that the P1 peak is also substantially more red-shifted for C12-PBTTT than for P3HT. Positions of polaronic features are determined by the extent of delocalization and/or crystalline perfection. [40,41] Hence the red-shifted polaron bands in C12-PBTTT suggest a larger polaron delocalization length relatively to P3HT. This is consistent with the larger and C12-PBTTT is also doped with FeCl3. Indeed, upon doping with increasing [FeCl3], the absorption band of the amorphous phase (at 526 nm in C12-PBTTT and 506 nm in P3HT, (see figure 3.7 A&B) progressively disappears. This result is at variance with the case of F4TCNQ-doped films for which no evidence for doping of the amorphous domains was seen by polarized UV-vis spectroscopy (see figure 3.7 D&E). [24] The fact that both crystalline and amorphous phases are doped with FeCl3 explains, at least in part, the higher overall conductivity relatively to F4TCNQ. [24] Comparison with F4TCNQ-doped films indicates that the doping level achieved with FeCl3 is substantially larger. In particular, the remaining proportion of neutral C12-PBTTT and P3HT is much lower (only a few percent) for FeCl3 with respect to F4TCNQ. [18] Following the work by Miller and coworkers, the high oxidation levels at saturation of C12-PBTTT and P3HT upon doping with FeCl3 could reflect higher dopant miscibility as compared to F4TCNQ. [29] Moreover, the spherical shape of the FeCl4 -anions may also favor closer proximity of the dopant to the polymer backbone as compared to the rod-like F4TCNQ whose long axis tends to align parallel to the alkyl side chains i.e. in a plane perpendicular to the polymer Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 135 backbone. [24,25] As seen hereafter, the higher overall oxidation of P3HT and C12-PBTTT chains achieved upon FeCl3 doping accounts for the improved charge conductivities.

Comparison of the polarized UV-vis-NIR spectra of oriented C12-PBTTT films D) and P3HT E) after doping with F4TCNQ and FeCl3 for the incident light polarization perpendicular to the rubbing direction. Polaronic band P2 is well visible in the films doped with FeCl3 whereas in the films doped with F4TCNQ only the characteristic bands of the F4TCNQ -anion are visible (marked with arrows). The absorption of amorphous C12-PBTTT and P3HT is labeled

Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2

136

A and is well visible in the films after doping with F4TCNQ but not in the films doped with FeCl3 indicating that the amorphous phase is only doped in the case of FeCl3.

Anisotropy of charge conductivity.

As demonstrated in the previous section by TEM and UV-vis-NIR spectroscopy, the proportion of oxidized versus neutral chains in the oriented P3HT and C12-PBTTT films can be tuned by adjusting [FeCl3] while maintaining the high level of in-plane orientation.

This gives a handle to control the charge carrier density in the oriented films and to determine its impact on anisotropic charge transport and TE properties. Therefore, a sequential doping protocol was used to probe the changes in charge conductivity and thermopower on a given oriented film sample: the same oriented film was doped in successive steps with a solution of FeCl3/nitromethane of increasing concentration. This alleviates all issues related to differences in morphology and structure of the rubbed samples that would result in a scattering of the charge conductivity values  and Seebeck coefficients S. Figure 3.8 shows the evolution of  and S versus FeCl3 concentration for P3HT and C12-PBTTT oriented films that were sequentially doped with FeCl3 solutions of increasing concentration. Table 3.1 collects the highest values of conductivity and power factors for both doped polymers and corresponding values from the literature for iodinedoped PA.

Upon doping, both P3HT and C12-PBTTT show a strong increase of conductivity with a saturation at higher FeCl3 concentration (for both // and ⊥ directions). In the aligned C12-PBTTT films, the conductivity increases by several orders of magnitude from 150 S/cm at 0.1 mM to 2.2 10 5 S/cm at 5 mM. The highest conductivities of C12-PBTTT surpass the 10 5 S/cm reported for stretch-oriented bulk films of polyacetylene doped with Iodine. [4] However, it is worth to note that it is necessary to realize a progressive doping with increasing dopant concentration to reach such large conductivities. Indeed, C12-PBTTT Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 137 films doped directly with a 10 mM solution reach only a 10 4 S/cm conductivity. As a matter of fact, PBTTTs with shorter or longer alkyl side chains showed also very high conductivities (see table 3.2) albeit below those of C12-PBTTT. This specificity of C12-PBTTT was previously noted in the case of F4TCNQ-doped PBTTTs and reflected a higher doping level for this polymer. [25] For the highly conducting C12-PBTTT films, one can estimate the charge carrier density based on the structural data indicating that the unit cell contains two FeCl4 -anions per unit cell, which translates to a maximum of two charge carriers per unit cell i.e. a carrier density of approx. 9 . 10 20 cm -3 . [42] This value is larger than the 0.3-0.95 . 10 20 cm -3 reported for FTS-doped PBTTT. [43] However, the present result appears consistent in view of the UV-vis-NIR results indicating full oxidation of the C12-PBTTT backbone contrary to the FTS-doped PBTTT that showed a substantial amount of neutral polymer chains and an oxidation level of only 3.5-11%. [43] As for P3HT, alignment improves substantially the conductivity from 63 S/cm in non-oriented FeCl3-doped P3HT films to 570 S/cm along the rubbing direction in oriented films (for [FeCl3] = 5mM). [38] The anisotropies of charge conductivity are much larger in C12-PBTTT than in P3HT (///⊥=70 for C12-PBTTT and 12 for P3HT).

This difference is partly due to different film structures. As shown by TEM, the rubbed C12-PBTTT films have a dominant face-on orientation whereas P3HT has a mixture of edge-on and face-on oriented domains. As a consequence, ⊥ is essentially probed along the insulating C12 side chains of C12-PBTTT whereas transport along -stacking and side chains coexist in P3HT films. As an alternative explanation, the reduced anisotropy of the charge conductivity in P3HT could be due to the loss of orientation/crystallinity and/or to the doping of amorphous zones for increasing [FeCl3] as suggested by polarized UV-Vis spectroscopy. As a rule, the maximum conductivity of doped P3HT films is always below that of C12-PBTTT. It does not reflect a difference in the degree of in-plane orientation since both polymers have similar high dichroic ratios prior to doping. It is also not due to different Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 139 doping levels since the spectroscopic features for 10 mM doped films of P3HT and C12-PBTTT are almost identical. Even in the direction perpendicular to the rubbing, C12-PBTTT films display particularly high conductivities of  2100 S/cm that are close to those observed in non-oriented FTSdoped C14-PBTTT. [21] This indicates that the conductivity measured in non-oriented PBTTT samples is essentially limited by the charge transport perpendicular to the chains.

Achieving such high conductivities in doped C12-PBTTT on a macroscopic scale implies also that charges must be transported quite effectively between the well doped and ordered domains seen in HRTEM. Charge tunnelling between highly conducting domains is the mechanism invoked in a granular metal model of Sheng to account for comparably high conductivities in heavily-doped PA. [44] In strong contrast, in P3HT, the large extension of amorphous interlamellar zones will hamper tunnelling between doped crystalline regions and limit the macroscopic conductivity. This difference in the Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 141 electrical conductivity can be also related to a different hopping mechanism in highly conducting PBTTT versus P3HT. Based on temperature-dependent conductivity measurements on FTS doped PBTTT and P3HT, Tanaka et. al proposed that a 2D variable range hopping accounts for transport in PBTTT films while a 1D process is at play in doped P3HT films. [43] The same situation might be observed in the present case for FeCl3 doped films.

Anisotropy of the Seebeck coefficient.

An enhancement of electrical conductivity along the direction of alignment was reported on most studies on stretch aligned polyacetylenes, polyanilines and polythiophenes. [3,45,46] However, there was no clear trend on the effect of solid-state drawing or stretching on Seebeck coefficient and its anisotropy. For instance, recently stretch aligned thin films of P3HT upon doping with molybdenum tris(dithiolene)

showed an anisotropy in electrical conductivity. However, the Seebeck coefficient was not affected by stretching. [47] The situation was different in the case of oriented thin films aligned by high-temperature rubbing. We have seen that the Seebeck coefficients and electrical conductivities were enhanced in the direction of rubbing and they were highly anisotropic. [24,25] Unlike the other studies, the simultaneous enhancement of both electrical conductivity and Seebeck coefficient along the rubbing direction enabled us to obtain improved power factors parallel to the chain direction.

As seen in figure 3.8 B and E, the thermopower variation with dopant concentration is inverse to that of charge conductivity: both Seebeck coefficients S// and S⊥ decrease with [FeCl3]. In all samples, the thermopower probed along the chain direction is larger than that measured perpendicular: S// > S⊥. [23] The anisotropy factor of the Seebeck coefficient S///S⊥ tends to decrease with increasing FeCl3 concentration for both polymers. The anisotropies are higher for C12-PBTTT (4.6) than for P3HT (2.4). Overall, as observed Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 142 previously for F4TCNQ-doped P3HT, the anisotropy in thermopower is always smaller than that of the charge conductivity. [24,25] The samples with the highest conductivities have very small thermopowers (see table 3.1). For instance, the rubbed C12-PBTTT sequentially doped with 5 mM

FeCl3/nitromethane have //=2.2 10 5 S/cm and S//=9.4±0.5 µV/K. Their Seebeck coefficients measured perpendicular to the alignment are even smaller with S⊥=1.5 ± 0.5 µV/K. Such low thermopowers are characteristic of metals. [48] For comparison, oriented films of PA doped with iodine have S//=15 µV/K. [3,4] At this point, we focus on the origin of the anisotropy of the thermopower S in oriented doped films of P3HT and C12-PBTTT. The present situation bears similarity with the case of doped oriented PA fibres that show S///S⊥>1. [49][50] Kaiser proposed a heterogeneous model of metallic fibrils separated by electrical barriers to explain this observation. [50] In doped PA, semi-conducting barriers alternate with metallic domains and modulate the thermopower and its anisotropy. The situation is similar for doped oriented P3HT and C12-PBTTT (see figure 3.9). Indeed, HRTEM evidences the semi-crystalline morphology of doped P3HT and C12-PBTTT with a clear coexistence of highly ordered domains and amorphous interlamellar zones. Moreover, it is known that the charge conductivity is higher in crystalline versus amorphous domains of doped P3HT (typically a few S/cm are obtained in F4TCNQ-doped regio-regular P3HT versus 0.01 S/cm in regio-random amorphous P3HT). [22] This difference in charge transport between amorphous and crystalline domains must translate into a corresponding difference of the thermopower Sa and S for amorphous and crystalline domains, respectively, with Sa > S. Accordingly, the semi-crystalline morphology of oriented conducting polymer films can be analyzed by using a heterogeneous material model similar to that proposed for doped PA.

Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 143

When the temperature gradient ∆T is oriented parallel to the rubbing 9 (see figure 3.9

A), it is distributed over a succession of crystalline and disordered (amorphous) domains connected in series. We consider that the sample consists of a fraction  of crystalline domains. The crystalline domains are characterized by an average temperature gradient ∆T (along the chain direction) and disordered/amorphous domains by an average temperature gradient ∆Ta. We also assume that the heat current flows along the same path as the charges, so that the temperature gradients across crystalline and amorphous domains are proportional to their thermal resistances (∆TW and ∆Ta  Wa with W and Wa the thermal resistances of crystalline and amorphous domains, respectively).

Hence, the Seebeck coefficient in the rubbing direction S// can be written as: [46] S//=S +(1-) Sa ( 1)

Where  is given by:

𝛼 = 𝜒𝑊 𝜒 𝜒𝑊 𝜒 +(1-𝜒)𝑊 𝑎 (2)
In the case of P3HT, the crystallinity  is determined by the rubbing temperature and is close to 50% for TR= 200°C. [26] In semi-crystalline polymers such as polyolefins, the thermal resistance of crystalline lamellae along the chain direction W is small whereas the one perpendicular to the chains is significantly larger and comparable to that of the amorphous phase i.e. W<<Wa. [51] This implies S//Sa i.e. the Seebeck coefficient S// is dominated by the contribution from the doped amorphous and less conducting phase.

Let us now consider what happens when the temperature gradient is oriented perpendicular to the chain direction (see figure 3.9 B). In this situation, crystalline and disordered zones are put in parallel along the temperature gradient and charges are Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 144 short-circuited through the most conductive domains i.e. along the crystalline domains.

As seen in TEM of the rubbed 180°C P3HT sample, the crystalline lamellae make a quite continuous conduction path given their lateral extension perpendicular to the rubbing direction (several hundreds of nm) and the way such lamellae are interconnected.

Accordingly, charges will be tunnelled out of the sample within the crystalline domains in the direction perpendicular to the rubbing and this implies that the observed Seebeck coefficient S⊥ is dominated by that of the most conductive lamellae and not by the disordered (amorphous) domains. This implies that S⊥Sa. As a consequence, the Seebeck coefficient anisotropy becomes S///S⊥Sa/S>1. The present analysis indicates that a nanostructured morphology in conducting polymers such as doped P3HT or PBTTT that results in a modulation of electronic properties can be beneficial in oriented films since both the thermopower S// and the conductivity // are increased in the direction of polymer chains. It is important to recall that one major strategy to improve TE properties in inorganic systems is also based on material nano-structuring. Indeed, the objective of nano-structuring inorganic materials such as Bi2Te3 is to break the heat transport in the materials through enhanced scattering of phonons at multiple interfaces between domains without breaking charge transport. [52] The resulting power factors for the most conductive C12-PBTTT films are very high and can reach values close to 1.94 mW . m -1. K -2 , a value slightly larger than for I2-doped PA (1.35 mW . m -1. K -2 ) (see figure 3.8 C and F). Such power factors are quite large compared to those of non-oriented films of C12-PBTTT doped with F4TCNQ from the vapor phase (0.1 mW . m -1. K -2 ). The combination of orientation and higher oxidation of the C12-PBTTT polymer after FeCl3 doping accounts for the remarkable power factors obtained in rubbed thin films. Previous UV-vis studies indicate that a high proportion of neutral polymer chains is present in oriented P3HT and C12-PBTTT films after doping with F4TCNQ. This implies a high proportion of semi-conducting chain segments and domains that prevents reaching similarly high conductivities as for FeCl3 doping.

Correlations between thermopower and charge conductivity in oriented conducting polymer films.

Kang and Snyder have recently demonstrated that details of microscopic charge transport can appear through the correlations between S and . [53] A general theoretical frame relating to the S=f() scaling law was recently proposed by these authors. They predicted a scaling behavior with an exponent -1/s where s is a function of the charge Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 146 transport mechanism. [53] In PEDOT:TOS the S- curve is well fit with s=1 pointing at a transport mechanism dominated by acoustic phonon scattering. The exponent -1/3 with s=3 is obtained for a transport mechanism involving scattering of carriers by ionized . [54,55] As a consequence, the power factor PF  1/2 (see figure 3.

A).This relationship indicates that the thermoelectric performance of a polymer can be enhanced by improving the conductivity

In our experiments, the set of S, data obtained over a large range of conductivities (1.0-10 5 S/cm) helps further probe the S=f() correlations and its anisotropy especially in a regime of high conductivities ( > 10 3 S/cm). This methodology alleviates the experimental difficulties associated with charge carrier determinations in doped films e.g. using AC Hall measurements, especially for such aligned films. The S- plots for directions parallel and perpendicular to the chains are shown in fig. 3.8.B and C for C12-PBTTT and P3HT, respectively. For both polymers, the S = f() plots show different behavior parallel and normal to the chain directions. In the direction parallel to the chains (R), the correlation between S// and // is well described by a power-law with an exponent -1/4. This is consistent with previous results by Chabinyc and coworkers for F4TCNQ-doped C14-PBTTT. [54,55] In the direction perpendicular to R, both polymers show a deviation from the power law, especially for high values of the conductivity. The conductivity ⊥ is observed to saturate close to 200 S/cm in P3HT and 2000 S/cm in C12-Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2

147

PBTTT. The departure from the scaling law in the case of S⊥-⊥ can have different origins.

First, it could be related to some intrinsic structural change in the films when the doping concentration becomes important. We observe that the dichroic ratio for the P1 band as well as the intensity of the reflections in the ED pattern of doped P3HT films tend to decrease when [FeCl3] approaches 10 mM. A disordering of the layers of -stacked backbones may be induced by the introduction of increasing amounts of FeCl3 in the layers of side chains for both P3HT and C12-PBTTT. This disordering may account for a "saturation" of the conductivity along the chain direction for high FeCl3 concentration. Alternatively, the departure from the power-law may indicate a different transport mechanism in the direction perpendicular to the polymer chains. As a matter of fact, attempts to fit the S⊥-⊥ plot indicate that S⊥ scales rather well with ln(⊥). A similar dependence was reported by Mateeva et al. for oriented polyaniline films for which the S- relation was fitted with the equation: [56] 𝑆

𝑖 = -1 𝛽 ( 𝑘 𝐵 𝑒 ) 𝑙𝑛 ( 𝜎 𝑖 𝜎 𝑖 𝑚𝑎𝑥 ) (3)
with kB the Boltzmann constant, e the charge of the electron and β a dimensionless coefficient. i refers to the direction either parallel (//) or perpendicular (⊥) to the chain direction. [56] When fitted with equation ( 3 direction and perpendicular to it, supporting different charge transport mechanisms. It is worth to mention that Kroon [57] et. al observed a departure from the S -1/4 law for  >1S/cm in non-oriented films of a oligo ethylated -polythiophene. We attribute this behavior to a mixing of the transport phenomena parallel and perpendicular to the polymer chains.

Interestingly, the same scaling relations (S -1/4 ) is observed for F4TCNQ-doped C12-PBTTT (figure3.8 D), indicating that they are apparently independent of the chemical nature of the dopant. The similarity of the behavior for FeCl3 and F4TCNQ further supports the fact that the high conductivity in FeCl3-doped films is not related to a contribution from ionic conductivity that should result in a departure from the S -1/4 scaling law. To the best of our knowledge, it is the first time that such a difference between the S- correlations is evidenced along and normal to the chain direction. As a matter of fact, different S- correlations along and normal to the chain direction would further explain, at least in part, the large spread in experimental data in reference 48 and 49 as the measurements concern isotropic films for which both contributions parallel and normal to the chain direction can combine to different extents.

Conclusion

Controlling crystallization and orientation of polymer semiconductors prior to doping gives a unique handle over the charge transport and thermoelectric properties of the obtained conducting polymer films. This strategy is versatile and can be applied to numerous doped polymer semiconductors including n-type polymers [58] or low-bandgap polymers [59,60] that can be oriented by different methods such as blade-coating, high temperature rubbing or epitaxy and subsequently doped using different types of dopants and doping methods. Central to this approach is the preservation of high order present in the pristine polymer semiconductors upon doping. This is possible when the Two different charge transport behaviors are evidenced in the directions parallel and normal to the polymer chains: S// -1/4 prevails in the direction parallel and S⊥ln() in the direction normal to the polymer chains. Regarding TE applications, the combination of effective chain alignment and controlled oxidation by sequential doping helps to reach a record power factor of 2 mW . m -1. K -2 that could result in high ZT factors provided the thermal conductivity is well contained along the chain direction. Such a result is important as it sets a new upper limit for the ultimate TE power factor of such conducting polymers, but, it will be essential to further address the important issue of thermal conductivity and its correlation to the charge conductivity to obtain TE polymers with outstanding performances.

Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 151

In the next chapter, we will focus on doped C12-PBTTT films with F6TCNNQ that is leading to a more stable doped state of the polymer films.

Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2

Chapter 4. Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F 6 TCNNQ and F 4 TCNQ F4TCNQ doping of organic semiconductors such as P3HT and PBTTT has been widely documented in the recent literature. This chapter focuses mainly on doping of C12-PBTTT using a new doping agent, namely F6TCNNQ. Unlike F4TCNQ, F6TCNNQ doping of polythiophenes has been marginally studied. This dopant is chosen because of its deeper LUMO energy level of -5.3 eV, versus -5.1 eV for F4TCNQ that can enhance charge transfer between C12-PBTTT and F6TCNNQ. Oriented C12-PBTTT thin films are prepared by high temperature rubbing and sequentially doped with F6TCNNQ using a new doping method, called "incremental concentration doping" (ICD). The ICD method implies that a given sample is doped by dipping it successively in dopant solutions of increasing concentration. ICD of C12-PBTTT using F6TCNNQ leads to record electrical conductivities of 2430 S/cm and thermoelectric power factors (PF) 530±200 µWm -1 K -2 . Moreover, this method leads to a higher electrical conductivity than the conventional doping method (direct doping). This study demonstrates that a progressive intercalation of dopants in PBTTT crystals using the ICD method helps preserve the high level of order initially present in the aligned C12-PBTTT films.

The diffusion coefficients of both F6TCNNQ and F4TCNQ anions have been determined using the doping kinetics followed by UV-Vis NIR spectroscopy. A lower diffusion coefficient is observed for F6TCNNQ anions, demonstrating the size dependence of diffusion coefficients into the polymer matrix. The longer and bulkier the anion, the lower its diffusion coefficient. Polarized UV-Vis-NIR spectroscopy on these doped thin

Introduction

Doped polymers are of central interest in plastic electronics not only for their charge transport properties used for instance in polymer injection layers (PEDOT: PSS) but also for their thermoelectric properties as demonstrated by Crispin and coworkers. [1] Doping polymer semi-conductors such as regioregular poly(3-hexylthiophene) or PBTTT with acceptor molecules such as F4TCNQ is a common strategy to transform semi-conducting films into conducting ones. [2][3][4][5][6] However, to control their TE properties, it is essential to tune the charge carrier density that is determined by the method of doping and the type of dopant for a given polymer. In particular, the power factor PF that determines the effectivity of a doped polymer for TE properties has been shown to scale with conductivity such as PF  . Enhancing TE properties is therefore achieved by improving charge conductivity of conducting polymers. Various groups have therefore focused on the doping method as a handle to reach improved TE performances in doped polymer films. [7][8][9][10][11][12] Sequential doping, i.e. the doping of a highly crystalline polymer semiconductor films with a dopant solution in an orthogonal solvent is an effective method to enhance the TE properties as compared to doped polymer films prepared from mixed polymer and dopant solution in a common solvent. [9,10,13,14] Vapor phase doping leads to even higher charge conductivities since it seems to preserve order along the chain direction in the doped polymer crystals. As demonstrated in the previous chapters, dopants are usually intercalated in the layers of alkyl side chains of the polymers [15,16] . Therefore, the amount of dopants incorporated in a thin polymer film depends strongly on the length and packing of alkyl side chains as demonstrated in chapter 2 for PBTTT. The final thermoelectric properties of the doped films can further be enhanced if the films are oriented. Both, thermopower and charge conductivity are enhanced in aligned P3HT and PBTTT films along the chain direction, making it possible to improve substantially the TE power factor of the layers.

Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ

161

The choice of the dopant is also of paramount importance. The LUMO level of the acceptor molecules must be well situated with respect to the HOMO of the p-type polymer. The higher electrical conductivity observed in F4TCNQ doped polythiophenes is mainly due to the favourable energetical offset between the energy levels of the polymer and F4TCNQ. For instance, in F4TCNQ-doped P3HT and PBTTT, the HOMO of these polymers is in the range -4.8 eV -5.2 eV and the LUMO of the dopant is around -5.24 eV. Besides F4TCNQ, various strong inorganic dopants such as I2, FeCl3, WO3 have also been used to dope the conducting polymers. In chapter 3, we have shown that FeCl3-doping of C12-PBTTT leads to very high conductivities above 10 5 S/cm. However, FeCl3 and the other cited dopants are difficult to use because of both their high instability in air and corrosive nature. [14,17,18] . Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ

162 It is therefore essential to find alternative dopants to both F4TCNQ and FeCl3. Instead of using inorganic based doping agents, electron-deficient molecular dopants such as TCNQ and derivatives such as 2-(perfluoronaphthalene-2,6-diylidene)-dimalonitrile (F6TCNNQ) (see figure 4.1C) were recently introduced for the doping of conducting interlayers in electronic devices [19,20] .

Karpov [21] and coworkers first reported the use of F6TCNNQ to dope effectively P3HT.

They reported a high electrical conductivity of  7 S/cm, which was three times higher than that of F4TCNQ-doped P3HT. The reason for the higher electrical conductivity was not demonstrated with clarity. They attributed the higher electrical conductivity to a combination of higher charge carrier mobilities and higher doping strength of F6TCNNQ (-5.37 eV for F6TCNNQ vs -5.24 eV for F4TCNQ) leading to a higher charge carrier concentration (see figure 4.1). Moreover, they proposed that similar to F4TCNQ doping of P3HT, the charge transfer between P3HT and F6TCNNQ is of integer charge transfer type. [22] Given the proposed context and state-of-the-art (chapter 1), the present chapter aims at investigating the following important issues:

-To what extent does the doping of C12-PBTTT with F6TCNNQ differ from that with F4TCNQ in terms of structure, doping kinetics and resulting TE properties?

-How different is the intercalation of F4TCNQ and F6TCNNQ in the layers of alkyl side chains of C12-PBTTT?

-Is there any relationship between the degree of order of the dopant anions in the polymer semiconductor and the resulting electrical conductivities?

-Is there a way to introduce bulky dopants such as F6TCNNQ in the crystals of C12-PBTTT so as to preserve best the pristine structure of the undoped polymer films?

Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 163 So far, there is no clear evidence of the effect of F6TCNNQ doping on the crystal structure of polythiophenes. Moreover, the position and the orientation of the F6TCNNQ -anion in doped polythiophenes such as PBTTT and P3HT is not well understood. In this chapter, we investigate the effect of F6TCNNQ -and F4TCNQ-doping on the crystal structure and TE properties of oriented C12-PBTTT. Polarized UV-Vis-NIR spectroscopy and low dose electron diffraction were used to uncover the location and orientation of the dopant anions. We demonstrate that F6TCNNQ -anions orient perpendicular to aligned C12-PBTTT chains, as observed for F4TCNQ -. The angle-dependent polarized UV-Vis-NIR spectroscopy helped determine the angular distributions of both the dopant anions and polarons in oriented C12-PBTTT and uncover the degree of orientation of dopant anions in the oriented C12-PBTTT films.

It is expected that the dopant size is an important factor that determines the intercalation of the dopant anions into the interdigitated layers of the alkyl side chains of PBTTT. The perpendicular orientation of the dopant anions enabled us to follow their individual diffusion kinetics into the layers of the interdigitated alkyl side chains of C12-PBTTT [16] . A correlation is observed between the dopant size and diffusion coefficient.

Finally, we propose a general method to prepare highly conducting oriented and crystalline C12-PBTTT thin films by a combination of high temperature rubbing followed a modified doping method i.e. incremental concentration doping (ICD). The ICD method implies that a given sample is doped by dipping it successively in dopant solutions of increasing concentration (see figure 4.2). On the other hand, in the conventional doping process (single doping), a film is directly dipped into the solution of at the final concentration of dopant. The changes in the crystal structure of the thin films were followed by low dose electron diffraction for both direct and incremental concentration doping. To the best of our knowledge, this is the first evidence on the impact of two Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 164 different doping methods with the same doping agent on the crystal structure and TE properties of PBTTT thin films. Promising thermoelectric performances were observed in oriented C12-PBTTT thin films after incremental concentration doping and record conductivities of 2430±500 S/cm were observed along the chain direction in F6TCNNQ doped PBTTT thin films.

2. Results and discussion.

Fabrication of highly oriented conducting PBTTT thin films

It has been shown in the previous chapters that polymers such as PBTTT can be easily oriented by rubbing at high temperature [23] . Highly oriented crystalline polymer films of C12-PBTTT with dichroic ratio7 were obtained by rubbing at 125 C. Herein we introduced a new doping procedure, which is called "incremental concentration doping" (ICD). It is a modified doping version of classical sequential doping (see figure 4.2). In ICD, the oriented thin films are doped by dipping a given sample successively in solutions of dopants of increasing concentration. This is different from the conventional sequential doping (we call it direct doping), that implies that each thin film is dipped a single time in a solution of given concentration (see figure 4.2). In this study, all the thin films are doped by using ICD.

Polarized UV-Vis-NIR spectroscopy

Polarized UV-vis-NIR spectroscopy is very helpful in determining the amount and orientation of dopant molecules such as F4TCNQ and F6TCNNQ in the doped and oriented thin films of C12-PBTTT. Let us first present the spectroscopic signatures of the two dopants. Both F4TCNQ -and F6TCNNQ -anions display distinct signatures in the form of vibronic structures as seen in figure 4.3. In order to confirm the presence of the F6TCNNQ -anions in the thin films of PBTTT, we prepared a solution of pure F6TCNNQ - anions by reacting F6TCNNQ with Ferrocene in DCM, the corresponding UV-Vis-NIR spectrum is shown in figure 4.3, that shows the spectra of both neutral and F4TCNQ -.

F6TCNNQ -anions show characteristic bands at 1143, 976 and 987 nm and there are no corresponding absorption bands in that spectral range for neutral F6TCNNQ. The latter is characterized by two bands at 481 and 454 nm that do overlap partially with some Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 166 bands of the F6TCNNQ -anion. As seen in figure 4.3B, the vibronic structure of the F6TCNNQ -anions in the doped C12-PBTTT matches quite well that of the anion in solution.

There is only a small redshift in the thin films (1157 nm versus 1143 nm in solution).

Having identified the spectral signatures of the F6TCNNQ -anion, we can now analyse in more details the polarized UV-vis-NIR spectra of the doped films. The polarized UV-Vis-NIR spectra of both F6TCNNQ and F4TCNQ doped C12-PBTTT are highly informative. As expected, the polaronic bands P1 and P2 are polarized along the rubbing direction and their intensities increase with doping concentration and saturate once the concentration reaches 0.5 mg/ml. Even for very low doping concentrations such as 0.01 mg/ml, the P1 and P2 bands were present. Contrary to the polaronic bands, the absorption bands of the neutral polymer (N) decreased with increasing dopant concentration. As expected, a higher reduction in the absorption of neutral polymer in F6TCNNQ doped C12-PBTTT indicates higher oxidation of polymer chains. This observation further confirms that the deeper LUMO level of F6TCNNQ (-5.37 eV) compared to -5.24 eV in F4TCNQ is more favourable for charge transfer.

The polarized UV-Vis-NIR spectra along the perpendicular direction are different from the spectra along the parallel direction. The highly polarized spectra along the perpendicular direction provide fruitful information regarding the location and orientation of the dopant anions in the C12-PBTTT crystals. As observed in our previous works [15,16] and previous chapter 2, both F4TCNQ -and F6TCNNQ -anion bands are polarized perpendicular to the rubbing direction. For the films doped with F6TCNNQ, the anion bands are similar to those observed for the anions in dichloromethane (DCM) (see figure 4.3).

For both F4TCNQ and F6TCNNQ dopants, the intensity of the anion bands increased with doping concentration and saturated for C= 0.5-1.0 mg/ml. The polaron bands P1 and P2

Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 167 show a similar saturation for POL//R. Interestingly, no absorption from the polarons is seen for POL ⊥R.

Angular distribution of dopant anions and polarons

To further understand the degree of order and distribution of the dopant anions and polarons in the oriented C12-PBTTT, angle-dependent polarized UV-Vis-NIR spectroscopy was used. Oriented C12-PBTTT thin films were prepared and doped by dipping in 1 mg/ml 2 mg/ml 1 mg/ml 0.5 mg/ml 0.1 mg/ml 0.01 mg/ml As rubbed

F 4 TCNQ - A) B) C) D)
Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 169 of the corresponding dopant solutions. Then the angle-dependent UV-Vis-NIR spectroscopy was recorded by changing the polarization angle of the incident light with respect to the direction of rubbing from 0 to 90 with an interval of 15 for both F4TCNQ and F6TCNNQ doped PBTTT thin films (see figure 4.5C and 4.5D).

angles (every 15°) between the direction of polarization and the direction of rubbing of the film. At 0°, the incident light is parallel to the direction of rubbing (POL // R) whereas at 90° it is perpendicular to the rubbing direction (POL ⊥ R). E) Example of deconvolution of the UV-Vis-NIR spectra for C12-PBTTT thin films doped with

F6TCNNQ when the incident polarization angle is 45. The spectra on top (blue) result from the overlap of the anionic contribution at 1158 nm, with the polaronic (P2) and N components. The spectra in red contain the sum of all components beneath the anion bands. The actual absorbance of the anion is obtained by subtraction of the red baseline from the initial spectra.

The angular distribution studies on oriented polymers before and after doping helps to determine the in-plane orientation of the aligned polymer chains, polarons and the dopant anions. Therefore, the angle-dependent UV-Vis-NIR spectroscopy on oriented polymer films provides a better understanding of the degree of orientation of both F4TCNQ -and F6TCNNQ -in the layers of the alkyl side chains of PBTTT.

For an incident polarization of 0(POL // R), a very small contribution of both dopant anions F4TCNQ -and F6TCNNQ -bands is observed. For increasing polarization angle, their intensities increase whereas the contributions from the polarons (P1 and P2) decreased.

Both F4TCNQ -and F6TCNNQ -anions showed their maximum absorption when the incident light is polarized perpendicular to the rubbing direction (90) while the polaronic bands have a maximum at 0. The plot of the intensities of absorbance of the neutral polymer, polarons and the dopant anions as a function of incident polarization angle help determine the in-plane angular distribution of the polarons and anions in oriented C12-PBTTT. The overlap of the anionic bands (873 nm for F4TCNQ and 1158 nm for F6TCNNQ) with the broad polaronic contribution made it difficult to evaluate the true absorbance of the anions. Hence, it was necessary to deconvolute the as-obtained spectra in order to extract the actual absorbance of the anion by subtracting the contribution from the polaron P2 (see figure 4.5E). The angular dependence of the Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 171 F4TCNQ -and F6TCNNQ -absorption after the subtraction of the P2 contribution is represented in figure 4.6 C and D, respectively. The figures 4.6C and D depict the angular distributions of the absorption for the polarons, F4TCNQ -and F6TCNNQ -anions in oriented C12-PBTTT. The resulting curves are fitted using the equation. [24] 0.8

A) B) C) D)
Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ

172

𝐴(𝜆, 𝜃) = 𝐴 0 -log [1 - 𝛽 𝛽+1 (1 + 3𝑐𝑜𝑠2𝜃)] (1)
Where,

𝛽 = 1 3 tanh(Δ𝐴 𝑙𝑛10 2 ) (2)
Here A0 is the unpolarised offset absorption contribution related to the non-aligned fraction of C12-PBTTT and ΔA= Amax -Amin, Amax (0°, //) and Amin (90°, ⏊) for P1 and Amax (90°, ⏊) and Amin (0°, //) for both F4TCNQ -and F6TCNNQ -.

The very small value of the absorption offset (A0 = 0.07) before doping indicates that most of the chains are aligned and there is almost no contribution from a non-oriented fraction of C12-PBTTT. The absorption offsets for both polarons and anions are also very small ( 0.05), which further confirms that the high degree of alignment of C12-PBTTT is maintained after doping. The same observation was also obtained by Untilova (ref) and coworkers for the angular distribution measurements of F4TCNQ-doped P3HT films. In addition to this, the maximum of the angular distributions of both F4TCNQ -and F6TCNNQ - anions are clearly centred at 90. This implies that the long axis of the dopants is indeed oriented around a plane orthogonal to the average polymer chain direction and not tilted. A clearer idea about the degree of orientation of both polarons and anions in the oriented C12-PBTTT matrix was obtained from the comparison of the dichroic ratio (DR) and 3D order parameters (OP), for the polaron, the anion and the pristine polymer.

Those values are collected in table 4.1. The 3D order parameter is calculated from the dichroic ratio using equation (3):

𝑂𝑃 = 𝐷𝑅-1 𝐷𝑅+2 (3)
𝐷𝑅 = 𝐴𝑏𝑠// 𝐴𝑏𝑠 ⊥ (4)
Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 173

Where Abs// is the absorption parallel to the rubbing and Abs⊥ is the absorption perpendicular to the rubbing direction. DR is the dichroic ratio. In the case of F4TCNQ doping, the calculated order parameters (OP) was 0.7 for the oriented C12-PBTTT and it was equal to 0.7 for the polarons. The similar values of the order parameter for both polarons and the oriented PBTTT chains indicates that doping does not alter the alignment of the chains. Surprisingly, the order parameter calculated for the F4TCNQ -anions (0.56) is substantially lower as compared to that of the polarons (0.7). Similar trends were obtained by Untilova and coworkers regarding the angular distribution of F4TCNQ in oriented P3HT films (ref) . It suggests that the F4TCNQ dopants are not strictly located in a plane orthogonal to the polymer chains, but are distributed around this average orientation.

The case of F6TCNNQ doping is different. In this case, the polymer chains, polarons and the dopant F6TCNNQ -anions showed similar OP 0.7. This result supports the idea that, unlike F4TCNQ -anions, the long axis of F6TCNNQ -lies essentially in a plane that is perpendicular to the oriented C12-PBTTT chains with little angular spread around this average plane. In other words, the orientation of the F6TCNNQ dopants is defined in a better way with respect to the polythiophene chains as compared to F4TCNQ. orientation of the F6TCNNQ -molecules in the layers of alkyl side chains of C12-PBTTT might be due to stronger Coulombic interactions between the F6TCNNQ -anion and the polaronic charges on the polymer backbone possibly because of shorter inter-charge distances. Indeed, the longer molecular axis of the F6TCNNQ -anion compared to F4TCNQ - (1.47 nm versus 1.1 nm) may bring the negatively charged F6TCNNQ -anions in closer proximity to the polaron. This is true since the lattice expansion experienced for both F4TCNQ and F6TCNNQ are of comparable magnitude (see below). However, such an effect should result in a blue-shift of the polaronic bands in F6TCNNQ-doped films and it may also result in a stronger localization of the polaronic charges, which could be detrimental for charge transport properties. A precise comprehension of this effect would require simulations of the position and orientation of F6TCNNQ when intercalated into the side chain layers of C12-PBTTT. In the following section, we will evaluate the impact of dopant geometry on the diffusion coefficient into the C12-PBTTT layers.

Doping kinetics

As seen in chapter 2, the fact that the spectra for POL ⊥ R are dominated by the absorption of the dopant anions helped determine the precise diffusion coefficients of F4TCNQ -and F6TCNNQ -anions into the layers of the alkyl side chains of the thin films. As shown in chapter 2 [15,16] , polarized UV-Vis-NIR spectroscopy can be used to follow the time dependence of the absorbance of both F4TCNQ -(875 nm) and F6TCNNQ -(1158 nm) anions (see figure 4.7) versus doping time. To measure the kinetics of dopant diffusion into the C12-PBTTT thin films, the oriented films were dipped in 1 mg/ml solution of dopant in acetonitrile for various time intervals.

The contact time of the oriented polymer with the dopant solution is increased and the corresponding UV-Vis-NIR spectra along both perpendicular and parallel directions are to saturate with doping time. The spectra along the parallel direction are similar for both dopants. As expected, the F6TCNNQ doped C12-PBTTT showed a higher degree of polymer chain oxidation than F4TCNQ doped PBTTT since the remaining contribution of neutral polymer N is lower for this dopant (see figure 4.4.A&B). The higher degree of chain oxidation in F6TCNNQ doped C12-PBTTT indicates a higher degree of charge transfer between the polymer and dopant anion.

Moreover, the integer charge transfer between polymer and dopant is verified from the linear dependence between the absorbance of the polarons and the anions of F4TCNQ - and F6TCNNQ -(see figure 4.8B&C). A perfectly linear relationship between the intensity of absorbance of the anions and polarons observed for both dopants supports the integer charge transfer in oriented PBTTT for both F4TCNQ -and F6TCNNQ -. The relation between F6TCNNQ -and polaron absorbance is remarkably linear due to the reduced error in absorbance measurement for the 1158 nm band that does not overlap with the PBTTT polaron bands.

Influence of the method of sequential doping on the charge conductivity.

The electrical conductivity of the oriented C12-PBTTT film has been measured as a function of dopant concentration. Two different procedures were used to dope the oriented films i.e. incremental concentration doping (ICD) and direct doping. In ICD, each sample is doped by dipping it successively in the solutions of dopants of increasing concentration up to a given concentration while in the direct doping procedure, each sample is dipped a single time in a solution of a given concentration. Typically for ICD, we consider the doping in successive solutions of 0.01, 0.1, 0.5, 1, 2 and 5 mg/ml.

As observed previously [14][15][16]27] , the doping of all oriented thin films produced higher electrical conductivity along the rubbing direction than perpendicular to it. The conductivity in the direction perpendicular to the rubbing at saturation is close to that Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 179 of non-oriented thin films (for F6TCNNQ and F4TCNQ doped C12-PBTTT we measure 65±4 S/cm and 120±94 S/cm, respectively). More interestingly, the incremental doping on the same device results in higher electrical conductivities as compared to single doping.

Incremental addition of F4TCNQ to PBTTT (5 mg/ml) produces a remarkably high electrical conductivity of  1380±700 along the rubbing direction. Surprisingly, the directly doped sample at the same concentration showed a conductivity of  920±60 S/cm on average. ICD produced 1.5 times higher electrical conductivity than direct doping in the case of F4TCNQ doping. Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 180 values of samples obtained for direct doping at high doping concentration (2 and 5 mg/ml for F4TCNQ and 5 mg/ml for F6TCNNQ). The conductivity of non-oriented samples were also measured at maximum doping concentration.

The same trend is observed for F6TCNNQ. ICD of oriented C12-PBTTT (2 mg/ml) results in record values of electrical conductivities of  2430±500 S/cm versus 600±200 S/cm for direct doping. The electrical conductivities measured perpendicular to the rubbing direction are similar to those observed for the non-oriented films of C12-PBTTT (65±4 S/cm and 120±94 S/cm for F6TCNNQ and F4TCNQ, respectively). This further confirms that the poor electrical conductivity in non-oriented samples is limited by the poor transport perpendicular to the polymer chains.

To better understand the reason for such an effect, we performed a test experiment that rules out the possible role of the solvent (ACN) on the film structure in the case of ICD. Indeed, it might be that the improved conductivity observed in ICD samples is related to some structural reorganization in the films upon multiple dips in ACN that could be beneficial for the charge transport properties. Therefore, an oriented C12-PBTTT thin film was dipped 4 times in pure ACN before single doping (2 mg/ml of F6TCNNQ in acetonitrile) (4ACN+DD). The two samples were compared using UV-Vis NIR spectra and electrical conductivity measurements.

As seen in figure 4.9, the polarized UV-Vis-NIR of both ICD and 4ACN+DD films show the same intensities of polaronic bands P1 and P2 and almost similar contribution from the neutral polymer. This implies that there is no clear difference in the doping level between a sample experiencing ICD or 4ACN+DD. The observed electrical conductivities of 4ACN+DD and DD films were  970+200 S/cm, and 600±200 S/cm, respectively. On the other hand, ICD produced samples with a conductivity of 2430 S/cm, which is 3-4 times higher than for the DD method. Since polarized UV-Vis-NIR spectroscopy shows Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 181 that both ICD and single doping (ICD and 4ACN+DD) lead to similar charge polaron concentrations, then the observed higher electrical conductivities for ICD films must be more related to a difference in the charge carrier mobility. For both dopants, slow and progressive intercalation of dopants leads to a higher electrical conductivity in oriented C12-PBTTT. This underlines the importance to preserve the crystallinity and initial orientation of the original semiconducting polymer films.

Clearly, the way the molecular dopants are introduced into the polymer matrix is also an important parameter that influences the TE performances. As observed in previous investigations, the maximum electrical conductivity from a conducting polymer can be reached only if the molecular dopant is introduced without affecting the initial structure of the thin film. Incremental solution doping bears some similarity with vapour phase doping in the sense that the dopants can be incorporated in the polymer matrix in a Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 182 progressive and non-destructive manner [28] . Interestingly, the oriented C12-PBTTT films doped by ICD show a much higher electrical conductivity than non-oriented vapourdoped films of PBTTT (670 S/cm) and P3HT (12.7 S/cm) as observed by Patel and

Hynynen respectively. [28,29] The observed high electrical conductivity in the F4TCNQ and F6TCNQ doped PBTTT thin film can be attributed to various reasons. First of all, high degree of in-plane chain alignment in the oriented PBTTT must be one important factor that can contribute to the observed high electrical conductivity in incrementally doped C12 PBTTT thin films.

For instance, Chabinyc and coworkers proposed that the electrical conductivity in conjugated polymers such as PBTTT and P3HT mainly depends on the processing methods and the conductivity scales with the length scale of oriented backbones. [28] This was evidenced by soft -X-ray scattering measurements on vapor doped thin films of PBTTT. They showed that, higher orientation along the polymer backbone followed by doping from the vapor phase without affecting the initial thin film microstructure can produce high electrical conductivity of 670 S/cm and a PF of 120 µWm -1 K -2 in PBTTT thin films doped with F4TCNQ. In our case, rubbed C12 PBTTT shows a high degree of chain orientation along the direction of rubbing. The electron diffraction measurements on oriented PBTTT show strong reflections (see figure 4.12 and 4.13) along the direction of the rubbing (mainly 003 and 004 represents the chain direction) and confirms the high degree of chain orientation in our samples. This high degree of chain alignment followed by incremental doping can lead to an apparent higher electrical conductivities and charge carrier mobilities (see table 4.2 for the calculated mobilities) similar to vapor doped PBTTT thin films.

The higher electrical conductivity observed in F6TCNNQ doped C12 PBTTT can have different origins. The angular distribution measurements on the doped thin films Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 183 indicate that F6TCNNQ -anions are better oriented in the layers of the alkyl side-chains of PBTTT with an order parameter of 0.7. In strong contrast, the dopant F4TCNQ -anions are less ordered in the layers of the side chains. Hence a high degree of chain alignment and better ordering of the dopant F6TCNNQ can produce higher charge carrier mobilities using ICD. This high mobility could be the reason for higher electrical conductivity incrementally doped F6TCNNQ/ C12-PBTTT (see table 4.2 for the calculated mobilities).

Influence of doping method on Seebeck coefficients and correlations with charge conductivity.

Seebeck coefficients were also anisotropic and higher Seebeck coefficients were measured along the rubbing direction for both F4TCNQ and F6TCNNQ doped C12-PBTTT.

The values of the Seebeck coefficients were little affected by the doping method. For instance, ICD with F4TCNQ (5 mg/ml) leads to S//= 50.7±4 µV/K while for DD S//= 57±3 µV/K (see figure 4.8 C and D). This difference is consistent with the higher conductivity seen for ICD versus DD. Similarly, ICD with F6TCNNQ (2 mg/ml) leads to S//=47±7 µV/K versus 64.5±3 µV/K for DD at the same concentration. It is worth to mention that the Seebeck coefficients of non-oriented samples were similar to the values measured along the chain direction (52.6±3 µV/K for F4TCNQ and 64.5±4 µV/K F6TCNNQ), indicating that alignment is not detrimental for the Seebeck coefficient while it is highly beneficial to enhance charge conductivity.

As expected, the increase of electrical conductivity in oriented C12-PBTTT thin films enhances the power factors along the chain direction. Incremental concentration doping (ICD) of C12-PBTTT films with F4TCNQ and F6TCNNQ produced very high-power factors of 355±200 µWm

-1 K -2 and 530±200 µWm -1 K -2
, respectively. The PF obtained for FeCl3 doped C12-PBTTT is 1.9 µW . m -1 . K -2 (see chapter 3) 3 times higher than F6TCNNQ doped PBTTTT.

As expected, ICD provided higher power factors than DD for both F4TCNQ and F6TCNNQ.

Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ

184

The poor electrical conductivities of the non-oriented C12-PBTTT thin film lead to a PF that is lower by more than one order of magnitude compared to the aligned films: only 33 µW . m -1 . K -2 for F4TCNQ and 27±3 µW . m -1 . K -2 for F6TCNNQ. The dependence of PF with doping concentration (figure 4.10) shows the same characteristic increase and saturation for concentrations in the range 2-5 mg/ml for both dopants. The ICD method is particularly useful to probe S- correlations because a single sample of well-defined initial structure and orientation is progressively doped at higher dopant concentration, which alleviates all issues related to sample-to-sample statistics due to structural variations. In the previous chapter 3, we have demonstrated that the S- correlation is well described as proposed by Chabinyc [4] and coworkers by the relation S//// -1/4 for F4TCNQ and FeCl3 dopants in the direction parallel to the polymer chains.

In oriented samples, we have observed a different correlation in the direction perpendicular to the backbone, namely S⊥-ln (⊥). The open question relative to these correlations is whether or not they depend on the chemical nature of the dopant. This Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ

185 is why we collected all S, data for films doped by ICD with the dopants F4TCNQ, F6TCNNQ and FeCl3 on the same plot (see figure 4.11). The correlations were plotted for both parallel and perpendicular orientations with respect to the chain direction.

F4TCNQ

Vapor Phase [8] 670 42 120

Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 186 limit correlates with the maximum doping/oxidation levels reached for each dopant that follows the sequence F4TCNQ<F6TCNNQ<FeCl3. The full lines correspond to the results of the fit using the scaling laws S//// -1/4 and S⊥-ln (⊥).

Impact of doping methods on the crystal structure

As seen in the previous chapters, doping results in the intercalation of dopant molecules in the layers of alkyl side chains. The extension of lattice expansion is correlated with the initial length of side chains: the shorter the side chains, the larger the lattice expansion along the alkyl side chains. Hereafter, we want to investigate the impact of dopant size on the lattice of C12-PBTTT upon doping. [8,10,15,30] Similarly, to the previous chapter, we followed the structural changes in both F4TCNQ The TEM grids were doped by ICD using dopant solutions (F4TCNQ or F6TCNNQ) in acetonitrile in the glove box. The dopant solution was allowed to stay on the TEM grid for 20-30 s and the excess solution was removed by blotting. The structural variation in the doped thin films was followed by electron diffraction using a transmission electron microscope. In the direct doping method, carbon-coated TEM grids were doped directly Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 189 at a given concentration of dopants. In the case of ICD, the grids were doped by dipping in solutions of increasing concentration up to the desired concentration.

This study report one of the first investigations on how the crystal structure of C12-PBTTT is affected by different doping procedures using the same dopant. As shown hereafter, interesting results are obtained by the two different doping procedures.

The thin films of C12-PBTTT prepared by high temperature rubbing show a sequence of h00 (h=1,2,3,4) reflections along the equatorial direction. The 003 reflection along the rubbing direction represents the periodicity along the backbone direction. The thin films were mainly in face-on orientation as indicated by the absence of the equatorial diffraction peak corresponding to the π-stacking (d020) (see figure 4.12 and 4.13). for ICD and DD, respectively). Surprisingly, ICD with the longer and bulkier F6TCNNQ produced a lower d100 spacing of 20.28Å than for F4TCNQ-doped films. Incidentally, the lattice expansion for both F4TCNQ and F6TCNNQ is significantly reduced with respect to FeCl3 (d100 =23.7 Å). Accordingly, the lattice expansion in C12-PBTTT is not correlated to the dimensions of the dopants: F4TCNQ has a long axis of 12 Å versus 14.4 Å for F6TCNNQ. UV-Vis-NIR spectroscopy has shown that the doping level for ICD at saturation are 14% and 20% (dopant per thiophene cycle), respectively. Hence, the lower lattice Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 190 expansion for F6TCNNQ is not associated with a lower doping level. The study of the doping kinetics indicates solely that the diffusion coefficient of F6TCNNQ into the structure of C12-PBTTT is slightly lower as compared to F4TCNQ. Accordingly, it seems that steric hindrance cannot be considered as the reason for a lower lattice expansion with F6TCNNQ. Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 191

Estimation of doping level and mobilities

Of importance for charge transport properties are the charge carrier density and the charge mobility that determine the resulting charge conductivity. The exact determination of charge carrier density is difficult and would require to use specific tools such as the Hall effect. [31] A rough estimation of the carrier density can be obtained from the measurement of absorbance of the polaron or of the anion since we have shown that integer charge transfer is involved in the doping process of C12-PBTTT with F4TCNQ and F6TCNNQ. However, doing so, one assumes that each anion generates one polaronic charge that can contribute to charge transport. To extract charge mobilities, the major assumption is that all polaronic charges seen by UV-Vis-NIR spectroscopy contribute to the charge transport process. This may not be true in absolute value as demonstrated by Pingel and Neher who showed that only 5% of charges generated by doping correspond to free charge carriers by doping P3HT with F4TCNQ. [22] To determine the anion concentration in the films, we determined the molar extinction coefficient of F6TCNNQ -anion in solution (in DCM). We find F6TCNNQ-= 53300 L . Mol -1. cm - 1 at 1153 nm. From the absorbance and the film thickness, one can extract a density of anions, hence of polarons, of 8.1×10 20 cm -3 . The polaron density calculated for the F4TCNQ doped C12 PBTTT is  5.7×10 20 cm -3 . Based on AC Hall effect measurements, Scholes and coworkers also reported similar charge carrier densities in non-oriented P3HT films doped from F4TCNQ [31] . This implies that the determined polaron densities using the extinction coefficients of the anions are reasonable.

A second approach implies to know the extinction coefficient of the polarons (P2). Here, we propose two approaches. First, we make use of the extinction coefficient determined by Charge Modulated Spectroscopy (CMS). We use the Beer-Lambert law as proposed Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 192 by Tanaka [32] and coworkers for FTS-doped PBTTT thin films. According to Beer-Lambert law,

𝐴 𝑃𝑜𝑙 = 𝜀 𝑃2 𝑛𝑡 (5)
Where APOL is the absorbance of the PBTTT polarons at 1.67 eV, P2 is the effective absorption cross-section of the polaron, t is the film thickness and n is the polaron density. CMS yields the absorption coefficients of C12-PBTTT polarons pol = 6×10 -16 cm 2 [33,34] . Only the contribution to the absorption of the free charge carriers involved in charge transport is determined using CMS whereas in conventional UV-Vis-NIR spectroscopy the absorption stems from both trapped and mobile polarons. By considering a film thickness of (40 nm F6TCNNQ doped (ICD) C12-PBTTT and 80 nm for F4TCNNQ doped (ICD) C12-PBTTT) and the relation P2=(log e) pol, the polaron density is in the range 2.83 -3×10 20 cm -3 for F6TCNNQ doped C12 PBTTT (2mg/ml) and 3.2×10 20 cm - 3 . Since the DD film shows the same polaronic absorption as the ICD film (see figure 4.9), a similar polaron density of 3.3×10 20 cm -3 is obtained.

Finally, we can also use a different value for P1 using the results obtained by polarized UV-Vis-NIR for the FeCl3-doped C12-PBTTT. In chapter 3, we have observed that C12-PBTTT is fully oxidized after doping with 10 mM FeCl3. The corresponding spectrum (see Figure 3.6 in chapter 3) shows no neutral form of the neutral polymer and only P1 and P2 absorption bands. Therefore, knowing the stoichiometry of the doped films (from EDX), we can show that a unit cell contains two FeCl4 -charges i.e. 2 polarons for 8 thiophene cycles. Knowing the absorption at 2500 nm and the film thickness, we extract P2=9.6×10 -17 cm 2 . This value is significantly different from the previous one obtained by CMS. The difference is related to the fact that our calibration with the FeCl3-doped films takes into account the UV-Vis-NIR absorption that stems from both mobile and trapped polarons. Using this value of P1, we extract a charge carrier density of 6.2 . 10 20 cm -3 for Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 193 F6TCNNQ doped C12-PBTTT (ICD) thin films and  6.7 . 10 20 cm -3 for F4TCNQ (ICD) doped C12-PBTTT. As expected, the carrier density extracted from the polaron absorption using P2 is lower than the value obtained using P1, since P1 corresponds to free and trapped charge carriers. As a matter of facts, the polaron densities extracted from P1 and F6TCNNQ are fairly similar, which is consistent with integer charge transfer and suggests that the extinction coefficients determined in solution for F6TCNNQ are consistent.

An estimate of the charge carrier mobility was calculated for both incremental and singly doped thin films using the charge carrier densities extracted based on the P1 and F6TCNNQ

and charge modulated spectroscopy (based on reference 32 Neher for rr-P3HT, would result in even larger charge mobilities. [22] Such values of charge mobilities are 1-2 orders of magnitude larger than those measured in OFETs of C14-PBTTT with oriented chains. Lee et al. observed mobilities of up to 0.4 cm 2 /V . s in oriented thin films. [35] Xue and coworkers also reported high mobilities of 1.67 cm 2 /V . s in stretch aligned films of C14-PBTTT. [36] More recently, Scholes and coworkers also reported similar mobilities of 0.12 cm 2 /V.s in non-oriented thin films of P3HT, sequentially doped with F4TCNQ from a charge carrier concentration of 4.9×10 20 cm -3 . [31] The difference between charge mobilities extracted from OFET and from highly doped system has multiple origins. First, the carrier density is much smaller than in our doped films. Second, the transport of charges is 3D in doped films whereas it is more 2D in OFETS. The estimated mobilities in our films are >10 cm 2 /V.s, which is assumed to be a limit to observe band-like transport. [37] There are several reasons why mobilities can be so high in our doped films. First, doping induces a stiffening and planarization of the polymer backbone whose structure becomes more quinoidal. At low doping concentration, the counterions tend to trap and localize the polarons and bipolarons because the Coulomb potential of the counterions extends quite far in the polymer (due to a low dielectric constant). Because of this, counterions act as traps for polarons and charge mobility is reduced. When the doping Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ

195 level increases, one observes a progressive overlap between the electric fields produced by the counterions, which results in a reduced hopping barrier and the charge carrier mobility tends to increase. [38][39] When the Coulomb traps fully overlap, then they form a band that tends to boost the charge carrier mobility. This is most likely why the charge mobilities can be so high in C12-PBTTT doped with F6TCNNQ or FeCl3. Second, we may also consider that doping influences the polymer microstructure in thin films. As shown in a recent work by Untilova et al., doping can promote some ordering in the smecticlike phase of rr-P3HT (ref) . Doping allows polythiophene backbones to reorganize towards a more ordered state upon intercalation of F4TCNQ dopants in the layers of alkyl side chains. We may suggest a similar phenomenon to occur in the present case in the more disordered parts of the PBTTT films. This would be similar to the observation by Chabynic and coworkers of doping-induced order in the amorphous phase of ra-P3HT. To conclude, we propose that ICD can produce such high mobility samples because i) the progressive intercalation of dopants preserves the pristine orientation order of PBTTTT, ii) dopants are introduced in a regular manner in the lattice of C12-PBTTT.

7.Conclusion

High temperature rubbing is a simple and versatile method, which can be applied to a wide variety of polymers. Similar to our previous works, we have applied high temperature rubbing to C12-PBTTT films and doped them with F6TCNNQ and F4TCNQ to produce highly conducting thin films. A combination of Polarized UV-Vis-NIR spectroscopy and electron diffraction measurements demonstrates that, similar to F4TCNQ -anions, F6TCNNQ -anions are also oriented perpendicular to the polymer backbone and prefer to reside in the alkyl side chains of C12-PBTTT crystals. [15,16] Moreover, we show that F6TCNNQ -anion are better oriented and lie in a plane strictly perpendicular to the C12-PBTTT backbone whereas F4TCNQ -anions are more distributed Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 196 around this direction. Similarly, to rr-P3HT, doping C12-PBTTT with F6TCNNQ and F4TCNQ produces an expansion of the unit cell along the d100 and a compression along the π stacking direction but the extent of lattice expansion along the alkyl side chain is not correlated to the long axis dimension of the dopant. The kinetics of doping was followed by UV-Vis-NIR spectroscopy and confirms that the longer and bulkier F6TCNNQ -anion has a lower diffusion coefficient (2×10 -12 cm 2 /s) than F4TCNQ -anions (9×10 -12 cm 2 /s).

In addition to this, we introduce a new doping method called incremental concentration doping (ICD). Incremental concentration doping leads to higher electrical conductivities and higher power factors in oriented C12-PBTTT than direct doping. ICD produced 4 times higher electrical conductivities (2430±500 S/cm) than direct doping (600±500 S/cm) and produced a record power factor of 530±200 µWm -1 K -2 . Also, similar to our previous results, non-oriented PBTTT sample limited electrical conductivities of 120±90 S/cm and 65±4 S/cm. A combination of alignment and incremental concentration doping produced 37 times higher electrical conductivities in F4TCNQ doped C12-PBTTT and 11 times higher electrical conductivities for F4TCNQ-doped films. It is shown that ICD and DD result in similar doping levels using UV-Vis-NIR spectroscopy. Therefore, the observed difference in charge conductivity is ascribed to the preservation of pristine order and higher mobilities observed in ICD doped C12 PBTTT. The enhanced electrical conductivities observed in F6TCNNQ doped C12-PBTTT is further pointing that the degree of ordering of the dopant anion is an important factor that determines the electrical conductivity in semiconducting polymers. Higher electrical conductivities are always observed in the oriented C12-PBTTT films when the dopants are introduced in a progressive manner in the films by ICD. This result demonstrates how important the doping process is in determining the resulting TE properties of the polymer thin films. From that perspective, ICD is more similar to vapor phase doping whereby the dopants enter the polymer films Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 197 progressively during evaporation. Our recent results on electrical conductivities on other semiconducting polymers further demonstrate that ICD is a versatile method that can be applied to different semi-crystalline polymers for better thermoelectric performance.

Conclusions and perspectives

The aim of my thesis was to investigate the effect of polymer orientation on thermoelectric properties of doped polythiophenes such as PBTTT and P3HT. We have used a simple and elegant method for aligning polymer semiconductors, which is known as high temperature rubbing. High temperature rubbing is a versatile method, used for aligning different classes of organic semiconductors. [1][2][3] This method produces highly anisotropic crystalline polymer films as evidenced by polarized UV-Vis-NIR spectroscopy and electron diffraction. Sequential doping of these oriented thin films using suitable doping agents such as F4TCNQ, F6TCNNQ and FeCl3 produced highly conducting polymer films. [4][5][6] . The structural, spectroscopic and thermoelectric characterization of these doped thin films shows a high degree of anisotropy of TE properties.

chapter 2 mainly discusses the relationship between the length of the alkyl side chain of PBTTT families and their thermoelectric properties. In addition to this, we have clearly demonstrated the relationship between the side chain length of PBTTT and the diffusion coefficients of the F4TCNQ -anions into the alkyl side chain layers. The perpendicular orientation of the F4TCNQ -anion in the alkyl side chains of PBTTT enabled us to follow the doping kinetics. The idea was simple and to the best of our knowledge, it was the first time, the diffusion of the dopant molecule was determined by using such a simple experimental method (polarized UV-Vis-NIR spectroscopy). We have seen that there is an optimum side chain length, (C12-PBTTT) which allows easy diffusion of dopants into the alkyl side chain layers. This lead to higher electrical conductivity.

The electrical conductivities and Seebeck coefficients of these oriented thin films are Some of the recent research work on organic thermoelectrics proposed that future efforts to improve the thermoelectric properties of semiconducting polymers must be focused on improving electrical conductivity (). [7] This can eventually lead to very high power factors according to the relation PF 1/2 as proposed by Glaudell [7] and coworkers. The third chapter demonstrates the use of strong Lewis acid-based doping agents such as FeCl3 on oriented polymers such as C12-PBTTT and P3HT. Sequential doping of liquid crystalline polymers such as C12-PBTTT using FeCl3 leads to record electrical conductivities of 2×10 5 S/cm and power factors of  2 mWm -1 K -2 along the chain direction. The measured electrical conductivities were higher than for I2 doped polyacetylene. [8][9][10] This higher electrical conductivity was obtained by using incremental concentration doping method i.e., each sample is doped by dipping it successively in the solutions of dopants of increasing concentration.

To the best of our knowledge, incremental concentration doping of semiconducting polymers has not been reported yet. The observed high electrical conductivity is mainly related to the regular intercalation of dopant molecules into the polymer structure.

Most importantly, we have seen that the relationships between the Seebeck coefficients and electrical conductivitiy along the direction // and ⊥ to the chains are different. Along the parallel direction, the Seebeck coefficient and the electrical conductivity follow the power-law S//  // -1/4 . Along the perpendicular direction, S⊥ln⊥, which is a fingerprint Conclusions and perspectives 206 for different charge transport mechanisms. A more detailed investigation is necessary to further understand the underlying charge transport mechanism.

In the last chapter we investigated the following issues: i) How does the doping of C12-PBTTT with F6TCNNQ differ from that with F4TCNQ in terms of structure, doping kinetics and resulting TE properties?

ii) How different are the intercalation of F4TCNQ and F6TCNNQ in the layers of alkyl side chains of C12-PBTTT?

iii) Is there any difference in electrical conductivity in thin films doped by different doping methods such as incremental concentration doping and direct doping, which was previously observed in the case of FeCl3 doped samples.

Similar to our previous results, the diffusion kinetics on oriented PBTTT thin films demonstrated the size dependence of the dopant anions and their diffusion coefficients.

As expected, longer and bulkier F6TCNNQ -anions showed slightly slower diffusion coefficients than F4TCNQ -. Moreover, angle-dependent polarized UV-Vis-NIR spectroscopy suggested that longer and bulkier F6TCNNQ -anions are better oriented in the layers of alkyl side chains of C12-PBTTT.

Most importantly, we observed that incremental concentration doping of C12-PBTTT thin films always produced higher conductivity than singly doped films, irrespective of the doping agent. The mobility calculation based on the extinction coefficients of the dopant anions and polarons suggested that the higher electrical conductivity in ICD might be related to the enhanced mobility of charge carriers due to the progressive intercalation of the dopant into the alkyl side chain layers. The higher electrical conductivity observed in incrementally doped C12-PBTTT thin films using F6TCNNQ -is Conclusions and perspectives 207 attributed to a better ordering of F6TCNNQ -anions in the side chain layers of the polymer.

In summary, the final outcome of the thesis is the following.

1) We proposed a simple method based on polarized UV-Vis-NIR spectroscopy to follow the diffusion of dopant anions such as F4TCNQ -and F6TCNNQ -in the oriented polymer matrix of various PBTTTs.

2) We were able to show that, polymer orientation through high temperature rubbing leads to highly conducting polymer thin films with enhanced electrical conductivities along the direction of rubbing. Moreover, these anisotropic thin films enabled us to see a difference in the S vs  relation along and perpendicular to the chain direction. This is of high importance since such an observation is not accessible from non-oriented semiconductor thin films. This is due to the fact that the measured electrical conductivity and Seebeck coefficient in non-oriented films is a mixture of contributions along c, b and a axes. Hence it is impossible to characterize the transport mechanism in one direction of the polymer crystal. It is worth to mention that, other polymer orientation techniques such as stretching and tensile drawing do not show a clear trend and anisotropy in the Seebeck coefficients but only showed a clear enhancement of electrical conductivity along the direction of applied stress. [9,11,12][13] In our case, S// is always larger than S⊥.

3) We proposed a simple method to characterize the orientation of dopant molecules such as F6TCNNQ and F4TCNQ in the layers of the alkyl side chains using angle-dependent polarized UV -Vis-NIR spectroscopy.

4) We introduced a new doping method called incremental concentration doping (ICD). Incremental concentration doping leads to higher electrical conductivities and higher power factors in oriented C12-PBTTT than direct doping.

Conclusions and perspectives 208

The ongoing project on incremental doping on other polymers further confirms that ICD is versatile and can be applied to a large palette of polymer semiconductors and dopants to enhance the TE properties. We believe that some of the results and publications from this thesis might be of high interest for the experimental and theoretical physicist to investigate further the underlying physical processes associated with TE properties. Also, the fundamental understandings and the results obtained from this thesis might be a small but important contribution to the field of plastic electronics.

However, there is always room for future developments. There is no doubt that the future of TE research on p-doped polymers must be based on enhancing electrical conductivity by introducing high crystallinity and long-range order along the polymer backbone. Preliminary results by electron diffraction measurements suggest that the introduction of ethylene glycol and ether-based side chains on PBTTT enhanced the order along the chain direction. Incremental concentration doping of these polymers by F6TCNNQ can lead to electrical conductivities  2×10 4 S/cm. This further confirms that long-range order along the polymer backbone and suitable doping methods such us ICD can provide enhanced TE properties. Side chain engineering can be combined to the methods of doping and orientation proposed in this thesis to further boost TE properties.

The measurement of ultimate thermal conductivity is also an important parameter, which can affect the ZT value of a TE material. Hence, it is important to measure the thermal conductivity of these oriented polymers. Our preliminary results on thermal conductivity measurements on oriented P3HT suggest that the rubbing process does not substantially increase the in-plane thermal conductivity and still results in an enhancement by a factor of 25 of the ZT value compared to non-oriented films. A thermoelectric device is an alternative array of p and n-type organic materials. Oriented TE devices would be possible only after extending the high temperature rubbing method Conclusions and perspectives 209 to n-type polymers. Gross and coworkers demonstrated that n-type polymers such as P(NDI2OD-T2) can also align using high temperature rubbing. [14] Vapor doping of oriented P(NDI2OD-T2) with tetrakis(dimethylamino)ethylene (TDAE) leads to a higher electrical conductivity of 8.3×10 -3 S/cm along the chain direction, which was 6 times higher than the direction perpendicular to the chains direction. This demonstrates the versatility of the high-T-rubbing method, that can be also applied to n-type polymers.

Combining alignment with side-chain engineering of polymers, dopant engineering and doping methods might be an interesting strategy forward to make oriented TE devices with enhanced electrical conductivities.

Experimental Details

Polymer synthesis and characterization

C8-PBTTT, C12-PBTTT, C14-PBTTT, and C18-PBTTT were synthesized by N. Leclerc and E.

Zaborova following the detailed procedure according to the references. [1][2][3][4][5] P3HT was purchased from Merck. 2,3,5,5-tetrafluoro(tetracyanoquinodimethane) (F4TCNQ) was purchased from TCI and solvents such as anhydrous acetonitrile, anhydrous nitromethane, Sodium polystyrenesulfonate (NaPSS), anhydrous FeCl3 and were obtained from Aldrich. The F6TCNNQ dopant has been synthesized [6] by P.Durand according to the reference (6). All the chemicals were used as received. The macromolecular parameters are collected in Table 1.

preparation of conducting polymer films 2.1 Preparation of glass slides

The glass slides for the thin film preparation are cleaned by sonicating in different solvents such as acetone, isopropanol, helmanex (diluted to 10% in Milli-Q water) for approximately 15 minutes. These glass slides are transferred to Milli-Q water and sonicated for 15 minutes to remove excess helmanex solution and kept in isopropanol.

The isopropanol is dried under nitrogen flush before doctor blading the films. ). [7,8] Experimental Details 217

The rubbing machine is composed of a metallic cylinder (diameter 4 cm) covered by a microfiber tissue (see figure 1). The cylinder can be rotated at a speed of approximately 300 RPM during the rubbing process. The cylinder is applied with a 2-3 bar pressure on the thin film. Most importantly, the temperature of the film during rubbing can be controlled. Calibration of the temperature was performed to account for the difference between the set temperature and actual temperature on the surface of the glass substrate. The films were rubbed at different rubbing temperatures (TR) depending on the polymer (125°C for C8 PBTTT, C12 PBTTT, C18 PBTTT, and 100°C for C14 PBTTT). The films of P3HT

were rubbed at 180 C to achieve a high degree of orientation as proposed by Hamidi-Sakr and coworkers [9,10] . The rubbing was done inside a glove box under nitrogen atmosphere to avoid oxidation of the polymer. To determine the film thickness after rubbing, the films were melt-annealed at 300 C in nitrogen in a Linkam hot stage (HS-95 controller and LTS420 stage) to randomize the in-plane chain direction and the thickness was extracted from the UV-vis absorbance.

Doping methods

A) F4TCNQ doping of polythiophenes

A 1 mg/ml solution of F4TCNQ was prepared by dissolving F4TCNQ in anhydrous acetonitrile in a glass vial. Doped PBTTT films were prepared by dipping the polymer films in the dopant solution for different time intervals (for doping kinetics).Films were dried using nitrogen to remove the excess F4TCNQ solution. Two different procedures were used to dope the oriented PBTTT i.e., incremental doping (ICD) and direct doping (DD). In ICD each sample is doped by dipping it successively in the solutions of dopants of increasing concentration. The excess F6TCNNQ or F4TCNQ crystals on the backside of the glass slides were cleaned using small amounts of acetone before measuring the polarized UV-Vis-NIR spectra. In the DD procedure, each sample is dipped a single time in a solution of a given concentration. All doping experiments, electrical conductivity measurements and Seebeck coefficients measurements were carried out in a glove box with PO2< 1 ppm and PH2O<1 ppm to avoid dedoping. The polarized UV-Vis-NIR measurements were done within 2 to 3 minutes after doping to avoid dedoping.

Sample preparation for TEM A) F4TCNQ doped PBTTT

Samples for transmission electron microscopy were prepared by spin coating (3000 rpm) a thin layer of NaPSS (10 mg/ml aq) on a precleaned glass substrate and this substrate was used for the preparation of oriented polymer film as described above. Doping was done by dipping the oriented polymer film in the dopant solution for the different time intervals. Thew were subsequently removed from the glass substrate by floating on distilled water and recovered on TEM copper grids. The films were coated with a thin amorphous carbon layer using an auto 306 Edwards evaporator.

TEM was performed in the bright field and diffraction modes using a CM12 Philips microscope equipped with an MVIII (Soft Imaging System) Charge Coupled Device camera. Beam exposure was set to a minimum using the low dose system to avoid de-doping under the electron beam that is observed when the same zone is exposed for a prolonged period of time.

B) FeCl3 doped P3HT and C12 PBTTT

The preparation of the samples for TEM analysis involved three steps: i) Floating the oriented polymer films (P3HT or PBTTT) on distilled water and recover them on TEM copper grids ii) Coat the films with a thin amorphous carbon layer iii) Dope the TEM grids directly in the glove box by dropping 20 µl of FeCl3/nitromethane solution on the grid. After 30s, the grids were blotted with a piece of absorbing paper to remove the excess of dopant solution.

To avoid sample ageing upon transfer in ambient atmosphere, the TEM grids were mounted on the TEM sample holder inside the glove box and the holder was then transferred to the TEM in an air-tight container. The samples were only swiftly exposed to air (1-2s) just before introducing the sample holder into the TEM.

C) Incremental concentration doping using F6TCNNQ and F4TCNQ

In order to verify the impact of the doping method, i.e. incremental concentration doping and single doping low dose electron diffraction were performed on doped PBTTT thin films. For TEM investigations, the oriented PBTTT thin films prepared on NaPSS by high-temperature rubbing were carbon-coated and recovered on copper TEM grids. The TEM grids were doped by incremental concentration doping using dopant solutions (F4TCNQ or F6TCNNQ) in acetonitrile in the glove box. The dopant solution was allowed to stay on the TEM grid for 20 to 30 s and the excess solution was removed by blotting.

In the direct doping method, carbon-coated TEM grids were doped directly at a given concentration of dopants.

Scanning electron microscopy

Scanning electron microscopy done by A. Carvalho. The experements were performed at 1kV with a HITACHI SU8010 FEG-SEM on as-prepared films (no coating) to verify the uniformity of the rubbed layers.

Scanning transmission electron microscopy-dispersive X-ray spectroscopy (STEM-EDX)

STEM-EDX experiments were performed using a Cs-corrected JEOL JEM-2100F microscope operated at 100 keV to avoid dedoping of the films under the beam. energy dispersive X-Ray spectroscopy (EDX) mapping was performed using a JEOL Silicon Drift Detector (DrySD60GV: sensor size 60 mm 2 ) with a solid angle of approximately 0.5 rad".

The EDX measurements were done at IPCMS by M. Bahri and O. Ersen.

Cyclic voltammetry

Cyclic voltammetric measurements were done by N. Leclerc at ICPESS. Oxidation potentials were determined by voltammetry with a conventional 3-electrode system using a voltammetric analyser equipped with a platinum microdisk (2 mm 2 ) working electrode and a platinum wire counter electrode. The reference electrode is constituted of a non-aqueous silver electrode including the following electrolyte solution: 0.01 M silver nitrate + 0.1 M Tetrabutylammonium perchlorate in Acetonitrile. Potentials were calibrated versus the saturated calomel electrode (SCE), using the ferrocene/ferrocenium (Fc/Fc + ) couple as an internal reference and a conventional scan rate of 100 mV/s. Recrystallized tetrabutylammonium hexafluorophosphate (Bu4NPF6) was used as the supporting electrolyte (0.1 M) in distilled and anhydrous acetonitrile (ACN). All potentials are referred to the SCE electrode that was calibrated at 0.41 V vs Fc/Fc + system. Following the work of Jenekhe et al. [11] , we estimated the ionization

Differential Scanning Calorimetry

DSC measurements were performed on a DSC Q2000 from TA Instruments under nitrogen atmosphere using a sample mass in the range 1.5-2.5 mg. The samples were heated to 250°C with a heating rate of 10°C.min -1 , then cooled to 20°C at 5°C/min .

Electrical conductivity and thermopower measurements

Rather than evaporating gold contacts on top of the oriented and doped films, we prepared all oriented films on NaPSS/glass substrate and recovered the films by floating On the same substrate, two devices are oriented along the rubbing direction) and two others in the perpendicular direction, which allows us to determine the charge transport and thermoelectric anisotropy on the same substrate (see figure 4 and5). Four-pointprobe measurements of electrical conductivity were performed using a Keithley 4200-SCS and a lab assistant semi probe station in a Jacomex glovebox under N2 atmosphere.

To derive the resistivity  from the sheet resistance R measured on the device geometry Thermopower measurements were conducted on the same devices in a glove box. The thermopower was measured via the differential temperature method: a temperature gradient is established across the sample either along the rubbing direction or perpendicular to it.

The measurements were performed by using a homemade setup made of one heating and one cooling Peltier cells (3 mm gap) (see figure 5) providing a controllable temperature difference ∆T. [9] Temperatures of the cold and warm sides were measured in a non-contact mode using two IR sensors avoiding thus thermal contact issues. The

Seebeck tension was measured using a Keithley 2634B source meter and a semiprobe lab assistant probe station in ambient conditions. The Seebeck coefficient is calculated from the slope of Vtherm versus temperature difference T for ∆T varying in a range ±10K around T=23±2°C.

Résumé de la thèse

Cette thèse traite de la fabrication de films minces orientés de polymères conducteurs pour applications en thermoélectricité. Plus, précisément, nous nous sommes penchés sur une série de polythiophènes, les PBTTTs dopés séquentiellement avec différents dopants (F4TCNQ, F6TCNNQ et FeCl3). Il s'agissait de comprendre comment contrôler et améliorer les propriétés TE de cette famille de polythiophène en tirant parti de leur alignement par une technique de brossage mécanique à haute température. La thèse est organisée autour de quatre chapitres et une conclusion. Une analyse quantitative de la cinétique de dopage a été réalisée en traçant l'absorbance de l'anion F4TCNQ -, A875(t) (à 875nm) en fonction du temps de dopage t (Figure 3). Afin de s'affranchir des problèmes de temps de diffusion différents pour des épaisseurs de films différentes, nous avons tracé les cinétiques en fonction de t 1/2 /l avec l, l'épaisseur des films. Les polymères ayant des chaînes alkyls différentes montrent des cinétiques de dopage différentes. Un ajustement des données a été réalisé à l'aide de l'équation phénoménologique suivante:

(1)

Où A0 est l'absorbance finale à saturation et D est le coefficient de diffusion caractéristique du F4TCNQ et l est l'épaisseur des films. Le dernier chapitre vise à comprendre l'effet de la taille et de la géométrie du dopant sur les propriétés TE des couches minces dopées de C12-PBTTT et sur la structure des couches. On a déterminé la distribution d'orientation précise des dopants dans les couches de chaînes alkyls par spectroscopie UV-vis-NIR polarisée. Nous avons aussi évalué comment la méthode de dopage influence les propriétés TE obtenues. En particulier, on a considéré deux modes de dopages : i) par augmentation progressive de alkyls et ont leur grand axe orienté dans un plan perpendiculaire aux chaînes de C12-PBTTT. La distribution d'orientation des dopants autour du plan perpendiculaire au chaînes de polythiophène dans les couches alkyls a été déterminée à partir de la dépendance angulaire des bandes d'absorption des deux anions et en la comparant à celle des chaînes polythiophènes (Figure 12). L'analyse a permis de mesurer précisément les paramètres d'ordre traduisant l'orientation des anions, du polaron et des chaînes du polymère dans les films avant dopage (tableau 1).

Tableau 1. Caractéristiques issues de l'analyse de la distribution angulaire des films minces de C12-PBTTT dopés avec F4TCNQ et F6TCNNQ.

L'orientation des dopants de type F6TCNNQ est plus proche du plan orthogonal aux chaînes de PBTTT que pour F4TCNQ dont l'orientation moyenne s'écarte davantage de cette direction. L'impact de l'intercalation sur la structure a aussi été déterminée par diffraction électronique et on a ainsi pu déterminer l'influence de l'intercalation sur les paramètres de maille de C12-PBTTT. Malgré une taille plus grande, la maille du C12-PBTTT ne présente pas une augmentation de volume supérieure. En étudiant par UV-vis-NIR la cinétique de diffusion de F4TCNQ et F6TCNQ dans les couches orientées de C12-PBTTT, nous avons montré que les coefficients de diffusion du F4TCNQ est supérieur à celui du F6TCNNQ (2.5×10 -12 cm 2 /s contre 9 . 10 -12 cm 2 /s pour F4TCNQ). Ceci .

Chapter 1 .

 1 Introduction .................................................................................................................................. Fundamental concepts & state of the art ........................................................................... 1.General aspects of conjugated polymers ............................................................................................ 1.1. Energy level design of conjugated polymers. .............................................................................. 1.2. P3HT: The working horse among organic semiconductors ......................................................... 2. Semi-crystalline P3HT-lamellar morphology ..................................................................................... 2.1. Crystal structure of P3HT ............................................................................................................. 3.Liquid -crystalline polymers with enhanced structural order............................................................ 3.1. Crystal structure of PBTTT ............................................................................................................ 4. Controlling orientation of conjugated polymers ................................................................................ 4.1. Orientation by mechanical rubbing ............................................................................................. 4.2. Orientation by friction transfer .................................................................................................... 4.3. Orientation by strain alignment ................................................................................................... 4.4. Orientation by using an orienting substrate (Epitaxy) ................................................................. 5. Anisotropy of charge transport .......................................................................................................... 6. Thermoelectricity ................................................................................................................................ 6.1. Basic principles of thermoelectricity............................................................................................ 6.2. Thermoelectricity in conducting polymers .................................................................................. 7. Doping of conjugated polymers ......................................................................................................... a) Electronic Process involved upon P-type doping ........................................................................ b) Charge carriers in conducting polymers ..................................................................................... 7.1. Factors influencing doping and thermoelectric properties of polythiophenes ........................... a) Regioregularity, solid-state order and crystallinity ..................................................................... b) Thermal conductivity .................................................................................................................. c) Size and electron affinity of dopant and IE of polymer ............................................................... d) Polymer orientation techniques ................................................................................................. e) Doping methods and resulting thermoelectric properties of polythiophenes ........................... i) Soution doping of polythiophenes…………………………………………………………………………………….58 ii) Sequential doping of polythiophenes……………………………………………………………………………….61 iii) Vapor phase doping…………………………………………………………………………………………………………62

Figure 1 . 1 .

 11 Figure 1.1. Schematic illustration of bonding and hybridization in ethylene molecule. Reproduced from (2)

Figure 1 . 2 .

 12 Figure 1.2. Evolution of HOMO and LUMO levels and the bandgap (Eg) with an increasing number of thiophenerepeating units and the resulting valence and conduction bands in polythiophene. Reproduced from(3)

Figure 1 . 3 .

 13 Figure 1.3. Chemical structures of most commonly used conjugated polymers. Reproduced from (3)

Figure 1 . 4 .

 14 Figure 1.4. A) Structures of possible coupling in 3-alkylthiophene dimers. B) Solid-state packing of HT coupled P3HT. C) Steric repulsion in P3HT as a result of the irregular coupling of monomers. Reproduced from(8)

Figure 1 . 5 .

 15 Figure 1.5. A) Schematic representation of the semi-crystalline morphology observed in an oriented P3HT film.B) Corresponding bright-field transmission electron microscopy image showing alternating crystalline and amorphous regions. A lamellar periodicity of 28 nm was obtained from the fast Fourier transform (FFT), which isshown in the inset. Reproduced from(14)

Chapter 1 :

 1 Fundamental concepts and state of the art 20 spin coating gives rise to less crystalline films than a slow drop-casting. The crystallinity and morphology of the thin films can be tuned by changing the molecular weight, processing conditions, solvent.

Figure 1 . 6 .

 16 Figure 1.6. A&B) STM images showing chain folding in regioregular poly(alkyl thiophenes) epitaxially oriented on HOPG. The 2D crystalline packing C) and D) are obtained after modelling using semi-empirical calculations.

Figure 1 . 7 .

 17 Figure 1.7. A) Representation of the crystal structure of P3HT form I proposed by Kayunkid and coworkers. Fig B) represents the evolution of lattice parameters as a function of molecular weight. Reproduced from (11&18)

Figure 1 . 8 .Figure 1 . 9 .

 1819 Figure 1.8. Chemical structure of Poly(2,5-bis(3-alkylthiophene-2-yl)thieno[3,2-b]thiophene)

  impressive hole mobilities of 1 cm 2 /V.s in undoped Poly(2, 5-bis(3-tetradecylthiophene-2-yl)thieno[3,2-b]thiophene) (C14-PBTTT). Moreover in 2008, Kumada and coworkers reported mobilities of 1 cm 2 /V.s in Poly(2,5-bis(3-hexadecyldecylthiophene-2yl)thieno[3,2-b]thiophene) (C14-PBTTT

Figure 1 . 10 .

 110 Figure 1.10. A&B) out of plane and in-plane scattering from C12-PBTTT. Figure (C&D) grazing incident X-ray scattering from C12-PBTTT.C) as spun C12-PBTTT D) annealed (180 °C) C12-PBTTT on OTS/SiO2 substrate.

Figure 1 . 11 .

 111 Figure 1.11. Simple packing model of C12-PBTTT based on X-ray scattering data. The lamellar stacking due to the alkyl side chains can be seen along the a-axis, π-stacking occurs along b-axis, and c is the chain direction.

Chapter 1 :

 1 Fundamental concepts and state of the art 27 observed for C14-PBTTT than for C12-PBTTT. Based on X-ray scattering studies, a simple packing model was proposed for PBTTT. This model suggests interdigitation of alkyl side chains (see figure 1.11).

Figure 1 .

 1 Figure 1.12. A) Simulated model representing the packing in C14-PBTTT viewed along the b and c-axes. The polymer backbone is shown in orange. The alkyl side-chains are shown in blue and green. B) Structural parameters as a function of side-chain length in the simulated PBTTT unit cells. Reproduced from(33)

Chapter 1 :

 1 Fundamental concepts and state of the art 28 a triclinic unit cell for C14-PBTTT with lattice parameters a= 21.5Å, b=5.4 Å, c=13.5 Å, α=137°, β=86°=ϒ=89°. These parameters were consistent with the previous results from McCulloch on C12-PBTTT. In addition to this, they demonstrated that the PBTTT backbone is planar and tilted by 14 ° to the substrate normal. The π-stacking distance was calculated as 3.58 Å. This short π -stacking distance leads to strong π overlaps. Alkyl side chains are highly interdigitated and form a lamellar structure with a d spacing of 21.3 Å.

Chapter 1 :

 1 Fundamental concepts and state of the art 30 oriented P3HT films with a dichroic ratio of 5.1 with 800 % enhancement of mobility with respect to non-oriented samples.

Figure 1 . 13 .

 113 Figure 1.13. Schematic illustration of the rubbing method. Reproduced from(36)

Figure 1 .

 1 Figure 1.14. A) UV-Vis polarized optical absorbance spectra of an oriented P3HT thin film (17.2 kDa). Parallel and perpendicular absorbance are shown in the figure. B) & D) High-resolution TEM image and electron diffraction pattern of as-rubbed P3HT thin films (17.2 kDa). C) Evolution of the dichroic ratio for rubbed P3HT films of different molecular weights. E) Schematic illustration of edge on to face-on transformation duringrubbing. All films were aligned at room temperature. Reproduced from(49)

Figure 1 .

 1 Figure 1.15. A) UV-Vis polarized optical absorbance spectra of an oriented PBTTT thin-film rubbed at 125 °C (Parallel and perpendicular absorbance are shown in the figure). B) Evolution of the dichroic ratio as a function of the rubbing temperature (TR). C) and D) are the electron diffraction and high-resolution electron microscopy images of oriented PBTTT rubbed at 125 °C. E) Schematic illustration of edge-on to face-on transformation duringrubbing. Reproduced from(44)

Figure 1 . 16 .

 116 Figure 1.16. Temperature dependence of the dichroic ratio in A) P3HT and B) PBTTT for samples with different molecular weights (Mw). Reproduced from(36)

Chapter 1 : 35 Figure 1 . 17 .

 135117 Figure 1.17. A) Schematic representation of the friction transfer deposition method. B) SEM image of an aligned P3HT film-oriented by friction transfer. C) and D) represent corresponding polarized Uv-Vis spectra along the parallel and perpendicular direction and the section profile from Grazing incidence X-ray diffraction respectively.

Figure 1 .

 1 Figure 1.18. A) Evolution of the dichroic ratio as a function of applied stress. B) Anisotropy of hole mobilities instrain aligned P3HT films in OFETs. Reproduced from(52)

Figure 1 . 19 .

 119 Figure 1.19. A) Schematic representation of epitaxial growth of P3HT on potassium 4-bromobenzoate. B) TEM bright-field image showing an oriented and nanostructured P3HT film after removal of the KBrBz. C) Electron diffraction pattern of crystallized P3HT thin films at 180 C on KBrBz. Reproduced from(55)

Chapter 1 : 40 Figure 1 . 20 .

 140120 Figure 1.20. Schematic representation (top) of the semi-crystalline morphology of P3HT. The green colour represents the crystalline regions. The intra and intermolecular charge transport in the crystalline region are represented. The amorphous zones and the corresponding mobilities are shown in the bottom. Reproduced from(57)

Chapter 1 :

 1 Fundamental concepts and state of the art 45 S2-S10. Generally, thermoelectric devices of inorganic materials are made by a combination of n-type and p-type materials. Each leg contributes to the total opencircuit voltage of the TE device (Voc).

Figure 1 .

 1 Figure 1.21. A) Diagram of a thermocouple made of n-type and p-type materials. Power generation or refrigeration modes are possible depending on the configuration. I is the current. B) A conventional thermoelectric module that comprises an array of elements, which are connected electrically in series but thermally in parallel. Reproduced from (3&66)

Figure 1 . 22 .

 122 Figure 1.22. The basic principle of acid-base doping (Top) and redox doping (bottom) that involves the transfer of an electron from the LUMO or from the HOMO. Reproduced from (3)

Figure 1 . 23 .

 123 Figure 1.23. Different models for explaining electron transport from P3HT to the dopant F4TCNQ Reproduced from(75)

Figure 1 . 24 .

 124 Figure 1.24. A) Common P-type dopants and B) Polymers used for thermoelectric applications. Reproduced from(5)

Figure 1 . 25 .

 125 Figure 1.25. Chemical structure of P3HT A) and the corresponding polaronic B) and bipolaronic C) structures. D) Energy levels of P3HT corresponding to the structural change. E) Doping induced absorption bands of RR P3HT doped with Iodine. Reproduced from (83&84).

Chapter 1 :

 1 Fundamental concepts and state of the art 54 new state called polaron. It is associated with a new energy level below the bottom of the CB and a level above the VB (see figure 1.25). This reduces the energy gap. These new energy levels produced within the bandgap lead to additional Mid and NIR transitions (labelled as P1 and P2 transitions) in the absorption spectra of doped P3HT

Chapter 1 :

 1 Fundamental concepts and state of the art 55 improve the PF. The doping and electrical conductivities of conjugated polymers such as polythiophenes can be affected by parameters such as regio-regularity, solid-state order, initial crystallinity of polymer films, IE of polymers, size and EA of the dopants, different doping methods, alkyl side chain length of the polymers and polymer orientation methods. The next section discusses these parameters in detail. a) Regioregularity, solid-state order and crystallinity: The influence of regioregularity of P3HT on electrical conductivity was studied by Hynynen and coworkers on vapor doped F4TCNQ: P3HT thin films. The study showed that doping of highly regio-regular P3HT leads to electrical conductivities of  5.3 S/cm while the amorphous P3HT domains are poorly conducting. Moreover, the electrical conductivity of the regio-regular P3HT was further improved to 12.7 S/cm by increasing the solid-state order of P3HT thin films by processing from p-xylene. The enhanced electrical conductivity in regio-regular and p-

Chapter 1 :

 1 Fundamental concepts and state of the art 59 the observed conductivity at this strong doping regime might be due to the intercalation of dopant into the crystalline regions of the polymer (see figure 1.27B).

Figure 1 . 26 .

 126 Figure 1.26. Different doping methods used for polythiophenes. A) Mixed solution doping of P3HT. The dopant and polymer are mixed in a common solvent before casting thin films. B) Sequential doping method. The polymer is spin-coated on a suitable substrate and then the dopant dissolved in an orthogonal solvent is spin-coated on

Chapter 1 :

 1 Fundamental concepts and state of the art 62 morphology and crystal structure of the polymer. Moreover, it is important to intercalate the dopant into the polymer matrix without modifying too much its initial crystal structure.

Figure 1 . 28 .

 128 Figure 1.28. Morphological changes in mixed solution and sequential doping. In the mixed solution doping A)Initially, solvated P3HT tend to aggregate in solution upon addition of F4TCNQ. These aggregates form disconnected domains with few tie chains. This prevents the delocalization of polarons. In the sequential doping B) the initial morphology of P3HT is maintained after sequential doping with Acetonitrile-F4TCNQ and the F4TCNQ -anions reside in the amorphous regions of P3HT. Incorporation of dopants into crystalline parts occurs by sequential doping from chlorobenzene -F4TCNQ mixture. Reproduced from(96)

Figure 1 .

 1 Figure 1.29. A) Out of plane and B) In-plane scattering profiles of F4TCNQ doped P3HT thin films extracted from GIXD e. C) 2D GIXD pattern of F4TCNQ doped P3HT thin film. Reproduced from(97)

Figure 1 .

 1 Figure 1.30. A) Method of preparation of oriented conducting polymer films by a combination of hightemperature rubbing and doping from a solution of F4TCNQ in acetonitrile. B) Polarized UV-Vis-NIR spectroscopy of oriented P3HT thin films doped by F4TCNQ in acetonitrile. The polaronic bands (P1, P2 and P3) are polarized along the rubbing direction while the F4TCNQ -anions are polarized perpendicular to the direction of rubbing. C) and D) Electron diffraction pattern of as-rubbed and doped P3HT thin films upon sequential doping with F4TCNQ in acetonitrile. E) Illustrates the corresponding structural changes in the P3HT unit cell as a result of the intercalation of F4TCNQ into the alkyl side chain layers. Reproduced from(46).

Figure 1 .

 1 Figure 1.31. A) and B) GIXD images of the solution and vapor doped F4TCNQ/C14-PBTTT thin films respectively.

Chapter 2 .

 2 Impact of alkyl side chain length on doping kinetics, crystal structure & thermoelectric properties of oriented PBTTTDoping of polymer semiconductors such as PBTTT with acceptor molecules such as F4TCNQ is widely used to tune the charge transport and thermoelectric (TE) properties in thin films. However, the mechanism of dopant insertion in the polymer matrix, the insertion kinetics and the ultimate doping levels have been investigated only marginally. This chapter addresses the impact of alkyl side chain length on the doping mechanism of a series of PBTTTs with linear side chains ranging from n-octyl to n-octyldecyl. The study focuses on thin films oriented by high temperature rubbing and sequentially doped in F4TCNQ solution. Structure-property correlations are established as a function of side chain length by a combination of transmission electron microscopy, polarized UV-Vis-NIR spectroscopy, charge transport and thermopower measurements. Intercalation of F4TCNQ into the layers of the alkyl side chains results in the expansion of the lattice along the side chains and the contraction along the π-stacking direction for all polymers. The extent of lattice expansion decreases with increasing side chain length. UV-Vis-NIR spectroscopy demonstrates integer charge transfer for all investigated PBTTTs. The doping kinetics and final doping level depend on both the side chain length and packing. Highly disordered n-octyl and crystalline n-octyldecyl side chain layers tend to hamper dopant diffusion in the side chain layers contrary to n-dodecyl side chains that can host the highest proportion of dopants. Consequently, the best TE properties are observed for C12-PBTTT films. Alignment of the polymers enhances significantly the TE performance by increasing the charge conductivity and the thermopower along the 1. Introduction Doping of semi-conducting polymers is a simple yet powerful way to control their charge transport properties by tuning the charge carrier densities. Depending on the doping level, different applications can be envisioned in plastic electronics. Low dopant concentrations can be used in blends of semi-conducting and insulating polymers to

Figure 2 . 1 .

 21 Figure 2.1. Chemical structures of A) C8-PBTTT B) C12-PBTTT C) C14-PBTTT and D) C18-PBTTT. E) Schematic illustration of high temperature rubbing and doping of PBTTT thin films by dipping in the dopant F4TCNQ solution in acetonitrile.

  exhibit high orientation prior to and after doping with F4TCNQ. The kinetics of F4TCNQ doping in solution is measured for four PBTTTs, namely C8-PBTTT, C12-PBTTT, C14-PBTTT and C18-PBTTT. The structure of the films with maximum doping level is investigated by low dose electron diffraction and polarized UV-Vis spectroscopy. Finally, the anisotropic Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 88 charge transport and Seebeck coefficients are measured in the rubbed thin films.

Figure 2 .

 2 Figure 2.2 illustrates the high birefringence in the doped films and the change in polarized UV-Vis absorption for two PBTTTs with C12 and C18 side chains. The polarized optical microscopy (POM) images show that the birefringence and thus orientation of the films are preserved after doping. Polarized UV-Vis-NIR spectroscopy was performed with the light polarization either parallel or perpendicular to the rubbing direction. It evidences the presence of highly polarized signals from the polaron (P1 and P2 bands), the neutral polymer (N) and from the F4TCNQ -anion.

Figure 2 . 2 .

 22 Figure 2.2. Example of alignment induced by high temperature rubbing (TR=125°C) for C12-PBTTT and C18-PBTTT thin films and evolution upon doping in solution with F4TCNQ/ACN (1mg/ml). A) and D) Polarized UV-Vis spectra prior to doping. B) and E) POM images under crossed polarizers showing a high birefringence of the rubbed filmswhen the films are oriented with the rubbing direction (R) at 45° to the polarizer and full extinction when R is parallel to the polarizer. The rubbing direction is given by the arrow R. C) and F) Polarized UV-Vis-NIR spectra of the doped PBTTT films at saturation, in red when the light polarization is parallel to the rubbing direction (POL//R), in black when the light polarization is perpendicular to the rubbing direction (POL⊥R). The main spectral features are highlighted. P1 and P2 correspond to polaronic bands of doped PBTTT. The F4TCNQ -bands are highlighted as well as the neutral polymer N and the amorphous non-oriented fraction of the polymer 'a'.

2. 2 .

 2 Doping kinetics using polarized UV-Vis-NIR spectroscopy Sequential doping of polymer semi-conductors is a complex mechanism based on the diffusion of dopants in the semi-crystalline polymer matrix coupled to a redox reaction between polymer and dopant. So far, little attention has been paid to the doping kinetics of semiconducting polymers such as PBTTT. It was mostly assumed that the doping is fast and almost instantaneous. Hereafter, UV-Vis-NIR spectroscopy is used to follow the kinetics of doping versus doping time. Oriented films of various PBTTTs are doped sequentially by increasing the time of contact between the film and the dopant solution in acetonitrile (1mg/ml). The evolution of the UV-Vis spectrum in both parallel (POL//R) and perpendicular (POL⊥R) polarizations vs rubbing direction are shown in figures 2.3 and 2.4 for four PBTTTs with different side chains, respectively.

Chapter 2 :

 2 Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 91 peak of the neutral polymer decreases substantially for C8-PBTTT, C12-PBTTT and C14-PBTTT whereas a smaller decrease of absorbance is evidenced in the case of C18-PBTTT.

Figure 2 . 3 .

 23 Figure 2.3. Kinetics of PBTTT doping with F4TCNQ in acetonitrile (1mg/ml) for POL//R as a function of side chain length. Polaronic bands are labelled P1 and P2. Note that the P2 band overlaps with a small contribution from the F4TCNQ -anion. The spectra for the saturation of doping are highlighted in red. The absorption spectra of the undoped films are shown in black.

Figure 2 . 4 .

 24 Figure 2.4. Kinetics of PBTTT doping with F4TCNQ in acetonitrile (1mg/ml) for POL⊥R as a function of side chain length. Polaronic bands are labelled P1 and P2. The spectra after saturation of doping are highlighted in red. The absorption spectra of the undoped films are shown in black.

Figure 2 . 5 .

 25 Figure 2.5. A) Doping kinetics of PBTTTs with F4TCNQ in acetonitrile (1mg/ml) for various lengths of alkyl side chains as indicated by the time dependence of the F4TCNQ -anion absorption peak (0-0 vibronic band in the range 870-883 nm). The representation against t 1/2 /l is independent of the nominal film thickness l and allows to extract the diffusion coefficient D for each polymer. For clarity, the kinetics curves are shifted along the ordinate axis.

Figure 2 . 6 .

 26 Figure 2.6. Differential Scanning Calorimetry on PBTTTs powder. Heating and cooling DSC curves showing the side chain melting/crystallization for C8-PBTTT, C12-PBTTT, C14-PBTTT and C18-PBTTT. B) Comparison of the UV-Vis-NIR spectra of F4TCNQ-doped C18-PBTTT films at room temperature and 70°C i.e. above the melting of C18 side chains. Note the higher intensity of polaronic bands and the decreased absorbance of the neutral polymer when doped at 70°C.

Chapter 2 :

 2 Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 98 1060 nm to extract the absorbance of the F4TCNQ -anion at 875-880 nm. The volume concentration of thiophene in the films (Cthiophene) was determined from the volume of the unit cell of the doped phases (we assume orthorhombic unit cells with two chains per unit cell). The cell parameters of the doped phases of the PBTTTs are determined from the values of d003, d100, and d020 of the diffraction patterns at saturation.

Figure 2 . 7 .

 27 Figure 2.7. Cyclic voltammograms, in the anodic region, of PBTTT thin-films in ACN containing 0.1 M tetra-nbutylammonium hexafluorophosphate at 20 °C; scan rate 100 mVs -1 . All peaks were calibrated against ferrocene (Fc + /Fc) at E½ = 0.41 V vs SCE.

Figure 2 . 8 illustratesChapter 2 : 101 Figure 2 . 8 . 2 . 4 .

 2821012824 Figure 2.8. Correlation between the absorbance of the polaron peak P1 at 2000 nm and the intensity of the F4TCNQ -anion peak at 875 nm as observed for different PBTTTs doped in solution of F4TCNQ/ACN (1mg/ml) at various time intervals.

Figure 2 . 9 .

 29 Figure 2.9. Evolution of the electron diffraction patterns of oriented thin films of various PBTTTs in as-rubbed films (TR=125°C) and after doping with F4TCNQ at different doping times. The reticular distances for the interlayer spacing d100 and for the π-stacking period 020 are given in orange. The arrow R indicates the chain direction induced upon rubbing.

Figure 2 . 10 .

 210 Figure 2.10. Equatorial section profiles of the ED patterns for thin films oriented by rubbing of PBTTTs with different side chains upon doping at different time intervals in a solution of F4TCNQ in acetonitrile for: A) C8-PBTTT, B) C12-PBTTT and C) C14-PBTTT

Figure 2 . 11 .

 211 Figure 2.11. Kinetics of the structural changes associated to the interlayer spacing d100 and to the -stacking d020 observed upon doping of the PBTTT thin films in a solution of F4TCNQ in acetonitrile for different side chain lengths.

Chapter 2 :

 2 Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 105 agreement with the kinetics of doping evidenced by UV-Vis spectroscopy. For C18-PBTTT, the evolution of the unit cell parameters is slower and similar to that observed in UV-Vis-NIR spectroscopy. More generally, for all PBTTTs the kinetics of the structural changes observed by TEM agree with the doping kinetics evidenced by UV-Vis spectroscopy. More precisely, for C8-PBTTT, a clear expansion of the interlayer spacing is observed from 15.1 Å to 16.4 Å and a contraction of π-stacking from 3.75 Å to 3.60 Å.

Chapter 2 :Figure 2 . 12 .

 2212 figure 2.11 D). In other words, the shorter the alkyl side chains, the larger the variation of the unit cell parameters upon doping.

107 Figure 2 . 13 .

 107213 Figure 2.13. Evolution of the charge conductivity, Seebeck coefficient and power factor as a function of increasing doping concentration of F4TCNQ for oriented thin films of C12-PBTTT. All values are measured along the directions parallel (//) and perpendicular (⊥) to the rubbing.

  3). The lowest conductivities are observed for C18-PBTTT with the longest side chains (// of a few S/cm), which is consistent with the fact that it has the smallest doping level at saturation (3.3%). The polymers C14-PBTTT and C8-PBTTT show intermediate doping levels around 7% and their conductivity // lies in the 15-40 S/cm range.

Figure 2 . 14 .

 214 Figure 2.14. Comparison of the TEM bright field images of C8-PBTTT films oriented by high temperature rubbing prior to doping (A) and after doping (B) with F4TCNQ/ACN (0.5 mg/ml). The insets correspond to the Fast Fourier Transforms (FFTs) and indicate a lamellar periodicity of 22±2 nm in both cases.

  whereas the anisotropy of conductivity /// ⊥ was in the range 50-100. Kaiser as well as Pukaci et. al proposed that the anisotropy of Seebeck coefficient is related to the heterogenous nature of the films where both metallic and semi-conducting domains coexist. This interpretation would also apply in the present case since the UV-Vis spectra of all doped films show the coexistence of doped (metal-like) and undoped (semiconducting) domains in the thin films (see figure2.3 and 2.4). It can be anticipated that Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 110 such a co-existence of doped and undoped phases could be beneficial in oriented films to observe simultaneously a high charge conductivity and a high Seebeck coefficient. In chapter 3, we develop in more details the heterogenous system model to account for the observed differences in Seebeck coefficients.

Figure 2 . 15 . 111 (

 215111 Figure 2.15. (A-C) Evolution of the charge conductivity versus doping time for C8-PBTTT, C14-PBTTT and C18-PBTTT respectively. (D-F) Evolution of the Seebeck coefficient versus doping time for C8-PBTTT, C14-PBTTT and C18-PBTTT respectively. All values are measured along the directions parallel (//) and perpendicular (⊥) to the rubbing.

  of S// with doping time parallels the kinetics of doping evidenced by UV-Vis spectroscopy versus side chain length. Both C8-PBTTT and C18-PBTTT have the slowest doping kinetics. Hence the Seebeck coefficients observed at low doping times are particularly high, especially for C18-PBTTT with S//> 400 µV/K for t= 5 s. Conversely, for C12-PBTTT and its fast doping kinetics, the variation of S with doping time is limited.

Figure 2 . 16 .

 216 Figure 2.16. Evolution of the maximum conductivity (//) and power factor (PF) along the rubbing direction (a) and of the average charge conductivity anisotropy ///⊥ (b) versus side chain length in F4TCNQ-doped PBTTT films oriented by high temperature rubbing.

Chapter 3 . 2

 32 PBTTTs was investigated from various points of view including structural variation, UV-Vis-NIR spectroscopic and thermoelectric properties. UV-Vis-NIR spectroscopy and DSC evidence the key role of side chain length and packing on the doping mechanism and helped determine the dopant diffusion coefficients and doping levels at saturation. Both, disordered C8 and crystalline C18 side chain layers hamper efficient dopant diffusion in the polymer films whereas loosely packed side chain layers of C12 and C14-PBTTT allow for fast dopant diffusion. The observed charge conductivity of the oriented PBTTT films correlates well with the doping level at saturation: an optimum in TE properties is observed for C12-PBTTT films. For all PBTTTs, alignment enhances substantially the TE performances by increasing simultaneously the charge conductivity and the thermopower along the chain direction. Aligned films of C12-PBTTT show charge conductivities of 193 S/cm along the rubbing direction and power factors of Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 113 approximately 100 µW . m -2. K -1 . More generally, this study underlines the importance to perform a controlled crystallization and orientation of the polymer semiconductors prior to doping from solution in order to enhance their TE properties. On-going studies indicate that the same trends observed for solution-doped PBTTTs are also valid when doping from the vapor phase. Bringing conducting polymers to high order: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW . m -1. K -This chapter presents an effective design strategy of polymer thermoelectric materials based on structural control in doped polymer semiconductors. The strategy is illustrated for two archetypical polythiophenes e.g. poly(2,5-bis(3-dodecyl-2-thienyl)thieno[3,2b]thiophene) (C12-PBTTT) and regioregular poly(3-hexylthiophene) (P3HT). FeCl3 doping of aligned films results in charge conductivities up to 2 . 10 5 S/cm and metallic-like thermopowers similar to iodine-doped polyacetylene. The films are almost optically transparent and show strongly polarized near-infra-red polaronic bands (dichroic ratio > 10). The comparative study of structure-property correlations in P3HT and C12-PBTTT identifies three conditions to obtain conductivities beyond 10 5 S/cm: I. Achieve high in-plane orientation of conjugated polymers with high persistence length II. Ensure uniform chain oxidation of the polymer backbones by regular intercalation of dopant molecules in the polymer structure without disrupting the alignment of π-stacked layers. III. Maintain a percolating nano-morphology along the chain direction. IV. The highly anisotropic conducting polymer films are ideal model systems to investigate the correlations between thermopower S and charge conductivity . A scaling law S -1/4 prevails along the chain direction but a different Sln relation is observed perpendicular to the chains, suggesting different charge transport mechanisms. The simultaneous increase of charge conductivity and Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 1. Introduction Conducting polymers were at the forefront of research in plastic electronics several decades ago and opened the new and vast research field of plastic electronics. A. Heeger, A. G. McDiarmid and H. Shirakawa were among the pioneers in this domain and their success on conducting polymers were recognized by the award of a Nobel Prize in 2000

Figure 3 . 1 .Chapter 3 :

 313 Figure 3.1. A) Chemical structure of P3HT and C12-PBTTT. B) Schematic illustration of the two-step process to prepare oriented conducting thin films combining high temperature rubbing and subsequent doping in a solution of FeCl3 in nitromethane. C) Illustration of the color change in rubbed C12-PBTTT films (left) upon doping with 1 mM and 10 mM FeCl3/nitromethane solutions. D) Polarized optical microscope images under crossed polarizers showing the high birefringence in the as-rubbed and doped C12-PBTTT films.

figure 3 . 1 )

 31 figure 3.1) yields highly oriented and crystalline conducting polymer films. The same strategy was applied to P3HT and C12-PBTTT doped with FeCl3. A drastic change of the film's color from purple to light blue (see figure 3.1 C) is observed after doping the polymer films with 10mM FeCl3 in nitromethane. polarized optical microscopy (POM)

Figure 3 . 2 (

 32 A to D) compares the electron diffraction patterns of the pristine and FeCl3-doped P3HT and PBTTT oriented films. The preserved orientation in the doped films is also demonstrated by transmission electron microscopy (TEM). The electron diffraction (ED) patterns of pristine oriented films indicate that the chain direction is parallel to the rubbing direction R and that crystalline domains are oriented mainly face-on the substrate in C12-PBTTT whereas face-on and edge-on crystals coexist in oriented P3HT films. Both the in-plane orientation and crystal contact plane are maintained after doping for P3HT and C12-PBTTT for the whole range of [FeCl3]. However, on closer inspection, the ED patterns after doping for [FeCl3] = 10 mM reveals a loss of crystallinity characterized by a partial loss of higher-order equatorial h00 reflections, a broadening of the remaining h00 and a reduction in the intensity of the meridional 002 reflection. As seen in figure 3.2.C&D and doping modifies the unit cell parameters of C12-PBTTT and P3HT but to different extents (see also figure 3.3 A and B). For both polymers, the unit cell expands along the alkyl side chain direction (d100) and the π-stacking periodicity Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 126 d(020) is reduced, demonstrating that dopant molecules penetrate inside the layers of alkyl side chains of the crystalline phase and that the change in lattice parameters correlate with the amount of incorporated dopants. For both polymers, the interlayer spacing d100 increases with [FeCl3] and saturates for [FeCl3] approaching 10 mM. A maximum of doping has been reached. d100 increases from 19.7 Å (pristine) to 23.6 Å (10 mM) in C12-PBTTT and from 16.4 Å (pristine) to 18.1 Å (10 mM) for P3HT. In the case of C12-PBTTT, the increase in lattice spacing is much larger than for F4TCNQ.  d100 is equal to 4 Å for FeCl3 doped C12-PBTTT and 0.9 Å for F4TCNQ.

Figure 3 . 2 .

 32 Figure 3.2. Structure evolution in doped P3HT and C12-PBTTT upon doping with FeCl3 as obtained by low dose TEM. Electron diffraction patterns of C12-PBTTT before A) and after C) doping with 10 mM FeCl3. Electron diffraction patterns of P3HT before B) and after D) doping with 10 mM FeCl3. The asterisk marks a Scherrer ring from excess FeCl3 on the surface of the doped P3HT film.

Figure 3 . 4 (

 34 Figure 3.4 (A to F) shows the elemental analysis in terms of EDX spectra and 2D maps for 5 mM FeCl3-doped C12-PBTTT and P3HT. First, one can see that there is a uniform

Figure 3 . 3 .Chapter 3 : 2 129Figure 3 . 4 .

 333234 Figure 3.3. A) Interlayer spacing d100 and π-stacking distance versus FeCl3 concentration in nitromethane for A) C12-PBTTT films and B) P3HT films. C) Equatorial section profiles of the ED patterns of oriented C12-PBTTT films upon doping with FeCl3 in nitromethane at different concentrations. The profiles have been shifted along the ordinate axis for clarity. D) Illustration of the structural change induced by FeCl3 doping for C12-PBTTT ([FeCl3] =10 mM) that results in the formation of a bimolecular crystal.

. 5 A 2 131Figure 3 .

 523 Figure 3.5 Low dose HRTEM image of FeCl3-doped P3HT A) and C12-PBTTT B) films aligned by high temperature

Chapter 3 :

 3 Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 132

2. 4 .

 4 Spectroscopic signatures of doping: Polarized UV-Vis-NIR spectroscopy TEM has shown that the amount of FeCl3 incorporated in the crystals of C12-PBTTT and P3HT depends on the concentration [FeCl3]. This trend is further supported by polarized UV-vis-NIR spectroscopy that helps visualize the level of chain oxidation versus [FeCl3] (see figure 3.6). Increasing the doping concentration results in a corresponding increase of the polaronic features P1 and P2 and a decrease of the neutral polymer absorption (N). For both polymers, the NIR absorption bands P1 and P2 are polarized parallel to the alignment direction R and the films are strongly birefringent under the polarizing microscope. Most remarkably, the P3HT and PBTTT films doped at 10 mM exhibit a spectrum without significant contribution of the original neutral polymer (1-2%),

Chapter 3 :

 3 Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 133 extension of aligned and planarized chain segments observed by HR-TEM in C12-PBTTT (see figure3.5 E).

Figure 3 . 6 .

 36 Figure 3.6. Evolution of the polarized UV-vis-NIR spectrum in oriented P3HT and C12-PBTTT thin films prepared by high temperature rubbing and subsequently doped in solution of FeCl3/nitromethane of increasing concentration. The light polarization is parallel (//) to the rubbing direction or perpendicular to (⊥). The asterisk highlights the absorption peak of the excess FeCl3 present on the film surface.

Chapter 3 :

 3 Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 134 Typically, C12-PBTTT films doped with 10 mM FeCl3 have a dichroic ratio of the P2 and P1 band of 10.6, which is very close to 10.8 of undoped C12-PBTTT. These similar values confirm that FeCl3 doping does not alter substantially the orientation of the C12-PBTTT chains. For oriented P3HT films, the situation is different: the dichroic ratios of the P1 and P2 bands decreases with increasing [FeCl3] (see figure3.7 C). This loss of orientation is possibly related to the loss of both alignment and crystallinity upon intercalation of dopant molecules. This is consistent with the decreased intensity of ED reflections and the loss of vibronic structure in the absorption of the non-doped phase (N).Interestingly, UV-vis spectra for POL ⊥ R (see figure3.6) indicate that the amorphous fraction of P3HT

Figure 3 . 7 .

 37 Figure 3.7. Polarized UV-Vis spectra of rubbed thin films of C12-PBTTT A) and regioregular P3HT (B). C12-PBTTTwas rubbed at 125°C and P3HT at 200°C. The absorption spectra recorded for incident light parallel (//) to the rubbing show a vibronic structure and correspond to planarized polymer chains forming -stacks. The featureless spectra obtained for incident light oriented perpendicular (⊥) to the rubbing correspond to the amorphous phase of the polymers. For P3HT, the 0-0 and 0-1 bands are centred at 605 nm and 560 nm whereas the amorphous phase shows a maximum at 506 nm. Likewise, for C12-PBTTT the 0-0 and 0-1 bands are centred at 526 nm and 552 nm whereas the amorphous phase shows a maximum at 526. C) Evolution of the dichroic ratio at 2500 nm of the polaronic P1 band in oriented P3HT films (TR=180°C) as a function of increasing FeCl3 concentration

Chapter 3 : 2 138Figure 3 . 8 .

 3238 Figure 3.8. Variation of the charge conductivity  (A and D), Seebeck coefficient S (B and E) and power factor PF (C and F) as a function of increasing FeCl3 concentration for doped P3HT and C12-PBTTT oriented films. , S and PF are measured along the rubbing direction (//) and perpendicular to it (⊥). The lines are guide to the eye. The blue dotted lines in D) and F) show the typical values of conductivity and power factors reached for PA doped with iodine, respectively.

Chapter 3 :

 3 figure 3.5 D) results in a larger conductivity //.

Figure 3 . 9 .Chapter 3 :

 393 Figure 3.9. Schematic representation for the observed differences in the Seebeck coefficient of oriented polymer films. A) When the rubbing direction is // to the applied temperature gradient and B) Rubbing direction is perpendicular to the applied temperature gradient.

  impurities. From an experimental point of view, Chabinyc and coworkers made a detailed examination of electrical conductivity measurements of polymers such as PBTTTC14, P3HT and P2TDC17FT4 (poly(2,5-bis(thiophen-2-yl)-(3,7diheptadecantyltetrathienoacene) by varying the dopant concentration and obtained a striking relationship between the thermopower and electrical conductivity for a wide range of conductivity for these P-doped polymers. Chabinyc et al. showed that S and  are related via the relation S -1/4

Figure 3 . 8 .

 38 Figure 3.8. Correlation between thermopower S and charge conductivity. A) F4TCNQ and FTS doped non-oriented films of PBTTT, P3HT, P2TDC17 -FT4, which follows the relation S//  -1/4 . B) Highly oriented thin films of C12-PBTTT (TR=125°C) and C) P3HT (TR=180°C) doped with FeCl3. // and ⊥ refer to the direction of the measurements

Figure A is reproduced from reference 52 .

 52 Figure A is reproduced from reference 52.

  ), one obtains 𝜎 ⊥ 𝑚𝑎𝑥 =3809 S/cm and β=26.7 for C12-PBTTT and 𝜎 ⊥ 𝑚𝑎𝑥 =211 S/cm and β=17.5 for P3HT. These results clearly indicate different intrinsic limitations in charge transport perpendicular to the chain direction for both polymers: charge transport perpendicular to the chains is larger in C12-PBTTT than in P3HT. Regarding the factor β, previous studies on polyaniline reported β  9 whereas much higher values are obtained in the present case for FeCl3-doped P3HT and C12-PBTTT. In oriented polyaniline samples, Mateeva et al. observed that equation (3) is verified both parallel and normal to the chain direction (albeit for a reduced domain of conductivities well below that measured in the present case). This is clearly not true for C12-PBTTT and P3HT that both show a power-law dependence along the chain direction and an apparent logarithmic dependence perpendicular to the chains. Altogether, these results indicate that the S- correlations in C12-PBTTT are different along the chain Chapter 3: Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 149

Chapter 3 :

 3 Bringing conducting polymers to high orders: towards conductivities beyond 10 5 S/cm and thermoelectric power factors of 2 mW.m -1 . K -2 150 dopants intercalate progressively into the layers of side chains by incremental increase of the dopant concentration in such a way that dopant intercalation preserves orientation and -stacking of the conjugated polymer backbones. Three conditions are identified to obtain conductivities of 10 5 S/cm: i)Achieve high in-plane orientation of conjugated polymers with a high persistence length ii) Ensure regular intercalation of dopant molecules in the polymer structure to form bimolecular polymer-dopant crystals and iii) Maintain a percolating nano-morphology along the chain direction e.g. via a liquid-crystal-like morphology (semi-crystalline lamellar morphologies are detrimental). iv) Moreover, using a heterogeneous model approach based on the semi-crystalline structure of the doped and aligned polymer films, we are able to account for the anisotropy of the Seebeck coefficient that is enhanced in the direction of alignment.

Figure 4 . 1 .

 41 Figure 4.1. A) Two-step preparation of highly oriented and crystalline C12-PBTTT films. The first step consists of chain alignment and crystallization by high-T rubbing followed by sequential doping in solutions of F4TCNQ or F6TCNNQ in acetonitrile. B) Polarized optical microscopic images of F 6TCNNQ doped C12-PBTTT showing high birefringence (TR =125C and CF4TCNQ= 1mg/ml) C) Schematic representation of HOMO and LUMO levels F4TCNQ, F6TCNNQ and PBTTT.

Figure 4 . 2 .DDChapter 4 :

 424 Figure 4.2. Schematic illustration of incremental concentration doping (ICD) and direct doping.

Figure 4 . 3 .Chapter 4 :

 434 Figure 4.3. A) Absorption bands of F4TCNQ -anions in doped PBTTT thin films and absorption of neutral F4TCNQ in solution. B) Absorption in solution of neutral F6TCNNQ (blue)andF6TCNNQ -anion (brown) after treating withFerrocene. The absorption of F6TCNNQ -anion in thin films is shown in red.

Figure 4 . 4 .

 44 Figure 4.4. Concentration dependent doping of C12-PBTTT with F4TCNQ and F6TCNNQ in acetonitrile. A) and B) represent the evolution of the UV-Vis-NIR spectra for POL//R as a function of doping concentration for F6TCNNQ - and F4TCNQ-doped C12-PBTTT, respectively. C) and D) show the evolution of the F6TCNNQ -and F4TCNQ -bands as a function of increasing doping concentration for POL ⊥R. P1 and P2 correspond to polaronic bands of doped C12-PBTTT and the neutral polymer absorption is labelled as N.

Figure 4 . 5 . 4 :

 454 Figure 4.5. A) and B) represents UV-Vis-NIR absorption spectra of oriented C12-PBTTT (before doping) plotted for different angles (every 15°) between the direction of polarization and the rubbing direction. C) and D)Polarized UV-Vis-NIR absorption spectra of F6TCNNQ doped PBTTT and F4TCNQ doped PBTTT plotted for

Figure 4 . 6 .

 46 Figure 4.6. Absorbance of the dopant anions after subtracting the background absorbance of the polaronic contribution P2 for A) F4TCNQ -and B) for F6TCNNQ -. C) Angular distribution of absorption bands of pristine C12-PBTTT before doping, the F4TCNQ -anion (= 873nm) and the P1 band (= 2500 nm). D) Angular distribution of the absorption bands of oriented pristine C12-PBTTT before doping, the F6TCNNQ -anion (= 1158 nm) and the polaronic P1 band (= 2500 nm).

Figure 4 . 8 .

 48 Figure 4.8. A) Plot of normalized absorptions of both F 6TCNNQ (λ=1158 nm) and F4TCNQ (λ=1158 nm) vs t 1/2 /l to determine the diffusion coefficients. B)-C) Plot of absorbance of the dopant anions vs that of the polarons (2500 nm). The linear fit supports integer charge transfer between C12-PBTTT and both F4TCNQ and F 6TCNNQ.

Figure 4 . 8 .

 48 Figure 4.8. Evolution of the charge conductivity (A and B) and of the Seebeck coefficient (C and D) as a function of increasing doping concentration of F4TCNQ and F6TCNNQ in oriented films of C12 -PBTTT, respectively. The curves correspond to the sample doped by incremental increase of doping concentration (ICD). Both S and  are measured parallel (red) and perpendicular (black) to the rubbing directions. For comparison, we also show the

Figure 4 . 9 .

 49 Figure 4.9. Polarized UV-Vis-NIR spectra of doped PBTTT thin films doped by ICD and dipping in acetonitrile. In the ICD, the thin film is dipped in increasing doping concentration of 0.01, 0.1, 0.5, 1 and 2 mg/ml F6TCNNQ in ACN. To exclude the effect of solvent, another thin film dipped 4 times in ACN and then directly doped by dipping in 2mg/ml F6TCNNQ in ACN

Figure 4 . 10 .

 410 Figure 4.10. Evolution of the power factor as a function of increasing doping concentration of A) F4TCNQ and B) F6TCNNQ doped films of C12-PBTTT. Films were oriented by high temperature rubbing at 125°C and doped by incremental increase of doping concentration.

Figure 4 . 11 .

 411 Figure 4.11. Correlation between Seebeck coefficient S and electrical conductivity  in highly oriented thin films of C12-PBTTT (TR=125°C) doped with F4TCNQ (black squares), F6TCNNQ (red squares) and FeCl3 (green squares).

  and F6TCNNQ doped C12-PBTTT thin films by low dose electron diffraction. Given the different TE properties evidenced for samples doped using the ICD or DD method, we also investigate to what extent the two doping methods affect the structure of the C12-PBTTT films.

Figure 4 . 12 .

 412 Figure 4.12. Evolution of the electron diffraction patterns in doped C12-PBTTT upon incremental concentration doping (A-C) and with single doping with F4TCNQ (D-F). R represents the rubbing direction.

Chapter 4 :

 4 Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 188 For TEM investigations, the oriented C12-PBTTT thin films prepared on NaPSS by high temperature rubbing were carbon-coated and the films were recovered on copper TEM grids by floating on water.

Figure 4 . 13 .

 413 Figure 4.13. Evolution of the electron diffraction patterns of F6TCNNQ doped C12-PBTTT (A-C) sequential addition (D-F) single doping as a function of dopant concentration represent the rubbing direction

Figure 4 .

 4 Figure 4.14 depicts the unit cell parameter variation as a function of dopant concentration for both ICD and DD methods. In both cases, doping induces an expansion along the side chain direction of the C12-PBTTT unit cell. The ICD thin films show a larger interlayer spacing variation up to 20.57 Å versus 20.27 Å for DD films. Also, a slight broadening of the diffraction peaks along the equatorial direction is observed for ICD samples (see figure 4.13C) suggesting some disordering in the lamellar packing upon doping. For F6TCNNQ, the same trends are observed as for F4TCNQ: the final lattice expansion at saturation is larger for ICD than for DD (d100 spacing of 20.28 Å vs 20.16 Å

Figure 4 . 14 .

 414 Figure 4.14. Evolution of d100 spacing of the sequentially and directly doped C12-PBTTT. A) Variation of d100 spacing of sequentially and singly doped F6TCNNQ/C12-PBTTT. A) Variation of d100 spacing of sequentially and singly doped F4TCNQ/C12-PBTTT

  highly anisotropic and enhanced along the rubbing direction. The enhancement of the electrical conductivities along the rubbing direction was observed for all PBTTTs. It leads to high PFs along the rubbing direction. Thin films of C12-PBTTT show very high electrical Conclusions and perspectives 205 conductivities of 193 S/cm along the rubbing direction and power factors of approximately 100 µW . m -2. K -1 . This study demonstrates the importance of side-chain engineering, which can be used as a tool to enhance the TE properties. Most importantly, we demonstrated that a combination of polymer orientation and sequential doping can be used to improve the thermoelectric properties of semiconducting polymers.

2. 2

 2 Thin film preparation, orientation by high-temperature rubbing. Thin films of PBTTT were prepared by doctor blading from a hot solution of 10 mg/ml PBTTT in ortho-dichlorobenzene (o-DCB) on a glass substrate at 160°C. Films are then aligned using a homemade rubbing machine. The rubbing machine was originally developed in the Sycommor group (L. Herrmann, P. Allgayer, J.-C. Wittmann and M. Brinkmann) for the preparation of alignment layers and further improved for the orientation of polymer semiconductors used in OFET (L. Biniek et al.

Figure 1 .

 1 Figure 1. A) Front view of rubbing machine B) side view of the rubbing machine.

Figure 2 .

 2 Figure 2. Schematic illustration of the orientation of PBTTT thin films by high-temperature rubbing and doping.

7 .

 7 potential (IP) or highest occupied molecular orbital (HOMO) from the redox data. The Experimental Details 222 HOMO levels were calculated from the following equations, using the oxidation and reduction onset potentials: HOMO (eV) = -[Eonsetox (vs SCE) + 4.4], based on an SCE energy level of 4.4 eV relative to the vacuum. Polarized UV-Vis-NIR spectroscopy The orientation of the polymer films was probed by UV-Vis-NIR absorption (350-2500 nm) using a Cary 5000 spectrometer with polarized incident light and spectral resolution of 1 nm. The UV-Vis-NIR spectra of the doped polymer film were measured along the parallel and perpendicular to the direction of rubbing. The orientation of the film with respect to the polarization of the incident light is controlled by a goniometer. The parallel configuration is the position where the cPBTTT axis is parallel to the light polarization. For each sample, a spectrum in the parallel and perpendicular configuration were measured as a function of doping time. Angle dependent polarized UV-Vis-NIR spectroscopic measurements were performed by rotating the angle of light polarization for different polarization angle from 0 to 90for with an interval of 15.

  them on distilled water and transferred them on the pre-patterned devices. Doing so, we avoided the problem of film dedoping when submitting the films to a secondary vacuum during evaporation. All devices were fabricated on glass substrates, cleaned by Experimental Details 223 ultrasonication in acetone, ethanol, hellmanex, deionized water, and isopropanol. The device fabrication was done by N. Zimmermann at ICUBE. The cleaned substrates were dried under nitrogen prior to use. Gold electrical contacts (40 nm thick) in a four-point probe geometry (1 mm spacing between electrodes, 5 mm length) were deposited via controlled thermal evaporation through a shadow mask, at an average rate of 4-6 Å/s. The first layer of Chromium (2.5 nm thick) was deposited prior to the gold to promote a good adhesion on the glass substrates (evaporation rate 0.5-1 Å/s). The contact geometry used for the electrical conductivity and thermopower measurements is shown in figure3 and 4.

Figure 3 .

 3 Figure 3. Photographic images of the glove box, electrical conductivity and Seebeck coefficient measurement setup.

figure 3 ,

 3 we have first determined the geometrical correction factor C such that  = C R t where t is the film thickness.To that aim, we have used a classical four-point probe system on a non-oriented doped P3HT film to obtain a reference value of the resistivity given by  =4.53Rt. Using this value of the resistivity we determined the geometrical correction factor for our four-line electrode geometry by measuring the sheet resistance on the same sample and obtained C=1.81 i.e. =1.81Rt. The average conductivity value for a given doping condition and polymer was taken as the average of four devices.

Figure 4 .

 4 Figure 4. Schematic representation of high temperature rubbing and recovery of oriented thin films on patterned devices with oriented electrodes // and ⊥ to the rubbing direction. The TE device is dipped in the corresponding dopant solution for successful doping.

Figure 5 .

 5 Figure 5. Schematic illustration of the experimental setup used for measuring the Seebeck coefficient of oriented PBTTT thin films along the direction parallel and perpendicular to the rubbing direction. The two curves represent the Seebeck coefficients measured along the parallel (black) and perpendicular (red) directions. Figure 4A is reproduced from reference 9.

Figure 1 .

 1 Figure 1. Structure chimique des PBTTTs étudiés et illustration du procédé de fabrication des couches minces par brossage mécanique et dopage séquentiel avec F4TCNQ/ACN.

Figure 2 .

 2 Figure 2. Exemple d'alignement induit par brossage à haute température (TR=125°C) pour des films de C12-PBTTT et de C18-PBTTT. (a) et (d) spectres UV-vis polarisés avant dopage et (c) et (f) après dopage avec F4TCNQ/ACN (1mg/ml) (// et ⊥ désignent l'orientation de la polarisation de la lumière par rapport à la direction de brossage). (b) and (e) Images sous microscope optique polarisé montrant la forte biréfringence des films après brossage et dopage.

Figure 4 .

 4 Figure 4. Evolution des clichés de diffraction électronique de films minces de différents PBTTTs orientés par brossage (TR=125°C) et dopés avec F4TCNQ pour différents temps de dopage. Les distances réticulaires correspondant à la périodicité lamellaire (le long des chaînes alkyls) d100 et selon le -stacking d020 sont notées en bleu. La flèche R indique la direction de brossage des films minces.

Figure 5 .

 5 Figure 5. Variation de la conductivité mesurée dans le sens du brossage et du facteur de puissance PF en fonction de la longueur des chaînes alkyls d'une famille de PBTTT.

Figure 6 .

 6 Figure 6. a) Structures chimique du P3HT et du PBTTT étudiés. b) Méthode de préparation des couches orientées conductrices par brossage à haute température et dopage en solution de FeCl3/nitrométhane. c) Evolution de la couleur des couches minces en fonction du dopage à 1mM et 10mM. d) Comparaison des biréfringences des films brossés de C12-PBTTT avant et après dopage avec FeCl3.

Figure 7 .

 7 Figure 7. Comparaison des structures liquide cristalline et semi-cristalline du C12-PBTTT et du P3HT, respectivement. a) et c) images obtenues en mode low dose et haute résolution des films minces alignés et dopés de C12-PBTTT. b) et d) illustration des structures liquide cristalline et semi-cristalline du C12-PBTTT et du P3HT, respectivement.

Figure 11 .

 11 Figure 11. Correlation entre le coefficient Seebeck S et la conductivité  dans les films très orientés de a) C12-PBTTT (TR=125°C) et b) P3HT (TR=180°C) dopé avec FeCl3. Les indices // et ⊥ font référence à la direction de mesure du transport de charge par rapport à la direction de brossage. Les courbes continues font correspondent à l'ajustement des points expérimentaux à l'aide des lois d'échelle S//  -1/4 et S⊥-ln ().

Figure 13 .

 13 Figure 13. Evolution de la conductivité de charge (A et B) et du coefficient Seebeck (C et D) en fonction de la concentration de dopant F4TCNQ et F6TCNNQ dans les films orientés de C12-PBTTT, respectivement. S et  sont mesurés parallèlement and perpendiculairement à la direction de brossage. Pour comparaison, nous montrons aussi les valeurs des échantillons dopés de manière directe à forte concentration (2 et 5 mg/ml pour F4TCNQ et 5 mg/ml pour F6TCNNQ). On indique aussi la conductivité des échantillons non orientés pour comparaison.

Table 1 . 1

 11 Representative thermoelectric performances in p-doped polythiophenes PBTTT and P3HT

	Semicon-ductor	Dopant	Doping method	Filler and binder	σ (S cm -1 )	α (µVK -1 )	∝ 𝟐 𝝈 (µWmK -2 )	Reference
	P3HT	Fe3-tos3.6H2O	Sequential Solution		55.4	64	30	Hwang et.

Table 2

 2

.1). The continuous lines are the result of the fitting following the phenomenological equation is given in the text. B) Doping kinetics of two oriented C14-PBTTT films with different thicknesses showing that the curves

Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 94 fall on a common master curve when plotted against t 1/2 /l showing that both films are characterized by a common diffusion coefficient of F4TCNQ -in C14-PBTTT.

Table 2 .

 2

	Polymer	C8-PBTTT C12-PBTTT C14-PBTTT	C18-PBTTT
	Characteristic doping time  (s)	50±5	5±1	8+1	59±5
	Diffusion constant of dopant 10 -12 cm 2. s -1	0.58±0.3	9.0±4.5	8.8±3.2	0.61±0.25

1. Estimated diffusion coefficients of F4TCNQ -anions from doping kinetics in oriented thin films of PBTTT with various linear alkyl side chains

Table 2 .

 2

	PBTTT-C8	-	-	-	-
	PBTTT-C12	71.3		4.3
	PBTTT-C14	62.2	62.7	14.9
	PBTTT-C18	54.3	82.2, 62.5	33.9	47.5
	The difference in diffusion coefficients of F4TCNQ -into the different PBTTTs underlines
	the essential role of side chain length and packing on the diffusion mechanism of F4TCNQ
	in the side chain layers. This assumption is further supported by differential scanning
	calorimetry (DSC) results that show well-defined side chain crystallization/melting peaks

2. Melting (Tm) and crystallization (Tc) temperatures and their corresponding enthalpy ΔHm and ΔHc,

attributed to side chains of four PBTTTs. The temperatures have been measured at the peak maximum, the enthalpy has been measured from the area of the entire peak.

Tc (°C)

Tm (°C) Δ Hc sc (J/g) Δ Hm sc (J/g) for C18-PBTTT, whereas no such peaks are seen for C8-PBTTT and only broad and weak peaks for C12-PBTTT and C14-PBTTT (see figure

2

.6 and table 2.2). The important Chapter 2: Impact of alkyl side chain length on doping kinetics, crystal structure and thermoelectric properties of oriented PBTTT 96

Table 2 . 3 .

 23 Estimated values of the number of F4TCNQ-anions per thiophene cycles for various PBTTT side chains.

	Polymer	C8 PBTTT C12 PBTTT	C14 PBTTT	C18 PBTTT
	F4TCNQ -anion per thiophene cycle * (%)	7.6±2	14±3	6.9±2	3.3±1
	The obtained values are reported in table 2.3. C12-PBTTT shows the highest doping
	concentration of 14±3 % of F4TCNQ -anion per thiophene cycle whereas the lowest
	doping level is observed for C18-PBTTT (3.3%). Both C14-PBTTT and C8-PBTTT show
	intermediate doping levels around 7 %.

Table 2 .4. Oxidation

 2

	Eox (V)	HOMO level (eV)

potentials Eox and HOMO energy levels of PBTTT polymers.

Table 3 .

 3

	1. Thermoelectric characteristics of the FeCl3-doped oriented films of C12-PBTTT and P3HT and
	comparison with stretch-oriented iodine-doped PA.

Table 3 . 2 .

 32 Maximum conductivity, Seebeck coefficients and power factors obtained for oriented films of variousPBTTTs aligned by rubbing (TR in the range 100-125°C) and doped with FeCl3 as measured in the direction parallel to the rubbing. The details of the film preparation and macromolecular parameters of the PBTTTs are found in chapter 2 and in the experimental section.

	PBTTT side chain	Conductivity	Seebeck	Power factor
	length	// (S/cm)	Coefficient	PF// (W/mK 2 )
			S// (V/K)
	C8	1.2±0.2 . 10 4	9.1±0.5	99±25
	C12	2.2±0.5 . 10 5	9.4±0.5	1944±626
	C14	1.5±0.3 . 10 4	14.1±1.0	132±36
	C18	1.7±0.4 . 10 4	11.9±1.0	241±77

Table 4 . 1 .

 41 Characteristics of angular distribution curves of F4TCNQ and F6TCNNQ doped PBTTT C12

Table 4 .

 4 It is remarkable to see that all the data points align on the same master curves over 5 decades of charge conductivity, regardless of the dopant and for both directions parallel and perpendicular to the C12-PBTTT chains.

	2. Summary of maximum electrical conductivities, Seebeck coefficient and power factors observed upon
	incremental and single doping methods.DD: direct doping. ICD: Incremental Concentration Doping. TE properties
	obtained for vapor doped F4TCNQ/ PBTTT thin film are also included from reference 8.

This demonstrates that the S- correlation in C12-PBTTT is not dependent on the chemical nature of the dopant but mainly on the structure, packing and electronic structure of the oxidized polymer.

  Accordingly, this correlation is characteristic of the charge transport mechanism of the polymer C12-PBTTT that is only marginally affected by the type of dopant. The only notable difference related to the dopants is the limit in conductivity that is reached for each dopant. This

	Dopant	Doping method	 (S/cm)	S (µV/K)	PF µWm -1 K -2
		ICD	2430±500	47±7	530±200
	F6TCNNQ	DD	600±200	64.5±3	247±74
		4ACN+DD	970±200	55±1	293±63
		ICD	1380±700 50.7±4	355±200
	F4TCNQ	DD	920±60	57±3	300±20

Table 4 . 3 .

 43 Calculated charge carrier densities as obtained from the UV-Vis-NIR absorption signal of the polaron and dopant anion.

		Density of charge carriers (cm -3 )		Mobility (cm 2 /V. s)
		F6TCNNQ/C12 PBTTT	F4TCNQ/C12 PBTTT	F6TCNNQ/C12 PBTTT ICD single	F4TCNQ/C12 PBTTT ICD single
	Polaron
	density based
	on P1 (our	6.2×10 20	6.7×10 20	24	6	13	8.5
	calibration
	with FeCl3)
	Polaron
	density based	8.1×10 20	5.7×10 20	18.8	4.6	15.3	10
	on F6TCNNQ
	Polaron
	density based on P2	2.8 to 3×10 20	3.2×10 20	53.6	13.25	26	17.4
	(Tanaka et.al)

  ). The values are summarized in table 4.3. The average mobilities calculated for ICD with F6TCNNQ and Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 194 F4TCNQ are higher than for DD films. i.e., µICD > uDD. As an example, the average mobility calculated for ICD is  32 cm 2 /V. s and 8 cm 2 /V. s for DD with F6TCNNQ. Average mobilities of 27 cm 2 /V. s and 11.9 cm 2 /V. s are obtained for F4TCNQ with ICD and DD respectively. Since both ICD and single doped thin films gave the same intensities of the P1 band, it is assumed that both doping methods lead to same polaron density (see figure 4.9). It is worth to mention that a lower charge carrier density corresponding to a few percent of the density extracted from UV-Vis-NIR data, as suggested by Pingel and

Table 1 .

 1 Macromolecular parameters of the semiconducting polymers used in this study. (Gel permeation chromatography was performed in hot tri-chlorobenzene versus PS standard)

	Polymer	Mn (g/mol)	Mw (g/mol)	PDI
	PBTTT C8	13000	24000	1.84
	PBTTT C12	26000	45000	1.73
	PBTTT C14	17000	51000	1.64
	PBTTT C18	31000	51000	3
	P3HT	24200	43600	1.8
	1.1 General procedure for the synthesis of polymers
	Poly(2,5-bis(3-alkylthiophen-2-yl)thieno[3,2-b]thiophene):

A Schlenk tube was charged with a stirrer bar, 5,5'-dibromo-4,4'-dialkylbithiophene (0.537 mmol), 2,5-bis(trimethylstannyl)-thieno

[3,2-b]

thiophene (250 mg, 0.537 mmol),

The organic layer was separated, dried on sodium sulphate, filtered and partially concentrated under reduced pressure. The resulting solution was precipitated in methanol. The polymer was collected by filtration and dried under vacuum to afford the product.

rubbing direction. Aligned films of C12-PBTTT show charge conductivities of 193 S/cm along the rubbing direction and power factors of approximately 100 µW m -2 K -1 versus a few µW . m -2 K -1 for non-oriented films.

thermopower along the chain direction results in a substantial improvement of thermoelectric power factors up to 2 mW . m -1. K -2 in C12-PBTTT.

films showed that F6TCNNQ -anions are better oriented in the layers of the alkyl side chains as compared to F4TCNQ. The estimation of charge carrier mobilities indicates that ICD leads to higher mobilities than DD. The lower degree of ordering of F4TCNQ -anions produced lower electrical conductivity than F6TCNNQ doped C12-PBTTT thin films.

Acknowledgements

First of all, It is my genuine pleasure to express my sincere gratitude to my PhD supervisor Dr. Martin Brinkmann, who was always there to help me with great patience and interest throughout these years. Over the course of the PhD, we had great scientific and personal discussions. Several ideas in this thesis were the outcomes of these scientific discussions. Joining his lab was the best decision of my life. Thank you so much Martin for giving me this opportunity and without your help, I would not accomplish this milestone.

Abbreviations used in the text

For short doping times, the absorption of both dopant anions scales like √ 𝑡 . This is an indication for a diffusion-limited doping mechanism. The saturation observed for longer doping times is due to the finite film thickness: the whole film is doped after 6-10 min (see Figure 4.7.A). Each polymer/dopant system is characterized by a typical doping time Chapter 4: Influence of dopant size and doping method on the thermoelectric properties of PBTTT films doped with F6TCNNQ and F4TCNQ 176 . However, this time constant increases with film thickness as previously observed in PBTTT polymers doped with F4TCNQ (see chapter 2). [16] The plot of normalized intensity of anions vs t 1/2 /l helped determine the diffusion coefficients of F4TCNQ -and F6TCNNQ - and this representation is independent of the polymer film thickness. The following phenomenological equation was used to fit the absorbance of the anion as a function of the doping time:

Where A is the absorption at wavelength  of the anion ( = 873 nm for F4TCNQ -and 1158 for F6TCNNQ -). A0 is the absorbance of the anion at saturation, l is the film thickness and D is the diffusion coefficient of F4TCNQ -or F6TCNNQ -in the matrix of C12-PBTTT.

The larger F6TCNNQ -anion shows a lower 2×10 -12 cm 2 /s diffusion coefficient with respect to 9×10 -12 cm 2 /s for F4TCNQ -(see figure 4.8A) . Both diffusion coefficients were close to the diffusion coefficients of iodine into the P3HT films reported by Maliakal (2.5×10 -11 cm 2 /s) [25] . This result points at the impact of the size of the dopant anions on the diffusion into the highly interdigitated side chain layers. [26] It suggests that the smaller anions such as F4TCNQ -can diffuse and intercalate into the alkyl side chains faster than the longer and bulkier anions of F6TCNNQ -. However, as seen hereafter, the amount of dopants at saturation in the layers of C12-PBTTT is not related to their diffusion kinetics.

The spectra measured along the parallel direction are also really instructive. The polaronic absorption bands in the doped PBTTT namely P1 and P2 are polarized along the // direction for both F4TCNQ and F6TCNNQ doped films as observed previously by

Hamid-Sakr et.al [15] . Similar to the anion spectra along the perpendicular direction, the absorption of the polarons can be seen already after 5s and continues to increase and

Résumé de la thèse Vishnu Vijayakumar

Films minces de polythiophènes orientés et conducteurs pour applications en thermoélectricité.

Directeur de la thèse : Martin Brinkmann l'intercalation des dopants dans les couches de chaînes alkyls des PBTTTs, quel que soit la longueur des chaînes alkyls. 9).

La conductivité des films de P3HT est toujours largement inférieure à celle du C12-PBTTT.

Cette différence n'est pas liée à une différence de densité de porteurs de charge à en juger par la spectroscopie UV-Vis-NIR, à une différence d'alignement des films. La faible conductivité observée pour P3HT est attribuée à la structure semi-cristalline de ce polymère. Dans le P3HT, les charges doivent passer par des zones désordonnées 241 amorphes (cf Figure 7). En revanche, la structure liquide-cristalline du C12-PBTTT, donne lieu à une meilleure percolation entre domaines ordonnés, ce qui permet d'atteindre des conductivités très élevées (cf images HRTEM de la Figure 7). Il faut noter que, même dans la direction perpendiculaire aux chaînes de PBTTT, le transport reste très important (conductivités de  2100 S/cm proches de celles rapportées pour des films de C14-PBTTT dopé avec des fluoro-trichlorosilanes). 2 mg/ml 1 mg/ml 0.5 mg/ml 0.1 mg/ml 0.01 mg/ml As rubbed

significativement plus élevées que par dopage directe. Ainsi, par APC on atteint 2430 S/cm et un facteur de puissance de 583 ±180 uWm

Summary

The aim of this thesis is to develop new polymeric thermoelectric (TE) materials based on oriented polythiophene (PBTTT) films. High-temperature rubbing produces oriented films of controlled orientation and crystallinity. Various doping methods with suitable dopants (F4TCNQ, F6TCNNQ and