Thermoelectric conjugated polymers

Abstract

The present invention relates to a semi-crystalline polymer material, preferably for a thermoelectric device, wherein the material comprises a first conjugated polymer, a second conjugated polymer mixed with the first conjugated polymer, and a dopant, the second conjugated polymer having a number-average molecular weight that is 25% or more greater than that of the first conjugated polymer. A method for producing a semi-crystalline polymer material, the use of the semi-crystalline polymer material, and a thermometric device comprising the semi-crystalline polymer material are also described.

Description

INTRODUCTION The invention relates to the field of thermoelectric materials. In particular, the invention relates to a semi-crystalline polymer material for use in a thermoelectric device. The use of the semi-crystalline polymer material as a thermoelectric material and a method for producing the semi-crystalline polymer material are also described. BACKGROUND Due to their relatively low thermal conductivity and the ability to tailor their molecular structure and charge carrier concentrations, conjugated polymers are increasingly being investigated as thermoelectric materials for operation at ambient temperature.1 However, their thermoelectric performance generally lags behind that of the most advanced inorganic thermoelectrics such as Bi2Te3(2). To achieve high performance with thermoelectric materials, strategies are needed that increase the electrical conductivity and Seebeck coefficients of the materials while minimizing the increase in thermal conductivity. This leads to an improved power factor (PF), defined by Equation 1 below, and an improved thermoelectric coefficient of performance (ZT), defined by Equation 2. In Equations 1 and 2, σ is the electrical conductivity, κ is the thermal conductivity, S is the Seebeck coefficient, and T is the temperature. A well-known approach to improving electrical conductivity, Seebeck coefficients (S), and performance factors is to induce uniaxial alignment of the polymer chains along the thermoelectric transport direction.3 Transport within the chain along the polymer backbone is much faster than transport between polymer chains, so conductivity along the polymer chain alignment direction is generally significantly higher than perpendicular to the alignment direction or than in polymer films without chain orientation.3c,4 However, this strategy is not always effective in increasing overall thermoelectric performance, as the improved electrical conductivity is often accompanied by an increase in thermal conductivity along the chain alignment direction. Furthermore, the semi-crystalline nature of these materials means that the aligned polymer chains form discrete crystalline domains separated by well-defined grain boundaries where the chain ends reside. These grain boundaries can then dominate the material's performance. For example, in the ribbon phase of poly[2,5-bis(3-dodecylthiophen-2-yl)thieno[3,2-b]thiophene] (PBTTT), one of the most highly ordered semi-crystalline conjugated polymers with a microstructure reminiscent of semi-crystalline polyethylene, scanning Kelvin probe microscopy (SKPM) studies have shown that these grain boundaries between the ribbons represent bottlenecks for overall charge transport.5 SUMMARY OF THE INVENTION In one aspect of the invention, a semi-crystalline polymer material, preferably for a thermoelectric device, is provided, wherein the material comprises: a first conjugated polymer; a second conjugated polymer mixed with the first conjugated polymer; and a dopant; wherein the second conjugated polymer has a number-average molecular weight that is 25% or more greater than the number-average molecular weight of the first conjugated polymer. In a further aspect of the invention, a semi-crystalline polymer material, preferably for a thermoelectric device, is provided, wherein the material comprises: a first conjugated polymer; a second conjugated polymer mixed with the first conjugated polymer; and a dopant, wherein the second conjugated polymer has a number-average molecular weight greater than the number-average molecular weight of the first conjugated polymer; and wherein the HOMO energy level of the second conjugated polymer is within 0.1 eV of the HOMO energy level of the first conjugated polymer; or wherein the LUMO energy level of the second conjugated polymer is within 0.1 eV of the LUMO energy level of the first conjugated polymer. In a further aspect of the invention, a semi-crystalline polymer material is provided, preferably for a thermoelectric device, wherein the material comprises: a first conjugated polymer; a second conjugated polymer mixed with the first conjugated polymer; and a dopant, wherein the second conjugated polymer has a number-average molecular weight greater than the number-average molecular weight of the first conjugated polymer; and the first conjugated polymer and the second conjugated polymer form crystalline domains comprising both the first and the second conjugated polymer, and preferably the first conjugated polymer and the second conjugated polymer comprise and preferably consist of the same conjugation system or the same delocalized pi system. In a further aspect of the invention, a method for producing a semi-crystalline polymer material, preferably for a thermoelectric device, is provided, wherein the method comprises: mixing a first conjugated polymer with a second conjugated polymer; applying the mixed polymers to a substrate; and doping the first and second conjugated polymers with a dopant; wherein the second conjugated polymer has a number-average molecular weight that is 25% or more greater than the number-average molecular weight of the first conjugated polymer. In a further aspect of the invention, a method for producing a semi-crystalline polymer material is provided, comprising: separately doping each first conjugated polymer and each second conjugated polymer with a dopant; mixing the first doped conjugated polymer with the second doped conjugated polymer; and applying the mixed polymers to a substrate; wherein the second conjugated polymer has a number-average molecular weight that is 25% or more greater than the number-average molecular weight of the first conjugated polymer. In a further aspect of the invention, a method for producing a semi-crystalline polymer material, preferably for a thermoelectric device, is provided, wherein the method comprises: mixing a first conjugated polymer with a second conjugated polymer; applying the mixed polymers to a substrate; and doping the first and second conjugated polymers with a dopant; wherein the second conjugated polymer has a number-average molecular weight greater than the number-average molecular weight of the first conjugated polymer; and wherein: a) the HOMO energy level of the second conjugated polymer is within 0.1 eV of the HOMO energy level of the first conjugated polymer, or the LUMO energy level of the second conjugated polymer is within 0.1 eV of the LUMO energy level of the first conjugated polymer;or b) the first conjugated polymer and the second conjugated polymer form crystalline domains comprising both the first and the second conjugated polymer, and preferably the first conjugated polymer and the second conjugated polymer comprise and preferably consist of the same conjugation system or the same delocalized pi system. In a further aspect of the invention, a method for producing a semi-crystalline polymer material, preferably for a thermoelectric device, is provided, the method comprising: separately doping each first conjugated polymer and each second conjugated polymer with a dopant; mixing the first doped conjugated polymer with the second doped conjugated polymer; and applying the mixed polymers to a substrate; wherein the second conjugated polymer has a number-average molecular weight greater than the number-average molecular weight of the first conjugated polymer; and wherein: a) the HOMO energy level of the second conjugated polymer is within 0.1 eV of the HOMO energy level of the first conjugated polymer, or the LUMO energy level of the second conjugated polymer is within 0.1 eV of the LUMO energy level of the first conjugated polymer;or b) the first conjugated polymer and the second conjugated polymer form crystalline domains comprising both the first and the second conjugated polymer, and preferably the first conjugated polymer and the second conjugated polymer comprise and preferably consist of the same conjugation system or the same delocalized pi system. The invention also relates to the use of the semi-crystalline polymer material described above for forming a layer on a substrate. The invention also relates to the use of the semi-crystalline polymer material described above as a thermoelectric material. The invention also relates to the use of the semi-crystalline polymer material described above for forming a layer of a thermoelectric device. The invention also relates to a layer on a substrate, wherein the layer comprises the semi-crystalline polymer material described above. The invention also relates to a thermoelectric device comprising the semi-crystalline polymer material described above. DEFINITIONS As used here, the term "semi-crystalline" refers to a material that includes discrete crystalline and amorphous regions. Specifically, a material that comprises crystalline domains separated by amorphous domains. As used here, the term “crystalline domains” refers to an area of ​​order within a polymer material due to the alignment, or at least partial alignment, of the polymers, i.e., the polymer chains. As used here, the term "conjugated" refers to polymers that include areas or regions of delocalization along their chain axis. For example, polymers that comprise a delocalized pi system. As used here, the term "mixed" refers to a combination or mixture of two polymers, where the two polymers are dispersed within the mixture, i.e., they are intimately mixed. It is understood that in semi-crystalline polymer materials comprising a mixture of two polymers, the polymer material includes crystalline domains containing both mixed polymers, as opposed to separate or discrete crystalline domains of each polymer. As used here, the term "doping agent" refers to a substance added to a polymer for the purpose of modulating its electrical properties. As used here, the term "HOMO energy level" refers to the energy level of the highest occupied molecular orbital. The difference in HOMO energy levels between two conjugated polymers can be calculated by determining the HOMO levels of the polymers separately using techniques such as cyclic voltammetry or ultraviolet photoelectron spectroscopy, or other techniques known to a person skilled in the art. As used here, the term "LUMO energy level" refers to the energy level of the lowest unoccupied molecular orbital. The difference in LUMO energy levels between two conjugated polymers can be calculated by determining the LUMO levels of the polymers separately using techniques such as cyclic voltammetry or ultraviolet photoelectron spectroscopy in combination with band gap measurements by optical spectroscopy or other techniques known to a person skilled in the art. Unless otherwise stated, the term "average molecular weight" here refers to the number-mean molecular weight determined by gel permeation chromatography. DETAILED DESCRIPTION OF THE INVENTION As explained above, in one aspect of the invention a semi-crystalline polymer material, preferably for a thermoelectric device, is provided, wherein the material comprises: a first conjugated polymer; a second conjugated polymer mixed with the first conjugated polymer; and a dopant; wherein the second conjugated polymer has a number-average molecular weight that is 25% or more greater than the number-average molecular weight of the first conjugated polymer. In general, the present invention proposes a new mechanism for increasing the thermoelectric performance of conjugated polymers by using a mixture of conjugated polymers with different number-mean molecular weights. Advantageously, it has been observed that this leads to an increase in electrical conductivity without significantly affecting the Seebeck coefficient and without a proportional increase in thermal conductivity, resulting in an improved thermoelectric coefficient of performance (QT). Without committing to specific theories, it is assumed that the incorporation of a conjugated polymer with a higher number-average molecular weight into a conjugated polymer with a lower number-average molecular weight leads to an increase in "linking chains." It is hypothesized that these linking chains connect adjacent crystalline domains and facilitate charge transfer between them without requiring chain transitions within the disordered domains. It was entirely unexpected that this arrangement would improve the electrical conductivity and Seebeck coefficient of the material without a proportional increase in thermal conductivity. Embodiments of the various aspects of the invention are described below. For the avoidance of doubt, it should be noted that all embodiments described herein with respect to one aspect of the present invention may also be applicable to the other aspects of the present invention. Connection chains As mentioned above, the materials of the present invention have been found to advantageously exhibit good thermoelectric performance. This is due to the mixture of a conjugated polymer with a lower number-average molecular weight and a conjugated polymer with a higher number-average molecular weight. Preferably, the first conjugated polymer and the second conjugated polymer form crystalline domains that encompass both the first and the second conjugated polymer. Without being bound to any specific theory, it is assumed that the improvement in thermoelectric properties is due to the conjugated polymer with the higher number-average molecular weight forming bonding chains that link discrete crystalline domains. Accordingly, in preferred embodiments of the invention, the semi-crystalline polymer material comprises discrete crystalline domains, and the second conjugated polymer links at least two of the crystalline domains. The verification of these bonding chains is discussed in detail in the following section, “Experiments.” Polymer alignment Preferably, the first and second conjugated polymers are oriented in a single direction. This can be advantageous for improving charge transport along the orientation direction. It can also allow the charge transport properties to be optimized in the direction of current flow or in the direction of the temperature gradient within a thermoelectric device. Accordingly, for thermoelectric devices comprising the semi-crystalline polymer material described herein, it is advantageous if the first and second conjugated polymers are preferably oriented in the direction of current flow or in the direction of a temperature gradient within the device. Relationship between the two polymers It is understood that the second conjugated polymer may be a different polymer than the first conjugated polymer, or that the second conjugated polymer may be the same polymer as the first conjugated polymer, only with a different number-mean molecular weight. Preferably, the second conjugated polymer is the same as the first conjugated polymer or closely related to it, i.e., with respect to its structure. Using the same or very similar conjugated polymers is advantageous to ensure that the electrical conductivity is not significantly impaired by the presence of the second conjugated polymer and that charge carriers can easily "jump" between the chains of both conjugated polymers. In particular, it is believed that by using two conjugated polymers whose energy levels are within a few kBTs, significant impairment of charge transport can be avoided. However, it should also be noted that a small difference between the conjugated polymers, and thus their energy levels, can be advantageous in order to achieve an increase in the Seebeck coefficient. Accordingly, in preferred embodiments, the HOMO energy level of the second conjugated polymer lies within 0.1 eV of the HOMO energy level of the first conjugated polymer. In other preferred embodiments, the LUMO energy level of the second conjugated polymer lies within 0.1 eV of the LUMO energy level of the first conjugated polymer. It is understood that the term "within 0.1 eV" refers to the situation in which the HOMO / LUMO energy level of the second conjugated polymer is less than or equal to 0.1 eV above or 0.1 eV below the HOMO / LUMO energy level of the first conjugated polymer, i.e., ±0.1 eV. Preferably, the HOMO energy level of the second conjugated polymer lies within 0.05 eV of the HOMO energy level of the first conjugated polymer, for example within 0.02 eV or within 0.01 eV. Preferably, the LUMO energy level of the second conjugated polymer lies within 0.05 eV of the LUMO energy level of the first conjugated polymer, for example within 0.02 eV or within 0.01 eV. It should be noted that a closely, but not exactly, matching HOMO level can occur in a situation where the second conjugated polymer has an essentially similar backbone structure to the first conjugated polymer. This similarity can also allow the second polymer to be incorporated into the crystalline domains without causing disruption. Accordingly, in optional embodiments, the first conjugated polymer and the second conjugated polymer essentially comprise the same polymeric backbone. Optionally, the first conjugated polymer and the second conjugated polymer comprise the same conjugation system and preferably consist of it. Optionally, the first conjugated polymer and the second conjugated polymer comprise the same delocalized pi system and preferably consist of it. Optionally, the first conjugated polymer and the second conjugated polymer differ only by single-atom substitutions in the polymer repeat unit, preferably while the conjugated polymer retains the same bonding pattern of the polymer backbone. For example, the first conjugated polymer can be PBTTT and the second polymer can be PBSTT. Optionally, the first conjugated polymer and the second conjugated polymer differ only by the addition of electron-withdrawing or electron-donating groups to the polymer backbone. Further differences between the first and second conjugated polymers (in addition to their different number-average molecular weights) may include different end groups and different types and / or numbers of side chains. For example, the first and second conjugated polymers may have side chains of different lengths. In particular, the first and second conjugated polymers may have the same (or a similar) backbone but different side chain lengths. It is assumed that the differences that can be tolerated are those that do not substantially alter the crystal structure or energy levels of the polymer backbone. However, it is thought that it could be advantageous for the second conjugated polymer to differ slightly from the first conjugated polymer so that the bonding chains can act as energy filters.When charges are forced to move across a conjugated polymer (i.e., a bonding chain) with a slightly higher energy, the Seebeck coefficient increases, i.e., the thermoelectric performance is improved. Optionally, the first conjugated polymer and the second conjugated polymer have the same side chains or at least side chains of the same length. Optionally, the first conjugated polymer and the second conjugated polymer preferably comprise the same repeating unit or units. Optionally, the first conjugated polymer and the second conjugated polymer are identical, except for their number-mean molecular weight and possibly their end groups. Molecular weights Preferably, the number-average molecular weight of the second conjugated polymer is 35% or more greater than the number-average molecular weight of the first conjugated polymer. For example, by 40% or more, or by 45% or more. It is further preferred that the number-mean molecular weight of the second conjugated polymer is 50% or more greater than the number-mean molecular weight of the first conjugated polymer. For example, by 60% or more, or by 70% or more. It is further preferred that the number-mean molecular weight of the second conjugated polymer is 75% or more greater than the number-mean molecular weight of the first conjugated polymer. For example, 80% or more greater. Without being bound to any particular theory, it is assumed that increasing the number-mean molecular weight (and thus the length of the polymer chain) of the second conjugated polymer increases the probability that the second conjugated polymer will be associated with two discrete crystalline domains, and thus increases the probability of forming a bonding chain. Preferably, the number-average molecular weight of the second conjugated polymer is 500% greater or less than the number-average molecular weight of the first conjugated polymer. It is further preferred that the number-mean molecular weight of the second conjugated polymer be 350% larger or smaller than the number-mean molecular weight of the first conjugated polymer. It is further preferred that the number-mean molecular weight of the second conjugated polymer be 250% larger or smaller than the number-mean molecular weight of the first conjugated polymer. In some preferred embodiments, the number-average molecular weight of the second conjugated polymer is 50% to 250% of the number-average molecular weight of the first conjugated polymer. More preferably, the number-average molecular weight of the second conjugated polymer is 75% to 150% of the number-average molecular weight of the first conjugated polymer. Preferably, the first conjugated polymer has a number-average molecular weight of at least 5 kDa. For example, at least 10 kDa or at least 12.5 kDa. Preferably, the first conjugated polymer has a number-average molecular weight of at least 15 kDa. For example, at least 20 kDa, at least 25 kDa, or at least 30 kDa. Preferably, the first conjugated polymer has a number-average molecular weight of less than 60 kDa, for example less than 50 kDa. Preferably, the second conjugated polymer has a number-average molecular weight of at least 30 kDa. Preferably, the second conjugated polymer has a number-average molecular weight of at least 50 kDa, for example at least 60 kDa. Preferably, the second conjugated polymer has a number-mean molecular weight of less than 120 kDa, for example less than 100 kDa. Preferably, the first conjugated polymer and / or the second conjugated polymer has a polydispersity index of 2 or less. Preferably, the first conjugated polymer and / or the second conjugated polymer has a polydispersity index of 1.5 or less. The first conjugated polymer may have a polydispersity index of 1.2 or less. Advantageously, low polydispersity can provide a well-ordered structure with discrete crystalline domains that can be effectively linked by bonding chains. Relative composition As shown in the examples below and in Fig. 2, the ratio of the two mixed conjugated polymers can have a significant influence on the overall thermoelectric performance. Without being bound to any specific theory, it is assumed that the first conjugated polymer is primarily responsible for the formation of ordered crystalline domains, while the second conjugated component is primarily responsible for the formation of the linking chains that connect these domains. Therefore, it is advantageous for the first conjugated polymer to be present in excess relative to the second conjugated polymer. Accordingly, in preferred embodiments, the molar ratio of the second conjugated polymer in the mixture of the first and second conjugated polymers is less than 50%, for example, less than 40%. Preferably, the molar ratio of the second conjugated polymer in the mixture of the first conjugated polymer and the second conjugated polymer is less than 30%, for example less than 25% or less than 20%. Preferably, the molar ratio of the second conjugated polymer in the mixture of the first conjugated polymer and the second conjugated polymer is greater than 5%. For example, greater than 7% or greater than 8%. Preferably, the molar ratio of the second conjugated polymer in the mixture of the first conjugated polymer and the second conjugated polymer is greater than 10%. For example, greater than 12.5% ​​or greater than 15%. In some embodiments, the molar ratio of the second conjugated polymer in the mixture of the first conjugated polymer and the second conjugated polymer is between 5 and 50%, for example between 7.5 and 40%. Preferably, the molar ratio of the second conjugated polymer in the mixture of the first conjugated polymer and the second conjugated polymer is between 10 and 30%, for example between 12.5 and 25%. Polymer structure Preferably, the first conjugated polymer and / or the second conjugated polymer is selected from homopolymers such as polythiophenes, polyselenophenes, polyfurans or polythiazoles; donor-donor copolymers such as polymers comprising thiophene, thienothiophene and / or selenophene units; and / or donor-acceptor copolymers such as polymers comprising diketopyrrolopyrrole units or copolymers comprising thiophene and fluorinated benzothiadiazole. Preferably, the first conjugated polymer and / or the second conjugated polymer comprises a thiophene unit within the polymer backbone. Preferably, the first conjugated polymer and / or the second conjugated polymer is selected from poly(2,5-bis(3-alkylthiophen-2-yl)thieno(3,2-b)thiophene (PBTTT), poly(3-hexylthiophene) (P3HT) or poly[[2,5-bis(2-octadecyl)-2,3,5,6-tetrahydro-3,6-diketopyrrolo[3,4-c]pyrrol-1,4-diyl]-alt-(2-octylnonyl)-2,1,3-benzotriazole] (DPPBTz). Preferably, the first conjugated polymer and / or the second conjugated polymer is poly(2,5-bis(3-alkylthiophen-2-yl)thieno(3,2-b)thiophene) (PBTTT). The first conjugated polymer and / or the second conjugated polymer can be poly(2,5-bis(3-alkylselenophen-2-yl)thieno(3,2-b)thiophene) (PBSTT). For example, the first conjugated polymer can be poly(2,5-bis(3-alkylthiophen-2-yl)thieno(3,2-b)thiophene) (PBTTT) and the second conjugated polymer can be poly(2,5-bis(3-alkylselenophen-2-yl)thieno(3,2-b)thiophene) (PBSTT), or vice versa. Preferably, the semi-crystalline polymer material comprises ribbons of aligned polymers. Preferably, the semi-crystalline polymer material comprises discrete crystalline domains, and the polymer chains of the first conjugated polymer adopt a substantially straight conformation within the crystalline domains. Advantageously, this "chain-extended" structure optimizes charge transport through the crystalline domains. It should be noted that the straight conformation can be achieved by optimizing the number-mean molecular weight of the first conjugated polymer. Preferably, the dopant is selected from a neutral molecular dopant, an ionic dopant, a closed shell ion introduced by an ion exchange process, a dopant that initiates an electrophilic or nucleophilic attack, or a dopant that induces a proton transfer. Preferably, the dopant is an ion-exchange dopant. The dopant preferably comprises FeCl3 and / or 1-butyl-1-methylpyrrolidinium bis(trifluoromethylsulfonyl)imide (BMP-TFSI). Preferably, the semi-crystalline polymeric material comprises a molar doping ratio of 0.01 to 2 charge carriers per monomer. A molar doping ratio of 0.1 to 1 charge carrier per monomer is further preferred. Method for producing the semi-crystalline polymer material As described in the aspects above, the method of the present invention can either comprise mixing two conjugated polymers and subsequently doping the mixed conjugated polymers, or it can comprise doping unmixed conjugated polymers and subsequently mixing the doped conjugated polymers together. Preferably, the doping is carried out after mixing, since the conjugated polymers are more soluble before doping, resulting in an improved mixing step. Preferably, the application of the mixed conjugated polymers further comprises aligning the mixed conjugated polymers. Advantageously, aligning the conjugated polymers can lead to improved charge transport through the material, particularly along the alignment direction. Preferably, aligning the mixed conjugated polymers includes coating the mixed conjugated polymers on the substrate with a doctor blade. Optionally, alignment includes rubbing, stretching, or fiber spinning of the mixed conjugated polymers. Preferably, the doping of a first conjugated polymer and a second conjugated polymer is carried out after the polymers have been aligned. Preferably, mixing the first conjugated polymer with the second conjugated polymer comprises mixing a solution comprising the first conjugated polymer with a solution comprising the second conjugated polymer. Preferably, doping a first conjugated polymer and a second conjugated polymer with a dopant comprises ion-exchange doping, molecular doping, electrochemical doping, solid-state diffusion doping, gas-phase doping, mixed-solution doping, sequential doping, and acid doping. For example, ion-exchange doping. Preferably, the doping of a first conjugated polymer and a second conjugated polymer with a dopant comprises ion exchange doping with a solution in which the oxidizing agent is present in a concentration range of 0.01 to 10 mM, for example 0.05 to 5 mM or 0.1 to 2 mM. BRIEF DESCRIPTION OF THE IMAGES Fig. 1 shows differential scanning calorimetry (DSC) measurements on undoped films containing exclusively the first conjugated polymer P1, exclusively the second conjugated polymer P2, or a mixture of P1 and P2 (see the "Experimental Part" section below for details). The measurements were performed at a cooling rate of 10 °C / min. Fig. 2 shows the electrical conductivity of doped mixed films as a function of the molar ratio of P2 introduced into the mixture. Fig. 3 shows the temperature-dependent electrical conductivity of doped films consisting exclusively of P1, exclusively of P2, and of a mixture of P1 and P2. The symbols represent experimental data, and the lines are fits to a Kaiser model. Fig. 4 shows the temperature-dependent Seebeck coefficient of doped films consisting exclusively of P1, exclusively of P2, and of a mixture of P1 and P2.The inset shows the distribution of Seebeck coefficient values ​​measured on several samples at room temperature. Fig. 5 shows the temperature-dependent thermal conductivity of doped and undoped layers consisting solely of P1, solely of P2, and of a mixture of P1 and P2. Fig. 6 shows the electrical conductivity (gray) and the Seebeck coefficient (black) as a function of exposure time to the dopant FeCl3 at a concentration of 0.33 mM. Fig. 7 shows the power factor as a function of exposure time to the dopant FeCl3 at a concentration of 0.33 mM. Fig. 8 is a graphical representation of an aligned PBTTT polymer in the ribbon phase with a single low molecular weight component (left) and the same material mixed with a higher molecular weight component (right). Fig. 9 shows a comparison of the surface morphology and charge carrier concentration.Figures (a)-(c) show AFM images and the corresponding Fast Fourier Transform (FFT) images (insertions) of aligned and undoped P1, mixture, and P2 films. Figures (d)-(f) show the corresponding sulfur 2p XPS spectra of aligned and undoped P1, mixture, and P2 films. Figure 10 shows a comparison of the crystalline order and paracrystallinity of films using GIWAXS: X-ray measurements along the chain alignment direction (a) P1, (b) mixture, (c) P2 and perpendicular to the chain alignment direction (d) P1, (e) mixture, (f) P2. Figure 11 shows the values ​​of Q(010) and the extracted paracrystallinity of aligned, doped PBTTT for X-rays incident parallel to the chain alignment direction.Figure 12 shows the crystalline order across individual grain boundaries, investigated by 4D-STEM: (a) Ring-shaped dark field (ADF) and the electron diffraction patterns (insertion) of P1; (b) Corresponding variance image highlighting the contrast of the ADF image in (a); (c) Fitted chain orientation map of P1; (d) ADF and the electron diffraction patterns (insertion) of a mixed film. (e) Corresponding variance image highlighting the contrast of the ADF image in (d); (f) Fitted chain orientation map of the mixed films. Only the dominant order is shown in the orientation maps. Figure 13 shows spectroscopic measurements of films consisting exclusively of P1, exclusively of P2, and of a mixture of P1 and P2. a) Polarized ultraviolet visible spectroscopy of undoped films. (b) Polarized ultraviolet visible spectroscopy of doped films.(c) Photothermal deflection spectroscopy (PDS) of the polymers, (d) polarized mid-infrared spectra of the doped films for incident light polarized parallel and perpendicular to the chain orientation. (e) Enlarged view of the spectra in (d) for incident light polarized perpendicular to the chain orientation. Fig. 14 shows the grain boundary conductivities determined by Kaiser model fitting of the experimental data shown in Fig. 3. Fig. 15 shows an AFM image of a doped film composed of a mixture of two different polymers, PBTTT and PBSTT. EXAMPLES Unless otherwise stated, all materials used in the examples were commercially available. Unless otherwise stated, anhydrous solvents (<20 ppm water) were used. Average molecular weight The number-mean molecular weight (Mn) and weight-mean molecular weight (Mw) were determined using an Agilent Technologies Series 1200 gel permeation chromatography (GPC) system at 80 °C in chlorobenzene using two series-connected PL mixed B columns and calibrated against low polydispersity polystyrene standards. Polydispersity Index (PDI) The polydispersity index was calculated by dividing the weight-mean molecular weight by the number-mean molecular weight (Mw / Mn). Synthesis of PBTTT (Poly(2,5-bis(3-alkylthiophen-2-yl)thieno(3,2-b)thiophene) A first batch of PBTTT (Poly(2,5-bis(3-alkylthiophen-2-yl)thieno(3,2-b)thiophene), here designated P1, was synthesized as described in I. McCulloch et al., Nature Materials 5, 328 (2006), the synthetic procedure being adopted hereby by reference. P1 had a number-mean molecular weight of 35 kDa and a polydispersity index (PDI) of 1.5. A second batch of PBTTT, designated P2, was sourced from 1-Material Inc. P2 had a number-average molecular weight of 65 kDa and a polydispersity index (PDI) of 2.5. Formation of mixed PBTTT films Solutions of the PBTTT polymers P1 and P2 were prepared at a concentration of 10 g / L in 1,2-dichlorobenzene. The solutions were then heated to 120 °C for 1 hour. Subsequently, defined volumes of the two solutions were mixed according to the desired ratio of P1 and P2 in the final mixture. The mixture was then heated to 120 °C for 10 minutes to ensure that the two polymers were thoroughly blended. Chain alignment was performed using a doctor blade recoater in a nitrogen glovebox (MBraun, <1 ppm O2, <1 ppm H2O). A PTFE rod was used as the coating blade. 10 µl of the mixed polymer solution was injected into the gap between the blade and a glass substrate. The temperature-controlled substrate (80 °C) was then moved under the PTFE coating blade at a constant speed of 110 µm / s. After coating with the blade, the polymer mixture deposited on the substrate was annealed at 270 °C for 20 minutes and then cooled to room temperature on a hot plate. Reference films consisting exclusively of P1 and P2 were produced analogously using polymer solutions containing only P1 or P2. Funding for PBTTT films BMP-TFSI solutions were prepared at a concentration of 1 M in acetonitrile. FeCl3 solutions were prepared at a concentration of 10 mM in acetonitrile. The two solutions were combined to obtain a 1 mM / 100 mM FeCl3 / BMP-TFSI solution in acetonitrile. The doping was carried out on the spin coater by dropping 140 ml of the 1 mM / 100 mM FeCl3 / BMP-TFSI solution in acetonitrile onto the prepared films, leaving the solution on the substrates for 100 seconds, then spinning off the excess at 2000 rpm and washing with 1 ml of acetonitrile. Lower doping levels were achieved by following the same protocol, but using 0.33 mM FeCl3 and leaving the solution on the substrate for a controlled period between 40 and 160 seconds. Film thickness The film thickness on a substrate was measured using surface profilometry with a Bruker DekTak XT. Conductivity measurements Thin-film conductivity measurements at room temperature were performed using an Agilent 4155B semiconductor parameter analyzer in a nitrogen atmosphere (Belle Technology, <10 ppm O2, <20 ppm H2O). Each instrument was isolated prior to measurement by scraping off the film outside the active area. Differential scanning calorimetry Differential scanning calorimetry experiments were performed using a TA Instruments Q2000 DSC and a heat flux method to measure heat flux. The instrument was calibrated with an indium standard. Samples were encapsulated in aluminum sample vials and heated from 0 °C to 300 °C at a constant rate of 10 °C / min using nitrogen as the purge gas. Temperature-dependent thermoelectric measurements Four-point conductivity and Seebeck measurements were performed using a microfabricated on-chip device architecture with integrated heating elements and calibrated temperature sensors. PBTTT films were structured by photolithography and oxygen plasma etching techniques as described in J. Chang et al., Adv. Funct. Mater. 2010, 20, 2825-283. Evaluation of the temperature-dependent electrical conductivity and Seebeck coefficient was performed in a closed LakeShore Cryotronics CRX-4K sampling station under high vacuum (10⁻⁷ to 10⁻⁶ mbar). A Keithley 2182A nanovoltmeter (for thermoelectric voltage measurements) and a Keithley 2612B source meter unit (for voltage supply and four-probe conductivity measurements) were used for electrical measurements. Failure analysis was performed according to the report in M. Statz et al., Comm. Phy. Conducted on January 1, 2018. AFM measurements Tapping-mode AFM measurements were performed using an Oxford Instruments Cypher ES Environmental AFM and an RTESPA-150 probe (resonance frequency 150 kHz, spring constant 0.4 N m⁻¹, Bruker) to characterize the surface morphology of the samples at a scan rate of 0.5 Hz and 512 samples per line. The data were analyzed using Gwyddion 2.63 (open-source software), including analysis of PBTTT bands by fast Fourier transform (FFT). Measurements to determine the melting points of pristine PBTTT bands were performed in an inert atmosphere protected by a continuous stream of dry nitrogen, using an RTESPA-150 probe at a scan rate of 0.5 Hz and 512 samples per line. Sample temperature was controlled by a programmable sample stage (provided by Cypher ES Environmental AFM) and simultaneously measured by an external K-type thermocouple attached to the sample. The melting point of the samples was determined by observing the dissociation of the band phase. XPS measurements XPS spectra were acquired using a Thermo Scientific Escalab 250xi. A pass energy of 20 eV, a step size of 0.1 eV, a spot size of 400 µm, and 30-fold scan averaging were employed. The data were processed using CasaXPS software, employing a Shirley background for all adjustments. Sulfur 2p spectra are characterized by a doublet (2p³ / ² and 2p¹ / ²) with an area ratio of 2:1 and a spin-orbit coupling Δ = 1.18 eV. The error bars were estimated using a Monte Carlo process in CasaXPS. GIWAXS characterization The GIWAXS measurements were performed at Brookhaven National Lab on the 12-ID Soft Matter Interfaces (SMI) beamline of the National Synchrotron Light Source II (NSLS-II) at a beam energy of 16.1 keV. The 2D scattering patterns were acquired at an X-ray incidence angle of 0.10° using a Pilatus 900 kW detector with a pixel size of 172 µm, positioned 279 mm from the sample. The sample and detector were enclosed in a vacuum chamber to suppress air scattering. To cover the desired scattering angle range, the vertically oriented, elongated detector was moved horizontally along a fixed arc. The images were subsequently visualized in Xi-CAM software and assembled and radially integrated using a custom code. Electron microscopy The PBTTT thin films were investigated using scanning electron diffraction (SED), a 4D STEM technique with a low convergence angle (< 1 mrad). SED data were acquired using a TS Spectra 300 at 300 kV with a convergence angle of approximately 0.5 mrad and a diffraction-limited beam diameter of approximately 6 nm. The current was approximately 20 pA with an image duration of 1 ms, resulting in an electron flux of approximately 140 e⁻ / Ų. The diffraction patterns for the SED were acquired using a single-chip Medipix3 direct electron detector (256 x 256 pixels) and a camera length of 145 mm, resulting in a Δs of 0.0094 Å⁻¹ and a maximum scattering vector of 1.1 Å(⁻¹). Additional data were collected under the same conditions at 200 kV. Prior to the measurements, thin layers were deposited by doctor blade coating onto substrates coated with polystyrene sulfonate (PSS), dissolved in water, transferred to TEM grids and then heated to the ribbon phase on the TEM grids. Polarized photothermal deflection spectroscopy The PDS was performed with a tunable light source consisting of a 250 W quartz-tungsten halogen lamp coupled with a 250 mm focal length monochromator. The monochromatic excitation beam was modulated at 13 Hz. The sub-gap absorption data A(E) were fitted to an exponential curve to extract the Urbach energy Eµ. Polarized ultraviolet visible spectroscopy UV-Vis-NIR spectra were acquired using a Shimadzu UV-3600i dual-beam spectrometer with a slit width of 3 nm and a data interval of 1 nm. Background spectra of the empty glass substrate were acquired separately. Results Mixing polymers Mixed and doped PBTTT films were prepared as described above using different ratios of P1 and P2. These were compared with reference films containing either doped P1 or P2 exclusively. Undoped pure and mixed samples were also prepared as references. Thin-film differential scanning calorimetry (DSC) was used to investigate the thermal transitions of the undoped reference films compared to an undoped mixed film. The results are shown in Fig. 1. The films composed exclusively of P1 and P2 both exhibited two exotherms upon cooling, but at significantly different temperatures of 156 °C and 244 °C (P1) and 139 °C and 227 °C (P2), respectively. The mixed film, consisting of 17.1 mol% P2 (and 82.9 mol% P1), also showed two corresponding transitions, occurring at temperatures of 149 °C and 234 °C. The different DSC transition temperatures were also confirmed by temperature-dependent AFM measurements.The DSC of the mixture was not observed as a combination of the DSCs of the films formed exclusively from P1 and P2, but showed distinct transition temperatures, indicating that no phase separation into P1 and P2 domains takes place in the mixed films, but that the two components are closely mixed within the mixed films. Improvements in electrical conductivity The electrical conductivity of the doped mixed films and the doped reference films was measured as described above. The development of the conductivity of the mixed films as a function of the molar ratio of P2 is shown in Fig. 2. Pure P1 and P2 films (0% and 100% P2) exhibited relatively high electrical conductivities of 1650 S / cm and 3170 S / cm, respectively. Surprisingly, however, many of the mixed films showed better performance than the pure P1 and P2 films. In particular, the electrical conductivity increased with increasing proportion of P2 in the mixed films from 0% to a peak value of 4810 S / cm at a composition of 17.1 mol% P2. This maximum electrical conductivity is more than 2.5 times higher than the conductivity of the layer formed from doped pure P1. With a further increase in the P2 content, the conductivity initially decreased and then remained at an approximately constant value in mixtures with P2 concentrations of ≈ 35 mol% or more. All further experimental data refer to the optimal mixture composition with 17.1 mol-% P2. Improvements in thermoelectric properties Fig. 3 shows the temperature-dependent electrical conductivity of the doped reference films P1 and P2 compared to a doped mixed film with 17.1 mol% P2 (and 82.9 mol% P1). For all three films, the conductivity exhibits a largely similar temperature dependence, typical for high-conductivity PBTTT, with a flattening or even a decrease above 250 K. However, the mixed film shows a significantly improved electrical conductivity compared to the reference films across the entire temperature range. Fig. 4 shows the Seebeck coefficient of all doped layers. The Seebeck coefficient varies linearly with temperature, which agrees well with the Mott formula for metals and previous measurements on ion-exchange doped PBTTT. The Seebeck coefficient for the mixed layer is consistently higher than that of pure P2 and comparable to that of pure P1. Fig. 5 shows the thermal conductivity of the doped reference films P1 and P2 compared to a doped mixed film containing 17.1 mol% P2 (and 82.9 mol% P1) and to analogous undoped films. As shown in the figure, all doped samples exhibit a significantly higher thermal conductivity than the undoped samples, reflecting the contribution of the heat transported by the introduced charge carriers. The doped mixed film exhibits the highest thermal conductivity across the entire temperature range. Surprisingly, however, the increase in thermal conductivity is not as large as would have been expected based on the Wiedemann-Franz relation for the increase in electrical conductivity. As a result, the doped mixed film shows a significant improvement in electrical conductivity compared to the doped, unmixed reference films, while the Seebeck coefficient remains comparable and the thermal conductivity increases only relatively slightly. This improved electrical conductivity without a proportional increase in thermal conductivity leads to an improvement in the quality factor for thermoelectric power (ZT), as defined by Equation 1 above. Simply put, the blended film can be assumed to exhibit improved thermoelectric performance (higher ZT) compared to the reference materials. It is unexpected that the blend of polymers leads to this performance improvement, and in particular to the advantageous increase in electrical conductivity, without significantly affecting the other factors, such as a significant increase in thermal conductivity. Optimization of the doping concentration To investigate the mixtures as a function of the doping concentration, the films were exposed to lower concentrations of FeCl3 for defined doping times, as described above in the Methodology section. As shown in Fig. 6, the electrical conductivity of the mixture films increased with the doping time, while the Seebeck coefficient decreased. A maximum power factor of 173 µW m⁻¹K⁻² was achieved at an exposure time of 130 s (Fig. 7). This is a relatively high power factor for this particular polymer. The measurements yielded a power factor for the thermoelectric power (ZT) in the range of 0.04 to 0.05. While this is relatively low compared to some conjugated polymers used for thermoelectric materials, it represents a significant improvement over the reference materials and thus confirms the overall approach. Proposed mechanism Without committing to a specific theory, it is assumed that the improved performance, particularly the increased conductivity of the doped mixed films compared to the doped pure films P1 or P2, is due to the fact that the presence of the higher molecular weight polymer in a mixture with the lower molecular weight polymer increases the concentration of bonding chains long enough to form a bridge between the crystalline ribbons. These bonding chains are thought to facilitate charge transport between the crystalline domains (in this case, crystalline ribbons). This is illustrated schematically in Fig. 8. On the left is a representation of P1 films comprising discrete crystalline domains (101) of aligned polymers separated by amorphous regions (102) where there are no or very few bonds between the domains (i.e.,minimal linkage chains). On the right is a representation of a mixed film containing both P1 and P2. The mixed structure includes a number of P2 polymers extending as linkage chains across adjacent crystalline domains (103), as well as some that fold over and remain within the discrete crystalline domain (104). The presence of the P2 polymer extending across the crystalline domains (103) is thought to enable enhanced charge transport between the crystalline domains. Evidence of bonding chains It is difficult to directly observe linkage chains using microstructural measurements. Therefore, a series of structural and spectroscopic measurements were performed to investigate differences in the microstructure of the mixed and reference polymer films, which might provide alternative explanations for the observed behavior. The morphology was investigated using atomic force microscopy (AFM), revealing clear and uniform banding in the undoped P1 and P2 films as well as in the doped mixed film with 17.1 mol% P2 (Figs. 9a, 9c, and 9b, respectively). The bands are preferably oriented perpendicular to the doctor blade coating direction, as shown by the two clearly defined, off-center peaks corresponding to the band periodicity in the Fourier transform of the AFM images (insertions). The white arrows indicate the doctor blade coating direction, i.e., the direction of chain alignment. Statistical analysis of the bandwidths averaged over three different samples and scan ranges of each polymer (Fig. 9a-c) revealed a similar periodicity of 68.9 ± 0.4 nm (P1), 71.8 ± 0.1 nm (mixtures), and 68.9 ± 0.4 nm (P2), even though P2 has almost twice the molecular weight of P1. This suggests that the polymer chains in P1 films are almost completely extended (i.e.,i.e., that the chains are predominantly straight and exhibit only minimal folds), whereas in P2 films most chains contain hairpin chain folds (see Fig. 8), so they exhibit a similar bandwidth to P1 films. The small but statistically significant increase in bandwidth in mixed films compared to P1 and P2 films suggests that the incorporation of higher molecular weight chains leads to a slight elongation of the chain extension in the bands, which is consistent with the incorporation of linker chains, likely resulting in a straightening of the chains and chain ends in the grain boundary regions. X-ray photoemission spectroscopy (XPS) was used to quantify the doping concentrations. Both TFSI and PBTTT contain sulfur atoms with a well-resolved chemical shift, allowing the quantification of the molar ratio of TFSI to PBTTT from the peak fitting of the SN 2p signal (Fig. 9d-f). Prior to the XPS measurements, the surface of the films was washed in acetonitrile to remove surface residues. For the peak fitting, separate 2p doublet components for each of the neutral PBTTT chain segments and PBTTT radical cations were used, along with separate components for the two sulfur atoms on TFSI. The calculated molar doping ratio for all three samples (doped P1, doped mixture and doped P2 films) was approximately one ion per monomer, with only slight differences of 8.5×1020cm-3(mixture), 8.3×1020cm-3(P2) and 7.6×1020cm-3(P1), which are not statistically significant. As a result, there are only minimal changes in the charge carrier concentration, which could otherwise dominate the transport properties. To investigate the crystal order of the three films (doped P1, doped mixed, and doped P2), grazing-incidence wide-angle X-ray scattering (GIWAXS) was used. For X-rays incident along the blade's coating direction, a strong (010) π-π stacking diffraction pattern and only a weak (003) backbone diffraction peak in the plane were observed (Fig. 10a-c), whereas for X-rays incident perpendicular to the blade's coating direction, the (003) backbone is much stronger and the π-π stacking peak is weaker (Fig. 10d-f). This suggests that the mixed films exhibit a similarly high level of chain alignment in the plane as the pure films. The paracrystallinities gπ-π for the π-π stacking (Fig. 11) were extracted. All three films exhibit very small gπ-π values ​​between 6.3% and 7.1% in their doped state (Fig.11), with only minor differences between the three samples, reflecting the high degree of structural order in the doped band-phase films. These small differences in gπ-π are insufficient to explain the sharp increase in conductivity observed for the mixture. Paracrystallinity values ​​were also extracted from the (003) backbone diffraction, and only a slight decrease in paracrystallinity was observed for P1 compared to P2 and the mixtures. Four-dimensional scanning transmission electron microscopy (4D-STEM) was used to capture the diffraction patterns of individual bands and the grain boundaries between them. The local diffraction patterns (inset in Fig. 12a,d) provide evidence of a clear chain alignment with strong anisotropic π-π stacking and backbone diffraction. The band structure is clearly visible in the ring-field diffraction (ADF) images of both P1 (Fig. 12a,b) and the mixed films (Fig. 12d,e), with a bandwidth of approximately 70 nm, similar to that observed in the AFM images. Interestingly, strong interchain ordering and π-π stacking diffraction were observed at all points in the local diffraction patterns, with only a relatively small decrease in intensity at the center of the grain boundary from one band to the next. The grain boundaries in the mixed films appear particularly narrow in the ADF images and were estimated to have a width of less than 5 nm.Without committing to a theory, this suggests that even at the grain boundaries, a pronounced chain alignment is present. The exact values ​​for the grain boundary width are uncertain due to the finite probe size required for sufficient diffraction resolution. The grain boundary regions in P1 proved to be somewhat wider (2.0 nm) than in the mixed films. This could be a direct manifestation of the linking chains in the mixed film, leading to a straightening of the polymer chains at the grain boundary regions. Areas where the grain boundary contrast disappeared, i.e., where adjacent bands appeared to merge, were also discovered in the ADF images. These are highlighted in Fig. 12e. It appeared that these were more pronounced in the mixed films than in the P1 films, where they were essentially absent.Both observations are interpreted as an expression of the incorporation of bond chains and could be directly related to the conductivity behavior, i.e., narrower grain boundaries probably lead to faster transport from band to band and higher conductivity. Spectroscopic characterization of the films was also performed. Using polarized ultraviolet visible (UV-VIS) spectroscopy, the degree of chain alignment was quantified by the dichroic ratio R = A∥∥ / A⊥, where A∥∥ and A⊥ represent the peak absorption at either the 0-0 or 0-1 absorption band. The dichroic ratio reached a value of 12.7 for undoped mixture films at the neutral π-π* absorption band of 553.1 nm (Fig. 13a) and 8.9 for doped mixtures at the polaron-induced absorption band at 875.0 nm (Fig. 13b). Both transitions are known to be polarized along the polymer chain axis, and the dichroic ratio therefore provides a quantitative measure of the overall chain alignment.Very similar dichroic conditions were also observed in films consisting exclusively of P1 or P2, showing that a high and comparable degree of chain alignment was achieved in the reference films compared to the mixed films. To compare the degree of energetic disorder, photothermal deflection spectroscopy (PDS) was performed on the undoped films (Fig. 13c). This allows for a volume-sensitive investigation of the sub-bandgap tail of the excitonic joint density of states and permits the extraction of the Urbach energy Em from an exponential fit in the sub-bandgap region. Films consisting exclusively of P1 ( ) exhibited a slightly lower E(m) value of 50.3 meV than films consisting exclusively of P2 (51.6 meV) or the mixture (51.1 meV). This could indicate higher crystalline order with fewer torsional defects in P1, which is consistent with DSC, although the effect is too small to be reflected in GIWAXS. In fact, the differences in Em are due to the measurement accuracy of PDS.A more convincing piece of evidence for a somewhat higher degree of crystalline order in P1 could be the observation of a slight decrease in Emsein, obtained from the power-law fittings of the optical absorption. Polarized Fourier-transform infrared spectroscopy (FTIR) was used on doped films to measure the polaron-induced absorption in the mid-infrared range, which is a sensitive indicator of the degree of polaron delocalization. Typically, two peaks are observed, designated as the B-peak at approximately 3500–2500 cm⁻¹ and the A-peak at approximately 1000 cm⁻¹. A redshift of the B-peak and a more intense A-peak indicate more pronounced polaron delocalization along the chains and in the direction between the chains. The spectra recorded with the polarizer along the coating direction of the doctor blade (Fig.Figure 13d), which primarily responds to the polarons located on the highly aligned chains in the crystalline domains, shows that the polarons in the crystalline domains of P1 are somewhat more delocalized than those in P2 and mixed films, while P2 and mixed films exhibit similar polaron delocalization. This could indicate a lower degree of torsional defects and a longer conjugation length in the crystalline domains of P1, which is consistent with the DSC and PDS results. When the polarizer is oriented perpendicular to the doctor blade coating direction, the spectra preferentially examine polarons located in the less well-aligned grain boundary regions (Fig. 5e). These are significantly more localized than those in the crystalline regions, which is consistent with the more disordered chain conformation expected at the grain boundaries.Overall, the spectroscopic characterization shows that while there may be subtle differences in the crystalline order of the three samples, with P1 exhibiting longer conjugation lengths and more delocalized polarons, these cannot be responsible for the observed conductivity behavior, as they would imply that P1 should have the highest conductivity. Finally, the electrical conductivity data were fitted using a heterogeneous transport model by Kaiser and Graham. The fittings show that for all samples and temperatures, the grain boundary resistance is at least an order of magnitude greater than the metallic transport within the grain, and that therefore the grain boundaries (where bonding chains would be present) strongly restrict the conductivity. Fig. 14 shows the conductivity of the two parallel grain boundary conduction pathways: hopping / tunneling (dashed lines) and disordered metallic transport (solid lines). Based on the above findings, and without committing to a specific theory, it seems likely that the improved conductivity compared to films consisting solely of P1 is due to the inclusion of bonding chains. In P1 films, transport across grain boundaries is restricted by the absence (or low concentration) of bonding chains, whereas in mixed films, the bonding chains facilitate transport across grain boundaries. Regarding the improvements over P2, without committing to a theory, it is assumed that the high concentration of hairpin chain folds present in the P2 films could increase the degree of structural and energetic disorder in the grain boundaries and limit the ability of the polymer chains to incorporate themselves as bonding chains. In contrast, there are probably fewer chain folds in the mixed films, and the longer chains are likely to be able to maintain their chain extension more easily and transition into the next crystal regions. Mixing two different polymers Fig. 15 shows an AFM image illustrating that band formation can also be achieved with a mixture of two different polymers with similar electronic structures. Specifically, Fig. 15 shows an example of a mixture where the second conjugated polymer (i.e., the high molecular weight polymer) is poly((2,5-bis(3-alkylselenophen-2-yl)thieno(3,2-b)thiophene) (PBSTT) and the first conjugated polymer (i.e., the lower molecular weight polymer) is poly(2,5-bis(3-alkylthiophen-2-yl)thieno(3,2-b)thiophene) (PBTTT). The substitution of some thiophene atoms with selenium atoms leads to a subtle adjustment of the energy levels and electronic structure between the two polymers. Selenophene is considered more electron-rich than thiophene, which is why PBSTT is likely to have a slightly lower ionization potential than PBTTT.However, the difference is small, typically less than 100 meV, and is also influenced by solid-state effects such as intermolecular interactions. Another method to achieve the desired matching of the energy levels of the two polymers could involve attaching various electron-withdrawing or electron-donating groups to the polymer backbone, for example, by replacing some of the backbone protons with electron-withdrawing fluorine atoms (which tend to increase the ionization potential) or by modifying the alkyl side chains, preferably the segment directly bonded to the backbone, by replacing some or all of the CH2 / CH3 units with more electron-withdrawing CF2 / CF3 units or more electron-donating ethylene glycol units, preferably maintaining the same side chain length. It is understood that when designing such modifications to the side chains or backbone, care must be taken to ensure that similar solubility properties are maintained for both polymers. When selecting the two polymers, it is important that the polymer chain conformation and molecular shape remain sufficiently similar so that the mixture of the two polymers can form a common ribbon phase, as shown in the AFM image in Fig. 15. In the system shown in Fig. 15, 20 mol% of the selenophene-substituted polymer (PBSTT) with a C10H21 alkyl chain length and a molecular weight of Mn = 36.9 kDa was used together with 80 mol% of the PBTTT polymer with a C12H25 alkyl chain and a molecular weight of 7.6 kDa. The 20% mixture of PBSTT in PBTTT was found to have a high conductivity of 3000 S / cm, which was more than 20% higher than that of the pure PBTTT film. It is likely that this can be further optimized by choosing different alkyl chains.For example, the same alkyl chain length could be used for both polymers to minimize disorder and structural incompatibility when mixing the second polymer. It is indeed remarkable that, despite this structural difference, the two polymers were able to form a common ribbon phase, and the mixture exhibited higher conductivity than the pure polymer. This demonstrates the robustness of the approach and suggests that there could be considerable scope for chemical modification of the second polymer while maintaining the compatibility of the two polymers for ribbon phase formation. References [1] (a) O. Zapata-Arteaga, S. Marina, J. Martin, J.S. Reparaz,* and M. Campoy-Quiles*, ACS Energy Lett. 2020, 5, 2972-2978; (b) S. Wang, G. Zuo, J. Kim, H. Sirringhaus, Progress in Polymer Science 2022, 129, 101548 .[2] M. G. Kanatzidis*, Nat. Mater. 2022, 21, 503-513 .[3] (a) Y. Huang, D.H. Lukito Tjhe, X. Huang, I. McCulloch and H. Sirringhaus, Appl. Phys. Lett. 2021, 119; (b) V. Vijayakumar, V. Untilova, M. Bahri, L. Herrmann, L. Biniek,* N. Leclerc,* und M. Brinkmann*, Adv. Energy Mater. 2019, 9, 1900266; (c) G. Giri, J. Reinspach, DA Fischer, LJ Richter, J. Xu, S. Benight, A. Ayzner, M. He, L. Fang, G. Xue, MF Toney und Z. Bao*, Chem. Mater. 2015, 27, 2350-2359 .[4] (a) M. Xiao, SB Lee, LMA Perdigão, A. Luci, DA Warr, SP Senanayak, M. Nikolka, M. Statz, Y. Wu, A. Sadhanala, S. Schott, R. Carey, Q. Wang, M. Lee, C. Kim, A. Onwubiko, C. Jellett, H. Liao, W. Yue, K. Cho, G. Costantini, I. McCulloch und H. Sirringhaus*, Adv. Mater. 2020, 32, 2000063; (b) L.Shaw, Y. Diao, JA Reinspach, JWF To, MF Toney, RT Weitz and Z. Bao*, ACS Appl. Mater. Interfaces 2016, 8, 9285-9296; (c) L. Biniek, T. Heiser, R. Bechara and M. Brinkmann*, Macromolecules 2013, 46, 4014-4023 .[5] T. Hallam, N. Zhao, I. Nandhakumar, M. Kemerink, M. Heeney, I. McCulloch and H. Sirringhaus*, Phy. Rew. Lett. 2009, 103, 256803 .[6] RJ Pandolfi, DB Allan, E Arenholz, L Barroso-luque, SI Campbell, TA Caswell, A Blair, F de Carlo, S Fackler, AP Fournier, G Freychet, M Fukuto, RJ Kline, R Li, C Liman, S Marchesini. J. Synchrotron Radiat. 2018, 25, 1261-1270. QUOTES INCLUDED IN THE DESCRIPTION This list of documents cited by the applicant was automatically generated and is included solely for the reader's convenience. The list is not part of the German patent or utility model application. The DPMA accepts no liability for any errors or omissions. Cited non-patent literature J. Chang et al, Adv. Funct. Mater. 2010, 20, 2825-283

[0103] M. Statz et al., Comm. Phy. 2018, 1, 16

[0103] O. Zapata-Arteaga, S. Marina, J. Martin, JS Reparaz,* und M. Campoy-Quiles*, ACS Energy Lett. 2020, 5, 2972-2978; (b) S. Wang, G. Zuo, J. Kim, H. Sirringhaus, Progress in Polymer Science 2022, 129, 101548

[0138] MG Kanatzidis*, Nat. Mater. 2022, 21, 503-513

[0138] Y. Huang, DH Lukito Tjhe, X. Huang, I. McCulloch and H. Sirringhaus, Appl. Phys. Lett. 2021, 119; (b) V. Vijayakumar, V. Untilova, M. Bahri, L. Herrmann, L. Biniek,* N. Leclerc,* und M. Brinkmann*, Adv. Energy Mater. 2019, 9, 1900266; (c) G. Giri, J. Reinspach, DA Fischer, LJ Richter, J. Xu, S. Benight, A. Ayzner, M. He, L. Fang, G. Xue, MF Toney and Z. Bao*, Chem. Mater. 2015, 27, 2350-2359

[0138] M. Xiao, SB Lee, LMA Perdigão, A. Luci, DA Warr, SP Senanayak, M. Nikolka, M. Statz, Y. Wu, A. Sadhanala, S. Schott, R. Carey, Q. Wang, M. Lee, C. Kim, A. Onwubiko, C. Jellett, H. Liao, W. K. Yue.Cho, G. Costantini, I. McCulloch und H. Sirringhaus*, Adv. Mater. 2020, 32, 2000063; (b) L. Shaw, Y. Diao, JA Reinspach, JWF To, MF Toney, RT Weitz und Z. Bao*, ACS Appl. Mater. Interfaces 2016, 8, 9285-9296; (c) L. Biniek, T. Heiser, R. Bechara und M. Brinkmann*, Macromolecules 2013, 46, 4014-4023

[0138] T. Hallam, N. Zhao, I. Nandhakumar, M. Kemerink, M. Heeney, I. McCulloch und H. Sirringhaus*, Phy. Rev. Lett. 2009, 103, 256803

[0138] RJ Pandolfi, DB Allan, E. Arenholz, L. Barroso-luque, SI Campbell, TA Caswell, A. Blair, F. de Carlo, S. Fackler, AP Fournier, G. Freychet, M. Fukuto, RJ Kline, R. Li, C. Liman, S. Marchesini. J. Synchrotron Radiat. 2018, 25, 1261–1270

[0138] .

Claims

Semi-crystalline polymer material, preferably for a thermoelectric device, wherein the material comprises: a first conjugated polymer; a second conjugated polymer mixed with the first conjugated polymer; and a dopant; wherein the second conjugated polymer has a number-average molecular weight that is 25% or more greater than the number-average molecular weight of the first conjugated polymer. Semi-crystalline polymer material according to claim 1, wherein the first and the second conjugated polymer are preferably aligned in a certain direction. Semi-crystalline polymer material according to one of the preceding claims, wherein the HOMO energy level of the second conjugated polymer is within 0.1 eV of the HOMO energy level of the first conjugated polymer; or wherein the LUMO energy level of the second conjugated polymer is within 0.1 eV of the LUMO energy level of the first conjugated polymer. Semi-crystalline polymer material according to one of the preceding claims, wherein the number-average molecular weight of the second conjugated polymer is 50%, preferably 75%, greater or more than the number-average molecular weight of the first conjugated polymer. Semi-crystalline polymer material according to one of the preceding claims, wherein the number-average molecular weight of the second conjugated polymer is 500% greater or less than the number-average molecular weight of the first conjugated polymer, preferably 250% greater or less. Semi-crystalline polymer material according to one of the preceding claims, wherein the first conjugated polymer and the second conjugated polymer differ only by single-atom substitutions in the polymer repeating unit or by the addition of electron-withdrawing or electron-donating groups to the polymer backbone, while the bonding pattern of the polymer backbone remains the same; and / or wherein the first conjugated polymer and the second conjugated polymer comprise and preferably consist of the same conjugation system. Semi-crystalline polymer material according to one of the preceding claims, wherein the molar ratio of the second conjugated polymer in the mixture of the first conjugated polymer and the second conjugated polymer is less than 50%, preferably less than 30%. Semi-crystalline polymer material according to one of the preceding claims, wherein the molar ratio of the second conjugated polymer in the mixture of the first conjugated polymer and the second conjugated polymer is greater than 5%, preferably greater than 10%. Semi-crystalline polymer material according to one of the preceding claims, wherein the semi-crystalline polymer material comprises discrete crystalline domains and wherein the polymers of the first conjugated polymer within the crystalline domains assume a substantially rectilinear conformation. Semi-crystalline polymer material according to one of the preceding claims, wherein the dopant is selected from a neutral molecular dopant, an ionic dopant, an ion introduced by an ion exchange process, an electrochemically introduced ion, a dopant introducing an electrophilic or nucleophilic attack, or a dopant inducing a proton transfer, wherein the dopant is optionally FeCl3. Semi-crystalline polymer material according to one of the preceding claims, comprising a molar doping ratio of 0.01 to 2 charge carriers per monomer, preferably 0.1 to 1 charge carrier per monomer. Semi-crystalline polymer material according to any one of the preceding claims, wherein the first conjugated polymer and / or the second conjugated polymer is selected from homopolymers such as polythiophenes, polyselenophenes, polyfurans or polythiazoles; donor-donor copolymers such as polymers comprising thiophene, thienothiophene and / or selenophene units; and / or donor-acceptor copolymers such as polymers comprising diketopyrrolopyrrole units or copolymers comprising thiophene and fluorinated benzothiadiazole. Semi-crystalline polymeric material according to one of the preceding claims, wherein the first conjugated polymer and / or the second conjugated polymer comprises a thiophene unit within the polymeric backbone. Semi-crystalline polymeric material according to one of the preceding claims, wherein the first conjugated polymer has a number-average molecular weight of at least 15 kDa and / or wherein the second conjugated polymer has a number-average molecular weight of at least 30 kDa. A method for producing a semi-crystalline polymer material, preferably for a thermoelectric device, wherein the method comprises: mixing a first conjugated polymer with a second conjugated polymer; applying the mixed polymers to a substrate; and doping the first and second conjugated polymers with a dopant; wherein the second conjugated polymer has a number-average molecular weight that is 25% or more greater than the number-average molecular weight of the first conjugated polymer. A method for producing a semi-crystalline polymer material, comprising: separately doping a first conjugated polymer and a second conjugated polymer with a dopant; mixing the first doped conjugated polymer with the second doped conjugated polymer; applying the mixed polymers to a substrate; wherein the second conjugated polymer has a number-average molecular weight that is 25% or more greater than the number-average molecular weight of the first conjugated polymer. The method of claim 15 or 16, wherein the application of the mixed polymers further comprises the alignment of the mixed polymers; preferably wherein the alignment of the mixed polymers comprises coating the mixed polymers on a substrate with a doctor blade, or wherein the alignment of the mixed polymers comprises drawing the mixed polymers into fibers. Method according to any one of claims 15 to 17, wherein the mixing of the first conjugated polymer with the second conjugated polymer comprises mixing a solution comprising the first conjugated polymer with a solution comprising the second conjugated polymer. Method according to one of claims 15 to 18, wherein the doping of a first conjugated polymer and a second conjugated polymer with a dopant comprises ion exchange doping, molecular doping, electrochemical doping, solid-state diffusion doping, gas-phase doping, mixed-solution doping, sequential doping, acid doping, preferably ion exchange doping. Use of the polymeric material according to any one of claims 1 to 14 for forming a layer on a substrate. Use of the polymeric material according to any one of claims 1 to 14 as a thermoelectric material. Use of the polymer material according to any one of claims 1 to 14 for forming a layer of a thermoelectric device. layer on a substrate, wherein the layer comprises the polymeric material according to any one of claims 1 to 14. Thermoelectric device comprising the polymer material according to any one of claims 1 to 14. Thermoelectric device according to claim 24, depending on claim 2, wherein the chains of the first and second conjugated polymer are preferably aligned in the direction of the current flow or in the direction of a temperature gradient in the device.

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