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Morphology controls the thermoelectric power factor of a doped semiconducting polymer

Morphology controls the thermoelectric power factor of a doped semiconducting polymer SCIENCE ADVANCES RESEARCH ARTICLE MATERIALS SCIENCE Copyright © 2017 The Authors, some rights reserved; Morphology controls the thermoelectric power factor of exclusive licensee American Association a doped semiconducting polymer for the Advancement of Science. Distributed 1 1,2 2 2 Shrayesh N. Patel, * Anne M. Glaudell, Kelly A. Peterson, Elayne M. Thomas, under a Creative 2 2 † Kathryn A. O’Hara, Eunhee Lim, Michael L. Chabinyc Commons Attribution NonCommercial The electrical performance of doped semiconducting polymers is strongly governed by processing methods and License 4.0 (CC BY-NC). underlying thin-film microstructure. We report on the influence of different doping methods (solution versus vapor) on the thermoelectric power factor (PF) of PBTTT molecularly p-doped with F TCNQ (n = 2 or 4). The vapor-doped films have more than two orders of magnitude higher electronic conductivity (s) relative to solution-doped films. On the basis of resonant soft x-ray scattering, vapor-doped samples are shown to have a large orientational correlation length (OCL) (that is, length scale of aligned backbones) that correlates to a high apparent charge carrier mobility (m). The Seebeck coefficient (a) is largely independent of OCL. This reveals that, unlike s, leveraging strategies to improve m have a smaller impact on a. Our best-performing sample with the largest OCL, vapor-doped PBTTT: −1 −2 F TCNQ thin film, has a s of 670 S/cm and an a of 42 mV/K, which translates to a large PF of 120 mWm K . In addition, despite the unfavorable offset for charge transfer, doping by F TCNQ also leads to a large PF of −1 −2 70 mWm K , which reveals the potential utility of weak molecular dopants. Overall, our work introduces im- portant general processing guidelines for the continued development of doped semiconducting polymers for thermoelectrics. INTRODUCTION poly(2,5-bis(3-tetradecylthiophen-2-yl)thieno[3,2-b]thiophene) Controlling the electrical doping of organic semiconductors is critical (PBTTT). We focus on p-doping with organic acceptors [2,3,5,6- to the performance of organic electronic devices (1). Doped semi- tetrafluoro-7,7,8,8-tetracyanoquinodimethane (F TCNQ) and 2,5- conducting polymers can serve as conductive interlayers for organic difluoro-7,7,8,8-tetracyanoquinodimethane (F TCNQ)] introduced light-emitting diodes (OLEDs) (2) and solar cells (2, 3) and can im- either in solution or from the vapor phase (Fig. 1). The results of prove the performance of organic thin-film transistors (OTFTs) (4). our experiments demonstrate how different processing and doping One emerging application of doped semiconducting polymers in- methods affect the thermoelectric PF. In particular, we find that align- volves organic thermoelectrics—materials that interconvert heat and ment of ordered domains is the critical factor leading to higher s with- electricity (5–7). The solution processability of semiconducting poly- out lowering a, thereby leading to enhancements in the PF. Using −1 −2 mers provides the opportunity to use roll-to-roll processing and print- these methods, we have found a PF of 120 mWm K for PBTTT: ing technologies for new classes of thermoelectric modules where the F TCNQ, which is among the highest reported values for semi- legs are thin films in rolled or corrugated designs (8–13). To realize conducting polymers (7). the potential of semiconducting polymers for thermoelectrics, how- The s value of semiconducting polymers is related to the product ever, the relationship between processing and the resulting thermo- of the carrier concentration (n) and carrier mobility (m). However, electric properties must be better understood. because of electronic disorder, the apparent m of a material will de- All of the physical properties of a material that define its thermo- pend on n because of the occupancy of electronic states with varying electric performance depend on carrier density (n), including electri- mobility (16–19). Through advances in molecular design and pro- cal conductivity (s), Seebeck coefficient (or thermopower) (a), and cessing, solution-processable semiconducting polymers, such as PBTTT, thermal conductivity (k)(14). The thermal-to-electrical energy con- have high charge carrier mobilities (m >1cm /V s) in field-effect version efficiency is related to the dimensionless figure of merit, ZT = transistors (20). These studies have revealed that the degree of 2 2 a sT/k, where T is the temperature in Kelvin and a s is the power electronic and structural disorder strongly influences m (21, 22). In factor (PF). Optimizing ZT is quite challenging because as n increases, s field-effect measurements, conduction occurs very close (within and k increase while a decreases (14). Organic semiconductors fre- ~1 nm) to the polymer-dielectric interface. The microstructure is quently have imperfect ordering in thin films, leading to an electronic generally described as (para)crystalline p-stacked domains intercon- structure that depends strongly on their morphology (15). Because nected by tie chains (21). Electrically doped films may require high processing methods widely vary in many studies of thermoelectric concentrations of dopant in the bulk (for example, >1 dopant per performance, it is difficult to form clear connections between mor- 10 monomers), which lead to strong perturbations of the morphology phology and thermoelectric performance (7). and structure relative to pristine films. Whether processing methods Here, we elucidate the connection between thin-film micro- that lead to high field-effect m also lead to high-bulk s has not been structure and thermoelectric transport properties (s and a) of p-doped well studied. The ability to tune the electronic structure of small-molecule or- ganic acceptors (23) has resulted in versatile p-type dopants for Materials Research Laboratory, University of California, Santa Barbara, Santa Barbara, CA 93106, USA. Materials Department, University of California, Santa Barbara, Santa semiconducting polymers (16, 24–26). Electron acceptors have been Barbara, CA 93106, USA. traditionally co-deposited with molecular organic donors to generate *Present address: Institute for Molecular Engineering, University of Chicago, 5640 organic charge transfer salts and metals (27–29). Mixing an organic South Ellis Avenue, Chicago, IL 60637, USA. †Corresponding author. Email: mchabinyc@engineering.ucsb.edu acceptor into a polymer leads to an integer charge transfer if the Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 1of13 | SCIENCE ADVANCES RESEARCH ARTICLE higher s relative to films cast from a polymer/dopant solution. For example, vapor doping a predeposited PBTTT with F TCNQ results in a s of ~250 S/cm, which is nearly two orders of magnitude higher relative to a solution-doped film with a similar concentration of F TCNQ (37). The general explanation for the enhancement in s is related to an overall better macroscopic film quality establishing the underlying microstructure for more efficient charge transport (35–37). Although these explanations describe enhancement in s,itisunclear how more efficient charge transport influences the overall thermo- electric properties of semiconducting polymers. Here, we examine how morphology affects the bulk s and a of PBTTT. PBTTT is a solution-processable polymer, where the field- effect mobility, crystal structure, and morphology have been well characterized for neat films (40–45), providing a strong foundation to study how molecular doping influences the morphology and charge transport. In addition, PBTTT has an accessible liquid crys- talline transition temperature above ~140°C, which permits thermal processing to enhance local and long-range order (40–42). Using PBTTT and other thiophene-based polymers, we have previously discovered an empirical connection between electrical conductivity and thermopower across a range of doping methods (33, 34). This broad correlation has been modeled as a result of the electronic density of states (DOS) and energy-dependent mobility, but the connection with morphology has not been made clear. Here, we use PBTTT as a model system to demonstrate how the correlation of alignment in ordered domains at the nanoscale dominates the resulting s at high n. These results suggest a pathway to increase the thermoelectric PF of semiconducting polymers. Fig. 1. Chemical structure and doping process. (A) Chemical structure of PBTTT and F TCNQ (n = 2 or 4) and the corresponding IE or EA. (B) Solution and vapor RESULTS AND DISCUSSION doping routes used to achieve doped films. Processing doped films of PBTTT We prepared highly conductive thin films of PBTTT using different offset between HOMO (highest occupied molecular orbital) [ioniza- processing methods, with the dopant added in solution or infiltrated tion energy (IE)] of the polymer and LUMO (lowest unoccupied mo- from the vapor phase (Fig. 1). A detailed procedure can be found in lecular orbital) [electron affinity (EA)] of the acceptor is sufficient to Materials and Methods, and we outline the critical differences here. provide a thermodynamic driving force for electron transfer (Fig. 1) We specifically focused on the limit of high doping to determine the (16, 30–32). connection between morphology and thermoelectric transport prop- How a dopant is incorporated into a semiconducting polymer is erties. For solution doping, 10 wt % of F TCNQ relative to PBTTT critical in dictating the resulting charge transport properties (33–37). [molar ratio (MR) of ~1 dopant to 4 monomers] was added to a so- Achieving high s (>10 S/cm) requires relatively high charge carrier lution of PBTTT. We have previously found that this composition is 19 3 concentrations (>10 /cm ) due to the observed superlinear increase near the maximum possible to readily form continuous thin films in conductivity in many materials (16, 30). If a dopant is added to the during spin casting (38). The solution was spin-coated to obtain a casting solution, this concentration requires as much as 10 weight % doped thin film in a N environment and annealed at 150°C for (wt %) relative to the monomer in the solution, which can be difficult 10 min to remove the solvent. These conditions were used to minimize because of solubility limits of neutral organic acceptors (26). Although weight loss of F TCNQ from the film (38) while also being above the chain aggregation appears to aid in efficient charge transfer in some liquid crystalline (LC) transition temperature of the neat polymer. The cases (32), highly charged polymers can gel or precipitate from solution same processing conditions were used with solution-doped samples, (38). Consequently, one must take great care in determining optimal where the dopant was F TCNQ. The typical thicknesses of solution- casting concentrations and temperatures that lead to macroscopically doped films were 40 to 50 nm. To form films doped by infiltration of homogenous films (35, 38). Alternatively, first casting a neat film from F TCNQ and F TCNQ from the vapor phase, we exposed spin-coated 4 2 solution and then subsequently doping has emerged as a versatile route films of PBTTT (21 ± 4 nm) prepared using different thermal treat- to yield macroscopic homogenous films with high s (33, 35–37). For ments to the vapor of each compound. In a N -filled glove box, we example, depositing a thin layer of organic acceptor from the vapor deposited the dopants by placing the samples underneath the lid of a phase can lead to diffusion of the dopant into the organic semi- sealed jar containing a few milligrams of dopant. The bottom of the conductor (37, 39) and solid-state charge transfer to generate highly jar was heated to ~210°C, which led to a rise in temperature of the conductive films (37). Recently, it has been shown that doping spin- substrate, which was located underneath the lid for the jar, to 75° to coated neat thiophene-based polymers either from the vapor phase 85°C. These temperatures are below the LC transition of PBTTT and (33, 37) or through sequential solution casting (35, 36) can lead to are known not to cause substantial changes in the structural order Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 2of13 | SCIENCE ADVANCES RESEARCH ARTICLE and charge mobility in OTFTs (41). The samples were exposed to the vapor for 10 min, which was sufficient to reach concentrations of Table 1. Summary of electronic conductivity (s), Seebeck coefficient F TCNQ in PBTTT films comparable to the solution-doped films. (a), and PF of doped PBTTT films. For the sample on an OTS-treated −1 −2 substrate, s = 670 ± 4S/cm, a =42 ± 6 mV/K, and PF = 120 ± 30 mWm K . All other samples reported in this table are on untreated quartz substrates. Changes in processing increase electrical conductivity −1 −2 Efficient charge transfer occurs between PBTTT and F TCNQ using Dopant Condition s (S/cm) a (mV/K) PF (mWm K ) both vapor- and solution-based doping. Ultraviolet-visible near- F TCNQ Solution— 2.08 ± 0.01 45 ± 4 0.42 ± 0.09 infrared (UV-vis-NIR) spectroscopy shows that neat poly(2,5-bis(3- as-cast tetradecylthiophen-2-yl)thieno[3,2-b]thiophene) (PBTTT-C ) thin Solution— 3.51 ± 0.05 60 ± 9 1.3 ± 0.4 film has a main absorption peak at ∼2.2 eV and a shoulder at ∼2.1 eV annealed that is bleached upon doping. New absorption peaks appear at 1.41 Vapor— 114.1 ± 0.5 32 ± 4 12 ± 3 and 1.60 eV upon introduction of F TCNQ and are assigned to its as-cast anion radical (16, 25, 37). The spectral features for the F TCNQ rad- ical anion absorption are similar to those in poly(3-hexylthiophene) Vapor— 220.00 ± 0.02 39 ± 5 32 ± 9 annealed (P3HT):F TCNQ films and have comparable absorptivity for heavily doped films (35). A subband gap transition for positive polarons of F TCNQ Solution— 0.41 ± 0.02 111.7 ± 0.1 0.52 ± 0.03 PBTTT is observed at ∼0.5 eV (fig. S2), but the precise position of as-cast polaronic features between 1 and 2 eV is difficult to assign because −3 Solution— 2×10 ± 755 ± 100 0.11 ± 0.03 −4 of the strong absorption peaks of the F TCNQ anion radical. We annealed 2× 10 observe no significant differences between UV-vis-NIR spectra be- Vapor— 13.7 ± 0.2 130 ± 20 23 ± 6 tween vapor-doped as-cast and annealed films, indicating that the as-cast annealing step does not change the concentration of F TCNQ in Vapor— 36 ± 3 140 ± 20 70 ± 20 the film (fig. S3). A comparison of the main absorption of PBTTT annealed (∼2.2 eV) for the vapor-doped film relative to the solution-doped film reveals more bleaching in the former (a factor of 0.78 lower peak area) and also a slightly higher absorbance of F TCNQ radical anion. This lower peak area translates to a slightly higher MR of 0.3 relative to the MR of 0.25 in the solution-doped film (37). but only the vapor-doped film has a small red shift. In contrast to Strikingly large differences in s are found between vapor- and F TCNQ, because of the relatively high vapor pressure of F TCNQ, 4 2 solution-doped samples despite this small difference in MR of the we observe significant dedoping of solution-doped PBTTT:F TCNQ dopant. Electrical conductivity measurements (Table 1) indicate films when thermally annealing at 150°C in N environment (fig. 65 times higher s for the vapor-doped PBTTT:F TCNQ annealed S2). The MR in the casting solution is ~0.28 F TCNQ per monomer 4 2 film (s = 220 ± 0.02 S/cm) relative to the solution-doped PBTTT: of PBTTT, similar to the solution-doped samples with F TCNQ. F TCNQ annealed film (s of 3.51 ± 0.05 S/cm). We observe a similar Doping from vapor allows for a comparison of infiltration into as- trend in as-cast films, where s = 114.1 ± 0.5 S/cm for the vapor- cast and annealed films. When vapor-doping an annealed PBTTT doped film and s = 2.08 ± 0.01 S/cm for the solution-doped film. film with F TCNQ, the primary absorption peak is 30% higher rela- Knowing that both doping methods yield comparable carrier con- tive to the as-cast doped films. This difference is, in part, from the fact centrations, the large difference in s must be related to the apparent that annealing a neat film results in an increase in primary absorption m, which is calculated to be ~2.5 and ~0.040 cm /V s for the vapor- peak by about 20% (comparison between Fig. 2, A and B). Despite dopedannealedfilmand the solution-dopedannealedfilm, respec- the small difference in the primary absorption peak, F TCNQ radical tively (assuming F TCNQ is fully ionized and all charges are free anion absorption is comparable between vapor-doped films but carriers). The higher apparent m with vapor-doped films is consistent slightly less than the solution-doped as-cast film. We attribute the with Hall effect mobility measurements on vapor-doped PBTTT: higher absorption of the F TCNQ radical anion in solution-doped films relative to vapor-doped films to the elevated temperature of F TCNQ, which was revealed to be ~2 cm /V s (37). Two possible factors that can contribute to the difference in apparent m are differ- the sample during vapor deposition. In addition, the absorption curve ences in local energetic disorder or the long-range morphology. is quantitatively similar in the NIR regime (fig. S2) for the F TCNQ- To determine whether the enhancement in s was unique to doped films. F TCNQ, we also examined samples doped with F TCNQ. The Comparison of the conductivity measurements reveals that the 4 2 intermediate fluorination level of F TCNQ resultsinanEAof vapor-doped PBTTT:F TCNQ films yield a higher s than solution- 2 2 ∼4.59 eV (46), which is expected to result in an unfavorable offset doped films do. The annealed film has a s of 36 ± 3 S/cm, and the for charge transfer with PBTTT (IE, ∼5.10 eV) in isolated materials as-cast film has a s of 13.7 ± 0.2 S/cm. On the other hand, the solution- compared to F TCNQ (EA, ∼5.24 eV). Note that charge transfer is doped as-cast film has the lowest s of 0.41 ± 0.02 S/cm despite the dictated not only by the offset between EA and IE but also by the higher concentration of F TCNQ radical anion. The high vapor pres- electrostatic interaction in solids (47). Although the EA of F TCNQ sure of F TCNQ does not allow direct comparison of solution and 2 2 is lower than the IE of PBTTT-C , integer charge transfer is indicated vapor-doped annealed films; it leads to significant dedoping and thus −3 by the presence of F TCNQ radical anion peaks at 1.43 and 1.62 eV several orders of magnitude lower s of 2 × 10 S/cm. Nevertheless, the (Fig. 2). Similar observations have been made for structurally similar fact that the vapor F TCNQ-doped films yield higher s further reiter- weak dopants with P3HT (26). For both doping methods, PBTTT: ates that an underlying microstructural feature is at play, which leads F TCNQ as-cast films yield a large decrease in the primary absorption, to a higher apparent m. Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 3of13 | SCIENCE ADVANCES RESEARCH ARTICLE Fig. 2. UV-vis spectra of neat and doped PBTTT thin films. UV-vis spectra of (A) neat PBTTT and PBTTT:F TCNQ thin films and (B) PBTTT:F TCNQ thin films. Solution- 4 2 doped films are at a dopant concentration of 10 wt %, and vapor-doped films are for dopant exposure of 10 min (all spectra normalized by thickness). Neat PBTTT films were annealed at 180°C, whereas solution-doped films were annealed at 150°C. Comparison of absorption spectra of annealed neat PBTTT and as-cast neat PBTTT can be found in the Supplementary Materials. Processing does not change crystalline structure in Thein-planescattering profiles alongthe q axis of the 2D xy doped films GIWAXS images (Fig. 4B and fig. S5) reveal that all highly doped To determine whether the difference in s between vapor- and solution- films lead to an increase in the scattering vector for the (110) reflec- doped films correlates to local structural order, we performed grazing tion (tables S1 and S2). This increase indicates a compression of the incidence wide-angle x-ray scattering (GIWAXS) (Fig. 3). PBTTT characteristic p-p–stacking distance. The value of d decreases forms semicrystalline films with strong texturing. For neat PBTTT, from 0.367 nm for the neat film to 0.355 nm for the F TCNQ the out-of-plane scattering features (along the q axis) correspond to solution-doped film and 0.353 nm for the vapor-doped film. The the lamella-stacked side chains (h00). The in-plane scattering features F TCNQ-doped films also showed compression upon doping, where −1 (along the q axis) correspond to the (-11-3) reflection at q = 14.2 nm d = 0.353 nm for an as-cast solution-doped film, d = 0.357 nm xy xy 110 110 −1 related to the chain axis and the (110) reflection at q =17.1 nm for a vapor-doped as-cast film, and d = 0.361 nm for a vapor- xy 110 associated to the p-stacking direction (43). The scattering pattern does doped annealed film. When thermal annealing a F TCNQ solution- not qualitatively change upon doping; that is, the features are shifted doped film at 150°C, we observe rapid dedoping, where d essentially and broadened upon doping, but no new strong scattering features returns to the value of an annealed neat film. The vapor F TCNQ- emerge. A significant difference between doped films and the neat film doped annealed film yields the smallest decrease in d ,which is −1 is the significant blurring of the off-axis features at q = 14.1 nm , consistent with the low concentration of F TCNQ radical anion xy 2 suggesting disorder along the chain axis, as expected from dopants (Fig. 2). We have previously shown that TCNQ, which does not re- with disordered molecular orientation. sult in an integer charge transfer with PBTTT due to the large offset The out-of-plane scattering features of the two-dimensional (2D) between EA and IE, causes no measurable changes to the d value GIWAXS images show changes upon doping in the alkyl side-chain relative to the neat film (38). This result suggested that the presence stacking direction (h00) (Fig. 4A). For all doped films, we observe a of ionized polymers and dopants is partly responsible for the de- small increase relative to the neat d = 2.12 nm (summarized in ta- crease in the characteristic d spacing. Overall, the changes to in- 100 110 bles S1 and S2). Because the true out-of-plane scattering is in the in- plane scattering features are quite similar with both doping methods. accessible region in the grazing incidence geometry, we also obtained Previously, nuclear magnetic resonance (NMR) experiments on high-resolution specular x-ray scattering on the vapor-doped PBTTT: bulk samples of PBTTT:F TCNQ cast from solution were carried out F TCNQ annealed films. The specular scattering results indicate that to determine the location of F TCNQ relative to the backbone of 4 4 13 1 d = 2.37 nm, which is about a 0.25-nm increase relative to the neat PBTTT (38). 2D C{ H} heteronuclear (HETCOR) NMR data dem- film (d = 2.12 nm). In addition, the change in peak width in the (h00) onstrated a very close contact between carbon atoms on the cyano direction reveals the degree of induced disorder along the side-chain group of F TCNQ and aromatic protons on the PBTTT backbone. stacking direction. As shown in fig. S4, the introduction of the dopant This proximity and the calculated geometry for model structures of from the vapor phase introduces additional disorder relative to the neat charge transfer between polymers and F TCNQ from density func- film, which is consistent with our previous work with PBTTT:F TCNQ tional theory suggested that the F TCNQ was intercalated between 4 4 solution-doped films (38). Similar changes are observed for films doped the polymer backbones (38). Full structural modeling to compare the with F TCNQ. The alkyl stacking spacing d increases by ∼0.25 nm in NMR and x-ray data was not carried out in the initial study. In con- 2 100 the solution-doped PBTTT:F TCNQ as-cast film, whereas d in- trast, a recent report on vapor-doped PBTTT:F TCNQ films proposed 2 100 4 creases by ∼0.10 nm in the vapor-doped as-cast film and by only that F TCNQ resides in the side-chain region of PBTTT films, with ∼0.04 nm in the vapor-doped annealed film. This trend indicates the the rationale that the p-stacking spacing, separation between chains, varying extent of dopant incorporated within the aliphatic side chains, was unperturbed in their x-ray scattering data (37). As discussed above, which is consistent with the UV-vis spectra showing a slightly smaller we observe a compression in the p-stacking spacing relative to the neat concentration of the F TCNQ radical anion in vapor-doped films. polymer, as observed in heavily doped PBTTT films with a variety of Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 4of13 | SCIENCE ADVANCES RESEARCH ARTICLE Fig. 3. GIWAXS for PBTTT films as a function of processing. 2D GIWAXS images for (A) annealed neat PBTTT, (B)F TCNQ vapor-doped annealed film, (C)annealedF TCNQ 4 4 solution-doped film, (D)F TCNQ vapor-doped annealed film, and (E) as-cast F TCNQ solution-doped film. The (100) reflection in GIWAXS image in (E) was blocked off with lead 2 2 tape to allow longer exposure time without saturating the detector. Images are obtained at Stanford Synchrotron Radiation Lightsource (SSRL) beamline 11-3. dopants (33, 38) and is reversible upon thermal removal of the dopant. the cyano groups of F TCNQ; such a change in tilt occurs in cocrystals Independent of these differences, one can infer that F TCNQ resides of PBTTT and [6,6]-phenyl C61 butyric acid methyl ester (PCBM) (48). in the side chains for vapor-doped samples by a simple considera- tion. On the basis of the observed out-of-plane scattering (h00), we Correlation of domain orientation controls the would expect <11% thickness change upon doping crystalline electrical conductivity domains. However, for the structural model where the molecule is Small perturbations in the local structure observed by GIWAXS are intercalated between in-plane p-stacked chains, we would expect a sig- similar for both solution doping and vapor doping methods and do nificant expansion of crystallites to compensate for the volume occu- not explain the differences in s. Our conclusion is in contrast to the pied by the dopant molecule, leading to a large in-plane expansion or explanation of Kang et al.(37), where they attributed the higher s only significant increase in thickness to maintain the lateral size. Starting to minimal perturbation to local order. However, it is important to with a neat film with thickness of 25 ± 3 nm, the vapor doping process remember that GIWAXS only provides structural information for leads to a thickness of 27 ± 3 nm, which is comparable to the expected short-range crystalline domains (5 to 20 nm). For larger length scales, expansion of the crystallites (fig. S6). Because the film thickness does we require a small-angle scattering method. Here, we use polarized not markedly increase, F TCNQ resides within the aliphatic side resonant soft x-ray scattering (RSoXS) that can reveal the length scale chains (assuming a high degree of crystallinity, which is known for of molecular orientation in both the crystalline and amorphous do- PBTTT). At this point, we cannot reconcile these data with the previous mains over the range from ~10 to 1000 nm (42). NMR data without detailed modeling of the spectra, which is beyond RSoXS leverages soft x-ray absorption to increase the scattering the scope of the work here. However, it is possible that the structure of length and to control the scattering contrast (42, 49). In p-conjugated the doped films has tilted PBTTT chains that allow a closer contact with molecules, the transition dipole moment (TDM) for the excitation Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 5of13 | SCIENCE ADVANCES RESEARCH ARTICLE entational alignment of polymer backbones is the key factor influenc- ing the performance of OTFTs (42). We prepared samples to mimic, or replicate, films used for mea- surements of electrical conductivity. 2D RSoXS patterns were measured by transmission through polymer thin films spin-coated on top of un- treated silicon nitride windows. We assume minimal difference in mor- phology when comparing films on silicon nitride windows and untreated quartz substrate used for conductivity measurements. On the other hand, vapor-doped PBTTT films on octadecyltrichlorosilane (OTS)–treated substrates were floated onto silicon nitride windows. All the scattering images showed a diffuse isotropic ring (fig. S7) that was azimuthally averaged to obtain a 1D scattering profile [intensity (I) versus q] (Fig. 5). The OCL is one-half of the characteristic length scale (d* =2p/q*) obtained from the primary scattering peak (q*)of the Lorentz-corrected (I*q ) scattering profile. The peak position was determined from fits using log-normal function. For equivalent annealing times (10 min), the OCL of neat PBTTT is sensitive to the annealing temperature above the LC transition. The OCL is ∼140 and ∼180 nm when annealed at 150° and 180°C, respectively (fig. S8 and table S3). Note that neat films were annealed at 180°C, whereas solution-doped films were annealed at 150°C to minimize dedoping of F TCNQ. Doped films with the higher values of the OCL correlate to higher s and, thus, higher apparent m because of the comparable level of dop- ing. The OCL of doped films of PBTTT depends strongly on the pro- cessing method (Fig. 5 and table S3). F TCNQ vapor-doped annealed films have an OCL of ∼220 nm, whereas the OCL for an annealed doped film cast from solution is ∼44 nm. The vapor-doped film and neat annealed PBTTT films have relatively similar OCLs. An- nealed films vapor-doped with F TCNQ have an OCL of ~210 nm, suggesting that the introduction of the molecular dopant does not have alarge impact on the OCL.The OCLof ∼100 nm of solution-doped PBTTT:F TCNQ films after thermal annealing at 150°C for 10 min approaches the OCL of a neat film annealed at 150°C (OCL, ~140 nm) likely due to dedoping, which allows the morphology to change. There- fore, we assert that the improved orientational alignment of the backbone leads to a higher m in vapor-doped films, similar to previous Fig. 4. Out-of-plane and in-plane scattering profiles of annealed neat and observations for OTFTs. doped PBTTT thin films. GIWAXS line cuts of (A) out-of-plane scattering and The trends in the OCL for as-cast doped films further confirm the (B) in-plane scattering. Black, neat; orange, F TCNQ solution doping; purple, significant role of backbone alignment in controlling s. The OCL of F TCNQ vapor doping; green, F TCNQ solution doping; blue, F TCNQ vapor doping. 4 2 2 The F TCNQ-doped film corresponds to as-cast conditions. Dashed red lines are an as-cast neat film is ∼70 nm, whereas that of the as-cast film from guides to the eye relative to the peak positions for the neat film. All scattering the PBTTT:F TCNQ solution is ∼40 nm. This difference indicates profiles correspond to thermally annealed films, except for the F TCNQ solution that doping in solution slightly reduces the OCL relative to a neat doping, which is for the as-cast case. a.u., arbitrary units. film before annealing. Moreover, the OCL observed for as-cast films is essentially identical to the annealed solution-doped film (OCL, of a core electron to an unoccupied p-orbital is orthogonal to the ∼44 nm). This similarity shows that annealing a heavily doped film p-conjugated plane of the molecule (50). By using a linearly polarized has no effect on enhancing the backbone alignment, and confirms why soft x-ray with the electric field vector in the plane of the film and tuning the s of as-cast and annealed solution-doped films is nearly identical. The the photon energy to the C 1s to p* resonance (285.4 eV for PBTTT), fact that the OCL does not increase with annealing a solution-doped film scattering contrast arises from the variations in molecular orientation. is not surprising because the PBTTT is now heavily charged, which shifts For PBTTT thin films, strong resonant scattering is expected in trans- the LC transition temperature to higher values and thus does not enter an mission mode because the p-stacking orientation is primarily in the LC mesophase when annealing at 150°C. Attempts to anneal the films at plane of the film (that is, TDM of the 1s to p* resonance is in the plane higher temperatures result in significant dedoping (38). For as-cast of film). Using RSoXS measurements, we can quantify the backbone F TCNQ solution-doped films, we observe a qualitatively different alignment through the orientational correlation length (OCL). The scattering profile relative to the as-cast neat and F TCNQ-doped films. OCL is defined as the average length over which the polymeric back- The scattering peaks are broader where the primary peak is at a lower −1 bones (that is, the LC director) drift out of alignment with each oth- q ∼ 0.03 nm (OCL, ∼130 nm). This suggests that aggregation of er. The OCL was shown to have an empirical exponential relationship PBTTT in solution, which dictates the as-cast thin film morphology, with field-effect m values, providing the first direct evidence that ori- is different when using a weak molecular dopant like F TCNQ. Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 6of13 | SCIENCE ADVANCES RESEARCH ARTICLE Fig. 5. RSoXS of doped films. Lorentz-corrected scattering profiles (log-log scale) from azimuthally averaged RSoXS images at a photon energy of 285.4 eV for (A)annealed and (B) as-cast films. The curves were offset for clarity. Black, neat; purple, F TCNQ vapor doping; blue, F TCNQ vapor doping; orange, F TCNQ solution doping; green, F TCNQ 4 2 4 2 solution doping. RSoXS experiments were performed at Advanced Light Source (ALS) beamline 11.0.1.2. When vapor-doping an as-cast film, we observe a qualitatively dif- A significant assumption in previous work on the relationship ferent scattering profile relative to the as-cast neat film. Essentially, two between the m from OTFTs and the OCL was that the OCL, a bulk −1 broad scattering peaks are seen around q ∼ 0.04 nm (OCL, ~130 nm), property, correlates to an interface-dominated charge transport mea- −1 and another is seen at the experimental low q limit (<0.01 nm ). We surement. Here, we can eliminate that assumption because s is a bulk see a similar effect for a F TCNQ vapor-doped film, where we see a transport property, and conclusively show the correlation between −1 peak around q ∼ 0.05 nm (OCL, ∼60 nm) and another at experi- OCL and s and, thus, apparent m. To determine the sensitivity of s mental low q limit. The second scattering peak corresponds to an to the OCL, we plot log(s) versus OCL for various doped films in OCL that is larger than that of an as-cast neat film. The larger cor- Fig. 6A. The increase in s is most significant at lower OCL values and related domains provide better charge transport, leading to the high then approaches a plateau at higher OCL values. The data points in s = 114.1 ± 0.5 S/cm for F TCNQ vapor-doped as-cast films and s = Fig. 6A are at comparable doping levels with respect to F TCNQ samples, 4 4 13.7 ± 0.2 S/cm for F TCNQ vapor-doped as-cast films. We can rule thus indicating that the apparent m is the parameter increasing with out that this change is due to a thermal annealing process; the sample OCL. The correlation of the interfacial mobility from OTFTs and that underneath the lid is calibrated to be approximately 85°C, which is from the bulk conductivity likely holds in PBTTT because GIWAXS well below the LC transition temperature and has no effect on the typically shows highly oriented crystallites with a thickness equivalent OCL (fig. S8 and table S3). to the total film. In materials where the interfacial and bulk structures are dissimilar, we might not expect such a relationship to hold. Increasing electrical conductivity through processing Extrapolation of s to higher OCL values using Fig. 6A provides Comparison of s and the OCL reveals that long backbone correlation valuable information on the limits of electrical performance. The lengths allow for more efficient charge transport. It has been previously chain alignment process effectively increases the OCL, where an infinite shown that the OCL of PBTTT can be controlled depending on the po- OCL corresponds to perfect alignment. When approaching the limit of larity of the substrate surface (42). In particular, substrates function- an infinite OCL, one would expect s to plateau as the net increase in m alized with a nonpolar monolayer of OTS lead to the largest OCL of becomes smaller (51). For comparison, the mobility in OTFTs of ∼380 nm for an annealed neat PBTTT thin film (42). Vapor-doping PBTTT thin films prepared on an OTS-treated substrate and subse- a PBTTT thin film annealed on an OTS-functionalized substrate with quently strained aligned has been measured (52). The alignment pro- F TCNQ yielded a s of 670 ± 4 S/cm. This s is a factor of ∼3higher cess resulted in a factor of ~2 increase in m according to field-effect value than the vapor-doped film on a bare substrate. The corresponding transistor measurements (52). This observation suggests that s could OCLisalsohigherat ∼350 nm (versus 210 nm for the doped film on a increase to ~1300 S/cm relative to our highest-performing PBTTT: bare substrate). UV-vis-NIR absorption measurements (fig. S3) con- F TCNQ vapor-doped film if the bulk and interfacial mobilities re- firm that the doping level is similar for both cases. Therefore, the in- main correlated. crease in s is entirely from a higher apparent m and reiterates the The determination of the connection between the OCL and s significance of the OCL as a parameter to describe the trends in s. shows the significance of processing on the transport properties of Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 7of13 | SCIENCE ADVANCES RESEARCH ARTICLE of the two studies because of the question of the carrier concentration with different dopants. One can also imagine achieving high s when casting films from a polymer/dopant solution if the processing steps lead to a solid-state thin film with a large OCL. Comparison to previous work reported in the literature further re- iterates the critical role of controlling the morphology to achieve effi- cient charge transport. Sequential doping of P3HT with F TCNQ results in films with higher s relative to solution-doped films (35), which is driven by improved interconnectivity between ordered p-stacked do- mains (36). Sequentially doping P3HT with F TCNQ in nonpolar sol- vents yields a larger 5 to 10 factor increase in s (36). These films also have good interconnectivity, but phase segregation of dopants in the disordered domains was suggested as the reason for the higher s. Other factors such as the molecular structure of the dopant and its sol- ubility in the polymer can also influence morphology and thus s (26). Polar side chains on polymers have been found to increase the ther- mal stability of F TCNQ-doped films, which is particularly useful for thermoelectric energy conversion (53). The environmental stability of F TCNQ in thiophene-based polymers is a concern due to photo- chemical reactions, which will require investigation into appropriate encapsulation for devices (53–55). Role of morphology on Seebeck coefficient and PF Although the impact of the orientation of domains on s is relatively straightforward to understand, the impact on a is less clear. Unlike s, which describes the transport of charge carriers relative to an electric field, a describes the migration of charge carriers relative to temper- ature gradient at the open circuit condition. This is quantified by measuring the voltage drop (DV) relative to temperature difference (DT). Fundamentally, a is related to the population in the electronic DOSs of a material and carrier scattering processes. The expression for the thermopower as a function of electronic conductivity func- tion s(E)isgiven as k ðE  E Þ sðEÞ ∂f ðEÞ B F a ¼ dE ð1Þ e k T s ∂E where E is the Fermi energy, s is the total conductivity, f(E)isthe Fer- mi function, and k /e is a natural unit of thermopower of 86.17 mV/K (56). As the semiconductor is p-doped, the Fermi level shifts closer to the valence band, which results in a decrease in the value of a.The introduction of a molecular dopant into a polymer will also modify the local structure and morphology, making it difficult to model the thermoelectric properties with a constant electronic DOS. There have been significant efforts to model the thermoelectric properties of poly- Fig. 6. The relationship between OCL and thermoelectric material proper- mers (56–58), but these models do not, as yet, consider morphology. ties. (A) Measured electronic conductivity (s), (B) measured Seebeck coefficient The overall thermoelectric properties of PBTTT depend strong- (a), and (C) calculated PF versus the corresponding OCL values, as determined ly on processing methods (Table 1). For example, a = 60 ± 9 and −1 −2 from the RSoXS experiments (table S3). PF = 1.3 ± 0.4 mWm K for the solution-doped annealed film of PBTTT:F TCNQ, whereas annealing followed by vapor infiltration −1 −2 doped semiconducting polymers. Although the vapor doping pro- of F TCNQ leads to a =39±5 mV/K and PF = 32 ± 9 mWm K cess yields the highest values of s here, the vapor doping process with (Table 1). Despite the lower a, PF is higher for vapor-doped films F TCNQ is not the only method to achieve high s. In our previous due to the nearly 100-fold increase in s. The difference in a with va- work, doping annealed films of PBTTT through exposure to a vapor por doping can be attributed, in part, to the slightly higher concen- of a fluorinated trichlorosilane (FTS) or immersion in a 4-ethylbenzene tration of F TCNQ based on the UV-vis spectra, but the morphology sulfonic acid solution yielded s of around 1000 S. The likely origin is of these two films is also quite different. Furthermore, a =42±6 mV/K that the OCL was set before doping and not substantially perturbed by for the F TCNQ vapor-doped film on an OTS-treated substrate, the doping process. It is difficult to directly compare the conductivities which indicates that a is less sensitive to the substrate treatment than Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 8of13 | SCIENCE ADVANCES RESEARCH ARTICLE s. However, the near factor of 3 increase in s results in the high PF of performance of polyacetylene (~10 S/cm) (63). A question is wheth- −1 −2 120 ± 30 mWm K . er the higher conductivities of these materials are due to the electronic With F TCNQ as the dopant, we observe a values of 111.7 ± structure of the conjugated backbone or other factors. In comparison 0.1 mV/K for the as-cast solution-doped film, 130 ± 20 mV/K for the to these materials, the presence of alkyl side chains on semiconducting vapor-doped as-cast film, and 140 ± 20 mV/K for the vapor-doped polymers, such as PBTTT, results in a significant volume of the annealed film (Table 1). The higher a relative to F TCNQ is not material being insulating. Accounting for the insulating side chains surprising because of the likely lower concentration of F TCNQ in reveals an effective electronic conductivity (s ) of the conjugated core 2 eff the film and potentially lower efficiency of carrier formation. The representing densely packed polymer chains (37). For PBTTT with F TCNQ vapor-doped annealed film yields the highest a, consist- tetradecyl side chains, the conjugated core accounts for approximately ent with the slightly lower doping efficiency, as indicated by UV-vis 15% of the volume in a film (determined from the unit cell of PBTTT). (Fig. 2). Owing to the high s of 36 ± 3 S/cm, the corresponding PF is The s would translate to ~4000 S/cm for our highest-performing eff −1 −2 70 ± 20 mWm K .Despite the lower s relative to the F TCNQ film. If these structural changes can be achieved without changing −1 −2 vapor-doped film on a bare quartz substrate, the PF is greater because the a, the PF could reach values of ~500 mWm K at high levels of the large a value. This shows the importance of tuning a while not of doping similar to PEDOT. Removal, or shortening, of the side significantly sacrificing s. chains of a polymer leads to difficulties in processing, that is, PEDOT Comparison of a to the corresponding OCL reveals that, unlike s, is poorly soluble and is cast as a dispersion or directly grown on a sub- the value of a does not show marked changes (Fig. 6B). For F TCNQ- strate. The dopant itself also modifies the volume of the doped doped samples, a is in the range of 30 to 60 mV/K for the full OCL material. However, this comparison shows that conjugated backbones window. In addition, the F TCNQ samples are essentially the same other than PEDOT have significant promise if their structure can be (~130 mV/K) at different OCL values. There appears to be a small up- judiciously modified to maintain their processability and allow for in- ward trend for vapor-doped samples, but overall a is less sensitive to corporation of the dopant. polymer chain alignment. Some recent work on the anisotropy of s and a on PEDOT:PSS [poly(3,4-ethylenedioxythiophene)-poly(styrenesulfonate)] Potential impact of processing on thermal conductivity films (59) and P3HT films (60)alsosuggeststhat a is relatively un- Our study reveals the potential benefits that local chain alignment can affected by polymer chain alignment, and thus, changes in PF with play in further improving PF, but one should also consider the impact the OCL (Fig. 6C) are dominated by the increase in s. This observa- on ZT. Changes in the OCL likely modify the thermal conductivity (k) tion would be expected if the shape of the electronic DOS was rela- through both the phonon (lattice) contribution (k )(64) and, at suf- tively constant and the number of carriers was similar, that is, no ficiently high n and s, the electronic contribution to thermal conduc- large change in the Fermi level. Temperature-dependent a and s mea- tivity (k ). Note that experimental challenges still remain on the surements are needed to fully elucidate the transport mechanism from determination of k of thin films and particularly along the in-plane both doping processes. direction (6). Techniques such as suspended microdevices (65)or Previously, we determined an empirical correlation where a follows the membrane-based ac calorimetry (66) can help to determine the −1/4 a power-law dependence with s (a º s ) and PF follows a square in-plane k. The membrane-based ac calorimetry method revealed an 1/2 −1 −1 root dependence with s (PFº s ) for a variety of p-doped thiophene- in-plane k of 0.39 W m K for an as-cast neat PBTTT film (1 mmthick) based polymers. This correlation was primarily determined on solution- (66). Films with high s,for example, F TCNQ-doped film at 640 S/cm, doped films and thermal annealing conditions 150°C or below. In Fig. 7 , could result in a significant contribution from k . PEDOT:PSS, for exam- we plot the empirical correlations (dashed lines) along with a and PF ple, has a significant contribution from k at a s of ~500 S/cm and higher values reported in this study. The solution-doped samples follow the em- (67). Therefore, there may be an optimization process on the extent of pirical trends. On the other hand, vapor-doped samples deviate from alignment and doping level that leads to a minimal increase of k to the empirical trends, where the values are observed to be higher than achieve a high ZT. A potential optimization process can be seen through expected at corresponding s values. The positive deviation of a and the example of PBTTT:F TCNQ, which yields lower s of 36 S/cm rel- PF relative to the empirical trend line is most pronounced for vapor- ative to PBTTT:F TCNQ (220 S/cm) at a comparable OCL value of doped PBTTT:F TCNQ and also for both solution- and vapor-doped ∼200 nm (Fig. 5A). Despite the lower s, the PBTTT:F TCNQ film 4 2 −1 −2 PBTTT:F TCNQ films. The PF of our highest-performing PBTTT: yields a higher a and thus a larger PF of 70 mWm K relative to −1 −2 F TCNQ film is similar to our previous work on PBTTT vapor-doped the PBTTT:F TCNQ film (PF = 32 mWm K ) (Fig. 5, B and C). As 4 4 with an interfacial FTS (green diamond marker in Fig. 7 ) (33). It has re- a consequence, the lower s can be leveraged to minimize, in principle, cently been proposed that the power-law relationship is due to a change the electronic contributions to k while still being able to achieve a rel- in the energy-dependent conductivity function of the material [s(E)] atively high PF. (56). The higher PF here would suggest that this function is affected by processing conditions speculated in that work. Summary We have explored how solution- and vapor-doping a high-mobility Future routes to improve the PF of polymers p-type polymer, PBTTT-C , affect its thermoelectric transport prop- Although we have focused only at the limit of high doping level, system- erties and its underlying microstructure. Overall, vapor-doping with atically varying the doping level and the OCL provides a route to opti- either F TCNQ or F TCNQ yields higher s relative to solution-doped 4 2 mization of the PF. We can also consider the performance of PEDOT, films. The enhancement in s is not related to the local order because which has been the benchmark conducting polymer for thermoelectrics the perturbations to the local structure are minimal and similar with −1 −2 (7). A PF as high as ∼460 mWm K has been reported (61), and s either doping route. We determined using RSoXS that the alignment of ∼5000 S/cm has been measured for metallic template polymerized of ordered domains, quantified through the OCL, is a critical parameter PEDOT doped with sulfuric acid (62). This s is comparable to the in explaining trends in s. The larger OCL for vapor-doped films allows Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 9of13 | SCIENCE ADVANCES RESEARCH ARTICLE Fig. 7. Trends in Seebeck coefficient and power factor. Log-log scale plot showing the trends in (A) Seebeck coefficient (a) and (B)PF(a s) versus electronic conductivity (s) for solution- and vapor-doped films. Orange markers are for vapor-doped films, and blue markers are for solution-doped films. Circle markers are for F TCNQ-doped films, square markers are for F TCNQ-doped films, and triangle marker is for the F TCNQ-doped film on OTS-treated substrate. Open markers 4 2 4 correspond to thermally annealed films, and filled markers correspond to as-cast films. The open green diamond is our previously reported FTS-doped PBTTT thin −1/4 1/2 film (33). Dashed lines are empirical trends [a º s and PF º s ] we previously reported on various doped semiconducting polymers (34). for efficient charge transport and thus a higher s relative to solution- All chemicals were used as received. PBTTT-C was synthesized doped films, which yield a smaller OCL. Owing to better long-range using literature procedure (68) with a number-average molecular correlation length of backbones, a F TCNQ vapor-doped PBTTT-C weight (M ) of 18,000 or 24,000 g/mol. 4 14 n casted on an OTS-treated substrate yields a high s of 670 ± 4 S/cm and corresponding large a of 42 ± 6 mV/K. This translates to a PF of 120 ± Neat PBTTT and PBTTT:F TCNQ (n = 2 or 4) thin-film −1 −2 30 mWm K —the highest reported value for F TCNQ-doped semi- sample fabrication conducting polymers. In addition, using a weaker molecular dopant Neat PBTTT solution preparation like F TCNQ can lead to a large a (140 ± 20 mV/K) while not signif- Neat PBTTT solutions at 5 mg/ml were prepared by dissolving PBTTT icantly sacrificing s (36 ± 3 S/cm), and thus yield a large PF of 70 ± in either CB or 1:1 CB:ODCB and heated to 120°C. Approximately −1 −2 20 mWm K . 1 hour was needed to fully dissolved PBTTT. The heated neat PBTTT With a better understanding of processing effects on s and a,we solution was filtered using 0.45-mm polytetrafluoroethylene (PTFE) sy- can now outline some general processing guidelines to achieve high ringe filter. The PBTTT solution gels when cooled below 80°C. As a thermoelectric PF. First, casting a neat semiconducting polymer film consequence, the neat PBTTT solution was maintained at 120°C before that forms locally p-stacked domains with long-range correlation doping or spin coating. lengths of the conjugated backbones provides an ideal microstructure Solution doping of PBTTT with F TCNQ (n = 2 or 4) for efficient charge transport for high apparent m. Second, introducing F TCNQ (3 wt %; n = 2 or 4) solution was prepared by dissolving the the molecular dopant into the polymer film (that is, from the vapor dopant in ODCB and heated to 150°C. To achieve 10 wt % dopant phase) to increase n, and in turn s, should lead to minimal perturba- concentration, an aliquot of F TCNQ solution was added to the neat tion to the local order while maintaining, or enhancing, the long-range PBTTT solution and heated to 120°C. After the addition of the dopant, correlation lengths of conjugated backbones. By leveraging the high the heated polymer solution became more viscous and immediately apparent m and then precisely controlling the dopant concentration, transitioned from a red to black color. These changes were indicative or by choosing a weaker molecular dopant, one can obtain a large a while of charge transfer between the polymer and dopant in solution. To not significantly sacrificing s. Overall, developing better doping routes minimize gelation and precipitation of the charged polymer, the and advancing our fundamental understanding of structure-property re- PBTTT:F TCNQ (n = 2 or 4) solution was maintained at 120°C before lationships of semiconducting polymers will have far-reaching implica- spin coating. tions on the deployment of lightweight and low-cost organic Substrates thermoelectric modules for thermal energy conversion and management. Thin films were prepared on quartz substrates (1.5 cm × 1.5 cm; Uni- versity Wafers) for conductivity and Seebeck measurements, native oxide silicon substrates (1.5 cm × 1.5 cm; International Wafer Services) MATERIALS AND METHODS for GIWAXS experiments, and silicon nitride windows (window size, Materials 1.5 mm × 1.5 mm; window thickness, 100 nm; frame size, 5 mm × 5 mm; Anhydrous chlorobenzene (CB) and o-dichlorobenzene (ODCB) thick frame, 200 mm; NX5150C, Norcada Inc.) for RSoXS experiments. were purchased from Sigma-Aldrich. F TCNQ and F TCNQ were The quartz and silicon substrates were cleaned by sonicating first in 4 2 purchased from TCI Chemicals. OTS was purchased from Gelest. acetone and then in isopropanol. Samples of PBTTT thin films on Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 10 of 13 | SCIENCE ADVANCES RESEARCH ARTICLE OTS-treated substrates were prepared by plasma-treating a quartz Four-point probe conductivity measurements were performed using substrate with air for 2 min. Then, the substrate was immersed in a custom-designed probe station in a N glove box. Voltage and current OTS and anhydrous toluene (volume ratio, 0.10 to 1) for 10 min, heated measurements were performed using a Keithley 2400 source measure to 80°C, and subsequently rinsed with toluene to achieve an OTS self- unit and Keithley 6221 precision current source. A constant current was assembled monolayer. The OTS-treated quartz substrates were heated applied to the outer contacts, and the resultant steady-state voltage re- to 80°C in a N glove box to remove the residual solvent. sponse was recorded from the inside contacts. The resistance (R;ohms) Spin-coating conditions of the sample was extracted from the slope of the VI sweep using Ohm’s Neat PBTTT thins films were spin-coated from a CB solution (5 mg/ml) law (V = IR). or from a 1:1 CB:ODCB solution (5 mg/ml). The heated solution (120°C) The Seebeck coefficient (a) measurements were performed in a wasspin-coated first at 1000rpm for45sandthenat3000rpm for15s N glove box using a custom-built setup. A detailed description of under ambient conditions. The heated (120°C) PBTTT:F TCNQ solu- the Seebeck coefficient measurement setup can be found in the study tion was spin-coated at 1000 rpm for 45 s and then at 3000 rpm for 15 s, of Glaudell et al.(34). Peltier elements 5 mm apart provided the tem- which leads to the macroscopically uniform thin films with thicknesses in perature difference (DT = T – T ). A minimal amount of thermal H C the range of 40 to 60 nm. conductive paste was applied to the tips of the thermal couple to ensure Thermal annealing good thermal contact between the thermocouple and the gold pads. The The neat PBTTT thin films were annealed in a N -filled glove box for measurement system has systematic error of 15% due to thermal 10 min at 180°C and then slowly cooled to 80°C. The thermal anneal- anchoring issues. A delay of 100 s was used for voltage measurements ing conditions for solution-doped films were at 150°C for 10 min in a to ensure that a steady-state temperature gradient was reached. The N -filled glove box. The thickness of the annealed film was approx- Seebeck coefficient was calculated from the slope of a linear fit for the imately in the range of 15 to 40 nm according to atomic force micros- DV versus DT plot. The measurements were taken within an approx- copy (AFM; Asylum MFP-3D) measurements. imate DT of ±3 K around 300 K. Vapor-doping process To infer the apparent charge carrier mobility (m), we used the MR Neat PBTTT thin films were first fabricated using the procedure out- (dopant/monomer) values of the vapor- and solution-doped films and lined above. Subsequently, an as-cast or thermal annealed neat film assumed that the dopants are fully ionized and all the hole carriers was vapor-doped with F TCNQ (n =2or 4) in aN glove box. Approx- generated contribute to conductivity. Knowing the unit cell of PBTTT, n 2 imately 5 to 10 mg of dopant were placed in a glass jar (Qorpak with a we can calculate the carrier concentration (n). After which, we can PTFE lined cap; diameter, ~5 cm; height, ~4.5 cm). The polymer sam- calculate the mobility using the equation m = s/(qn), where s is the ple was placed underneath the cap (near the center) using double- electronic conductivity and q is the charge (+1). sided tape. The closed jar was heated on a hotplate set to ~210°C. The typical heating times were in the range of 2 to 10 min. This Synchrotron x-ray scattering heating process leads to a partial vapor pressure of the dopant in 2D GIWAXS images were obtained using beamline 11-3 at SSRL lo- the jar. Successful doping of a PBTTT thin film was confirmed when cated on the SLAC (Stanford Linear Accelerator Center) National the film has a nearly transparent appearance (typically achieved after Accelerator Laboratory campus. Thin-film samples for GIWAXS 5 to 10 min for 25-nm PBTTT thin film). Successful doping was con- experiments were prepared following the procedures outlined above. firmed through UV-vis-NIR measurements. The temperature of a The samples were exposed to x-rays with a wavelength of 0.9752 Å, sample underneath the cap was measured using a thermocouple and 2D scattering images were obtained using a MAR345 image while the jar was heated. The sample was around 75°C after 5 min plate detector or MarCCD detector, which was placed 400 mm from and equilibrates to around 95°C after 30 min. the sample. A LaB sample was used as a standard for calibration. All samples were placed in a He-filled chamber to reduce air scattering UV-vis-NIR spectroscopy and minimize beam damage to the sample. The reported GIWAXS UV-vis-NIR spectra of thin films on 0.5-mm-thick quartz substrates images were taken at a grazing incident x-ray angle of 0.10 or 0.12, (1.5 cm × 1.5 cm) were obtained using the Shimadzu UV-3600 UV- which is above the critical angle of the polymer film and below the Nir-NIR Spectrometer at the UC (University of California) Santa critical angle of the silicon substrate. Barbara Materials Research Laboratory TEMPO Facility. The doped X-ray specular scattering was collected on beamline 2-1 using the films were placed in a custom-built airtight holder to ensure doping setup with Soller slits and a photomultiplier tube. The incident x-ray stability. Measurements were taken within a wavelength (l) range of energy was 11.5 keV. 300 to 2300 nm. RSoXS samples were prepared by directly spin-coating onto the silicon nitride windows following thin-film fabrication process out- Conductivity and Seebeck measurements lined above. RSoXS samples from doped PBTTT thin films on an Gold contact layers (~100 nm thick) for electronic conductivity and See- OTS-treated quartz substrate were prepared by first scribing small beck coefficient measurements were thermally evaporated (Angstrom square grids using a razor blade. A solution of 15% hydrogen fluoride Engineering Amod) onto either neat PBTTT thin films or solution- in deionized water was used to partially etch the oxide layer. The films doped PBTTT:F TCNQ (n = 2 or 4) thin films through a shadow mask. were then lifted off the substrate by dipping them in deionized water. Four-point probe conductivity contacts had a channel length of 0.2 mm The pieces of freestanding films were then lifted out of the water using and a channel width of 1 mm. Seebeck measurement contacts consisted silicon nitride windows. of 1-mm gold pads adjacent to 0.2-mm × 1-mm gold bars. The distance RSoXS experiments were performed on beamline 11.0.1.2 at the between the gold pads (temperature probes) and gold bars (voltage ALS located on the Lawrence Berkeley National Laboratory campus. probes) was 3, 4, and 5 mm apart. A detailed schematic is provided 2D RSoXS scattering images were collected in transmission mode in fig. S9. using a charge-coupled device camera (Princeton Instrument PI-MTE) Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 11 of 13 | SCIENCE ADVANCES RESEARCH ARTICLE 16. P. Pingel, D. Neher, Comprehensive picture of p-type doping of P3HT with the molecular cooled to −45°C in a high-vacuum chamber (add pressure). The sample- acceptor F TCNQ. Phys. Rev. B 87, 115209 (2013). to-detector distance was set to 175 mm. The 2D scattering image was 17. V. I. Arkhipov, P. Heremans, E. V. Emelianova, G. J. Adriaenssens, H. Bässler, Charge carrier reduced by azimuthal integration over all values of q (scattering mobility in doped semiconducting polymers. Appl. Phys. Lett. 82, 3245–3247 (2003). vector). After subtraction of the dark image, the Lorentz-corrected 18. V. I. Arkhipov, E. V. Emelianova, P. Heremans, H. Bässler, Analytic model of carrier mobility in doped disordered organic semiconductors. Phys. Rev. B 72, 235202 (2005). profile (I*q versus q) was obtained. The data reduction was per- 19. B. Lüssem, C.-M. Keum, D. Kasemann, B. Naab, Z. Bao, K. Leo, Doped organic transistors. formed in Wavemetrics Igor Pro using NIKA macro developed by Chem. Rev. 116, 13714–13751 (2016). J. Ilvasky at the Advanced Photon Source (69). 20. H. Sirringhaus, 25th anniversary article: Organic field-effect transistors: The path beyond amorphous silicon. Adv. Mater. 26, 1319–1335 (2014). 21. R. Noriega, J. Rivnay, K. Vandewal, F. P. V. Koch, N. Stingelin, P. Smith, M. F. Toney, A. Salleo, A general relationship between disorder, aggregation and charge transport in SUPPLEMENTARY MATERIALS conjugated polymers. Nat. Mater. 12, 1038–1044 (2013). Supplementary material for this article is available at http://advances.sciencemag.org/cgi/ 22. D. Venkateshvaran, M. Nikolka, A. Sadhanala, V. Lemaur, M. Zelazny, M. Kepa, content/full/3/6/e1700434/DC1 M. Hurhangee, A. J. Kronemeijer, V. Pecunia, I. Nasrallah, I. Romanov, K. Broch, fig. S1. AFM height and phase images of neat annealed PBTTT and F TCNQ vapor-doped films I. McCulloch, D. Emin, Y. Olivier, J. Cornil, D. Beljonne, H. Sirringhaus, Approaching at 5 and 10 min. disorder-free transport in high-mobility conjugated polymers. Nature 515, 384–388 fig. S2. Absorption spectra showing the NIR regime for doped PBTTT films and the thermal (2014). stability of F TCNQ-doped films. 23. J. B. Torrance, An overview of organic charge-transfer solids: Insulators, metals, and the fig. S3. Additional UV-vis-NIR spectra of F TCNQ vapor-doped films relative to a neat film. neutral-ionic transition. Mol. Cryst. Liq. Cryst. 126,55–67 (1985). fig. S4. Williamson-Hall plot for neat (black circle) and F TCNQ vapor-doped film. 24. K.-H. Yim, G. L. Whiting, C. E. Murphy, J. J. M. Halls, J. H. Burroughes, R. H. Friend, J.-S. Kim, fig. S5. In-plane scattering profiles of as-cast neat and doped films. Controlling electrical properties of conjugated polymers via a solution-based p-type fig. S6. Thin-film thickness profile of neat and vapor-doped PBTTT:F TCNQ film. doping. Adv. Mater. 20, 3319–3324 (2008). fig. S7. Representative 2D RSoXS images for neat PBTTT, F TCNQ vapor-doped, and F TCNQ 4 4 25. C. Wang, D. T. Duong, K. Vandewal, J. Rivnay, A. Salleo, Optical measurement of doping solution-doped thin films (all thermally annealed). efficiency in poly(3-hexylthiophene) solutions and thin films. Phys. Rev. B 91, 85205 (2015). fig. S8. Lorentz-corrected scattering profiles of neat PBTTT for different annealing 26. J. Li, G. Zhang, D. M. Holm, I. E. Jacobs, B. Yin, P. Stroeve, M. Mascal, A. J. Moulé, temperatures. Introducing solubility control for improved organic p-type dopants. Chem. Mater. 27, fig. S9. Schematic of the geometry of the contacts for electronic conductivity and Seebeck 5765–5774 (2015). measurements on thin films of doped polymers. 27. J. B. Torrance, The difference between metallic and insulating salts of table S1. X-ray reflection peaks of annealed PBTTT thin films from GIWAXS. tetracyanoquinodimethone (TCNQ): How to design an organic metal. Acc. Chem. Res. 12, table S2. X-ray reflection peaks of as-cast PBTTT thin films from GIWAXS. 79–86 (1979). table S3. Summary of OCLs for doped films. 28. F. Wudl, From organic metals to superconductors: Managing conduction electrons in organic solids. Acc. Chem. Res. 17, 227–232 (1984). 29. M. R. Bryce, L. C. Murphy, Organic metals. Nature 309, 119–126 (1984). 30. P. Pingel, R. Schwarzl, D. Neher, Effect of molecular p-doping on hole density and REFERENCES AND NOTES mobility in poly(3-hexylthiophene). Appl. Phys. Lett. 100, 143303 (2012). 1. B. Lüssem, M. Riede, K. Leo, Doping of organic semiconductors. Phys. Status Solidi A 210, 31. I. Salzmann, G. Heimel, M. Oehzelt, S. Winkler, N. Koch, Molecular electrical doping of 9–43 (2013). organic semiconductors: Fundamental mechanisms and emerging dopant design rules. 2. H. Ma, H.-L. Yip, F. Huang, A. K.-Y. Jen, Interface engineering for organic electronics. Acc. Chem. Res. 49, 370–378 (2016). Adv. Funct. Mater. 20, 1371–1388 (2010). 32. J. Gao, E. T. Niles, J. K. Grey, Aggregates promote efficient charge transfer doping of 3. M.-C. Jung, S. R. Raga, L. K. Ono, Y. Qi, Substantial improvement of perovskite solar cells poly(3-hexylthiophene). J. Phys. Chem. Lett. 4, 2953–2957 (2013). stability by pinhole-free hole transport layer with doping engineering. Sci. Rep. 5, 33. S. N. Patel, A. M. Glaudell, D. Kiefer, M. L. Chabinyc, Increasing the thermoelectric 9863 (2015). power factor of a semiconducting polymer by doping from the vapor phase. ACS Macro 4. G. Lu, J. Blakesley, S. Himmelberger, P. Pingel, J. Frisch, I. Lieberwirth, I. Salzmann, Lett. 5, 268–272 (2016). M. Oehzelt, R. Di Pietro, A. Salleo, N. Koch, D. Neher, Moderate doping leads to high 34. A. M. Glaudell, J. E. Cochran, S. N. Patel, M. L. Chabinyc, Impact of the doping method performance of semiconductor/insulator polymer blend transistors. Nat. Commun. 4,1588 on conductivity and thermopower in semiconducting polythiophenes. Adv. Energy Mater. (2013). 5, 1401072 (2015). 5. O. Bubnova, X. Crispin, Towards polymer-based organic thermoelectric generators. 35. D. T. Scholes, S. A. Hawks, P. Y. Yee, H. Wu, J. R. Lindemuth, S. H. Tolbert, B. J. Schwartz, Energy Environ. Sci. 5, 9345–9362 (2012). Overcoming film quality issues for conjugated polymers doped with F TCNQ by solution 6. S. N. Patel, M. L. Chabinyc, Anisotropies and thermoelectric properties of semiconducting sequential processing: Hall effect, structural, and optical measurements. J. Phys. Chem. Lett. polymers. J. Appl. Polym. Sci. 134, 44403 (2016). 6,4786–4793 (2015). 7. B. Russ, A. Glaudell, J. J. Urban, M. L. Chabinyc, R. A. Segalman, Organic thermoelectric 36. I. E. Jacobs, E. W. Aasen, J. L. Oliveira, T. N. Fonseca, J. D. Roehling, J. Li, G. Zhang, materials for energy harvesting and temperature control. Nat. Rev. Mater. 1, 16050 (2016). M. P. Augustine, M. Mascal, A. J. Moulé, Comparison of solution-mixed and sequentially 8. O. Owoyele, S. Ferguson, B. T. O’Connor, Performance analysis of a thermoelectric processed P3HT:F4TCNQ films: Effect of doping-induced aggregation on film cooler with a corrugated architecture. Appl. Energy 147, 184–191 (2015). morphology. J. Mater. Chem. C 4, 3454–3466 (2016). 9. Q. Wei, M. Mukaida, K. Kirihara, Y. Naitoh, T. Ishida, Polymer thermoelectric modules 37. K. Kang, S. Watanabe, K. Broch, A. Sepe, A. Brown, I. Nasrallah, M. Nikolka, Z. Fei, screen-printed on paper. RSC Adv. 4, 28802 (2014). M. Heeney, D. Matsumoto, K. Marumoto, H. Tanaka, S.-i. Kuroda, H. Sirringhaus, 2D 10. J.-H. Bahk, H. Fang, K. Yazawa, A. Shakouri, Flexible thermoelectric materials and device coherent charge transport in highly ordered conducting polymers doped by solid state optimization for wearable energy harvesting. J. Mater. Chem. C 3, 10362–10374 (2015). diffusion. Nat. Mater. 15, 896–902 (2016). 11. C. Wan, R. Tian, A. B. Azizi, Y. Huang, Q. Wei, R. Sasai, S. Wasusate, T. Ishida, K. Koumoto, 38. J. E. Cochran, M. J. N. Junk, A. M. Glaudell, P. L. Miller, J. S. Cowart, M. F. Toney, C. J. Hawker, Flexible thermoelectric foil for wearable energy harvesting. Nano Energy 30, 840–845 B. F. Chmelka, M. L. Chabinyc, Molecular interactions and ordering in electrically (2016). doped polymers: Blends of PBTTT and F TCNQ. Macromolecules 47,6836–6846 (2014). 12. H. Fang, B. C. Popere, E. M. Thomas, C.-K. Mai, W. B. Chang, G. C. Bazan, M. L. Chabinyc, 39. J. Li, C. W. Rochester, I. E. Jacobs, S. Friedrich, P. Stroeve, M. Riede, A. J. Moulé, R. A. Segalman, Large-scale integration of flexible materials into rolled and corrugated Measurement of small molecular dopant F4TCNQ and C F diffusion in organic bilayer 60 36 thermoelectric modules. J. Appl. Polym. Sci. 134, 44208 (2017). architectures. ACS Appl. Mater. Interfaces 7, 28420–28428 (2015). 13. K. Kirihara, Q. Wei, M. Mukaida, T. Ishida, Thermoelectric power generation using 40. I. McCulloch, M. Heeney, C. Bailey, K. Genevicius, I. MacDonald, M. Shkunov, D. Sparrowe, nonwoven fabric module impregnated with conducting polymer PEDOT:PSS. Synth. Met. S. Tierney, R. Wagner, W. Zhang, M. L. Chabinyc, R. J. Kline, M. D. McGehee, M. F. Toney, 225,41–48 (2017). Liquid-crystalline semiconducting polymers with high charge-carrier mobility. Nat. Mater. 5, 14. G. J. Snyder, E. S. Toberer, Complex thermoelectric materials. Nat. Mater. 7, 105–114 328–333 (2006). (2008). 15. A. Salleo, R. J. Kline, D. M. DeLongchamp, M. L. Chabinyc, Microstructural characterization 41. M. L. Chabinyc, M. F. Toney, R. J. Kline, I. McCulloch, M. Heeney, X-ray scattering study of and charge transport in thin films of conjugated polymers. Adv. Mater. 22, 3812–3838 thin films of poly(2,5-bis(3-alkylthiophen-2-yl)thieno[3,2-b]thiophene). J. Am. Chem. Soc. (2010). 129,3226–3237 (2007). Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 12 of 13 | SCIENCE ADVANCES RESEARCH ARTICLE 42. B. A. Collins, J. E. Cochran, H. Yan, E. Gann, C. Hub, R. Fink, C. Wang, T. Schuettfort, 62. M. N. Gueye, A. Carella, N. Massonnet, E. Yvenou, S. Brenet, J. Faure-Vincent, S. Pouget, C. R. McNeill, M. L. Chabinyc, H. Ade, Polarized X-ray scattering reveals non-crystalline F. Rieutord, H. Okuno, A. Benayad, R. Demadrille, J.-P. Simonato, Structure and dopant orientational ordering in organic films. Nat. Mater. 11, 536–543 (2012). engineering in PEDOT thin films: Practical tools for a dramatic conductivity 43. P. Brocorens, A. Van Vooren, M. L. Chabinyc, M. F. Toney, M. Shkunov, M. Heeney, enhancement. Chem. Mater. 28, 3462–3468 (2016). I. McCulloch, J. Cornil, R. Lazzaroni, Solid-state supramolecular organization of 63. C. K. Chiang, C. R. Fincher Jr., Y. W. Park, A. J. Heeger, H. Shirakawa, E. J. Louis, S. C. Gau, polythiophene chains containing thienothiophene units. Adv. Mater. 21, 1193–1198 A. G. MacDiarmid, Electrical conductivity in doped polyacetylene. Phys. Rev. Lett. 39, (2009). 1098–1101 (1977). 44. T. Schuettfort, B. Watts, L. Thomsen, M. Lee, H. Sirringhaus, C. R. McNeill, Microstructure 64. X. Wang, V. Ho, R. A. Segalman, D. G. Cahill, Thermal conductivity of high-modulus of polycrystalline PBTTT films: Domain mapping and structure formation. ACS Nano 6, polymer fibers. Macromolecules 46, 4937–4943 (2013). 1849–1864 (2012). 65. A. Weathers, Z. U. Khan, R. Brooke, D. Evans, M. T. Pettes, J. W. Andreasen, X. Crispin, L. Shi, 45. D. M. DeLongchamp, R. J. Kline, Y. Jung, D. S. Germack, E. K. Lin, A. J. Moad, L. J. Richter, Significant electronic thermal transport in the conducting polymer poly(3,4- M. F. Toney, M. Heeney, I. McCulloch, Controlling the orientation of terraced nanoscale ethylenedioxythiophene). Adv. Mater. 27, 2101–2106 (2015). “ribbons” of a poly(thiophene) semiconductor. ACS Nano 3,780–787 (2009). 66. H. Ushirokita, H. Tada, In-plane thermal conductivity measurement of conjugated 46. K. Kanai, K. Akaike, K. Koyasu, K. Sakai, T. Nishi, Y. Kamizuru, T. Nishi, Y. Ouchi, K. Seki, polymer films by membrane-based AC calorimetry. Chem. Lett. 45,735–737 Determination of electron affinity of electron accepting molecules. Appl. Phys. A 95, (2016). 309–313 (2009). 67. J. Liu, X. Wang, D. Li, N. E. Coates, R. A. Segalman, D. G. Cahill, Thermal conductivity and 47. A. Mityashin, Y. Olivier, T. Van Regemorter, C. Rolin, S. Verlaak, N. G. Martinelli, D. Beljonne, elastic constants of PEDOT:PSS with high electrical conductivity. Macromolecules 48, J. Cornil, J. Genoe, P. Heremans, Unraveling the mechanism of molecular doping in 585–591 (2015). organic semiconductors. Adv. Mater. 24, 1535–1539 (2012). 68. I. Mcculloch, M. Heeney, M. L. Chabinyc, D. Delongchamp, R. J. Kline, M. Cölle, W. Duffy, 48. N. C. Miller, E. Cho, M. J. N. Junk, R. Gysel, C. Risko, D. Kim, S. Sweetnam, C. E. Miller, D. Fischer, D. Gundlach, B. Hamadani, R. Hamilton, L. Richter, A. Salleo, M. Shkunov, L. J. Richter, R. J. Kline, M. Heeney, I. McCulloch, A. Amassian, D. Acevedo-Feliz, C. Knox, D. Sparrowe, S. Tierney, W. Zhang, Semiconducting thienothiophene copolymers: Design, M. R. Hansen, D. Dudenko, B. F. Chmelka, M. F. Toney, J.-L. Brédas, M. D. McGehee, Use synthesis, morphology, and performance in thin-film organic transistors. Adv. Mater. of X-ray diffraction, molecular simulations, and spectroscopy to determine the molecular 21, 1091–1109 (2009). packing in a polymer-fullerene bimolecular crystal. Adv. Mater. 24, 6071–6079 (2012). 69. J. Ilavsky, Nika: Software for two-dimensional data reduction. J. Appl. Cryst. 45, 324–328 49. C. Wang, D. H. Lee, A. Hexemer, M. I. Kim, W. Zhao, H. Hasegawa, H. Ade, T. P. Russell, (2012). Defining the nanostructured morphology of triblock copolymers using resonant soft X-ray scattering. Nano Lett. 11,3906–3911 (2011). Acknowledgments: Use of the SSRL, SLAC National Accelerator Laboratory, is supported by 50. J. Stöhr, NEXAFS Spectroscopy, vol. 25 of Springer Series in Surface Sciences (Springer, 1992). the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under 51. D. S. Pearson, P. A. Pincus, G. W. Heffner, S. J. Dahman, Effect of molecular weight contract no. DE-AC02-76SF00515. This research used resources of the ALS, which is a U.S. and orientation on the conductivity of conjugated polymers. Macromolecules 26, Department of Energy Office of Science User Facility under contract no. DE-AC02-05CH11231. 1570–1575 (1993). The MRL Shared Experimental Facilities are supported by the MRSEC Program of the NSF 52. X. Xue, G. Chandler, X. Zhang, R. J. Kline, Z. Fei, M. Heeney, P. J. Diemer, O. D. Jurchescu, under award no. DMR 1121053, a member of the NSF-funded Materials Research Facilities B. T. O’Connor, Oriented liquid crystalline polymer semiconductor films with large Network. Funding: The authors acknowledge the support of the Air Force Office of Scientific ordered domains. ACS Appl. Mater. Interfaces 7, 26726–26734 (2015). Research through the Multidisciplinary University Research Initiative on Controlling Thermal 53. J. Li, C. W. Rochester, I. E. Jacobs, E. W. Aasen, S. Friedrich, P. Stroeve, A. J. Moulé, The and Electrical Transport in Organic and Hybrid Materials (AFOSR FA9550-12-1-0002). A.M.G. effect of thermal annealing on dopant site choice in conjugated polymers. Org. Electron. received partial support from the ConvEne IGERT Program of the NSF under NSF-DGE 0801627. 33,23–31 (2016). E.L. acknowledges support from the NSF Graduate Research Fellowship (DGE-1144085). 54. J. Fuzell, I. E. Jacobs, S. Ackling, T. F. Harrelson, D. M. Huang, D. Larsen, A. J. Moulé, Optical Author contributions: S.N.P. designed and performed all experiments related to thermoelectric dedoping mechanism for P3HT:F4TCNQ mixtures. J. Phys. Chem. Lett. 7, 4297–4303 (2016). measurements, GIWAXS, and RSoXS and wrote the manuscript with input from M.L.C. A.M.G. 55. I. E. Jacobs, F. Wang, N. Hafezi, C. Medina-Plaza, T. F. Harrelson, J. Li, M. P. Augustine, assisted in thermoelectric measurements and assisted in the analysis of the data. K.A.P. M. Mascal, A. J. Moulé, Quantitative dedoping of conductive polymers. Chem. Mater. 29, performed UV-vis-NIR experiments. E.M.T. performed AFM experiments. K.A.O. performed 832–841 (2017). high-resolution x-ray scattering experiments. E.L. assisted with RSoXS experiments. M.L.C. 56. S. D. Kang, G. J. Snyder, Charge-transport model for conducting polymers. Nat. Mater. 16, supervised all aspects of the project, designed experiments, and was involved in writing of the 252–257 (2017). manuscript. All coauthors assisted with writing and editing of the manuscript. Figure credit: 57. D. Wang, W. Shi, J. Chen, J. Xi, Z. Shuai, Modeling thermoelectric transport in organic S.N.P. for Figs. 1, 3, and 4; S.N.P., K.A.P., E.M.T., and A.M.G. for Fig. 2; S.N.P. and E.L. for Figs. 5 and materials. Phys. Chem. Chem. Phys. 14, 16505–16520 (2012). 6; S.N.P. and A.M.G. for Fig. 7. Competing interests: The authors declare that they have no 58. G. Zuo, H. Abdalla, M. Kemerink, Impact of doping on the density of states and the competing interests. Data and materials availability: All data needed to evaluate the mobility in organic semiconductors. Phys. Rev. B 93, 235203 (2016). conclusions in the paper are present in the paper and/or the Supplementary Materials. 59. Q. Wei, M. Mukaida, K. Kirihara, T. Ishida, Experimental studies on the anisotropic Additional data related to this paper may be requested from the authors. thermoelectric properties of conducting polymer films. ACS Macro Lett. 3, 948–952 (2014). Submitted 9 February 2017 60. S. Qu, Q. Yao, L. Wang, Z. Chen, K. Xu, H. Zeng, W. Shi, T. Zhang, C. Uher, L. Chen, Highly Accepted 28 April 2017 anisotropic P3HT films with enhanced thermoelectric performance via organic small Published 16 June 2017 molecule epitaxy. NPG Asia Mater. 8, e292 (2016). 10.1126/sciadv.1700434 61. O. Bubnova, Z. U. Khan, H. Wang, S. Braun, D. R. Evans, M. Fabretto, P. Hojati-Talemi, D. Dagnelund, J.-B. Arlin, Y. H. Geerts, S. Desbief, D. W. Breiby, J. W. Andreasen, Citation: S. N. Patel, A. M. Glaudell, K. A. Peterson, E. M. Thomas, K. A. O’Hara, E. Lim, R. Lazzaroni, W. M. Chen, I. Zozoulenko, M. Fahlman, P. J. Murphy, M. Berggren, X. Crispin, M. L. Chabinyc, Morphology controls the thermoelectric power factor of a doped Semi-metallic polymers. Nat. Mater. 13, 190–194 (2013). semiconducting polymer. Sci. Adv. 3, e1700434 (2017). Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 13 of 13 http://www.deepdyve.com/assets/images/DeepDyve-Logo-lg.png Science Advances Pubmed Central

Morphology controls the thermoelectric power factor of a doped semiconducting polymer

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SCIENCE ADVANCES RESEARCH ARTICLE MATERIALS SCIENCE Copyright © 2017 The Authors, some rights reserved; Morphology controls the thermoelectric power factor of exclusive licensee American Association a doped semiconducting polymer for the Advancement of Science. Distributed 1 1,2 2 2 Shrayesh N. Patel, * Anne M. Glaudell, Kelly A. Peterson, Elayne M. Thomas, under a Creative 2 2 † Kathryn A. O’Hara, Eunhee Lim, Michael L. Chabinyc Commons Attribution NonCommercial The electrical performance of doped semiconducting polymers is strongly governed by processing methods and License 4.0 (CC BY-NC). underlying thin-film microstructure. We report on the influence of different doping methods (solution versus vapor) on the thermoelectric power factor (PF) of PBTTT molecularly p-doped with F TCNQ (n = 2 or 4). The vapor-doped films have more than two orders of magnitude higher electronic conductivity (s) relative to solution-doped films. On the basis of resonant soft x-ray scattering, vapor-doped samples are shown to have a large orientational correlation length (OCL) (that is, length scale of aligned backbones) that correlates to a high apparent charge carrier mobility (m). The Seebeck coefficient (a) is largely independent of OCL. This reveals that, unlike s, leveraging strategies to improve m have a smaller impact on a. Our best-performing sample with the largest OCL, vapor-doped PBTTT: −1 −2 F TCNQ thin film, has a s of 670 S/cm and an a of 42 mV/K, which translates to a large PF of 120 mWm K . In addition, despite the unfavorable offset for charge transfer, doping by F TCNQ also leads to a large PF of −1 −2 70 mWm K , which reveals the potential utility of weak molecular dopants. Overall, our work introduces im- portant general processing guidelines for the continued development of doped semiconducting polymers for thermoelectrics. INTRODUCTION poly(2,5-bis(3-tetradecylthiophen-2-yl)thieno[3,2-b]thiophene) Controlling the electrical doping of organic semiconductors is critical (PBTTT). We focus on p-doping with organic acceptors [2,3,5,6- to the performance of organic electronic devices (1). Doped semi- tetrafluoro-7,7,8,8-tetracyanoquinodimethane (F TCNQ) and 2,5- conducting polymers can serve as conductive interlayers for organic difluoro-7,7,8,8-tetracyanoquinodimethane (F TCNQ)] introduced light-emitting diodes (OLEDs) (2) and solar cells (2, 3) and can im- either in solution or from the vapor phase (Fig. 1). The results of prove the performance of organic thin-film transistors (OTFTs) (4). our experiments demonstrate how different processing and doping One emerging application of doped semiconducting polymers in- methods affect the thermoelectric PF. In particular, we find that align- volves organic thermoelectrics—materials that interconvert heat and ment of ordered domains is the critical factor leading to higher s with- electricity (5–7). The solution processability of semiconducting poly- out lowering a, thereby leading to enhancements in the PF. Using −1 −2 mers provides the opportunity to use roll-to-roll processing and print- these methods, we have found a PF of 120 mWm K for PBTTT: ing technologies for new classes of thermoelectric modules where the F TCNQ, which is among the highest reported values for semi- legs are thin films in rolled or corrugated designs (8–13). To realize conducting polymers (7). the potential of semiconducting polymers for thermoelectrics, how- The s value of semiconducting polymers is related to the product ever, the relationship between processing and the resulting thermo- of the carrier concentration (n) and carrier mobility (m). However, electric properties must be better understood. because of electronic disorder, the apparent m of a material will de- All of the physical properties of a material that define its thermo- pend on n because of the occupancy of electronic states with varying electric performance depend on carrier density (n), including electri- mobility (16–19). Through advances in molecular design and pro- cal conductivity (s), Seebeck coefficient (or thermopower) (a), and cessing, solution-processable semiconducting polymers, such as PBTTT, thermal conductivity (k)(14). The thermal-to-electrical energy con- have high charge carrier mobilities (m >1cm /V s) in field-effect version efficiency is related to the dimensionless figure of merit, ZT = transistors (20). These studies have revealed that the degree of 2 2 a sT/k, where T is the temperature in Kelvin and a s is the power electronic and structural disorder strongly influences m (21, 22). In factor (PF). Optimizing ZT is quite challenging because as n increases, s field-effect measurements, conduction occurs very close (within and k increase while a decreases (14). Organic semiconductors fre- ~1 nm) to the polymer-dielectric interface. The microstructure is quently have imperfect ordering in thin films, leading to an electronic generally described as (para)crystalline p-stacked domains intercon- structure that depends strongly on their morphology (15). Because nected by tie chains (21). Electrically doped films may require high processing methods widely vary in many studies of thermoelectric concentrations of dopant in the bulk (for example, >1 dopant per performance, it is difficult to form clear connections between mor- 10 monomers), which lead to strong perturbations of the morphology phology and thermoelectric performance (7). and structure relative to pristine films. Whether processing methods Here, we elucidate the connection between thin-film micro- that lead to high field-effect m also lead to high-bulk s has not been structure and thermoelectric transport properties (s and a) of p-doped well studied. The ability to tune the electronic structure of small-molecule or- ganic acceptors (23) has resulted in versatile p-type dopants for Materials Research Laboratory, University of California, Santa Barbara, Santa Barbara, CA 93106, USA. Materials Department, University of California, Santa Barbara, Santa semiconducting polymers (16, 24–26). Electron acceptors have been Barbara, CA 93106, USA. traditionally co-deposited with molecular organic donors to generate *Present address: Institute for Molecular Engineering, University of Chicago, 5640 organic charge transfer salts and metals (27–29). Mixing an organic South Ellis Avenue, Chicago, IL 60637, USA. †Corresponding author. Email: mchabinyc@engineering.ucsb.edu acceptor into a polymer leads to an integer charge transfer if the Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 1of13 | SCIENCE ADVANCES RESEARCH ARTICLE higher s relative to films cast from a polymer/dopant solution. For example, vapor doping a predeposited PBTTT with F TCNQ results in a s of ~250 S/cm, which is nearly two orders of magnitude higher relative to a solution-doped film with a similar concentration of F TCNQ (37). The general explanation for the enhancement in s is related to an overall better macroscopic film quality establishing the underlying microstructure for more efficient charge transport (35–37). Although these explanations describe enhancement in s,itisunclear how more efficient charge transport influences the overall thermo- electric properties of semiconducting polymers. Here, we examine how morphology affects the bulk s and a of PBTTT. PBTTT is a solution-processable polymer, where the field- effect mobility, crystal structure, and morphology have been well characterized for neat films (40–45), providing a strong foundation to study how molecular doping influences the morphology and charge transport. In addition, PBTTT has an accessible liquid crys- talline transition temperature above ~140°C, which permits thermal processing to enhance local and long-range order (40–42). Using PBTTT and other thiophene-based polymers, we have previously discovered an empirical connection between electrical conductivity and thermopower across a range of doping methods (33, 34). This broad correlation has been modeled as a result of the electronic density of states (DOS) and energy-dependent mobility, but the connection with morphology has not been made clear. Here, we use PBTTT as a model system to demonstrate how the correlation of alignment in ordered domains at the nanoscale dominates the resulting s at high n. These results suggest a pathway to increase the thermoelectric PF of semiconducting polymers. Fig. 1. Chemical structure and doping process. (A) Chemical structure of PBTTT and F TCNQ (n = 2 or 4) and the corresponding IE or EA. (B) Solution and vapor RESULTS AND DISCUSSION doping routes used to achieve doped films. Processing doped films of PBTTT We prepared highly conductive thin films of PBTTT using different offset between HOMO (highest occupied molecular orbital) [ioniza- processing methods, with the dopant added in solution or infiltrated tion energy (IE)] of the polymer and LUMO (lowest unoccupied mo- from the vapor phase (Fig. 1). A detailed procedure can be found in lecular orbital) [electron affinity (EA)] of the acceptor is sufficient to Materials and Methods, and we outline the critical differences here. provide a thermodynamic driving force for electron transfer (Fig. 1) We specifically focused on the limit of high doping to determine the (16, 30–32). connection between morphology and thermoelectric transport prop- How a dopant is incorporated into a semiconducting polymer is erties. For solution doping, 10 wt % of F TCNQ relative to PBTTT critical in dictating the resulting charge transport properties (33–37). [molar ratio (MR) of ~1 dopant to 4 monomers] was added to a so- Achieving high s (>10 S/cm) requires relatively high charge carrier lution of PBTTT. We have previously found that this composition is 19 3 concentrations (>10 /cm ) due to the observed superlinear increase near the maximum possible to readily form continuous thin films in conductivity in many materials (16, 30). If a dopant is added to the during spin casting (38). The solution was spin-coated to obtain a casting solution, this concentration requires as much as 10 weight % doped thin film in a N environment and annealed at 150°C for (wt %) relative to the monomer in the solution, which can be difficult 10 min to remove the solvent. These conditions were used to minimize because of solubility limits of neutral organic acceptors (26). Although weight loss of F TCNQ from the film (38) while also being above the chain aggregation appears to aid in efficient charge transfer in some liquid crystalline (LC) transition temperature of the neat polymer. The cases (32), highly charged polymers can gel or precipitate from solution same processing conditions were used with solution-doped samples, (38). Consequently, one must take great care in determining optimal where the dopant was F TCNQ. The typical thicknesses of solution- casting concentrations and temperatures that lead to macroscopically doped films were 40 to 50 nm. To form films doped by infiltration of homogenous films (35, 38). Alternatively, first casting a neat film from F TCNQ and F TCNQ from the vapor phase, we exposed spin-coated 4 2 solution and then subsequently doping has emerged as a versatile route films of PBTTT (21 ± 4 nm) prepared using different thermal treat- to yield macroscopic homogenous films with high s (33, 35–37). For ments to the vapor of each compound. In a N -filled glove box, we example, depositing a thin layer of organic acceptor from the vapor deposited the dopants by placing the samples underneath the lid of a phase can lead to diffusion of the dopant into the organic semi- sealed jar containing a few milligrams of dopant. The bottom of the conductor (37, 39) and solid-state charge transfer to generate highly jar was heated to ~210°C, which led to a rise in temperature of the conductive films (37). Recently, it has been shown that doping spin- substrate, which was located underneath the lid for the jar, to 75° to coated neat thiophene-based polymers either from the vapor phase 85°C. These temperatures are below the LC transition of PBTTT and (33, 37) or through sequential solution casting (35, 36) can lead to are known not to cause substantial changes in the structural order Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 2of13 | SCIENCE ADVANCES RESEARCH ARTICLE and charge mobility in OTFTs (41). The samples were exposed to the vapor for 10 min, which was sufficient to reach concentrations of Table 1. Summary of electronic conductivity (s), Seebeck coefficient F TCNQ in PBTTT films comparable to the solution-doped films. (a), and PF of doped PBTTT films. For the sample on an OTS-treated −1 −2 substrate, s = 670 ± 4S/cm, a =42 ± 6 mV/K, and PF = 120 ± 30 mWm K . All other samples reported in this table are on untreated quartz substrates. Changes in processing increase electrical conductivity −1 −2 Efficient charge transfer occurs between PBTTT and F TCNQ using Dopant Condition s (S/cm) a (mV/K) PF (mWm K ) both vapor- and solution-based doping. Ultraviolet-visible near- F TCNQ Solution— 2.08 ± 0.01 45 ± 4 0.42 ± 0.09 infrared (UV-vis-NIR) spectroscopy shows that neat poly(2,5-bis(3- as-cast tetradecylthiophen-2-yl)thieno[3,2-b]thiophene) (PBTTT-C ) thin Solution— 3.51 ± 0.05 60 ± 9 1.3 ± 0.4 film has a main absorption peak at ∼2.2 eV and a shoulder at ∼2.1 eV annealed that is bleached upon doping. New absorption peaks appear at 1.41 Vapor— 114.1 ± 0.5 32 ± 4 12 ± 3 and 1.60 eV upon introduction of F TCNQ and are assigned to its as-cast anion radical (16, 25, 37). The spectral features for the F TCNQ rad- ical anion absorption are similar to those in poly(3-hexylthiophene) Vapor— 220.00 ± 0.02 39 ± 5 32 ± 9 annealed (P3HT):F TCNQ films and have comparable absorptivity for heavily doped films (35). A subband gap transition for positive polarons of F TCNQ Solution— 0.41 ± 0.02 111.7 ± 0.1 0.52 ± 0.03 PBTTT is observed at ∼0.5 eV (fig. S2), but the precise position of as-cast polaronic features between 1 and 2 eV is difficult to assign because −3 Solution— 2×10 ± 755 ± 100 0.11 ± 0.03 −4 of the strong absorption peaks of the F TCNQ anion radical. We annealed 2× 10 observe no significant differences between UV-vis-NIR spectra be- Vapor— 13.7 ± 0.2 130 ± 20 23 ± 6 tween vapor-doped as-cast and annealed films, indicating that the as-cast annealing step does not change the concentration of F TCNQ in Vapor— 36 ± 3 140 ± 20 70 ± 20 the film (fig. S3). A comparison of the main absorption of PBTTT annealed (∼2.2 eV) for the vapor-doped film relative to the solution-doped film reveals more bleaching in the former (a factor of 0.78 lower peak area) and also a slightly higher absorbance of F TCNQ radical anion. This lower peak area translates to a slightly higher MR of 0.3 relative to the MR of 0.25 in the solution-doped film (37). but only the vapor-doped film has a small red shift. In contrast to Strikingly large differences in s are found between vapor- and F TCNQ, because of the relatively high vapor pressure of F TCNQ, 4 2 solution-doped samples despite this small difference in MR of the we observe significant dedoping of solution-doped PBTTT:F TCNQ dopant. Electrical conductivity measurements (Table 1) indicate films when thermally annealing at 150°C in N environment (fig. 65 times higher s for the vapor-doped PBTTT:F TCNQ annealed S2). The MR in the casting solution is ~0.28 F TCNQ per monomer 4 2 film (s = 220 ± 0.02 S/cm) relative to the solution-doped PBTTT: of PBTTT, similar to the solution-doped samples with F TCNQ. F TCNQ annealed film (s of 3.51 ± 0.05 S/cm). We observe a similar Doping from vapor allows for a comparison of infiltration into as- trend in as-cast films, where s = 114.1 ± 0.5 S/cm for the vapor- cast and annealed films. When vapor-doping an annealed PBTTT doped film and s = 2.08 ± 0.01 S/cm for the solution-doped film. film with F TCNQ, the primary absorption peak is 30% higher rela- Knowing that both doping methods yield comparable carrier con- tive to the as-cast doped films. This difference is, in part, from the fact centrations, the large difference in s must be related to the apparent that annealing a neat film results in an increase in primary absorption m, which is calculated to be ~2.5 and ~0.040 cm /V s for the vapor- peak by about 20% (comparison between Fig. 2, A and B). Despite dopedannealedfilmand the solution-dopedannealedfilm, respec- the small difference in the primary absorption peak, F TCNQ radical tively (assuming F TCNQ is fully ionized and all charges are free anion absorption is comparable between vapor-doped films but carriers). The higher apparent m with vapor-doped films is consistent slightly less than the solution-doped as-cast film. We attribute the with Hall effect mobility measurements on vapor-doped PBTTT: higher absorption of the F TCNQ radical anion in solution-doped films relative to vapor-doped films to the elevated temperature of F TCNQ, which was revealed to be ~2 cm /V s (37). Two possible factors that can contribute to the difference in apparent m are differ- the sample during vapor deposition. In addition, the absorption curve ences in local energetic disorder or the long-range morphology. is quantitatively similar in the NIR regime (fig. S2) for the F TCNQ- To determine whether the enhancement in s was unique to doped films. F TCNQ, we also examined samples doped with F TCNQ. The Comparison of the conductivity measurements reveals that the 4 2 intermediate fluorination level of F TCNQ resultsinanEAof vapor-doped PBTTT:F TCNQ films yield a higher s than solution- 2 2 ∼4.59 eV (46), which is expected to result in an unfavorable offset doped films do. The annealed film has a s of 36 ± 3 S/cm, and the for charge transfer with PBTTT (IE, ∼5.10 eV) in isolated materials as-cast film has a s of 13.7 ± 0.2 S/cm. On the other hand, the solution- compared to F TCNQ (EA, ∼5.24 eV). Note that charge transfer is doped as-cast film has the lowest s of 0.41 ± 0.02 S/cm despite the dictated not only by the offset between EA and IE but also by the higher concentration of F TCNQ radical anion. The high vapor pres- electrostatic interaction in solids (47). Although the EA of F TCNQ sure of F TCNQ does not allow direct comparison of solution and 2 2 is lower than the IE of PBTTT-C , integer charge transfer is indicated vapor-doped annealed films; it leads to significant dedoping and thus −3 by the presence of F TCNQ radical anion peaks at 1.43 and 1.62 eV several orders of magnitude lower s of 2 × 10 S/cm. Nevertheless, the (Fig. 2). Similar observations have been made for structurally similar fact that the vapor F TCNQ-doped films yield higher s further reiter- weak dopants with P3HT (26). For both doping methods, PBTTT: ates that an underlying microstructural feature is at play, which leads F TCNQ as-cast films yield a large decrease in the primary absorption, to a higher apparent m. Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 3of13 | SCIENCE ADVANCES RESEARCH ARTICLE Fig. 2. UV-vis spectra of neat and doped PBTTT thin films. UV-vis spectra of (A) neat PBTTT and PBTTT:F TCNQ thin films and (B) PBTTT:F TCNQ thin films. Solution- 4 2 doped films are at a dopant concentration of 10 wt %, and vapor-doped films are for dopant exposure of 10 min (all spectra normalized by thickness). Neat PBTTT films were annealed at 180°C, whereas solution-doped films were annealed at 150°C. Comparison of absorption spectra of annealed neat PBTTT and as-cast neat PBTTT can be found in the Supplementary Materials. Processing does not change crystalline structure in Thein-planescattering profiles alongthe q axis of the 2D xy doped films GIWAXS images (Fig. 4B and fig. S5) reveal that all highly doped To determine whether the difference in s between vapor- and solution- films lead to an increase in the scattering vector for the (110) reflec- doped films correlates to local structural order, we performed grazing tion (tables S1 and S2). This increase indicates a compression of the incidence wide-angle x-ray scattering (GIWAXS) (Fig. 3). PBTTT characteristic p-p–stacking distance. The value of d decreases forms semicrystalline films with strong texturing. For neat PBTTT, from 0.367 nm for the neat film to 0.355 nm for the F TCNQ the out-of-plane scattering features (along the q axis) correspond to solution-doped film and 0.353 nm for the vapor-doped film. The the lamella-stacked side chains (h00). The in-plane scattering features F TCNQ-doped films also showed compression upon doping, where −1 (along the q axis) correspond to the (-11-3) reflection at q = 14.2 nm d = 0.353 nm for an as-cast solution-doped film, d = 0.357 nm xy xy 110 110 −1 related to the chain axis and the (110) reflection at q =17.1 nm for a vapor-doped as-cast film, and d = 0.361 nm for a vapor- xy 110 associated to the p-stacking direction (43). The scattering pattern does doped annealed film. When thermal annealing a F TCNQ solution- not qualitatively change upon doping; that is, the features are shifted doped film at 150°C, we observe rapid dedoping, where d essentially and broadened upon doping, but no new strong scattering features returns to the value of an annealed neat film. The vapor F TCNQ- emerge. A significant difference between doped films and the neat film doped annealed film yields the smallest decrease in d ,which is −1 is the significant blurring of the off-axis features at q = 14.1 nm , consistent with the low concentration of F TCNQ radical anion xy 2 suggesting disorder along the chain axis, as expected from dopants (Fig. 2). We have previously shown that TCNQ, which does not re- with disordered molecular orientation. sult in an integer charge transfer with PBTTT due to the large offset The out-of-plane scattering features of the two-dimensional (2D) between EA and IE, causes no measurable changes to the d value GIWAXS images show changes upon doping in the alkyl side-chain relative to the neat film (38). This result suggested that the presence stacking direction (h00) (Fig. 4A). For all doped films, we observe a of ionized polymers and dopants is partly responsible for the de- small increase relative to the neat d = 2.12 nm (summarized in ta- crease in the characteristic d spacing. Overall, the changes to in- 100 110 bles S1 and S2). Because the true out-of-plane scattering is in the in- plane scattering features are quite similar with both doping methods. accessible region in the grazing incidence geometry, we also obtained Previously, nuclear magnetic resonance (NMR) experiments on high-resolution specular x-ray scattering on the vapor-doped PBTTT: bulk samples of PBTTT:F TCNQ cast from solution were carried out F TCNQ annealed films. The specular scattering results indicate that to determine the location of F TCNQ relative to the backbone of 4 4 13 1 d = 2.37 nm, which is about a 0.25-nm increase relative to the neat PBTTT (38). 2D C{ H} heteronuclear (HETCOR) NMR data dem- film (d = 2.12 nm). In addition, the change in peak width in the (h00) onstrated a very close contact between carbon atoms on the cyano direction reveals the degree of induced disorder along the side-chain group of F TCNQ and aromatic protons on the PBTTT backbone. stacking direction. As shown in fig. S4, the introduction of the dopant This proximity and the calculated geometry for model structures of from the vapor phase introduces additional disorder relative to the neat charge transfer between polymers and F TCNQ from density func- film, which is consistent with our previous work with PBTTT:F TCNQ tional theory suggested that the F TCNQ was intercalated between 4 4 solution-doped films (38). Similar changes are observed for films doped the polymer backbones (38). Full structural modeling to compare the with F TCNQ. The alkyl stacking spacing d increases by ∼0.25 nm in NMR and x-ray data was not carried out in the initial study. In con- 2 100 the solution-doped PBTTT:F TCNQ as-cast film, whereas d in- trast, a recent report on vapor-doped PBTTT:F TCNQ films proposed 2 100 4 creases by ∼0.10 nm in the vapor-doped as-cast film and by only that F TCNQ resides in the side-chain region of PBTTT films, with ∼0.04 nm in the vapor-doped annealed film. This trend indicates the the rationale that the p-stacking spacing, separation between chains, varying extent of dopant incorporated within the aliphatic side chains, was unperturbed in their x-ray scattering data (37). As discussed above, which is consistent with the UV-vis spectra showing a slightly smaller we observe a compression in the p-stacking spacing relative to the neat concentration of the F TCNQ radical anion in vapor-doped films. polymer, as observed in heavily doped PBTTT films with a variety of Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 4of13 | SCIENCE ADVANCES RESEARCH ARTICLE Fig. 3. GIWAXS for PBTTT films as a function of processing. 2D GIWAXS images for (A) annealed neat PBTTT, (B)F TCNQ vapor-doped annealed film, (C)annealedF TCNQ 4 4 solution-doped film, (D)F TCNQ vapor-doped annealed film, and (E) as-cast F TCNQ solution-doped film. The (100) reflection in GIWAXS image in (E) was blocked off with lead 2 2 tape to allow longer exposure time without saturating the detector. Images are obtained at Stanford Synchrotron Radiation Lightsource (SSRL) beamline 11-3. dopants (33, 38) and is reversible upon thermal removal of the dopant. the cyano groups of F TCNQ; such a change in tilt occurs in cocrystals Independent of these differences, one can infer that F TCNQ resides of PBTTT and [6,6]-phenyl C61 butyric acid methyl ester (PCBM) (48). in the side chains for vapor-doped samples by a simple considera- tion. On the basis of the observed out-of-plane scattering (h00), we Correlation of domain orientation controls the would expect <11% thickness change upon doping crystalline electrical conductivity domains. However, for the structural model where the molecule is Small perturbations in the local structure observed by GIWAXS are intercalated between in-plane p-stacked chains, we would expect a sig- similar for both solution doping and vapor doping methods and do nificant expansion of crystallites to compensate for the volume occu- not explain the differences in s. Our conclusion is in contrast to the pied by the dopant molecule, leading to a large in-plane expansion or explanation of Kang et al.(37), where they attributed the higher s only significant increase in thickness to maintain the lateral size. Starting to minimal perturbation to local order. However, it is important to with a neat film with thickness of 25 ± 3 nm, the vapor doping process remember that GIWAXS only provides structural information for leads to a thickness of 27 ± 3 nm, which is comparable to the expected short-range crystalline domains (5 to 20 nm). For larger length scales, expansion of the crystallites (fig. S6). Because the film thickness does we require a small-angle scattering method. Here, we use polarized not markedly increase, F TCNQ resides within the aliphatic side resonant soft x-ray scattering (RSoXS) that can reveal the length scale chains (assuming a high degree of crystallinity, which is known for of molecular orientation in both the crystalline and amorphous do- PBTTT). At this point, we cannot reconcile these data with the previous mains over the range from ~10 to 1000 nm (42). NMR data without detailed modeling of the spectra, which is beyond RSoXS leverages soft x-ray absorption to increase the scattering the scope of the work here. However, it is possible that the structure of length and to control the scattering contrast (42, 49). In p-conjugated the doped films has tilted PBTTT chains that allow a closer contact with molecules, the transition dipole moment (TDM) for the excitation Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 5of13 | SCIENCE ADVANCES RESEARCH ARTICLE entational alignment of polymer backbones is the key factor influenc- ing the performance of OTFTs (42). We prepared samples to mimic, or replicate, films used for mea- surements of electrical conductivity. 2D RSoXS patterns were measured by transmission through polymer thin films spin-coated on top of un- treated silicon nitride windows. We assume minimal difference in mor- phology when comparing films on silicon nitride windows and untreated quartz substrate used for conductivity measurements. On the other hand, vapor-doped PBTTT films on octadecyltrichlorosilane (OTS)–treated substrates were floated onto silicon nitride windows. All the scattering images showed a diffuse isotropic ring (fig. S7) that was azimuthally averaged to obtain a 1D scattering profile [intensity (I) versus q] (Fig. 5). The OCL is one-half of the characteristic length scale (d* =2p/q*) obtained from the primary scattering peak (q*)of the Lorentz-corrected (I*q ) scattering profile. The peak position was determined from fits using log-normal function. For equivalent annealing times (10 min), the OCL of neat PBTTT is sensitive to the annealing temperature above the LC transition. The OCL is ∼140 and ∼180 nm when annealed at 150° and 180°C, respectively (fig. S8 and table S3). Note that neat films were annealed at 180°C, whereas solution-doped films were annealed at 150°C to minimize dedoping of F TCNQ. Doped films with the higher values of the OCL correlate to higher s and, thus, higher apparent m because of the comparable level of dop- ing. The OCL of doped films of PBTTT depends strongly on the pro- cessing method (Fig. 5 and table S3). F TCNQ vapor-doped annealed films have an OCL of ∼220 nm, whereas the OCL for an annealed doped film cast from solution is ∼44 nm. The vapor-doped film and neat annealed PBTTT films have relatively similar OCLs. An- nealed films vapor-doped with F TCNQ have an OCL of ~210 nm, suggesting that the introduction of the molecular dopant does not have alarge impact on the OCL.The OCLof ∼100 nm of solution-doped PBTTT:F TCNQ films after thermal annealing at 150°C for 10 min approaches the OCL of a neat film annealed at 150°C (OCL, ~140 nm) likely due to dedoping, which allows the morphology to change. There- fore, we assert that the improved orientational alignment of the backbone leads to a higher m in vapor-doped films, similar to previous Fig. 4. Out-of-plane and in-plane scattering profiles of annealed neat and observations for OTFTs. doped PBTTT thin films. GIWAXS line cuts of (A) out-of-plane scattering and The trends in the OCL for as-cast doped films further confirm the (B) in-plane scattering. Black, neat; orange, F TCNQ solution doping; purple, significant role of backbone alignment in controlling s. The OCL of F TCNQ vapor doping; green, F TCNQ solution doping; blue, F TCNQ vapor doping. 4 2 2 The F TCNQ-doped film corresponds to as-cast conditions. Dashed red lines are an as-cast neat film is ∼70 nm, whereas that of the as-cast film from guides to the eye relative to the peak positions for the neat film. All scattering the PBTTT:F TCNQ solution is ∼40 nm. This difference indicates profiles correspond to thermally annealed films, except for the F TCNQ solution that doping in solution slightly reduces the OCL relative to a neat doping, which is for the as-cast case. a.u., arbitrary units. film before annealing. Moreover, the OCL observed for as-cast films is essentially identical to the annealed solution-doped film (OCL, of a core electron to an unoccupied p-orbital is orthogonal to the ∼44 nm). This similarity shows that annealing a heavily doped film p-conjugated plane of the molecule (50). By using a linearly polarized has no effect on enhancing the backbone alignment, and confirms why soft x-ray with the electric field vector in the plane of the film and tuning the s of as-cast and annealed solution-doped films is nearly identical. The the photon energy to the C 1s to p* resonance (285.4 eV for PBTTT), fact that the OCL does not increase with annealing a solution-doped film scattering contrast arises from the variations in molecular orientation. is not surprising because the PBTTT is now heavily charged, which shifts For PBTTT thin films, strong resonant scattering is expected in trans- the LC transition temperature to higher values and thus does not enter an mission mode because the p-stacking orientation is primarily in the LC mesophase when annealing at 150°C. Attempts to anneal the films at plane of the film (that is, TDM of the 1s to p* resonance is in the plane higher temperatures result in significant dedoping (38). For as-cast of film). Using RSoXS measurements, we can quantify the backbone F TCNQ solution-doped films, we observe a qualitatively different alignment through the orientational correlation length (OCL). The scattering profile relative to the as-cast neat and F TCNQ-doped films. OCL is defined as the average length over which the polymeric back- The scattering peaks are broader where the primary peak is at a lower −1 bones (that is, the LC director) drift out of alignment with each oth- q ∼ 0.03 nm (OCL, ∼130 nm). This suggests that aggregation of er. The OCL was shown to have an empirical exponential relationship PBTTT in solution, which dictates the as-cast thin film morphology, with field-effect m values, providing the first direct evidence that ori- is different when using a weak molecular dopant like F TCNQ. Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 6of13 | SCIENCE ADVANCES RESEARCH ARTICLE Fig. 5. RSoXS of doped films. Lorentz-corrected scattering profiles (log-log scale) from azimuthally averaged RSoXS images at a photon energy of 285.4 eV for (A)annealed and (B) as-cast films. The curves were offset for clarity. Black, neat; purple, F TCNQ vapor doping; blue, F TCNQ vapor doping; orange, F TCNQ solution doping; green, F TCNQ 4 2 4 2 solution doping. RSoXS experiments were performed at Advanced Light Source (ALS) beamline 11.0.1.2. When vapor-doping an as-cast film, we observe a qualitatively dif- A significant assumption in previous work on the relationship ferent scattering profile relative to the as-cast neat film. Essentially, two between the m from OTFTs and the OCL was that the OCL, a bulk −1 broad scattering peaks are seen around q ∼ 0.04 nm (OCL, ~130 nm), property, correlates to an interface-dominated charge transport mea- −1 and another is seen at the experimental low q limit (<0.01 nm ). We surement. Here, we can eliminate that assumption because s is a bulk see a similar effect for a F TCNQ vapor-doped film, where we see a transport property, and conclusively show the correlation between −1 peak around q ∼ 0.05 nm (OCL, ∼60 nm) and another at experi- OCL and s and, thus, apparent m. To determine the sensitivity of s mental low q limit. The second scattering peak corresponds to an to the OCL, we plot log(s) versus OCL for various doped films in OCL that is larger than that of an as-cast neat film. The larger cor- Fig. 6A. The increase in s is most significant at lower OCL values and related domains provide better charge transport, leading to the high then approaches a plateau at higher OCL values. The data points in s = 114.1 ± 0.5 S/cm for F TCNQ vapor-doped as-cast films and s = Fig. 6A are at comparable doping levels with respect to F TCNQ samples, 4 4 13.7 ± 0.2 S/cm for F TCNQ vapor-doped as-cast films. We can rule thus indicating that the apparent m is the parameter increasing with out that this change is due to a thermal annealing process; the sample OCL. The correlation of the interfacial mobility from OTFTs and that underneath the lid is calibrated to be approximately 85°C, which is from the bulk conductivity likely holds in PBTTT because GIWAXS well below the LC transition temperature and has no effect on the typically shows highly oriented crystallites with a thickness equivalent OCL (fig. S8 and table S3). to the total film. In materials where the interfacial and bulk structures are dissimilar, we might not expect such a relationship to hold. Increasing electrical conductivity through processing Extrapolation of s to higher OCL values using Fig. 6A provides Comparison of s and the OCL reveals that long backbone correlation valuable information on the limits of electrical performance. The lengths allow for more efficient charge transport. It has been previously chain alignment process effectively increases the OCL, where an infinite shown that the OCL of PBTTT can be controlled depending on the po- OCL corresponds to perfect alignment. When approaching the limit of larity of the substrate surface (42). In particular, substrates function- an infinite OCL, one would expect s to plateau as the net increase in m alized with a nonpolar monolayer of OTS lead to the largest OCL of becomes smaller (51). For comparison, the mobility in OTFTs of ∼380 nm for an annealed neat PBTTT thin film (42). Vapor-doping PBTTT thin films prepared on an OTS-treated substrate and subse- a PBTTT thin film annealed on an OTS-functionalized substrate with quently strained aligned has been measured (52). The alignment pro- F TCNQ yielded a s of 670 ± 4 S/cm. This s is a factor of ∼3higher cess resulted in a factor of ~2 increase in m according to field-effect value than the vapor-doped film on a bare substrate. The corresponding transistor measurements (52). This observation suggests that s could OCLisalsohigherat ∼350 nm (versus 210 nm for the doped film on a increase to ~1300 S/cm relative to our highest-performing PBTTT: bare substrate). UV-vis-NIR absorption measurements (fig. S3) con- F TCNQ vapor-doped film if the bulk and interfacial mobilities re- firm that the doping level is similar for both cases. Therefore, the in- main correlated. crease in s is entirely from a higher apparent m and reiterates the The determination of the connection between the OCL and s significance of the OCL as a parameter to describe the trends in s. shows the significance of processing on the transport properties of Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 7of13 | SCIENCE ADVANCES RESEARCH ARTICLE of the two studies because of the question of the carrier concentration with different dopants. One can also imagine achieving high s when casting films from a polymer/dopant solution if the processing steps lead to a solid-state thin film with a large OCL. Comparison to previous work reported in the literature further re- iterates the critical role of controlling the morphology to achieve effi- cient charge transport. Sequential doping of P3HT with F TCNQ results in films with higher s relative to solution-doped films (35), which is driven by improved interconnectivity between ordered p-stacked do- mains (36). Sequentially doping P3HT with F TCNQ in nonpolar sol- vents yields a larger 5 to 10 factor increase in s (36). These films also have good interconnectivity, but phase segregation of dopants in the disordered domains was suggested as the reason for the higher s. Other factors such as the molecular structure of the dopant and its sol- ubility in the polymer can also influence morphology and thus s (26). Polar side chains on polymers have been found to increase the ther- mal stability of F TCNQ-doped films, which is particularly useful for thermoelectric energy conversion (53). The environmental stability of F TCNQ in thiophene-based polymers is a concern due to photo- chemical reactions, which will require investigation into appropriate encapsulation for devices (53–55). Role of morphology on Seebeck coefficient and PF Although the impact of the orientation of domains on s is relatively straightforward to understand, the impact on a is less clear. Unlike s, which describes the transport of charge carriers relative to an electric field, a describes the migration of charge carriers relative to temper- ature gradient at the open circuit condition. This is quantified by measuring the voltage drop (DV) relative to temperature difference (DT). Fundamentally, a is related to the population in the electronic DOSs of a material and carrier scattering processes. The expression for the thermopower as a function of electronic conductivity func- tion s(E)isgiven as k ðE  E Þ sðEÞ ∂f ðEÞ B F a ¼ dE ð1Þ e k T s ∂E where E is the Fermi energy, s is the total conductivity, f(E)isthe Fer- mi function, and k /e is a natural unit of thermopower of 86.17 mV/K (56). As the semiconductor is p-doped, the Fermi level shifts closer to the valence band, which results in a decrease in the value of a.The introduction of a molecular dopant into a polymer will also modify the local structure and morphology, making it difficult to model the thermoelectric properties with a constant electronic DOS. There have been significant efforts to model the thermoelectric properties of poly- Fig. 6. The relationship between OCL and thermoelectric material proper- mers (56–58), but these models do not, as yet, consider morphology. ties. (A) Measured electronic conductivity (s), (B) measured Seebeck coefficient The overall thermoelectric properties of PBTTT depend strong- (a), and (C) calculated PF versus the corresponding OCL values, as determined ly on processing methods (Table 1). For example, a = 60 ± 9 and −1 −2 from the RSoXS experiments (table S3). PF = 1.3 ± 0.4 mWm K for the solution-doped annealed film of PBTTT:F TCNQ, whereas annealing followed by vapor infiltration −1 −2 doped semiconducting polymers. Although the vapor doping pro- of F TCNQ leads to a =39±5 mV/K and PF = 32 ± 9 mWm K cess yields the highest values of s here, the vapor doping process with (Table 1). Despite the lower a, PF is higher for vapor-doped films F TCNQ is not the only method to achieve high s. In our previous due to the nearly 100-fold increase in s. The difference in a with va- work, doping annealed films of PBTTT through exposure to a vapor por doping can be attributed, in part, to the slightly higher concen- of a fluorinated trichlorosilane (FTS) or immersion in a 4-ethylbenzene tration of F TCNQ based on the UV-vis spectra, but the morphology sulfonic acid solution yielded s of around 1000 S. The likely origin is of these two films is also quite different. Furthermore, a =42±6 mV/K that the OCL was set before doping and not substantially perturbed by for the F TCNQ vapor-doped film on an OTS-treated substrate, the doping process. It is difficult to directly compare the conductivities which indicates that a is less sensitive to the substrate treatment than Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 8of13 | SCIENCE ADVANCES RESEARCH ARTICLE s. However, the near factor of 3 increase in s results in the high PF of performance of polyacetylene (~10 S/cm) (63). A question is wheth- −1 −2 120 ± 30 mWm K . er the higher conductivities of these materials are due to the electronic With F TCNQ as the dopant, we observe a values of 111.7 ± structure of the conjugated backbone or other factors. In comparison 0.1 mV/K for the as-cast solution-doped film, 130 ± 20 mV/K for the to these materials, the presence of alkyl side chains on semiconducting vapor-doped as-cast film, and 140 ± 20 mV/K for the vapor-doped polymers, such as PBTTT, results in a significant volume of the annealed film (Table 1). The higher a relative to F TCNQ is not material being insulating. Accounting for the insulating side chains surprising because of the likely lower concentration of F TCNQ in reveals an effective electronic conductivity (s ) of the conjugated core 2 eff the film and potentially lower efficiency of carrier formation. The representing densely packed polymer chains (37). For PBTTT with F TCNQ vapor-doped annealed film yields the highest a, consist- tetradecyl side chains, the conjugated core accounts for approximately ent with the slightly lower doping efficiency, as indicated by UV-vis 15% of the volume in a film (determined from the unit cell of PBTTT). (Fig. 2). Owing to the high s of 36 ± 3 S/cm, the corresponding PF is The s would translate to ~4000 S/cm for our highest-performing eff −1 −2 70 ± 20 mWm K .Despite the lower s relative to the F TCNQ film. If these structural changes can be achieved without changing −1 −2 vapor-doped film on a bare quartz substrate, the PF is greater because the a, the PF could reach values of ~500 mWm K at high levels of the large a value. This shows the importance of tuning a while not of doping similar to PEDOT. Removal, or shortening, of the side significantly sacrificing s. chains of a polymer leads to difficulties in processing, that is, PEDOT Comparison of a to the corresponding OCL reveals that, unlike s, is poorly soluble and is cast as a dispersion or directly grown on a sub- the value of a does not show marked changes (Fig. 6B). For F TCNQ- strate. The dopant itself also modifies the volume of the doped doped samples, a is in the range of 30 to 60 mV/K for the full OCL material. However, this comparison shows that conjugated backbones window. In addition, the F TCNQ samples are essentially the same other than PEDOT have significant promise if their structure can be (~130 mV/K) at different OCL values. There appears to be a small up- judiciously modified to maintain their processability and allow for in- ward trend for vapor-doped samples, but overall a is less sensitive to corporation of the dopant. polymer chain alignment. Some recent work on the anisotropy of s and a on PEDOT:PSS [poly(3,4-ethylenedioxythiophene)-poly(styrenesulfonate)] Potential impact of processing on thermal conductivity films (59) and P3HT films (60)alsosuggeststhat a is relatively un- Our study reveals the potential benefits that local chain alignment can affected by polymer chain alignment, and thus, changes in PF with play in further improving PF, but one should also consider the impact the OCL (Fig. 6C) are dominated by the increase in s. This observa- on ZT. Changes in the OCL likely modify the thermal conductivity (k) tion would be expected if the shape of the electronic DOS was rela- through both the phonon (lattice) contribution (k )(64) and, at suf- tively constant and the number of carriers was similar, that is, no ficiently high n and s, the electronic contribution to thermal conduc- large change in the Fermi level. Temperature-dependent a and s mea- tivity (k ). Note that experimental challenges still remain on the surements are needed to fully elucidate the transport mechanism from determination of k of thin films and particularly along the in-plane both doping processes. direction (6). Techniques such as suspended microdevices (65)or Previously, we determined an empirical correlation where a follows the membrane-based ac calorimetry (66) can help to determine the −1/4 a power-law dependence with s (a º s ) and PF follows a square in-plane k. The membrane-based ac calorimetry method revealed an 1/2 −1 −1 root dependence with s (PFº s ) for a variety of p-doped thiophene- in-plane k of 0.39 W m K for an as-cast neat PBTTT film (1 mmthick) based polymers. This correlation was primarily determined on solution- (66). Films with high s,for example, F TCNQ-doped film at 640 S/cm, doped films and thermal annealing conditions 150°C or below. In Fig. 7 , could result in a significant contribution from k . PEDOT:PSS, for exam- we plot the empirical correlations (dashed lines) along with a and PF ple, has a significant contribution from k at a s of ~500 S/cm and higher values reported in this study. The solution-doped samples follow the em- (67). Therefore, there may be an optimization process on the extent of pirical trends. On the other hand, vapor-doped samples deviate from alignment and doping level that leads to a minimal increase of k to the empirical trends, where the values are observed to be higher than achieve a high ZT. A potential optimization process can be seen through expected at corresponding s values. The positive deviation of a and the example of PBTTT:F TCNQ, which yields lower s of 36 S/cm rel- PF relative to the empirical trend line is most pronounced for vapor- ative to PBTTT:F TCNQ (220 S/cm) at a comparable OCL value of doped PBTTT:F TCNQ and also for both solution- and vapor-doped ∼200 nm (Fig. 5A). Despite the lower s, the PBTTT:F TCNQ film 4 2 −1 −2 PBTTT:F TCNQ films. The PF of our highest-performing PBTTT: yields a higher a and thus a larger PF of 70 mWm K relative to −1 −2 F TCNQ film is similar to our previous work on PBTTT vapor-doped the PBTTT:F TCNQ film (PF = 32 mWm K ) (Fig. 5, B and C). As 4 4 with an interfacial FTS (green diamond marker in Fig. 7 ) (33). It has re- a consequence, the lower s can be leveraged to minimize, in principle, cently been proposed that the power-law relationship is due to a change the electronic contributions to k while still being able to achieve a rel- in the energy-dependent conductivity function of the material [s(E)] atively high PF. (56). The higher PF here would suggest that this function is affected by processing conditions speculated in that work. Summary We have explored how solution- and vapor-doping a high-mobility Future routes to improve the PF of polymers p-type polymer, PBTTT-C , affect its thermoelectric transport prop- Although we have focused only at the limit of high doping level, system- erties and its underlying microstructure. Overall, vapor-doping with atically varying the doping level and the OCL provides a route to opti- either F TCNQ or F TCNQ yields higher s relative to solution-doped 4 2 mization of the PF. We can also consider the performance of PEDOT, films. The enhancement in s is not related to the local order because which has been the benchmark conducting polymer for thermoelectrics the perturbations to the local structure are minimal and similar with −1 −2 (7). A PF as high as ∼460 mWm K has been reported (61), and s either doping route. We determined using RSoXS that the alignment of ∼5000 S/cm has been measured for metallic template polymerized of ordered domains, quantified through the OCL, is a critical parameter PEDOT doped with sulfuric acid (62). This s is comparable to the in explaining trends in s. The larger OCL for vapor-doped films allows Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 9of13 | SCIENCE ADVANCES RESEARCH ARTICLE Fig. 7. Trends in Seebeck coefficient and power factor. Log-log scale plot showing the trends in (A) Seebeck coefficient (a) and (B)PF(a s) versus electronic conductivity (s) for solution- and vapor-doped films. Orange markers are for vapor-doped films, and blue markers are for solution-doped films. Circle markers are for F TCNQ-doped films, square markers are for F TCNQ-doped films, and triangle marker is for the F TCNQ-doped film on OTS-treated substrate. Open markers 4 2 4 correspond to thermally annealed films, and filled markers correspond to as-cast films. The open green diamond is our previously reported FTS-doped PBTTT thin −1/4 1/2 film (33). Dashed lines are empirical trends [a º s and PF º s ] we previously reported on various doped semiconducting polymers (34). for efficient charge transport and thus a higher s relative to solution- All chemicals were used as received. PBTTT-C was synthesized doped films, which yield a smaller OCL. Owing to better long-range using literature procedure (68) with a number-average molecular correlation length of backbones, a F TCNQ vapor-doped PBTTT-C weight (M ) of 18,000 or 24,000 g/mol. 4 14 n casted on an OTS-treated substrate yields a high s of 670 ± 4 S/cm and corresponding large a of 42 ± 6 mV/K. This translates to a PF of 120 ± Neat PBTTT and PBTTT:F TCNQ (n = 2 or 4) thin-film −1 −2 30 mWm K —the highest reported value for F TCNQ-doped semi- sample fabrication conducting polymers. In addition, using a weaker molecular dopant Neat PBTTT solution preparation like F TCNQ can lead to a large a (140 ± 20 mV/K) while not signif- Neat PBTTT solutions at 5 mg/ml were prepared by dissolving PBTTT icantly sacrificing s (36 ± 3 S/cm), and thus yield a large PF of 70 ± in either CB or 1:1 CB:ODCB and heated to 120°C. Approximately −1 −2 20 mWm K . 1 hour was needed to fully dissolved PBTTT. The heated neat PBTTT With a better understanding of processing effects on s and a,we solution was filtered using 0.45-mm polytetrafluoroethylene (PTFE) sy- can now outline some general processing guidelines to achieve high ringe filter. The PBTTT solution gels when cooled below 80°C. As a thermoelectric PF. First, casting a neat semiconducting polymer film consequence, the neat PBTTT solution was maintained at 120°C before that forms locally p-stacked domains with long-range correlation doping or spin coating. lengths of the conjugated backbones provides an ideal microstructure Solution doping of PBTTT with F TCNQ (n = 2 or 4) for efficient charge transport for high apparent m. Second, introducing F TCNQ (3 wt %; n = 2 or 4) solution was prepared by dissolving the the molecular dopant into the polymer film (that is, from the vapor dopant in ODCB and heated to 150°C. To achieve 10 wt % dopant phase) to increase n, and in turn s, should lead to minimal perturba- concentration, an aliquot of F TCNQ solution was added to the neat tion to the local order while maintaining, or enhancing, the long-range PBTTT solution and heated to 120°C. After the addition of the dopant, correlation lengths of conjugated backbones. By leveraging the high the heated polymer solution became more viscous and immediately apparent m and then precisely controlling the dopant concentration, transitioned from a red to black color. These changes were indicative or by choosing a weaker molecular dopant, one can obtain a large a while of charge transfer between the polymer and dopant in solution. To not significantly sacrificing s. Overall, developing better doping routes minimize gelation and precipitation of the charged polymer, the and advancing our fundamental understanding of structure-property re- PBTTT:F TCNQ (n = 2 or 4) solution was maintained at 120°C before lationships of semiconducting polymers will have far-reaching implica- spin coating. tions on the deployment of lightweight and low-cost organic Substrates thermoelectric modules for thermal energy conversion and management. Thin films were prepared on quartz substrates (1.5 cm × 1.5 cm; Uni- versity Wafers) for conductivity and Seebeck measurements, native oxide silicon substrates (1.5 cm × 1.5 cm; International Wafer Services) MATERIALS AND METHODS for GIWAXS experiments, and silicon nitride windows (window size, Materials 1.5 mm × 1.5 mm; window thickness, 100 nm; frame size, 5 mm × 5 mm; Anhydrous chlorobenzene (CB) and o-dichlorobenzene (ODCB) thick frame, 200 mm; NX5150C, Norcada Inc.) for RSoXS experiments. were purchased from Sigma-Aldrich. F TCNQ and F TCNQ were The quartz and silicon substrates were cleaned by sonicating first in 4 2 purchased from TCI Chemicals. OTS was purchased from Gelest. acetone and then in isopropanol. Samples of PBTTT thin films on Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 10 of 13 | SCIENCE ADVANCES RESEARCH ARTICLE OTS-treated substrates were prepared by plasma-treating a quartz Four-point probe conductivity measurements were performed using substrate with air for 2 min. Then, the substrate was immersed in a custom-designed probe station in a N glove box. Voltage and current OTS and anhydrous toluene (volume ratio, 0.10 to 1) for 10 min, heated measurements were performed using a Keithley 2400 source measure to 80°C, and subsequently rinsed with toluene to achieve an OTS self- unit and Keithley 6221 precision current source. A constant current was assembled monolayer. The OTS-treated quartz substrates were heated applied to the outer contacts, and the resultant steady-state voltage re- to 80°C in a N glove box to remove the residual solvent. sponse was recorded from the inside contacts. The resistance (R;ohms) Spin-coating conditions of the sample was extracted from the slope of the VI sweep using Ohm’s Neat PBTTT thins films were spin-coated from a CB solution (5 mg/ml) law (V = IR). or from a 1:1 CB:ODCB solution (5 mg/ml). The heated solution (120°C) The Seebeck coefficient (a) measurements were performed in a wasspin-coated first at 1000rpm for45sandthenat3000rpm for15s N glove box using a custom-built setup. A detailed description of under ambient conditions. The heated (120°C) PBTTT:F TCNQ solu- the Seebeck coefficient measurement setup can be found in the study tion was spin-coated at 1000 rpm for 45 s and then at 3000 rpm for 15 s, of Glaudell et al.(34). Peltier elements 5 mm apart provided the tem- which leads to the macroscopically uniform thin films with thicknesses in perature difference (DT = T – T ). A minimal amount of thermal H C the range of 40 to 60 nm. conductive paste was applied to the tips of the thermal couple to ensure Thermal annealing good thermal contact between the thermocouple and the gold pads. The The neat PBTTT thin films were annealed in a N -filled glove box for measurement system has systematic error of 15% due to thermal 10 min at 180°C and then slowly cooled to 80°C. The thermal anneal- anchoring issues. A delay of 100 s was used for voltage measurements ing conditions for solution-doped films were at 150°C for 10 min in a to ensure that a steady-state temperature gradient was reached. The N -filled glove box. The thickness of the annealed film was approx- Seebeck coefficient was calculated from the slope of a linear fit for the imately in the range of 15 to 40 nm according to atomic force micros- DV versus DT plot. The measurements were taken within an approx- copy (AFM; Asylum MFP-3D) measurements. imate DT of ±3 K around 300 K. Vapor-doping process To infer the apparent charge carrier mobility (m), we used the MR Neat PBTTT thin films were first fabricated using the procedure out- (dopant/monomer) values of the vapor- and solution-doped films and lined above. Subsequently, an as-cast or thermal annealed neat film assumed that the dopants are fully ionized and all the hole carriers was vapor-doped with F TCNQ (n =2or 4) in aN glove box. Approx- generated contribute to conductivity. Knowing the unit cell of PBTTT, n 2 imately 5 to 10 mg of dopant were placed in a glass jar (Qorpak with a we can calculate the carrier concentration (n). After which, we can PTFE lined cap; diameter, ~5 cm; height, ~4.5 cm). The polymer sam- calculate the mobility using the equation m = s/(qn), where s is the ple was placed underneath the cap (near the center) using double- electronic conductivity and q is the charge (+1). sided tape. The closed jar was heated on a hotplate set to ~210°C. The typical heating times were in the range of 2 to 10 min. This Synchrotron x-ray scattering heating process leads to a partial vapor pressure of the dopant in 2D GIWAXS images were obtained using beamline 11-3 at SSRL lo- the jar. Successful doping of a PBTTT thin film was confirmed when cated on the SLAC (Stanford Linear Accelerator Center) National the film has a nearly transparent appearance (typically achieved after Accelerator Laboratory campus. Thin-film samples for GIWAXS 5 to 10 min for 25-nm PBTTT thin film). Successful doping was con- experiments were prepared following the procedures outlined above. firmed through UV-vis-NIR measurements. The temperature of a The samples were exposed to x-rays with a wavelength of 0.9752 Å, sample underneath the cap was measured using a thermocouple and 2D scattering images were obtained using a MAR345 image while the jar was heated. The sample was around 75°C after 5 min plate detector or MarCCD detector, which was placed 400 mm from and equilibrates to around 95°C after 30 min. the sample. A LaB sample was used as a standard for calibration. All samples were placed in a He-filled chamber to reduce air scattering UV-vis-NIR spectroscopy and minimize beam damage to the sample. The reported GIWAXS UV-vis-NIR spectra of thin films on 0.5-mm-thick quartz substrates images were taken at a grazing incident x-ray angle of 0.10 or 0.12, (1.5 cm × 1.5 cm) were obtained using the Shimadzu UV-3600 UV- which is above the critical angle of the polymer film and below the Nir-NIR Spectrometer at the UC (University of California) Santa critical angle of the silicon substrate. Barbara Materials Research Laboratory TEMPO Facility. The doped X-ray specular scattering was collected on beamline 2-1 using the films were placed in a custom-built airtight holder to ensure doping setup with Soller slits and a photomultiplier tube. The incident x-ray stability. Measurements were taken within a wavelength (l) range of energy was 11.5 keV. 300 to 2300 nm. RSoXS samples were prepared by directly spin-coating onto the silicon nitride windows following thin-film fabrication process out- Conductivity and Seebeck measurements lined above. RSoXS samples from doped PBTTT thin films on an Gold contact layers (~100 nm thick) for electronic conductivity and See- OTS-treated quartz substrate were prepared by first scribing small beck coefficient measurements were thermally evaporated (Angstrom square grids using a razor blade. A solution of 15% hydrogen fluoride Engineering Amod) onto either neat PBTTT thin films or solution- in deionized water was used to partially etch the oxide layer. The films doped PBTTT:F TCNQ (n = 2 or 4) thin films through a shadow mask. were then lifted off the substrate by dipping them in deionized water. Four-point probe conductivity contacts had a channel length of 0.2 mm The pieces of freestanding films were then lifted out of the water using and a channel width of 1 mm. Seebeck measurement contacts consisted silicon nitride windows. of 1-mm gold pads adjacent to 0.2-mm × 1-mm gold bars. The distance RSoXS experiments were performed on beamline 11.0.1.2 at the between the gold pads (temperature probes) and gold bars (voltage ALS located on the Lawrence Berkeley National Laboratory campus. probes) was 3, 4, and 5 mm apart. A detailed schematic is provided 2D RSoXS scattering images were collected in transmission mode in fig. S9. using a charge-coupled device camera (Princeton Instrument PI-MTE) Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 11 of 13 | SCIENCE ADVANCES RESEARCH ARTICLE 16. P. Pingel, D. Neher, Comprehensive picture of p-type doping of P3HT with the molecular cooled to −45°C in a high-vacuum chamber (add pressure). The sample- acceptor F TCNQ. Phys. Rev. B 87, 115209 (2013). to-detector distance was set to 175 mm. The 2D scattering image was 17. V. I. Arkhipov, P. Heremans, E. V. Emelianova, G. J. Adriaenssens, H. Bässler, Charge carrier reduced by azimuthal integration over all values of q (scattering mobility in doped semiconducting polymers. Appl. Phys. Lett. 82, 3245–3247 (2003). vector). After subtraction of the dark image, the Lorentz-corrected 18. V. I. Arkhipov, E. V. Emelianova, P. Heremans, H. Bässler, Analytic model of carrier mobility in doped disordered organic semiconductors. Phys. Rev. B 72, 235202 (2005). profile (I*q versus q) was obtained. The data reduction was per- 19. B. Lüssem, C.-M. Keum, D. Kasemann, B. Naab, Z. Bao, K. Leo, Doped organic transistors. formed in Wavemetrics Igor Pro using NIKA macro developed by Chem. Rev. 116, 13714–13751 (2016). J. Ilvasky at the Advanced Photon Source (69). 20. H. Sirringhaus, 25th anniversary article: Organic field-effect transistors: The path beyond amorphous silicon. Adv. Mater. 26, 1319–1335 (2014). 21. R. Noriega, J. Rivnay, K. Vandewal, F. P. V. Koch, N. Stingelin, P. Smith, M. F. Toney, A. Salleo, A general relationship between disorder, aggregation and charge transport in SUPPLEMENTARY MATERIALS conjugated polymers. Nat. Mater. 12, 1038–1044 (2013). Supplementary material for this article is available at http://advances.sciencemag.org/cgi/ 22. D. Venkateshvaran, M. Nikolka, A. Sadhanala, V. Lemaur, M. Zelazny, M. Kepa, content/full/3/6/e1700434/DC1 M. Hurhangee, A. J. Kronemeijer, V. Pecunia, I. Nasrallah, I. Romanov, K. Broch, fig. S1. AFM height and phase images of neat annealed PBTTT and F TCNQ vapor-doped films I. McCulloch, D. Emin, Y. Olivier, J. Cornil, D. Beljonne, H. Sirringhaus, Approaching at 5 and 10 min. disorder-free transport in high-mobility conjugated polymers. Nature 515, 384–388 fig. S2. Absorption spectra showing the NIR regime for doped PBTTT films and the thermal (2014). stability of F TCNQ-doped films. 23. J. B. Torrance, An overview of organic charge-transfer solids: Insulators, metals, and the fig. S3. Additional UV-vis-NIR spectra of F TCNQ vapor-doped films relative to a neat film. neutral-ionic transition. Mol. Cryst. Liq. Cryst. 126,55–67 (1985). fig. S4. Williamson-Hall plot for neat (black circle) and F TCNQ vapor-doped film. 24. K.-H. Yim, G. L. Whiting, C. E. Murphy, J. J. M. Halls, J. H. Burroughes, R. H. Friend, J.-S. Kim, fig. S5. In-plane scattering profiles of as-cast neat and doped films. Controlling electrical properties of conjugated polymers via a solution-based p-type fig. S6. Thin-film thickness profile of neat and vapor-doped PBTTT:F TCNQ film. doping. Adv. Mater. 20, 3319–3324 (2008). fig. S7. Representative 2D RSoXS images for neat PBTTT, F TCNQ vapor-doped, and F TCNQ 4 4 25. C. Wang, D. T. Duong, K. Vandewal, J. Rivnay, A. Salleo, Optical measurement of doping solution-doped thin films (all thermally annealed). efficiency in poly(3-hexylthiophene) solutions and thin films. Phys. Rev. B 91, 85205 (2015). fig. S8. Lorentz-corrected scattering profiles of neat PBTTT for different annealing 26. J. Li, G. Zhang, D. M. Holm, I. E. Jacobs, B. Yin, P. Stroeve, M. Mascal, A. J. Moulé, temperatures. Introducing solubility control for improved organic p-type dopants. Chem. Mater. 27, fig. S9. Schematic of the geometry of the contacts for electronic conductivity and Seebeck 5765–5774 (2015). measurements on thin films of doped polymers. 27. J. B. Torrance, The difference between metallic and insulating salts of table S1. X-ray reflection peaks of annealed PBTTT thin films from GIWAXS. tetracyanoquinodimethone (TCNQ): How to design an organic metal. Acc. Chem. Res. 12, table S2. X-ray reflection peaks of as-cast PBTTT thin films from GIWAXS. 79–86 (1979). table S3. Summary of OCLs for doped films. 28. F. Wudl, From organic metals to superconductors: Managing conduction electrons in organic solids. Acc. Chem. Res. 17, 227–232 (1984). 29. M. R. Bryce, L. C. Murphy, Organic metals. Nature 309, 119–126 (1984). 30. P. Pingel, R. Schwarzl, D. Neher, Effect of molecular p-doping on hole density and REFERENCES AND NOTES mobility in poly(3-hexylthiophene). Appl. Phys. Lett. 100, 143303 (2012). 1. B. Lüssem, M. Riede, K. Leo, Doping of organic semiconductors. Phys. Status Solidi A 210, 31. I. Salzmann, G. Heimel, M. Oehzelt, S. Winkler, N. Koch, Molecular electrical doping of 9–43 (2013). organic semiconductors: Fundamental mechanisms and emerging dopant design rules. 2. H. Ma, H.-L. Yip, F. Huang, A. K.-Y. Jen, Interface engineering for organic electronics. Acc. Chem. Res. 49, 370–378 (2016). Adv. Funct. Mater. 20, 1371–1388 (2010). 32. J. Gao, E. T. Niles, J. K. Grey, Aggregates promote efficient charge transfer doping of 3. M.-C. Jung, S. R. Raga, L. K. Ono, Y. Qi, Substantial improvement of perovskite solar cells poly(3-hexylthiophene). J. Phys. Chem. Lett. 4, 2953–2957 (2013). stability by pinhole-free hole transport layer with doping engineering. Sci. Rep. 5, 33. S. N. Patel, A. M. Glaudell, D. Kiefer, M. L. Chabinyc, Increasing the thermoelectric 9863 (2015). power factor of a semiconducting polymer by doping from the vapor phase. ACS Macro 4. G. Lu, J. Blakesley, S. Himmelberger, P. Pingel, J. Frisch, I. Lieberwirth, I. Salzmann, Lett. 5, 268–272 (2016). M. Oehzelt, R. Di Pietro, A. Salleo, N. Koch, D. Neher, Moderate doping leads to high 34. A. M. Glaudell, J. E. Cochran, S. N. Patel, M. L. Chabinyc, Impact of the doping method performance of semiconductor/insulator polymer blend transistors. Nat. Commun. 4,1588 on conductivity and thermopower in semiconducting polythiophenes. Adv. Energy Mater. (2013). 5, 1401072 (2015). 5. O. Bubnova, X. Crispin, Towards polymer-based organic thermoelectric generators. 35. D. T. Scholes, S. A. Hawks, P. Y. Yee, H. Wu, J. R. Lindemuth, S. H. Tolbert, B. J. Schwartz, Energy Environ. Sci. 5, 9345–9362 (2012). Overcoming film quality issues for conjugated polymers doped with F TCNQ by solution 6. S. N. Patel, M. L. Chabinyc, Anisotropies and thermoelectric properties of semiconducting sequential processing: Hall effect, structural, and optical measurements. J. Phys. Chem. Lett. polymers. J. Appl. Polym. Sci. 134, 44403 (2016). 6,4786–4793 (2015). 7. B. Russ, A. Glaudell, J. J. Urban, M. L. Chabinyc, R. A. Segalman, Organic thermoelectric 36. I. E. Jacobs, E. W. Aasen, J. L. Oliveira, T. N. Fonseca, J. D. Roehling, J. Li, G. Zhang, materials for energy harvesting and temperature control. Nat. Rev. Mater. 1, 16050 (2016). M. P. Augustine, M. Mascal, A. J. Moulé, Comparison of solution-mixed and sequentially 8. O. Owoyele, S. Ferguson, B. T. O’Connor, Performance analysis of a thermoelectric processed P3HT:F4TCNQ films: Effect of doping-induced aggregation on film cooler with a corrugated architecture. Appl. Energy 147, 184–191 (2015). morphology. J. Mater. Chem. C 4, 3454–3466 (2016). 9. Q. Wei, M. Mukaida, K. Kirihara, Y. Naitoh, T. Ishida, Polymer thermoelectric modules 37. K. Kang, S. Watanabe, K. Broch, A. Sepe, A. Brown, I. Nasrallah, M. Nikolka, Z. Fei, screen-printed on paper. RSC Adv. 4, 28802 (2014). M. Heeney, D. Matsumoto, K. Marumoto, H. Tanaka, S.-i. Kuroda, H. Sirringhaus, 2D 10. J.-H. Bahk, H. Fang, K. Yazawa, A. Shakouri, Flexible thermoelectric materials and device coherent charge transport in highly ordered conducting polymers doped by solid state optimization for wearable energy harvesting. J. Mater. Chem. C 3, 10362–10374 (2015). diffusion. Nat. Mater. 15, 896–902 (2016). 11. C. Wan, R. Tian, A. B. Azizi, Y. Huang, Q. Wei, R. Sasai, S. Wasusate, T. Ishida, K. Koumoto, 38. J. E. Cochran, M. J. N. Junk, A. M. Glaudell, P. L. Miller, J. S. Cowart, M. F. Toney, C. J. Hawker, Flexible thermoelectric foil for wearable energy harvesting. Nano Energy 30, 840–845 B. F. Chmelka, M. L. Chabinyc, Molecular interactions and ordering in electrically (2016). doped polymers: Blends of PBTTT and F TCNQ. Macromolecules 47,6836–6846 (2014). 12. H. Fang, B. C. Popere, E. M. Thomas, C.-K. Mai, W. B. Chang, G. C. Bazan, M. L. Chabinyc, 39. J. Li, C. W. Rochester, I. E. Jacobs, S. Friedrich, P. Stroeve, M. Riede, A. J. Moulé, R. A. Segalman, Large-scale integration of flexible materials into rolled and corrugated Measurement of small molecular dopant F4TCNQ and C F diffusion in organic bilayer 60 36 thermoelectric modules. J. Appl. Polym. Sci. 134, 44208 (2017). architectures. ACS Appl. Mater. Interfaces 7, 28420–28428 (2015). 13. K. Kirihara, Q. Wei, M. Mukaida, T. Ishida, Thermoelectric power generation using 40. I. McCulloch, M. Heeney, C. Bailey, K. Genevicius, I. MacDonald, M. Shkunov, D. Sparrowe, nonwoven fabric module impregnated with conducting polymer PEDOT:PSS. Synth. Met. S. Tierney, R. Wagner, W. Zhang, M. L. Chabinyc, R. J. Kline, M. D. McGehee, M. F. Toney, 225,41–48 (2017). Liquid-crystalline semiconducting polymers with high charge-carrier mobility. Nat. Mater. 5, 14. G. J. Snyder, E. S. Toberer, Complex thermoelectric materials. Nat. Mater. 7, 105–114 328–333 (2006). (2008). 15. A. Salleo, R. J. Kline, D. M. DeLongchamp, M. L. Chabinyc, Microstructural characterization 41. M. L. Chabinyc, M. F. Toney, R. J. Kline, I. McCulloch, M. Heeney, X-ray scattering study of and charge transport in thin films of conjugated polymers. Adv. Mater. 22, 3812–3838 thin films of poly(2,5-bis(3-alkylthiophen-2-yl)thieno[3,2-b]thiophene). J. Am. Chem. Soc. (2010). 129,3226–3237 (2007). Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 12 of 13 | SCIENCE ADVANCES RESEARCH ARTICLE 42. B. A. Collins, J. E. Cochran, H. Yan, E. Gann, C. Hub, R. Fink, C. Wang, T. Schuettfort, 62. M. N. Gueye, A. Carella, N. Massonnet, E. Yvenou, S. Brenet, J. Faure-Vincent, S. Pouget, C. R. McNeill, M. L. Chabinyc, H. Ade, Polarized X-ray scattering reveals non-crystalline F. Rieutord, H. Okuno, A. Benayad, R. Demadrille, J.-P. Simonato, Structure and dopant orientational ordering in organic films. Nat. Mater. 11, 536–543 (2012). engineering in PEDOT thin films: Practical tools for a dramatic conductivity 43. P. Brocorens, A. Van Vooren, M. L. Chabinyc, M. F. Toney, M. Shkunov, M. Heeney, enhancement. Chem. Mater. 28, 3462–3468 (2016). I. McCulloch, J. Cornil, R. Lazzaroni, Solid-state supramolecular organization of 63. C. K. Chiang, C. R. Fincher Jr., Y. W. Park, A. J. Heeger, H. Shirakawa, E. J. Louis, S. C. Gau, polythiophene chains containing thienothiophene units. Adv. Mater. 21, 1193–1198 A. G. MacDiarmid, Electrical conductivity in doped polyacetylene. Phys. Rev. Lett. 39, (2009). 1098–1101 (1977). 44. T. Schuettfort, B. Watts, L. Thomsen, M. Lee, H. Sirringhaus, C. R. McNeill, Microstructure 64. X. Wang, V. Ho, R. A. Segalman, D. G. Cahill, Thermal conductivity of high-modulus of polycrystalline PBTTT films: Domain mapping and structure formation. ACS Nano 6, polymer fibers. Macromolecules 46, 4937–4943 (2013). 1849–1864 (2012). 65. A. Weathers, Z. U. Khan, R. Brooke, D. Evans, M. T. Pettes, J. W. Andreasen, X. Crispin, L. Shi, 45. D. M. DeLongchamp, R. J. Kline, Y. Jung, D. S. Germack, E. K. Lin, A. J. Moad, L. J. Richter, Significant electronic thermal transport in the conducting polymer poly(3,4- M. F. Toney, M. Heeney, I. McCulloch, Controlling the orientation of terraced nanoscale ethylenedioxythiophene). Adv. Mater. 27, 2101–2106 (2015). “ribbons” of a poly(thiophene) semiconductor. ACS Nano 3,780–787 (2009). 66. H. Ushirokita, H. Tada, In-plane thermal conductivity measurement of conjugated 46. K. Kanai, K. Akaike, K. Koyasu, K. Sakai, T. Nishi, Y. Kamizuru, T. Nishi, Y. Ouchi, K. Seki, polymer films by membrane-based AC calorimetry. Chem. Lett. 45,735–737 Determination of electron affinity of electron accepting molecules. Appl. Phys. A 95, (2016). 309–313 (2009). 67. J. Liu, X. Wang, D. Li, N. E. Coates, R. A. Segalman, D. G. Cahill, Thermal conductivity and 47. A. Mityashin, Y. Olivier, T. Van Regemorter, C. Rolin, S. Verlaak, N. G. Martinelli, D. Beljonne, elastic constants of PEDOT:PSS with high electrical conductivity. Macromolecules 48, J. Cornil, J. Genoe, P. Heremans, Unraveling the mechanism of molecular doping in 585–591 (2015). organic semiconductors. Adv. Mater. 24, 1535–1539 (2012). 68. I. Mcculloch, M. Heeney, M. L. Chabinyc, D. Delongchamp, R. J. Kline, M. Cölle, W. Duffy, 48. N. C. Miller, E. Cho, M. J. N. Junk, R. Gysel, C. Risko, D. Kim, S. Sweetnam, C. E. Miller, D. Fischer, D. Gundlach, B. Hamadani, R. Hamilton, L. Richter, A. Salleo, M. Shkunov, L. J. Richter, R. J. Kline, M. Heeney, I. McCulloch, A. Amassian, D. Acevedo-Feliz, C. Knox, D. Sparrowe, S. Tierney, W. Zhang, Semiconducting thienothiophene copolymers: Design, M. R. Hansen, D. Dudenko, B. F. Chmelka, M. F. Toney, J.-L. Brédas, M. D. McGehee, Use synthesis, morphology, and performance in thin-film organic transistors. Adv. Mater. of X-ray diffraction, molecular simulations, and spectroscopy to determine the molecular 21, 1091–1109 (2009). packing in a polymer-fullerene bimolecular crystal. Adv. Mater. 24, 6071–6079 (2012). 69. J. Ilavsky, Nika: Software for two-dimensional data reduction. J. Appl. Cryst. 45, 324–328 49. C. Wang, D. H. Lee, A. Hexemer, M. I. Kim, W. Zhao, H. Hasegawa, H. Ade, T. P. Russell, (2012). Defining the nanostructured morphology of triblock copolymers using resonant soft X-ray scattering. Nano Lett. 11,3906–3911 (2011). Acknowledgments: Use of the SSRL, SLAC National Accelerator Laboratory, is supported by 50. J. Stöhr, NEXAFS Spectroscopy, vol. 25 of Springer Series in Surface Sciences (Springer, 1992). the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under 51. D. S. Pearson, P. A. Pincus, G. W. Heffner, S. J. Dahman, Effect of molecular weight contract no. DE-AC02-76SF00515. This research used resources of the ALS, which is a U.S. and orientation on the conductivity of conjugated polymers. Macromolecules 26, Department of Energy Office of Science User Facility under contract no. DE-AC02-05CH11231. 1570–1575 (1993). The MRL Shared Experimental Facilities are supported by the MRSEC Program of the NSF 52. X. Xue, G. Chandler, X. Zhang, R. J. Kline, Z. Fei, M. Heeney, P. J. Diemer, O. D. Jurchescu, under award no. DMR 1121053, a member of the NSF-funded Materials Research Facilities B. T. O’Connor, Oriented liquid crystalline polymer semiconductor films with large Network. Funding: The authors acknowledge the support of the Air Force Office of Scientific ordered domains. ACS Appl. Mater. Interfaces 7, 26726–26734 (2015). Research through the Multidisciplinary University Research Initiative on Controlling Thermal 53. J. Li, C. W. Rochester, I. E. Jacobs, E. W. Aasen, S. Friedrich, P. Stroeve, A. J. Moulé, The and Electrical Transport in Organic and Hybrid Materials (AFOSR FA9550-12-1-0002). A.M.G. effect of thermal annealing on dopant site choice in conjugated polymers. Org. Electron. received partial support from the ConvEne IGERT Program of the NSF under NSF-DGE 0801627. 33,23–31 (2016). E.L. acknowledges support from the NSF Graduate Research Fellowship (DGE-1144085). 54. J. Fuzell, I. E. Jacobs, S. Ackling, T. F. Harrelson, D. M. Huang, D. Larsen, A. J. Moulé, Optical Author contributions: S.N.P. designed and performed all experiments related to thermoelectric dedoping mechanism for P3HT:F4TCNQ mixtures. J. Phys. Chem. Lett. 7, 4297–4303 (2016). measurements, GIWAXS, and RSoXS and wrote the manuscript with input from M.L.C. A.M.G. 55. I. E. Jacobs, F. Wang, N. Hafezi, C. Medina-Plaza, T. F. Harrelson, J. Li, M. P. Augustine, assisted in thermoelectric measurements and assisted in the analysis of the data. K.A.P. M. Mascal, A. J. Moulé, Quantitative dedoping of conductive polymers. Chem. Mater. 29, performed UV-vis-NIR experiments. E.M.T. performed AFM experiments. K.A.O. performed 832–841 (2017). high-resolution x-ray scattering experiments. E.L. assisted with RSoXS experiments. M.L.C. 56. S. D. Kang, G. J. Snyder, Charge-transport model for conducting polymers. Nat. Mater. 16, supervised all aspects of the project, designed experiments, and was involved in writing of the 252–257 (2017). manuscript. All coauthors assisted with writing and editing of the manuscript. Figure credit: 57. D. Wang, W. Shi, J. Chen, J. Xi, Z. Shuai, Modeling thermoelectric transport in organic S.N.P. for Figs. 1, 3, and 4; S.N.P., K.A.P., E.M.T., and A.M.G. for Fig. 2; S.N.P. and E.L. for Figs. 5 and materials. Phys. Chem. Chem. Phys. 14, 16505–16520 (2012). 6; S.N.P. and A.M.G. for Fig. 7. Competing interests: The authors declare that they have no 58. G. Zuo, H. Abdalla, M. Kemerink, Impact of doping on the density of states and the competing interests. Data and materials availability: All data needed to evaluate the mobility in organic semiconductors. Phys. Rev. B 93, 235203 (2016). conclusions in the paper are present in the paper and/or the Supplementary Materials. 59. Q. Wei, M. Mukaida, K. Kirihara, T. Ishida, Experimental studies on the anisotropic Additional data related to this paper may be requested from the authors. thermoelectric properties of conducting polymer films. ACS Macro Lett. 3, 948–952 (2014). Submitted 9 February 2017 60. S. Qu, Q. Yao, L. Wang, Z. Chen, K. Xu, H. Zeng, W. Shi, T. Zhang, C. Uher, L. Chen, Highly Accepted 28 April 2017 anisotropic P3HT films with enhanced thermoelectric performance via organic small Published 16 June 2017 molecule epitaxy. NPG Asia Mater. 8, e292 (2016). 10.1126/sciadv.1700434 61. O. Bubnova, Z. U. Khan, H. Wang, S. Braun, D. R. Evans, M. Fabretto, P. Hojati-Talemi, D. Dagnelund, J.-B. Arlin, Y. H. Geerts, S. Desbief, D. W. Breiby, J. W. Andreasen, Citation: S. N. Patel, A. M. Glaudell, K. A. Peterson, E. M. Thomas, K. A. O’Hara, E. Lim, R. Lazzaroni, W. M. Chen, I. Zozoulenko, M. Fahlman, P. J. Murphy, M. Berggren, X. Crispin, M. L. Chabinyc, Morphology controls the thermoelectric power factor of a doped Semi-metallic polymers. Nat. Mater. 13, 190–194 (2013). semiconducting polymer. Sci. Adv. 3, e1700434 (2017). Patel et al., Sci. Adv. 2017; 3 : e1700434 16 June 2017 13 of 13

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