Thermotropic Phase Transition of Benzodithiophene Copolymer Thin

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Thermotropic Phase Transition of Benzodithiophene Copolymer Thin Films and Its Impact on Electrical and Photovoltaic Characteristics Sangwon Ko,†,‡,§ Do Hwan Kim,†,⊥,△ Alexander L. Ayzner,⊥,#,○ Stefan C. B. Mannsfeld,# Eric Verploegen,⊥,# Alexander M. Nardes,∥ Nikos Kopidakis,∥ Michael F. Toney,# and Zhenan Bao*,⊥ ‡

Department of Chemistry, Stanford University, 333 Campus Drive, Stanford, California 94305-4401, United States Department of Chemical Engineering, Stanford University, 443 Via Ortega, Stanford, California 94305-4125, United States # Stanford Synchrotron Radiation Lightsource, SLAC National Accelerator Laboratory, Menlo Park, California 94025, United States ∥ National Renewable Energy Laboratory, Golden, Colorado 80401, United States § Transportation Environmental Research Team, Korea Railroad Research Institute, Uiwang-si 437-757, Korea △ Department of Organic Materials and Fiber Engineering, Soongsil University, Seoul 156-743, Korea ○ Department of Chemistry and Biochemistry, University of California Santa Cruz, 1156 High Street, Santa Cruz, California 95064, United States ⊥

S Supporting Information *

ABSTRACT: We observed a thermotropic phase transition in poly[3,4-dihexyl thiophene-2,2′:5,6′-benzo[1,2-b:4,5-b′]dithiophene] (PDHBDT) thin films accompanied by a transition from a random orientation to an ordered lamellar phase via a nearly hexagonal lattice upon annealing. We demonstrate the effect of temperature-dependent molecular packing on charge carrier mobility (μ) in organic field-effect transistors (OFETs) and photovoltaic characteristics, such as exciton diffusion length (LD) and power conversion efficiency (PCE), in organic solar cells (OSCs) using PDHBDT. The μ was continuously improved with increasing annealing temperature and PDHBDT films annealed at 270 °C resulted in a maximum μ up to 0.46 cm2/(V s) (μavg = 0.22 cm2/(V s)), which is attributed to the well-ordered lamellar structure with a closer π−π stacking distance of 3.5 Å as shown by grazing incidence-angle X-ray diffraction (GIXD). On the other hand, PDHBDT films with a random molecular orientation are more effective in photovoltaic devices than films with an ordered hexagonal or lamellar phase based on current−voltage characteristics of PDHBDT/C60 bilayer solar cells. This observation corresponds to an enhanced dark current density (JD) and a decreased LD upon annealing. This study provides insight into the dependence of charge transport and photovoltaic characteristics on molecular packing in polymer semiconductors, which is crucial for the management of charge and energy transport in a range of organic optoelectronic devices.

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INTRODUCTION

strongly depends on molecular design such as planarity of conjugated backbones, alkyl chain lengths, interfacial tension between the backbone and side chains, and side chain structures,5−8 as well as processing parameters such as thermal or solvent annealing. Regioregular P3HT, which is the most studied organic semiconductors,9,10 has a coplanar rigid polymer backbone in its thin film with high density of hexyl chains, facilitating the formation of a lamellar packing structure. Furthermore, a delocalized π-conjugation along the polymer backbone and a

The advantages of easy fabrication, low cost, and compatibility with flexible and lightweight plastic substrates have promoted the development of many promising conjugated polymers for applications, such as field-effect transistors (FETs) and solar cells.1−3 The solution processable conjugated polymers consisting of a rigid backbone with flexible side-chains tend to self-organize and form microphase-separated structures (e.g., lamellar phase for poly(3-hexylthiophene) (P3HT) and hexagonal phase for poly(9,9-dioctylfluorene-2,7-diyl)) because of a repulsive interaction between the polymer backbones and the side-chains.4 These characteristics often affect charge or exciton transport in optoelectronic devices. However, the molecular packing structure in these conjugated polymers © XXXX American Chemical Society

Received: October 31, 2014 Revised: January 31, 2015

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DOI: 10.1021/cm503773j Chem. Mater. XXXX, XXX, XXX−XXX

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Chemistry of Materials

The reaction mixture was kept at −78 °C for 30 min and warm up to −10 °C. After cooling to −50 °C, 5.50 mL (5.50 mmol) of trimethyltin chloride solution (1.0 M in THF) was added, and the reaction mixture was warmed to room temperature and stirred for 3 h. The mixture was quenched with 20 mL of water and extracted with diethyl ether. The organic extraction wad dried with anhydrous magnesium sulfate and evaporated in vacuo. Recrystallization of the residue from hexane yields the desired product. (0.65 g, yield 48%). 1H NMR (500 MHz, CDCl3) δ 0.44 (s, 18H), 7.41 (s, 2H), 8.27 (s, 2H). 13 C NMR (125 MHz, CDCl3) δ −8.27, 115.15, 131.03, 138.64, 141.38, 141.82. HRMS (ESI) m/z = 515.9195 (M+), calcd m/z = 515.9201. Poly(3,4-dihexyl thiophene-2,2′:5,6′-benzo[1,2-b:4,5-b′]dithiophene) (PDHBDT). Tris(dibenzylideneacetone)dipalladium(0) (13.7 mg, 0.015 mmol) and tri(o-tolyl)phosphine (18.3 mg, 0.06 mmol) was added to a mixture of 2,5-dibromo-3,4-dihexylthiophene (123.1 mg, 0.30 mmol) and 2,6-bis(trimethyltin)benzo[1,2-b:4,5b′]dithiophene (154.8 mg, 0.30 mmol) in 4.0 mL of chlorobenzene under argon. The reaction mixture was frozen and degassed under vacuum three times. The reaction mixture was stirred for 16 h at 90 °C. Crude polymers were precipitated in methanol and washed with acetone and hexane. The polymer was redissolved in 100 mL of chloroform along with a Pd(0) scavenger31 and stirred for 4 h at 70 °C. The polymer was collected by reprecipitation and filtration from methanol. After washing the final product using Soxhlet apparatus with methanol and acetone, the polymer PDHBDT was recovered with chloroform. The polymer was further purified with column chromatography on silica gel with chloroform to give PDHBDT as an orange solid (112 mg, 85%). GPC: Mn = 28 kg/mol, Mw = 107 kg/ mol, PDI = 3.82; 1H NMR (400 MHz, CDCl3) δ 0.91 (br, 6H), 1.35 (m, 12H), 1.66 (br, 4H), 2.85 (br, 4H), 7.07 (bs, 2H), 7.15 (bs, 2H); anal. calcd for C26H30S3: C 71.18, H 6.89; found C 70.90, H 6.63. Instrumentation. 1H and 13C NMR spectra were recorded using Merc 400 or INOVA 500 spectrometers in CDCl3 at 293 K. Gel permeation chromatography (GPC) was performed in tetrahydrofuran (THF). Molecular weight was determined by size exclusion chromatography with Viscotec GPC system (Model 350 HTGPC) in THF. The calibration curve was made with monodipersed polystyrene standards. UV−vis absorption spectra were recorded with a UV−vis spectrophotometer (Cary 6000i) at room temperature. A thin film for UV−vis in solid state was prepared by spin coating on glass from a polymer solution in 1,2-dichlorobenzene. An optical gap was calculated from the onset of the longest wavelength absorption band of the thin film. The photoluminescence excitation/emission measurements were performed by using the Nanolog spectrofluorometer for nanomaterials (Horiba Jobin Yvon). A 450-W broadband cw xenon lamp and a monochromator supplied the excitation light in the range of 450−850 nm with 5 nm steps. The HOMO level of PDHBDT thin films was evaluated by photoelectron spectroscopy (PES) measurements (AC-2, manufactured by Riken Keiki). Electrochemical analysis was carried out by cyclic voltammetry using a CHI411 instrument from CH Instruments, Inc. The experiment was performed under a stream of argon in a saturated solution of 0.1 M tetra-n-butylammonium hexafluorophosphate (nBu4NPF6, from Strem Chemicals, Inc., recrystallized from ethanol) as a supporting electrolyte in anhydrous acetonitrile. The experiment was carried out using platinum (Pt) electrodes at a scan rate of 10 mV·s−1 against silver wire as a pseudoreference electrode at ambient temperature. The Fc/Fc+ redox couple was used as a reference oxidation potential for the electrochemical measurement. The polymer film was drop cast onto a Pt disk electrode from a chloroform solution. Thickness was measured by a Dektak 150 profilometer (Veeco Metrology Group). AFM images were taken using tapping mode (light tapping regime) using a Multimode AFM (Veeco). Grazing incidence X-ray scattering (GIXS) measurements were performed at the Stanford Synchrotron Radiation Lightsource (SSRL; Menlo Park, CA) using beamline 11−3. The incidence angle was chosen in the range of 0.10°−0.12° to optimize the signal tobackground ratio. The data were distortion-corrected (angle-dependent image distortion introduced by planar detector surface) before

strong π−π interchain stacking result in effective charge and exciton transport within lamellar planes. In particular, this strong π−π stacking of the planar polymer backbone has been a key structural feature of many high charge carrier mobility conjugated polymers, leading to intensive studies of materials containing 3-alkylthiophene building blocks.11−16 On the other hand, the poly(3,4-dialkylthiophene) systems have been less explored.17,18 Recently, we demonstrated that 3,4-disubsituted poly(alkylthiophenes), which do not show distinct π−π stacking, can be used to make FETs and bulk heterojunction (BHJ) solar cells that rival benchmark P3HT devices.19 Specifically, we found that the incorporation of 3,4dialkylthiophene into the backbone is a critical method to finetune the packing structure of the corresponding polymer, since comonomers can be easily varied for Stille polymerization with a 3,4-dialkylthiophene. While the 3,4-dialkylthiophene unit may induce twists in the polymer chain, the strategy involving a fused aromatic ring structure is frequently used for extended πconjugation, resulting in high performance in optoelectronic devices.13−15,20−22 Thus, the combination of 3,4-dihexylthiophene and a rigid building block, benzo[1,2-b:4,5-b′]dithiophene (BDT), can expand the understanding of tuning polymer packing structures, optoelectronic properties, and device performance in thin films, specifically, less explored 3,4disubstituted polythiophene systems. Even though continuous development of rigid-rod conjugated polymers has improved their electrical performance in the devices, the relationship between the molecular packing structure and the optoelectrical properties has been less extensively studied. This is partly due to the lack of tunability in molecular packing structures and the availability of polymer systems with various easily accessible molecular packing structures. Herein, we report a study of the impact of temperaturedependent molecular packing structures on charge carrier mobility in FETs and power conversion efficiency (PCE) in solar cells using poly[3,4-dihexyl thiophene-2,2′:5,6′-benzo[1,2b:4,5-b′]dithiophene] (PDHBDT) comprising of benzo[1,2b:4,5-b′]dithiophene (BDT) and 3,4-dihexylthiophene units. In addition, exciton diffusion in organic solar cells has been suggested to depend on structural order in the conjugated molecules or polymers.23−29 Yet, the correlation between exciton diffusion length and molecular packing structure in rigid-rod conjugated polymers remains ambiguous. Here, we observed the exciton diffusion length of PDHBDT depends on the polymer molecular packing. This study provides new insights into the control and ultimately the tunability of the exciton diffusion length in conjugated polymer systems, which is crucial for the management of energy transport in a wide range of important organic electronic devices.

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EXPERIMENTAL SECTION

Materials. 2,5-Dibromo-dihexylthiophene was prepared according to the reported procedures.30 THF was purified through the Pure SolMD standard design solvent purification system, Innovative Technology Inc. Benzo[1,2-b:4,5-b′]dithiophene, trimethyltin chloride solution (1.0 M in THF), chlorobenzene, tris(dibenzylideneacetone)dipalladium(0), and tri(o-tolyl)phosphine were purchased from Aldrich and TCI America and used without further purification. C60 and PC71BM were purchased from Nano-C inc. Synthesis. . 2,6-Bis(trimethylstannyl)benzo[1,2-b:4,5-b′]dithiophene (1). Benzo[1,2-b:4,5-b′]dithiophene (0.50 g, 2.63 mmol) was dissolved in 10 mL of anhydrous THF and cooled in an acetone/dry ice bath under argon protection. t-Butyllithium solution (3.70 mL, 6.31 mmol) was added dropwise with stirring at −78 °C. B

DOI: 10.1021/cm503773j Chem. Mater. XXXX, XXX, XXX−XXX

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Chemistry of Materials Scheme 1. Synthetic Route to PDHBDTa

a

Reagents and conditions: (i) t-BuLi, SnMe3Cl, THF, −78 °C to room temperature; (ii) Pd2(dba)3, tris(o-tolyl)phosphine, chlorobenzene, 90 °C.

Figure 1. (a) Normalized UV−vis spectra of PDHBDT in ODCB solution and films. (b) Cyclic voltammetry of polymer films in 0.1 M n-Bu4NPF6 in acetonitrile at a scan rate of 10 mV s−1. The oxidation potential was determined from the onset of the cyclic voltammogram. performing quantitative analysis on the images. GIXS images were collected on polymer samples that were deposited on Si substrates with a native oxide layer. Although polymer FET devices were fabricated on OTS-coated Si substrates, we believe this has little effect on the bulk film crystal structure. This is because even for pentacene films in the thin film phase, changing from native oxide Si to OTS did not change the unit cell parameters, and the molecular orientation changed very slightly.32 We expect that the difference between two weakly interacting surfaces would be smaller still for semicrystalline polymers. Calibration of reciprocal space and all the data analysis was performed using WxDiff software written by Mannsfeld. Typical exposure time was 180 s. The GIXS samples were prepared by spincoating the same polymer solutions used for fabricating devices onto silicon wafers at 2000 rpm for 30 s. Device Fabrication and Characterization. Polymer solutions in chlorobenzene for PDHBDT (0.5 wt %) were spin-coated onto OTSY treated SiO2/Si substrates at 2000 rpm for 30 s inside a nitrogen glovebox for thin-film transistors. The films were thermal annealed at 80 °C for 15 min. Additional films were annealed at specified temperature for 5 min with a slow cooling (3−4 °C/min) inside a glovebox. Subsequently, gold electrodes (40 nm) were deposited through shadow masks with a channel length (L) of 50 μm and channel width (W) of 1000 μm. OTFT transfer and output characteristics were recorded under dry N2 atmosphere (glovebox) and under ambient conditions using a Keithley 4200 semiconductor parametric analyzer (Keithley Instruments, Cleveland OH). PDHBDT/C60 bilayer solar cell devices were fabricated on glass substrates with the architecture of ITO/MoO 3 (8 nm)/ PDHBDT:C60/Ca(7 nm)/Al(100 nm), where the 40 nm-thick C60 was sublimed under vacuum on top of spin-cast PDHBDT films of 20

nm. PDHBDT/PC71BM -based BHJ solar cells were also fabricated in the device structure with ITO/PEDOT:PSS(30 nm)/polymer:PC71BM/Ca(7 nm)/Al(100 nm). PDHBDT (10 mg/mL) was codissolved with PC71BM in chlorobenzene with weight ratios of 1:1 at 65 °C. The polymer:PC71BM blends were spun at 600 rpm for 60 s on the ITO/PEDOT:PSS substrates inside the glovebox. Current− voltage characteristics were recorded in the dark and under simulated 1 sun AM 1.5 radiation with a Keithley 2400 source meter. Illumination was achieved with a 91160 300 W Oriel solar simulator equipped with a 6258 ozone-free Xe lamp and an air mass AM 1.5 G filter. The light intensity was calibrated using an NREL calibrated silicon photodiode with a KG5 filter that had a spectral mismatch factor error of less than 2% for the devices analyzed. External quantum efficiency (EQE) measurements were taken at short circuit using monochromated white light from a tungsten lamp which was modulated by an optical chopper. The current from the devices were measured as a function of wavelength and compared to the current obtained from a photodiode with a NIST traceable calibration photocurrent action spectrum.

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RESULTS AND DISCUSSION The synthetic route to PDHBDT is shown in Scheme 1. Comonomer 1, bis(trimethylstannyl)benzo[1,2-b:4,5-b′]dithiophene, was prepared via lithiation followed by quenching with trimethyl tin chloride in a moderate yield (85%). Polymerization from Stille coupling was carried out using dibromo-3,4-dihexyl thiophene 2 and comonomer 1 to afford PDHBDT. The crude polymer was treated with Pd(0) C

DOI: 10.1021/cm503773j Chem. Mater. XXXX, XXX, XXX−XXX

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Figure 2. 2D GIXD pattern of PDHBDT films annealed at various temperatures. Films were annealed at (a) 80, (b) 215, (c) 250, and (d) 270 °C. (e) Schematic diagram of molecular structure as a function of annealing temperature.

−4.80 eV below vacuum level as the potential of Fc/Fc+. The lowest unoccupied molecular orbital (LUMO) energy level of −2.96 eV was estimated from the HOMO and the optical band gap (Egopt = 2.31 eV). The polymer shows a good thermal stability with a 50% weight loss (Td) at 432 °C for PDHBDT determined by thermogravimetric analysis (TGA) (see the Supporting Information, Figure S1a). The thermal behavior of PDHBDT was characterized by differential scanning calorimetry (DSC). During the heating scan, an exothermic transition temperature at 286 °C was observed, which is attributed to the melting of the polymer backbone, and crystallization temperature (Tc) was observed at Tc = 242 °C (Supporting Information Figure S1b). To elucidate the impact of thermotropic transition on molecular packing in PDHBDT thin films, we performed twodimensional (2D) grazing incidence-angle X-ray diffraction (GIXD). We thermally annealed our polymer films at 80, 215, 250, and 270 °C, respectively. These temperatures were chosen to be in the range between the nominal phase transition temperatures (Tg1 = 162 °C and Tg2 = 225 °C) and the melting temperature (Tm = 286 °C), as shown in the DSC thermogram (see the Supporting Information, Figure S1b). Figure 2 shows GIXD images of the polymer films annealed at different temperatures. As shown in Figure 2a, the scattering pattern for the film annealed at 80 °C contains two relatively prominent Bragg peaks in the low Q region. These Bragg reflections correspond to lamellar packing, as summarized in the cartoon in Figure 2e. Both the (100) and the (200) reflections are distributed along a constant radius arc, demonstrating that the crystallite orientation distribution is broad. That is, the underlying crystal lattice is the same, but the orientation of the lattice with respect to the substrate differs for different crystallites. The scattering vector corresponding to the

scavenger in chloroform at 70 °C for 4 h. After reprecipitation in methanol, the polymer was further purified with column chromatography on silica gel with chloroform eluent to remove any residual palladium catalyst. After washing the final product using Soxhlet apparatus with methanol and acetone, the polymer was recovered with chloroform. PDHBDT is soluble in tetrahydrofuran (THF), chloroform (CF), chlorobenzene (CB), and 1,2-dichlorobenzene (ODCB) solvents. The synthesized polymer has the number-averaged molecular weight (Mn) of 28 kg/mol with polydispersity index (PDI) of 3.82 as determined by gel permeation chromatography (GPC) using THF as an eluent and polystyrenes as the standards. PDHBDT displayed absorption maximum (λmax) at 461 nm both in solution (in ODCB) and thin films (Figure 1a). There was little red-shift from solution to thin film, indicating little change in aggregation state. An optical band gap of 2.31 eV was estimated from the onset absorption (536 nm) of the polymer thin film. In comparison with Egopt of PDHTT (1.96 eV) containing a bithiophene unit,19 the fused ring containing PDHBDT polymer exhibited a larger Egopt. It has been known that the delocalization of electrons from fused thienothiophenes along the backbone is less favorable than from a single thiophene ring, due to the larger resonance stabilization energy of the fused ring over the single thiophene ring.13 This reduced delocalization along the backbone may be the reason for deepening of the PDHBDT polymer HOMO and widening the band gap. The electrochemical properties of PDHBDT films were characterized with cyclic voltammetry (CV) in 0.1 M nBu4NPF 6−acetonitrile solution (Figure 1b). The onset oxidation potentials vs Fc/Fc+ is 0.47 V, and thus the highest occupied molecular orbital (HOMO) energy level of this polymer is estimated as −5.27 eV, which was calculated using D

DOI: 10.1021/cm503773j Chem. Mater. XXXX, XXX, XXX−XXX

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out-of-plane d spacing changes from 26 Å in the hexagonal case to 15 Å, which suggests, when compared to the larger lamellar spacing in the film annealed at 80 °C, that side-chains in adjacent lamellar may be interdigitated. The crystalline morphology of the films then shows classic edge-on, though PDHBDT has a relatively high PDI of 3.82.33 Second, there is a vertical streaking of the Bragg reflection that peak near the horizon at Qxy ≈ 1.8 Å−1, which corresponds to a d-spacing of ∼3.5 Å. From a comparison to scattering patterns of many other conjugated polymer films, this peak corresponds to the repeat distance along an interchain pistack−one of the faster charge transfer directions in a semicrystalline polymer films. This pi-stacking distance is among the smallest for conjugated polymers, especially for conjugated polymers without strong donor−acceptor repeating units, suggesting that in the limit of vanishing disorder, hole transfer along the pi-stack within well-ordered regions of a crystallite is relatively efficient. Several (likely mixed index) Bragg peaks are also present, to which we have as-yet been unable to assign Miller indices. This supports our conclusion that the side chains are interdigitated. Because the hexagonal lattice parameters in the film annealed at 250 °C are unchanged relative to the film annealed at 215 °C, we believe that the scattering pattern of the 250 °C film corresponds to a coexistence of hexagonal-packed and lamellarpacked solid phases. That is, there is no evidence for a superstructure, where large crystals are packed in a hexagonal lattice with the interior of the lattice itself containing lamellae. We would expect that if additional order developed within the hexagonally packed crystallites, the hexagonal lattice parameters would be significantly perturbed with respect to the lattice of the film annealed at 215 °C. Moreover, using time-resolved photoluminescence measurements, we found that the average excited state lifetimes at 550 nm for films annealed at 80, 215, and 270 °C are 127.7, 90.9, and 97.4 ps, respectively (see the Supporting Information, Figure S2). The decrease in lifetime is also consistent with increased intermolecular interaction between adjacent polymer chains, leading to larger rates of nonradiative relaxation. It is interesting to note that the absorption spectrum of the polymer film annealed at 270 °C is shifted to the red by 10 nm compared to the film annealed at 80 °C (Figure 3a). We attribute this to the fact that annealing at the higher temperature leads to planarization of the polymer backbone, which is responsible for giving rise to the lamellar packing and

lamellar (100) reflection yields a d-spacing of 21.5 Å (Table 1), which is consistent with the side-chain length of the polymer and shows that there is no side-chain interdigitation. Table 1. Summary of d-Spacing and π−π Stacking Distance of PDHBDT upon Annealing annealing temp. (°C)

disordered packing (Å)

80 215 250 270

21.5

hexagonal packing (Å) 26.0 26.0

lamellar packing (Å)

π−π stacking distance (Å)

15.0 15.0

3.5 3.5

In stark contrast, Figure 2b shows that annealing the polymer film at 215 °C results in the appearance of relatively sharp Bragg peaks with a well-defined orientation, which cannot be rationalized with lamellar molecular packing. We find that a nearly hexagonal crystal lattice is consistent with this scattering pattern; we have labeled the relevant Miller indices on the diffraction image. The orientation of the lattice is such that the (−001) direction is normal to the film surface, a packing arrangement that can be approximated by a lying-down closepacked cylinder array. The nearly hexagonal lattice constants a and b are 26.8 and 25.5 Å, respectively. At the moment we are unable to determine the repeat distance along the long cylinder axis, i.e. the length of the lattice vector normal to the a−b plane. We envision the polymer backbone extending along the long cylinder axis; therefore, the repeat distance along the long cylinder axis is related to the effective monomer length (roughly 9.92 Å). The difficulty associated with identifying Bragg peaks corresponding to this periodicity is likely due to a small electron density variation (contrast) along this direction and to the disorder that is frequently present in polymer films. Interestingly, annealing the film at 250 °C results in a scattering pattern that contains additional peaks that are absent from the 215 °C pattern. Annealing at 270 °C just short of the melting temperature results in the nearly complete disappearance of hexagonal peaks, and the peaks present in the 250 °C film that are not due to the hexagonal lattice are now significantly more intense. The peaks that are present are again consistent with lamellar packing, though the lamellar peak positions are larger than the film that was annealed at 80 °C. There are a number of significant diffraction changes that are visible in the scattering patterns for the 215 and 270 °C films. First, the position of Bragg peaks near the Qz axis differ: the

Figure 3. (a) UV−vis spectra of PDHBDT films annealed at 80, 215, 250, and 270 °C; and (b) in-plane GIXD curves upon annealing. E

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Figure 4. AFM topographs of PDHBDT films annealed at (a) 80, (b) 215, (c) 250, and (d) 270 °C.

improved order along the chain−both effects that are expected to alter the electronic structure of the crystalline polymer regions and lead to a lowering of the optical band gap. These results are consistent with the X-ray scattering images (Figure 3b), because the π-stacking Bragg reflection is absent in the film annealed at 80 °C. There have been several reports on thermotropic phase transitions of the rigid-rod polymers with flexible alkyl side chains.6−8,34 Poly(p-biphenylene) terephthalate with hexadecyloxy side chains exhibited a layered mesophase at a lower temperature, while the polymer showed a transition to hexagonal mesophase at a higher temperature.8 A theoretical modeling study on the microstructure formation also revealed the transition to hexagonal packing from lamellar as temperature increasing.7 In contrast to the reports, PDHBDT showed predominantly lamellar scattering patterns at higher annealing temperature. This may be due to the strong π−π interaction facilitated by the more planar polymer backbone induced by annealing at a higher temperature. The hexagonal packing structures have been observed mostly with polymers having bulky side chains that drive the backbone twisting, such as in poly(9,9-dialkylfluorene)35 and regioregular poly(3-ethylhexylthiophene).36 The rich molecular packing structures of PDHBDT make it a good candidate to gain insight into the influence of phase behavior on the device performance. As shown in Figure 4, AFM height images of the 80 °C annealed film showed many pinhole-like voids, which implies poor interdomain connectively, resulting in low charge carrier transport performance (vide infra). The PDHBDT film annealed at higher temperature (215 and 250 °C) exhibited nodular shape domains with reduced pinholes and gives better inter domain connectivity. The 270 °C annealed film showed more fibrillar shape domains, which form more interconnected crystalline grains from the substantially stronger intermolecular interactions, resulting in high charge carrier mobility (vide infra). As expected, the various packing motifs influence the charge transport in FET devices. The charge transport properties of the spin-coated PDHBDT films were evaluated using bottomgate, top-contact FETs built on n-doped SiO2 (300 nm)/Si wafers modified with a crystalline octadecyltrimethoxysilane (OTS-Y) monolayer.37 The PDHBDT thin films were prepared by spin-casting from 0.5 wt % solutions in chlorobenzene. Table 2 summarizes the average transistor characteristics at various annealing temperatures tested in an inert atmosphere. Representative I−V transfer and output curves are shown in Figure 5. The FET mobility (μ) monotonically improved with increasing annealing temperature and the more developing π−π stacking of polymer backbone (see the Supporting Information, Figure S3). After annealing at 80 °C, the initial mobility (μ) was 2.3 × 10−4 cm2/(V s) for PDHBDT with an on/off ratio in the order of 1 × 102. After annealing at 215 °C, 1 order of magnitude larger mobility (6.0

Table 2. Summary of Top-Contact OFETs (W/L = 20, L = 50 um) Tested in an Inert Atmosphere annealing temp. (°C) 80 215 250 270

μavg (cm2/(V s)) −4

on/off −5

2.3 × 10 ± 5.1 × 10 6.0 × 10−3 ± 3.4 × 10−4 7.7 × 10−2 ± 9.8 × 10−3 0.22 ± 0.11

1 1 1 1

× × × ×

2

10 103 104 106

VT (V) −2.0 −13 −22 −16

± ± ± ±

0.5 2.6 3.6 8.3

× 10−3 cm2/(V s)) was observed in the hexagonal films, compared to the initial μ. The average μ of 0.077 cm2/(V s) was observed for 250 °C annealed film, which started to show the (010) peak from π−π stacking. We conclude that hexagonal microstructure (annealed at 215 °C) exhibits a higher carrier mobility than that of the random microstructure due to better interdomain connectivity with reduced pinholes as shown in Figure 4a, b. Moreover, the 250 °C annealed films show the (010) peak from π−π stacking in Figure 2c and well-organized crystalline domains in Figure 4c, thereby allowing much higher carrier mobility compared to the random microstructure as depicted in Table 2. It is worth noticing that the annealed PDHBDT films at 270 °C resulted in maximum μ up to 0.46 cm2/(V s) (μavg = 0.22 cm2/(V s)) with an on/off ratio of 1 × 106. The improved μ is attributed to the ordered lamellar packing with closer π−π stacking distance of 3.5 Å upon annealing at 270 °C as shown in the GIXD measurement. To investigate the impact of molecular packing controlled by the phase behavior in PDHBDT thin films on solar cell performance of PDHBDT/C60 bilayer heterojunction films, we observed photovoltaic characteristics in an inert atmosphere. The current density−voltage (J−V) characteristics under AM1.5G are shown in Figure 6a as a function of annealing temperature in the range from 80 to 270 °C. To gain a better interfacial contact of the bilayer devices, we prepared the device structure ITO/MoO3(8 nm)/PDHBDT:C60/Ca(7 nm)/Al(100 nm), where the 40 nm-thick C60 was sublimed under vacuum on top of spin-cast and annealed PDHBDT films of 20 nm as seen in the schematic representation of Figure 6a. C60 was chosen to avoid dissolving PDHBDT films by the solvent. We found that the short-circuit current (Jsc) in PDHBDT/C60 bilayer solar cells decreased with an increase of annealing temperature from 80 to 215 °C, but there is no notable change in open circuit voltage (Voc) irrespective of annealing temperatures. In contrast, Jsc was unchanged for the devices annealed at 270 °C relative to those annealed at 215 °C, but Voc dramatically decreased to 0.59 V, which is lower than other devices. Based on these current−voltage characteristics of PDHBDT/C60 bilayer solar cells, the PCE monotonically decreased as a function of annealing temperature (Figure 6b), because of changes in Jsc and Voc. The maximum PCE of 1.25% was for devices annealed at 80 °C. This result shows the randomly oriented phase is more effective in photovoltaic F

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Figure 5. Representative transfer and output curves of organic thin film transistor fabricated with PDHBDT thin film annealed at 270 °C tested in the inert atmosphere (VD = −100 V).

Figure 6. (a) Current−voltage characteristics and (b) PCE of PDHBDT/C60 bilayer heterojunction solar cells as a function of annealing temperature.

Figure 7. (a) Change of Voc of PDHBDT/C60 bilayer heterojunction solar cells upon annealing and (b) energy diagram at the interface of PDHBDT/C60 bilayer proved by UV−vis spectrophotometer and photoelectron spectroscopy (see the Supporting Information, Figure S4).

UV−vis spectrophotometer and photoelectron spectroscopy (PES), respectively (see the Supporting Information, Figure S4). Interestingly, it was found that Eg gradually decreased from 2.33 to 2.26 eV with a difference of 0.07 eV upon annealing (Table 3 and the Supporting Information, Figure S4a), which comes from a change in the HOMO level (see the Supporting

devices than the hexagonal or lamellar phases with a higher molecular ordering. In particular, to get a better understanding of the difference in Voc induced by thermotropic phase transition (Figure 7a), we have also investigated the band gap (Eg) and HOMO level of PDHBDT thin films as a function of annealing temperature by G

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are able to give us a sense for how molecular packing structure affects the photovoltaic characteristics, this is insufficient to directly correlate device data with crystalline nanostructure in the PDHBDT thin film. Exciton diffusion is characterized by a diffusion length LD, which is the characteristic distance that excitons are able to diffuse during the excited-state lifetime. As such, LD largely determines the number of excitons that can contribute to the photocurrent and consequently to the device efficiency. To elucidate the impact of polymer crystal structure on the exciton diffusion length in the film, we performed commonly used thickness-dependent photoluminescence (PL) quenching experiments.23−25 For this experiment, as seen in Figure 9a, we

Table 3. Summary of Band Gap (Eg), HOMO, LUMO, Voc, and PCE for PDHBDT-Based Bilayer Heterojunction Solar Cells upon Annealing annealing temp. (°C)

Eg (eV)

HOMO (eV)

LUMO (eV)

Voc (V)

PCEavg (%)

80 150 215 250 270

2.33

5.20 5.16 5.16 5.16 5.14

2.87 2.87

0.67 0.65 0.64

1.25 1.15 1.0

2.88

0.59

0.9

2.29 2.26

Information, Figure S4b). On the basis of these findings, a proposed energy diagram at the interface PDHBDT/C60 bilayer can be depicted in Figure 7b, and we speculate that the deeper HOMO level of PDHBDT thin films annealed at 80 °C may be due to more severe backbone twisting and their resulting different interchain interactions. Further, as seen in GIXD results, this can be explained with randomly oriented molecular packing in the PDHBDT thin films. On the contrary, a red-shift in UV−vis spectra, i.e., lower Eg, in annealing for PDHBDT thin films reflects the strong intermolecular interaction between PDHBDT polymer chains. The minimum Eg is in the lamellar phase molecular packing in films annealed at 270 °C. This decrease of Eg with a fixed LUMO level results in the lower Voc in PDHBDT/C60 bilayer solar cell, as shown in Figure 7b. The relevant optoelectrical and photovoltaic parameters obtained from Figure 6 and Figure 7 are summarized in Table 3. It is also interesting to note that dark current density JD (measured via hole-only device as shown from inset in Figure 8) is enhanced with increased annealing temperature as shown

Figure 9. (a) Schematic geometries for getting exciton diffusion length (LD) using PDHBDT or PDHBDT/C60 bilayer thin films. (b) Relative PL intensity ratio of PDHBDT with vs film thickness for films grown on 10 nm of C60 on glass. The lines are fits yielding LD = 6.0 ± 0.8 nm for random phase PDHBDT (annealed at 80 °C) black line, LD =12.5 ± 1.1 nm for hexagonal phase PDHBDT (annealed at 215 °C) red line, and LD =8.6 ± 0.5 nm for lamellar phase PDHBDT (annealed at 270 °C) solid line. The excitation wavelength was = 440 nm.

fabricated two kinds of samples based on PDHBDT and PDHBDT/C60 bilayer thin films, where PDHBDT film with different thickness in the range from 8 to 20 nm was prepared on precleaned glass substrates, and annealed at various temperatures to develop different molecular packing structures. Subsequently, 10 nm of C60 was coated on top of PDHBDT films to produce the quenching interface. Finally, the PL spectra were obtained. Further, assuming that the exciton density is proportional to the PL signal, the experimental data were fitted to a model in which the steady-state exciton diffusion23,24 is described by

Figure 8. Experimental dark-current densities of PDHBDT films as a function of annealing temperature for a hole-only device.

in Figure 8, which is another factor that can lead to low Voc. Moreover, Jsc in bilayer solar cells prepared from PDHBDT film annealed at 80 °C, where this film represents randomly oriented phase, is higher than both the hexagonal phase and the lamellar phase with more ordered molecular packing. This suggests that exciton and/or charge transport in the PDHBDT films with the broad orientation distribution of polymer crystallites may be more efficient along out-of-plane direction.38 However, this is not the case for PDHBDT thin films showing hexagonal or lamellar phase molecular packing, because most crystallites pack edge-on: therefore, transport is more efficient primarily in-plane. Although the variable Eg and HOMO level

L (1 − e−2d / LD) PL1 =1− D LD = PL 2 d(1 + e−2d / LD)

DEτE

(1)

where PL1 is exciton density in the presence of C60 and PL2 is exciton density in the absence of C60. d, DE, and τE are the H

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of 3.33% was obtained in nonannealed film, whereas thermal annealed solar cells at 110 °C for 10 min showed decreased PCE of 2.78%. It has been reported that PCBM transition temperature (and cold crystallization) in P3HT/PCBM blends is usually around 110 °C.39 Thus, we have chosen 110 °C as an annealing temperature for BHJ solar cells for the thiophene based PDHBDT polymer. Because of the limited absorption (Egopt = 2.31 eV), Jsc was limited to 6.09 mA/cm2; however, a high Voc of 0.873 V was observed because of deep lying HOMO (−5.27 eV from CV measurement), resulting in an optimized PCE of 3.33%. Although the external quantum efficiency (EQE) spectrum shows charge collection in the limited absorption range; moderately high quantum efficiency (57% at 400 nm) was observed in PDHBDT-based solar cells.

thickness of the donor layer, the diffusion coefficient, and the exciton lifetime, respectively. This model assumes a constant exciton density as a function of vertical position in the film. We have used the thin film absorption coefficient of 0.0175 nm−1 and the Beer−Lambert law to estimate the difference in exciton density generated near the front surface, through which light is incident, and the back (exit) surface. For a 20 nm thin film at the PL excitation wavelength, we find a difference of 30%, and 16% at 10 nm thickness. We find very similar absorption coefficients for films annealed at different temperatures. Thus, the largest thickness accumulates the largest error as far as fitting to the model specified in the paper is concerned (eq 1). Although the exciton density is not entirely constant throughout the film, we believe that the relatively small difference−particularly at smaller thicknesses−makes the constancy of exciton generation a reasonable assumption for comparing different annealing conditions. We note that PDHBDT thin film displayed thickness-dependent PL quenching when forming a contact with C60 as a quenching layer. The PL quenching in PDHBDT thin film with isotropic crystallite orientation (annealed at 80 °C) is larger than that in hexagonal or lamellar phase molecular packing (see the Supporting Information, Figure S5). Moreover, the PL intensity of the PDHBDT films on glass exhibits a linear dependence with PDHBDT thickness and this suggests that to a first approximation, optical interference effects can be ignored.25 In the analysis that follows, we make two further assumptions: that (a) the complex refractive index at the excitation wavelength change little for different annealing temperatures, and that (b) if some amount of C60 interpenetration into the bottom polymer layers takes place, it does so to a similar extent for all of our semicrystalline polymer films. Since the bilayer film was not subjected to further annealing post-C60 deposition, we believe that the latter assumption is reasonable. However, it certainly makes a contribution to the error bars of the calculated exciton diffusion length. The ratio PL1/PL2 (which means quenching efficiency) is shown as a function of the thickness of the PDHBDT films in Figure 9b. Fitting of the ratio PL1/PL2 for randomly oriented PDHBDT films yields LD = 12.5 ± 1.1 nm, whereas the hexagonal phase and the lamellar phase PDHBDT films exhibit LD = 6.0 ± 0.8 nm and LD = 8.6 ± 0.5 nm, respectively. We note that the trend of exciton diffusion length is the same as that observed for the excited-state lifetime. This data suggest that an isotropic crystallite orientation may correlate with a larger exciton diffusion length. It is possible that a primarily face-on orientation (π-stacking direction normal to substrate) would lead to a further increase in the vertical diffusion length. However, without additional information on the crystallite volume fraction, the crystallite connectivity, and the extent of C60 interpenetration, we are unable to draw a stronger connection between the orientation distribution and mesoscopic exciton diffusion. The photovoltaic properties of the PDHBDT-based on bulk heterojunction (BHJ) solar cells were also investigated in the device structure ITO/PEDOT:PSS/polymer:PC71BM/Ca/Al, with the active layers spun from chlorobenzene. Optimized polymer blends were obtained with a weight ratio of 1:1 (PDHBDT:PC71BM) and polymer concentration of 10 mg/ mL. Although device characteristics are summarized in the Supporting Information, Table S1, J−V curves and EQE spectra are shown in Supporting Information, Figure S6. The best PCE

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CONCLUSION In conclusion, we have prepared disubstituted polythiophene, PDHBDT, and studied the impact of temperature-dependent molecular packing of PDHBDT thin films on electrical and photovoltaic characteristics. The mobilities (μ) showed a strong relationship to molecular packing structures derived by thermotropic transition. While we observed the maximum μ up to 0.46 cm2/(V s) (μavg = 0.22 cm2/(V s)) in an edge-on lamellar packed structure, the broad orientation distribution of polymer crystallites is more effective to PDHBDT/C60 bilayer photovoltaic characteristics than hexagonal or ordered lamellar phase. This observation corresponds to the enhanced dark current density (JD) and the decreased LD upon annealing, leading to lower Voc and Jsc in the bilayer devices. We believe that above findings provide insight into the tuning electrical and photovoltaic properties depending on polymer packing structures, which can be effective strategy for optimizing organic optoelectronic devices.

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ASSOCIATED CONTENT

S Supporting Information *

TGA, DSC, time-resolved PL dynamics, transfer/output curves, UV−vis, photoelectron spectroscopy, thickness-dependent photoluminescence quenching, and photovoltaic properties. This material is available free of charge via the Internet at http://pubs.acs.org

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Phone: 1-650-723-2419. Fax: 1650-723-9780. Author Contributions †

S.K. and D.H.K. contributed equally to this work.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Center for Advanced Molecular Photovoltaics, Award KUS-C1-015-21, made by King Abdullah University of Science and Technology (KAUST). GIXS measurements were carried out at the Stanford Synchrotron Radiation Lightsource, a national user facility operated by Stanford University on behalf of the U.S. Department of Energy, Office of Basic Energy Sciences. S.K. acknowledges financial support by Korea Railroad Research Institute through the project “Development of Improvement I

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Technology of Railroad Environment (PK1504C)” D.H.K. acknowledges financial support by the Center for Advanced Soft-Electronics under the Global Frontier Project (CASE2014M3A6A5060932) and the Basic Science Research Program (2014R1A1A1005933) of the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT, and Future Planning. The photoluminescence measurements (A.M.N. and N.K.) were carried out under funding from the Energy Frontier Research Center “Molecularly Engineered Energy Materials (MEEMs)” funded by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract DE-SC0001342:001.

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