Article pubs.acs.org/Macromolecules
Crystallization-Driven Enhancement in Photovoltaic Performance through Block Copolymer Incorporation into P3HT:PCBM Blends Dargie Deribew,† Eleni Pavlopoulou,† Guillaume Fleury,*,† Célia Nicolet,† Cedric Renaud,† Sébastien-Jun Mougnier,† Laurence Vignau,‡ Eric Cloutet,† Cyril Brochon,† Fabrice Cousin,§ Giuseppe Portale,∥ Mark Geoghegan,⊥ and Georges Hadziioannou*,† †
Laboratoire de Chimie des Polymères Organiques (LCPO), Université de Bordeaux, CNRS UMR 5629, 16 Avenue Pey-Berland, F-33607 Pessac Cedex, France ‡ Laboratoire de l’Intégration du Matériau au Système (IMS), Université de Bordeaux, CNRS UMR 5218, 16, Avenue Pey-Berland, F-33607 Pessac Cedex, France § Laboratoire Léon Brillouin, CEA-CNRS, CEA Saclay, F-91191 Gif sur Yvette Cedex, France ∥ Netherlands Organization for Scientific Research (NWO), DUBBLE-CRG at the ESRF, Grenoble, France ⊥ Department of Physics and Astronomy, The University of Sheffield, Hicks Building, Hounsfield Road, Sheffield S3 7RH, United Kingdom S Supporting Information *
ABSTRACT: We report the increased crystallization of poly(3-hexylthiophene) (P3HT) in the donor−acceptor mixture of [6,6]-phenyl-C61-butyric acid methyl ester (PCBM) with P3HT by the addition of a block copolymer, P3HT-b-PI, where PI refers to polyisoprene. The photovoltaic performance of devices created using this blend is markedly improved by the addition of the diblock copolymer. We have characterized the structure of thin films of the P3HT-b-PI containing mixtures using optical microscopy, scanning force microscopy, UV−vis absorption spectroscopy, neutron reflectometry, and grazing incidence X-ray diffraction (GIXD). The GIXD data provide the information on the crystallinity of the films, the absorption data were used to confirm that the addition of the diblock was responsible for the increase in crystallization, neutron reflectometry data reveal a PCBM-rich region near the hole injection layer, and the two microscopy techniques revealed the structural effect of the crystallization at the surface of the films.
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INTRODUCTION Solar cells made of a photoactive conjugated polymer as the electron donor and a soluble fullerene derivative as the electron acceptor have shown great potential in the field of renewable energy technologies.1−4 Since the first demonstration of the photoinduced electron transfer from a conducting polymer to fullerene moieties in 1992,5 tailoring the properties of the active layer through macromolecular design and/or process optimization has led to considerable achievements as regards to device performance.6−9 However several issues still require attention before organic solar cells become efficient enough for commercialization. For the bulk heterojunction (BHJ) configuration that is formed upon blending an electron-donor and an electron-acceptor material, the components should phase-separate and form an interpenetrating network that exhibits maximum donor/acceptor interfacial area and characteristic domain size of about 10 nm.10−12 Thus, exciton dissociation at the interfaces is enhanced, and high power conversion efficiencies, PCE, can be achieved. The BHJ’s mesostructure is strongly correlated to the processing parameters during device fabrication. In the poly(3-hexylthiophene):[6,6]-phenyl-C61-butyric acid methyl ester, © 2013 American Chemical Society
P3HT:PCBM, blendthe most widely studied BHJthe solvent and the casting method,13−15 the composition of the donor:acceptor blend,16−18 the molecular weight of the photoactive material,19−22 and the annealing temperature23−29 are reported to strongly affect the final morphology of the phase-separated blend. For example, annealing the films at an appropriate temperature is known to increase P3HT crystallinity; P3HT chains organize into an arrangement of crystalline domains, the extent and local orientation of which can result in improved charge percolation pathways that are essential for efficient charge carrier harvesting and high PCE. Besides, the phase segregation of the bulk blend can change drastically upon heating due to significant interdiffusion of both materials30 which is driven by the miscibility or PCBM in the disordered P3HT domains.30,31 Therefore, both the overall mesostructure (inherent to the P3HT:PCBM phase separation) and the local ordering (inherent to the P3HT crystallinity and PCBM Received: October 10, 2012 Revised: April 2, 2013 Published: April 12, 2013 3015
dx.doi.org/10.1021/ma302128h | Macromolecules 2013, 46, 3015−3024
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and colleagues added 5% of a P3HT-b-poly(styrene-co-acrylate) copolymer with fullerenes chemically linked to the acrylate unit into a P3HT:PCBM blend and observed an improvement in PCE from 2.6% for the pristine device to 3.5%.51 Lee et al. studied the addition of a C60-end-capped P3HT in the P3HT:PCBM blend and observed that a long-term thermal stability of device performance can be achieved.52 The incorporation of poly(3-hexylthiophene)-b-poly(ethylene oxide), P3HT-b-PEO, in pure P3HT and P3HT:PCBM blends has also been reported,53 and a change in P3HT domain size with respect to copolymer concentration was observed. Poly(3hexylthiophene)-b-poly(4-vinylpyridine), P3HT-b-P4VP, has been tested as well, and the noncovalent supramolecular interactions between P4VP and PCBM were exploited in order to hamper PCBM aggregation and, consequently, improve significantly the performance of annealing-free P3HT:PCBM devices.54 Finally, Xiao et al. recently presented an integrated study of the effect of a polystyrene-b-poly(3-hexylthiophene), PS-b-P3HT, copolymer as interfacial compatibilizer in the P3HT:PCBM blend.55 Their work revealed that adding an optimized weight fraction of PS-b-P3HT copolymer increases the crystallinity of P3HT and homogenizes the vertical distribution of P3HT and PCBM in the active layer. They attributed these outcomes to the favorable interactions of the P3HT block with the P3HT homopolymer and the strong affinity of PS block toward PCBM leading to enhanced hole transport and charge extraction. All the studies mentioned above confirm the ability of block copolymers to modify the structure of the active layer, leading to improved device performance. However, the additional functionalities that the coil-like block of a P3HT-b-coil copolymer can bring to the structuration of the P3HT:PCBM:P3HT-b-coil blend has not been fully exploited yet. In the quest for a P3HT-b-coil copolymer that can induce the optimum morphology in the archetypical P3HT:PCBM photoactive blend, we explore herein the use of P3HT-bpolyisoprene. Polyisoprene, PI, is a low-Tg polymer (Tg = −71 °C for the PI-block studied herein), which means that at room temperature the PI chains exhibit high mobility. We chose to link a low molecular weight, Mn, PI-block (Mn = 2000 g mol−1) to a relatively high molecular weight P3HT-block (Mn = 25 000 g mol−1) to form the P3HT-b-PI copolymer that was incorporated in the P3HT:PCBM blend in order to ensure miscibility of the PI blocks within the P3HT phase. Thus, we hypothesize that the PI block can act as a guiding factor that promotes heterogeneous nucleation in the P3HT matrix. Simultaneously, by acting as a plasticizer, the PI block can increase the segmental mobility of P3HT that may favor crystallization through promoting rearrangements of the P3HT chains to their equilibrium state.56 Low quantities of the P3HT-b-PI copolymer (up to 10 wt %) were added in a 1:1 w/w P3HT:PCBM blend, and 60 nm thick films were prepared. An integrated structural study of the copolymer-containing films using optical microscopy, scanning force microscopy (SFM), grazing-incidence X-ray diffraction (GIXD), neutron reflectivity (NR), and UV−vis spectroscopy is presented, and the implications of the copolymer addition in the active layer morphology are discussed. In addition, differential scanning calorimetry (DSC) has been employed to determine the effect of the block copolymer on the crystallization properties of the polymer and subsequently those nucleating properties. Finally, the photovoltaic cell performance characteristics were related to the structural properties of the
agglomeration) are directly correlated to the performance of organic solar cells. As already mentioned, the formation of micro- and nanoscale mesostructures in the BHJs begins with spontaneous phase separation between the donor and the acceptor components during solvent evaporation. Although phase separation in those films is governed by thermodynamics,32 it is difficult to control experimentally, and the final structure contains phases with a distribution of sizes (due to coarsening and nucleation effects) while surface segregation and wetting also complicate the spatial distribution of materials. Controlling phase separation using block copolymers has been demonstrated to be a promising approach for tuning microphase separation in photovoltaic active layers.33,34 The amphiphilic nature of block copolymers and their ability to self-assemble35−39 into highly ordered and thermodynamically stable mesostructures render them attractive for tailoring BHJ morphology. Block copolymers used for this purpose comprise a rod-like block usually a conjugated semiconducting polymerand a coil-like block which brings additional functionalities to the system. Block copolymers with rod−rod architectures can also lead to well-defined p−n junctions;40 however, their use in photovoltaic applications has been limited so far due to the modest processability they exhibit owing to the strong rod−rod interactions.41,42 There are three ways to incorporate a block copolymer in a BHJ system. First of all, a donor-b-acceptor copolymer can be used alone to form a single, self-structured, active layer upon self-assembly. Hadziioannou and colleagues43,44 demonstrated this concept using a donor-b-acceptor copolymer with poly(pphenylenevinylene), PPV, being the donor block and a C60functionalized polystyrene being the acceptor block. Other molecules have also been tested;45 however, PCE remained low in all cases due to poor charge mobility and carrier losses in the coil phase.46 In a second approach, a copolymer that contains a donor block can be blended with a fullerene to form the BHJ.34,35,47 Yet, the PCE was far from the state-of-the-art performance of the archetypical P3HT:PCBM BHJ due to the presence of an important volume fraction of insulating material in the blend. The most promising approach for the use of block copolymers for photovoltaic applications has proved to be their utilization as compatibilizers in polymer:fullerene blends. This methodology is inspired by the improved compatibility obtained in common A:B polymer blends through the incorporation of an AB block copolymer.48 In a photovoltaic active layer, one of the blocks is commonly of the same nature as the donor semiconducting polymer while the second block can vary depending on the extra functionalities that are desirable. In this regard, several coil blocks (e.g., polyacrylates, polystyrene, poly(acrylonitrile), poly(4-vinylpyridine), perylene diimides (PDI), poly(methyl methacrylate), polyisoprene, poly(tert-butyl acrylate), poly(L-lactic acid)) could be covalently attached to a P3HT block to form a copolymer.49 Fréchet and co-workers reported the use of a diblock copolymer containing a P3HT-based block and a fullerene-containing block as a compatibilizer in order to lower the interfacial energy between the polymer and fullerene components in photovoltaic blends and, thus, to improve the stability of the thin film morphology upon thermal annealing.11 Rajaram et al. used a diblock copolymer to compatibilize the P3HT:PDI blend, and they achieved a 3-fold improvement in PCE compared to previously reported P3HT:PDI BHJ devices.50 In another report, Wudl 3016
dx.doi.org/10.1021/ma302128h | Macromolecules 2013, 46, 3015−3024
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Figure 1. Synthetic route for the fabrication of P3HT-b-PI block copolymers. The 25 kg mol−1 P3HT was freeze-dried in toluene and then dissolved in 40 mL of toluene at 50 °C to improve solubility. Five equivalents of lithium 2-methoxyethanolate per active center were added to the living polyisoprene. Then the last mixture was added in excess (2 equiv) on the P3HT, and the reaction was stirred at 50 °C for 3 days. The reaction was quenched with degassed methanol and the block copolymer was recovered by precipitation in methanol. The block copolymer was washed in a Soxhlet with pentane for 48 h to eliminate the excess of polyisoprene, then dissolved in chloroform, and recovered by precipitation in methanol. After drying under reduced pressure, the block copolymer was analyzed by SEC and 1H NMR. All reactions were carried out under reduce pressure using flame-dried glassware. The synthesis route is stated in Figure 1. Characterization of Poly(3-hexylthiophene)-b-polyisoprene. 1 H nuclear magnetic resonance (NMR) spectra were recorded using a Bruker AC-400 NMR at room temperature, and the samples were dissolved in deuterated chloroform CDCl3. SEC was performed in THF at 40 °C at a flow rate of 1 mL min−1 using a differential refractometer (Varian) and a UV−vis spectrophotometer (Varian) operating at 254 nm and using three TSKgel Tosoh columns (G4000HXL, G3000HXL, and G2000HXL). The elution times were converted into molecular weights using a calibration curve based on low-dispersity polystyrene standards. The molecular weight Mn and the dispersity Đ were analyzed by SEC. The regioregularity was calculated by integration of the corresponding peaks on the 1H NMR spectrum. Thin Film Preparation and Device Fabrication. First, the P3HT-b-PI copolymer was dissolved in o-dichlorobenzene, and appropriate amounts of the solution were introduced in a 1:1 (by weight) P3HT:PCBM solution in o-dichlorobenzene to obtain the desired copolymer concentrations in the blend (0−10 wt %) with respect to the amount of P3HT:PCBM. The solutions were stirred overnight at 50 °C to promote complete dissolution. For device fabrication the conventional architecture was used i.e. glass/ITO/ PEDOT:PSS/active layer/Al. The ITO-coated glass substrates were cleaned in an ultrasonic bath for 15 min using sequentially three different solvents; namely acetone, ethanol, and isopropanol. The substrates were dried and placed in a UV-ozone chamber for 15 min. A 50 nm thin (measured by an Alpha-step IQ Surface Profiler) layer of PEDOT:PSS was spin-coated at 5000 rpm and dried in a vacuum oven at 110 °C for 30 min. Next, the P3HT:PCBM:P3HT-b-PI solutions were filtered and spin-coated on the PEDOT:PSS layers to form the active layers. The thickness of the active layer is typically in the range of ∼60 nm. Finally, an 80 nm thick aluminum cathode was thermally deposited through a shadow mask at a pressure of 10−7 mbar. Thermal annealing treatment was performed after Al deposition on a temperature-controlled hot plate at 165 °C for 20 min. The devices were left to cool down to room temperature before further characterizations. All procedures after PEDOT:PSS deposition were performed in a nitrogen glovebox with oxygen and moisture levels
