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Highly Crystalline Films of PCPDTBT with Branched Side Chains by Solvent Vapor Crystallization: Influence on Opto-Electronic Properties Florian S. U. Fischer, Daniel Trefz, Justus Back, Navaphun Kayunkid, Benjamin Tornow, Steve Albrecht, Kevin G. Yager, Gurpreet Singh, Alamgir Karim, Dieter Neher, Martin Brinkmann, and Sabine Ludwigs* Conjugated copolymers with alternating donor and acceptor units in the backbone have attracted great attention recently due to their superior charge transport properties and high solar cell efficiencies compared with conventional conjugated polymers such as poly(3-hexylthiophene) (P3HT). Understanding and controlling the semicrystalline morphology of these polymers remains a challenge. A highly interesting polymer class is based on the donor cyclopentadithiophene and the acceptor benzothiadiazol (CPDT and BT). While PCPDTBT (poly{[4,4bis-alkyl-cyclopenta-(2,1-b;3,4-b′)dithiophen]-2,6-diyl-alt-(2,1,3benzo-thiadiazole)-4,7-diyl}) based on linear C16-alkyl side chains shows strong crystallization and high mobilities[1–4] the crystallization tendency of PCPDTBT changes dramatically when the linear side chains are replaced with branched alkyl side groups. These systems are reported to be much less crystalline. PCPDTBT with ethylhexyl side chains was first introduced by Brabec and co-workers[5] in 2006. Bulk heterojunction solar cells of PCPDTBT have reached power conversion efficiencies of up to 4.5%–5.5%.[6–9] For high efficiencies the addition of high boiling point solvent additives such as diiodooctane (DIO) or alkanedithiols during solution processing F. S. U. Fischer, D. Trefz, J. Back, Prof. S. Ludwigs IPOC-Functional Polymers University of Stuttgart Pfaffenwaldring 55, 70569, Stuttgart, Germany E-mail: [email protected] Dr. N. Kayunkid,[+] Dr. M. Brinkmann Institut Charles Sadron (UPR22) CNRS, 23 Rue du Loess 67034, Strasbourg Cedex 2, France B. Tornow, S. Albrecht, Prof. D. Neher University of Potsdam Institute of Physics and Astronomy Karl-Liebknecht-Str. 24-25, 14476, Potsdam, Germany Dr. K. G. Yager Brookhaven National Laboratory Upton, New York, USA Dr. G. Singh,[†] Prof. A. Karim Department of Polymer Engineering University of Akron Ohio, USA [+]Present Address: College of Nanotechnology, King Mongkut’s Institute of Technology Ladkrabang (KMITL), Bangkok 10520, Thailand [†]Present Address: IBM Almaden Research Center, San Jose, California, USA
DOI: 10.1002/adma.201403475
Adv. Mater. 2014, DOI: 10.1002/adma.201403475
of the bulk heterojunction is reported to lead to an improved blend morphology as well as increased crystallinity of the polymer.[9–11] GIWAXS (grazing incidence wide angle x-ray scattering) experiments by Nelson and co-workers,[10] Bazan and co-workers,[12] and Russell co-workers[13] on the morphology of the homopolymer films display only very few reflections. In all cases, a broad reflection at a q value of 1.6 Å−1 corresponding to a distance of 3.8 Å is indexed as π-stacking (010) direction with an edge-on orientation. A strong out-of-plane reflection (OOP)-(100) at distances of 10.5 Å (q = 0.6 Å−1) without additive and 12 Å (q = 0.51 Å−1) with solvent additive were found by Nelson and co-workers,[10] and Russel and co-workers.[13] Bazan and co-workers[12] reported in both cases, i.e., with and without solvent additive, a mixture of the polymorph with (100) reflections at 11.4 Å (q = 0.55 Å−1) and edge-on orientation and an additional polymorph with (100) at 12.6 Å (q = 0.50 Å−1) and face-on orientation. Along the polymer chain axis Nelson and co-workers[10] reported a reflection at 11.8 Å (q = 0.53 Å−1). The available data, however, do not allow a detailed crystal structure analysis. Regarding surface morphology films deposited following the typical literature protocol, i.e., chlorobenzene (CB) solutions containing 2 w% DIO, can be described as randomly oriented aggregate structures in the size range of 35–45 nm, Figure 1E and Figure S1A (Supporting Information). This sample preparation will be abbreviated as CB/DIO-film in the following. Interestingly, adding of DIO to a number of different processing solvents leads to similar surface structures, compare Figure S2 (Supportingt Information). In a recent publication, we could show that the use of 1-chloronaphthalene as processing solvent leads to fiber-like patterns that also seem to be made of these 35–45 nm aggregates, Figure S1B (Supporting Information).[14] Though these aggregates appear to be crystalline, both GIWAXS and electron diffraction show very low crystallinity. Here, we show an approach based on solvent vapor annealing that allows crystallization of even a marginally crystallizable polymer such as PCPDTBT with bulky branched alkyl side chains into well-defined semicrystalline morphologies. A tentative crystal structure of PCPDTBT is proposed—that is to our knowledge the first determination of a crystal structure from highly crystalline films and electron diffraction of a p-type donor–acceptor copolymer in the literature. We compare our highly crystalline films with CB/DIO-films following the wellestablished literature protocol for solar cells. Both absorption
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Figure 1. Morphologies after different preparation methods of thin PCPDTBT films (annealing protocol see Figure S4, Supporting Information). A)–D) PCPDTBT films after CB-annealing. A) POM image of a sample with only spherulites (substrate: Si-wafer). B) POM image of a sample with spherulites and terrace-like structures, inset: magnification of a terrace-like structure (SiOx-wafer). C) SEM image of the dendritic structure of a spherulite (Au-substrate). D) AFM height image of a terrace-like structure (SiOx-wafer). E) AFM height image of CB/DIO-film, inset: chemical formula of PCPDTBT: F) AFM height image of a thin PCPDTBT film created by CS2-annealing at RT (Si-wafer).
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spectroscopy and structural data suggest that we have prepared a highly stable new crystal structure of PCPDTBT that does not give any evidence for classical π-stacking between chains and therefore differs from literature data. We were further able to compare the influence of different crystal structures on the performance of bilayer solar cells. Exposure of solvent-cast films to solvent vapor atmospheres has become a good alternative to conventional temperature annealing protocols both for conjugated polymers and blends of conjugated polymers with electron acceptors (e.g., PCBM) for bulk heterojunction solar cells.[15–17] Most research groups use saturated vapor, e.g., by putting the films in a closed vial filled with solvent. Some other groups, including ours have built setups that allow control of the atmosphere over a range of solvent vapor pressures.[17–19] These setups have proven to be highly successful for directing block copolymer self-assembly in thin films.[20–22] Recently, we used solvent vapor protocols for the controlled swelling of P3HT to erase any pre-history of the film and induce recrystallization into defined spherulitic structures.[19,23,24] In this manuscript, we used CB and carbon disulfide (CS2) for solvent vapor annealing of initially rather amorphous films spincoated from CHCl3, Figure S3 (Supporting Information).[14] In the following first CB-annealed films and then CS2-annealed films will be discussed. CB-Annealed Films: A typical polarized optical microscopy (POM) image of a film recrystallized at vapor pressures between 73%–70% after swelling to the solution like state at 80%–90%, is shown in Figure 1A (detailed protocol in Figure S4A, Supporting Information). These samples have a strong birefringence and the Maltese crosses indicate radial growth from specific nucleation points. SEM images, Figure 1C and Figure S5 (Supporting Information) suggest formation of dendritic spherulites. The center of the spherulite is made of parallel lamellar subunits, which further outside splay and branch radially.[25] Both temperature-dependent POM-measurements of spherulite samples, Figure S9A (Supporting Information), and DSCdata locate the melting transition around 280 °C, Figure S9B (Supporting Information). Varying the crystallization conditions allows further tuning of the nucleation density as is shown in Figure S6 (Supporting Information). A similar behavior was observed recently for P3HT spherulites.[19] In addition to the spherulites some of the films reveal deep blue almost circular regions in POM, Figure 1B. These features become especially prominent at low nucleation densities (compare for example Figure S6A bottom left corner, Supporting Information). Atomic force microscopy (AFM) suggests a terrace-like character with flat terraces and distinct height steps of 15 ± 3 nm, Figure 1D and corresponding height profile in Figure S7 (Supporting Information). It has to be pointed out that the crystallization behavior is substrate independent: the thin-film structures were found on silicon wafers with natural and 300 nm thick SiOx layers, Au- and transparent substrates such as ITO, glass, and quartz. The coexistence of mostly spherulites with lamellar substructure and fewer terrace-like features could also be visualized with transmission electron microscopy (TEM), Figure 2A. In particular the terrace-like features reveal very distinct reflections up to the sixth order and high symmetries, Figure 2B, which strongly suggests high crystallinity.[10,12,13] The sequence
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of reflections in Figure 2B along the equator can be indexed as (h 0 0) and the lengths are calculated to 9.7 Å (2 0 0), 4.8 Å (4 0 0), and 3.2 Å (6 0 0). Along the meridian only one reflection of 6.2 Å was found and indexed as (0 2 0). All other reflections are listed and indexed in Table S1 (Supporting Information). It should be noted that only the even reflections are present and that three selection rules for the crystal structure can be identified: h k 0: h + k = 2n, h 0 0: h = 2n, and 0 k 0: k = 2n. These selection rules support an orthorhombic unit cell with a Pn2n space group i.e., a high symmetry structure. Taking into account the defined height steps (≈15 nm) evident from the AFM images (Figure 1D) we suggest that the polymer chains are standing perpendicular on the substrate in these terracelike areas and that the a–b plane of the film (crystal structure) lies in the plane of the substrate, Figure 2E right sketch.
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Figure 2. TEM/ED studies of CB-annealed PCPDTBT thin films. A) TEM bright field image of a CB-annealed film, with terrace-like (red, dashed circle) and spherulite regions with lamellar substructure (green, solid circle). B) TEM/ED pattern of a terrace-like region as shown in A). C) and D) TEM/ED pattern of the sublamellar region of the spherulites (C) and after tilting the lamellar region around the fiber axis (020) by 45°–50° (D). E) Sketches of the proposed orientation of PCPDTBT chains with standing (terrace-like, red) and tilted (spherulites, green) chains with respect to the substrate. Unit cell parameters are indicated.
The electron diffraction (ED) pattern of the spherulites shows less distinct reflections, Figure 2C. The meridian reflection labeled as (0 2 0) in the ED points along the direction of the lamellar/fiber axis and is the same as for the terrace-like regions. Upon tilting of the sample relative to the electron beam around the (0 2 0) axis, the meridian reflection (0 2 0) remains unchanged, whereas additional peaks appear. Closer examination shows that the same ED pattern appears as for the terrace-like structures at a tilting angle of 45°–50°, compare Figure 2B and D. This means that after the tilting around the b-axis the a–b plane lies perpendicular to the e−-beam. The identical reflections prove that the spherulites crystallize in the same crystal structure as the terrace-like areas and all reflections visible in Figure 2C belong to the same polymorph. The polymer chains are just tilted by 45°–50° relative to the substrate as schematically shown in the left sketch in Figure 2E. GIWAXS measurements further support the findings from ED and are discussed in the supplementary information (Figures S10, 11, and Table S1, Supporting Information). CS2-Annealed Films: Furthermore, information on the crystal structure could be gained from the CS2-annealed samples that show fibers with widths in the 30–40 nm range, see Figure 1F. The fibers show no defined nucleation points and only weak birefringence is detected in POM, Figure S8 (Supporting Information), which strongly suggests that these samples are overall less crystalline. This is also supported by TEM/ED and GIWAXS showing a limited set of weak reflections (Figures S10C and S12, Supporting Information). However, since the reticular distances observed in these patterns are identical to those observed in the CB-annealed samples, it supports the same polymorph as the CB-annealed ones. One single reflection is detected OOP at q = 0.66 Å−1 (9.5 Å) that we index as (2 0 0) in accordance with the CB-annealed data, Figures S10C and S11 (Supporting Information). Additionally, we find in-plane a new reflection at q = 1.07 Å−1 (5.9 Å). This reflection can be indexed as (0 0 4) in the polymer chain direction and this fits also nicely with the data by Nelson and co-workers.[10] From polarized Raman experiments[37] of highly oriented CS2-annealed samples[24] we can identify an orientation of the polymer chains perpendicular to the fiber growth direction and parallel to the substrate, i.e., an in-plane orientation of the polymer chains on the substrate, see supplementary information (Figures S14–S17, Supporting Information). We note that while Russell and co-workers,[13] Bazan and co-workers,[12] and Nelson and co-workers[10,12,13] also report on an in-plane orientation of the PCPDTBT chains in CB/solvent additive films, their reflections do not coincide with the patterns obtained for the CB- and CS2-annealed films reported in this publication. Mainly, we do not see any evidence for efficient long-range π-stacking. This means that we have indeed prepared a new polymorph of PCPDTBT. Tentative Crystal Structure: Combining the diffraction data from CB- and CS2-annealed samples we can complement the unit cell parameters of this polymorph: a = 19.3 Å, b = 12.4 Å, c = 23.6 Å, and α = β = γ = 90° (Figure 2 and Figure S13C, Supporting Information). The ED pattern in Figure 2B is considered as a crystal with standing chains in accordance to AFM data. This pattern shows high symmetry and as such, is strikingly different from that observed for a polymer such as P3HT.[26] In particular, this pattern does not reveal a lattice
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characterized by a large lamellar period corresponding to layers of π-stacked chains separated by layers of side chains (usually in the range 15–25 Å), and a short 3.5–4.0 Å period corresponding to a regular π-stacking of the chains. In the new PCPDTBT polymorph, both the a and b cell parameters are large and difficult to reconcile with a short π-stacking distance. It is the fuzzy reflection indexed as (5 1 0) that could indicate a π-stacking period whereas in the present case its origin may as well be related to some characteristic stacking distance of alkyl side chains in the large unit cell of the new polymorph. Therefore, the structure of PCPDTBT cannot be analyzed similarly to the layered structure of P3HT and PBTTT.[26,27] The high symmetry of the observed ED pattern reminds that observed for much more rigid polymers such as polyfluorenes for which the chain packing is characterized by the absence of long-range π-stacking and the pairing of chains into dimer-like motives in the unit cell.[28] A tentative structural model of this new polymorph of PCPDTBT was therefore constructed along this line and is displayed in Figure S13 (Supporting Information). To obtain a reasonable crystal density of ≈1.3 g cm–3, four chains comprising two comonomers (cyclopentadithiophene–benzothiadiazole units) must be placed in the unit cell. Electron diffraction patterns corresponding to the [0 0 1] zone were calculated using a chain without the branched alkyl side chains. Figure S13 (Supporting Information) compares both calculated
and experimental ED patterns. Despite the absence of the branched side chains, one can reproduce the highly symmetric experimental ED pattern with the intense 2 0 0 and 0 2 0 reflections seen in Figure 2B. Additional structural studies are in progress to determine the role of the side chains on the packing of the PCPDTBT backbones in this polymorph and to obtain additional information on the exact backbone conformation. Optical and Opto-Electronic Properties: The different crystal structures that are obtained either from CB/DIO-films or from solvent-annealing (CB- and CS2-annealed) are also reflected by the different absorption behavior of the samples, Figure 3B and Figure S18 (Supporting Information). The rather prominent 800 nm band in the CB/DIO-films that is commonly attributed to π-stacking in the literature[10,13,29] is not obtained for the solvent-annealed samples. This is in accordance to the formation of a new polymorph. The films with the new polymorph even exhibit a new blue-shifted maximum at 680 nm, the origin of which is subject of ongoing studies. We made the interesting experimental observation that the 800 nm absorption in the film vanishes gradually when the CB/DIO-films are heated from room temperature up to 100 °C (Figure S18B, Supporting Information). This finding strongly supports the statement by Peet and co-workers[29] that these aggregates at 800 nm are not thermodynamically stable. GIWAXS experiments on temperature-annealed films further showed that upon heating a loss of
Figure 3. A) Schematic setup of the prepared bilayer solar cells and the in-plane chain orientation of the PCPDTBT films. B) Normalized absorption spectra of PCPDTBT thin films. C) and D) Bilayer solar cell results with polymer layers prepared by CS2-annealing and CB/DIO spincoating. C) J–V curves under AM 1.5 G illumination with 100 mW cm–2 (solid) are compared with J–V curves with pure polymer excitation using monochromatic LED light at 740 nm (dashed) with the intensity adjusted to give ca. the same Jsc as under AM 1.5G illumination. D) EQE spectra as function of applied bias ranging from –1.6 to 0.4 V (left scale) compared with the absorption spectra of the pure polymer layers (grey, right scale).
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ordering in the (010) occurred, which Russell and co-workers[13] attribute to a more equilibrium crystalline state of the PCPDTBT chains at higher temperatures. The solvent-annealed films on the other hand are much more stable and only show changes above 100 °C, Figures S18C and S18D (Supporting Information). Literature proposes that the morphology of conjugated polymers strongly affects the device performance in optoelectronic devices.[13,16,30] This is particularly true for solar cells, where the orientation and packing of the conjugated main chains was shown to have a huge influence on the efficiency of the generation and extraction of photogenerated charges. We have shown that PCPDTBT is a system in which we are able to fabricate layers of the very same material but with different structural parameters, tuned via solvent annealing. This provides us with the opportunity to study the performance of solar cells in relation to the polymer morphology. Here, a comparison of the CS2annealed and the CB/DIO-films is of great interest as both form homogenous films with an in-plane polymer chain orientation but different crystal structures, one without and one with classical π-stacking, respectively. In addition, layers with as-prepared low crystallinity CHCl3-films were included into the study. All studies were performed on bilayer solar cells (device design see Figure 3A), with evaporated C60 on top of the polymer films to prevent penetration of the fullerene into the polymer layers or changes in the morphology as could occur using solvent coating. The J–V curves of the bilayer solar cells are shown in Figure 3C (CS2-annealed and CB/DIO-film) and Figure S19 (CHCl3-film, Supporting Information) illuminated with either the entire solar spectrum (solid line) or by monochromatic 740 nm LED light that only excites the polymer (dashed line). Despite a large difference in the polymer morphology, the shape of the J–V characteristics is quite comparable. In particular, we see a similar field dependence of the photocurrent at reverse bias and comparable absolute photocurrents. The performance data are summarized in Table S2 (Supporting Information). The external quantum efficiency (EQE) spectra as function of applied bias are compared with the absorption spectra in Figure 3D and Figure S20 (Supporting Information). The high-energy region (400–540 nm) is dominated by the C60 absorption whereas the low-energy region (600–900 nm) mirrors the polymer absorption, with distinct changes of the onset and peak position depending on the polymer morphology (compare vertical dashed lines Figure 3D). This allows us to conclude that the polymer chains involved in the photogeneration of free charges exhibit the same packing and environment as the chains monitored by optical absorption spectroscopy and structural analysis. Importantly, EQEs in the polymer absorption region show almost no field dependence between –1.6 and 0.2 V for all morphologies tested. A stronger field dependence is, however, observed when exciting the fullerene layer, which we attribute to the field-induced split-up of CT excitations in pure fullerene domains.[31] Using time delayed collection field (TDCF) measurements, discrimination between the field dependence of charge generation and extraction is possible. In Figure S21 (Supporting Information) the J–V curve is compared with the EQE and the total extracted charge Qtot from the TDCF measurements with primary C60 (480 nm) or polymer (650 nm) excitation. Again, we find very similar properties, irrespective of the protocol to prepare
the polymer layer, with a very weak field dependence of generation under direct polymer excitation. This leads to the important conclusion that in these PCPDTBT/C60 bilayer devices, free charge generation is not controlled by the structural properties of the hole-transporting polymer phase. One might argue that this result is expected as the delocalization of electrons in pure and extended fullerene phases was proposed to lead to efficient split-up of CT excitons at the polymer: fullerene heterojunction. However, work by Köhler and co-workers[32,33] revealed a large variation in the field dependence of generation in polymer/C60 bilayer devices depending on the type of polymer used. These authors proposed the photocurrent field dependence to correlate mainly with the conjugation length (and with that hole delocalization) of the polymer backbone. Our work demonstrates that for bilayer solar cells with the very same donor polymer but different layer morphologies, the yield and field dependence of free charge generation is, surprisingly, the same. In summary, we have prepared unusual highly crystalline PCPDTBT films by controlled solvent vapor annealing protocols applied to precast thin films. Both CB- and CS2-annealing induce a new polymorph of PCPDTBT, only differing in the orientation of the polymer chains on the substrate. Our electron diffraction patterns and preliminary structural model of the new polymorph do not give evidence for any efficient π-stacking and differ from the reflections reported in literature for films deposited by the typical device protocol from CB in the presence of solvent additives (CB/DIO-films). In contrast to CB/DIO-films, our solvent-annealed films seem to be quasi-equilibrated and are highly stable up to the melting point that was determined to be at 280 °C. The difference in crystal structure is also reflected by differing absorption spectra between the solvent-annealed and the CB/DIO-films. Using the possibility to induce certain polymorphs we were able—to our knowledge for the first time—to prepare bilayer solar cells with evaporated C60 differing only in the polymer morphology. Our data suggest charge generation in such cells to be independent of the polymer morphology. This finding is intriguing as earlier theoretical and experimental work suggested that the CT split up is largely affected by the polymer conjugation length, the hole mobility, or the orientation of the polymer chains with respect to the fullerene phase.[32–36] Furthermore experiments with defined device architectures and full control of the polymer morphology are needed to arrive at a conclusive view of the parameters determining free charge generation across a polymer: fullerene heterojunction. Last but not the least, we want to highlight the potential of solvent vapor crystallization as an approach applicable to any semicrystalline polymer. The feasibility of using low boiling point solvents as well as high boiling solvents as demonstrated here offers significant control over the crystallization behavior. This will help the community to better understand and control crystallization of functional semicrystalline polymers and to establish structure function relationships.
Supporting Information Supporting Information is available from the Wiley Online Library or from the author.
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Acknowledgements Experimental work by M. Plachetta, the scientific discussions with C. Ruiz Delgado and B. Lotz and the assembly of the TDCF setup by J. Kurpiers are gratefully acknowledged. We thank the DFG for funding within the priority program “Elementary processes in organic solar cells” (SPP-1355), the Emmy Noether Program, the EU for funding within Smartonics, the BMBF for funding within the project PVCOMB (FKZ 03IS2151D) the European Community for funding via the Interreg IV-A program (C25 Rhin-Solar). Research carried out in part at the Center for the Functional Nanomaterials and the National Synchrotron Light Source, the Brookhaven National Laboratory, which are supported by the U.S. Department of Energy, Office of Basic Energy Sciences, under Contract No. DE-AC02–98CH10886. A. Karim and G. Singh acknowledge support by the U.S. Department of Energy, Division of Basic Energy Sciences under contract No. DE-FG02–10ER4779 for the research towards this work. Received: July 31, 2014 Revised: September 19, 2014 Published online:
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[11] M. Morana, M. Wegscheider, A. Bonanni, N. Kopidakis, S. Shaheen, M. Scharber, Z. Zhu, D. Waller, R. Gaudiana, C. J. Brabec, Adv. Funct. Mater. 2008, 18, 1757. [12] J. T. Rogers, K. Schmidt, M. F. Toney, E. J. Kramer, G. C. Bazan, Adv. Mater. 2011, 23, 2284. [13] Y. Gu, C. Wang, T. P. Russell, Adv. Energy Mater. 2012, 2, 683. [14] F. S. U. Fischer, K. Tremel, A.-K. Saur, S. Link, N. Kayunkid, M. Brinkmann, D. Herrero-Carvajal, J. T. L. Navarrete, M. C. R. Delgado, S. Ludwigs, Macromolecules 2013, 46, 4924. [15] S. Miller, G. Fanchini, Y.-Y. Lin, C. Li, C.-W. Chen, W.-F. Su, M. Chhowalla, J. Mater. Chem. 2008, 18, 306. [16] B. Gholamkhass, P. Servati, Org. Electron. 2013, 14, 2278. [17] E. Verploegen, C. E. Miller, K. Schmidt, Z. Bao, M. F. Toney, Chem. Mater. 2012, 24, 3923. [18] H. Chen, S. Hu, H. Zang, B. Hu, M. D. Dadmun, Adv. Funct. Mater. 2013, 23, 1701. [19] E. J. W. Crossland, K. Rahimi, G. Reiter, U. Steiner, S. Ludwigs, Adv. Funct. Mater. 2011, 21, 518. [20] B. Kim, S. W. Hong, S. Park, J. Xu, S.-K. Hong, T. P. Russell, Soft Matter 2011, 7, 443. [21] S. Park, B. Kim, J. Xu, T. Hofmann, B. M. Ocko, T. P. Russell, Macromolecules 2009, 42, 1278. [22] S. Ludwigs, A. Böker, A. Voronov, N. Rehse, R. Magerle, G. Krausch, Nat. Mater. 2003, 2, 744. [23] E. J. W. Crossland, K. Tremel, F. S. U. Fischer, K. Rahimi, G. Reiter, U. Steiner, S. Ludwigs, Adv. Mater. 2012, 24, 839. [24] F. S. U. Fischer, K. Tremel, M. Sommer, E. J. W. Crossland, S. Ludwigs, Nanoscale 2012, 4, 2138. [25] G. R. Strobl, The Physics of Polymers, Springer, Berlin 1996. [26] N. Kayunkid, S. Uttiya, M. Brinkmann, Macromolecules 2010, 43, 4961. [27] D. M. DeLongchamp, R. J. Kline, E. K. Lin, D. A. Fischer, L. J. Richter, L. A. Lucas, M. Heeney, I. McCulloch, J. E. Northrup, Adv. Mater. 2007, 19, 833. [28] M. Brinkmann, Macromolecules 2007, 40, 7532. [29] J. Peet, N. S. Cho, S. K. Lee, G. C. Bazan, Macromolecules 2008, 41, 8655. [30] H. Chen, Y.-C. Hsiao, B. Hu, M. D. Dadmun, Adv. Funct. Mater. 2014, 24, 5129. [31] G. F. Burkhard, E. T. Hoke, Z. M. Beiley, M. D. McGehee, J. Phys. Chem. C 2012, 116, 26674. [32] C. Schwarz, S. Tscheuschner, J. Frisch, S. Winkler, N. Koch, H. Bässler, A. Köhler, Phys. Rev. B 2013, 87, 155205. [33] C. Schwarz, H. Bässler, I. Bauer, J.-M. Koenen, E. Preis, U. Scherf, A. Köhler, Adv. Mater. 2012, 24, 922. [34] D. Caruso, A. Troisi, Proc. Natl. Acad. Sci. U. S. A. 2012, 109, 13498. [35] T. Offermans, S. C. J. Meskers, R. A. J. Janssen, Chem. Phys. 2005, 308, 125. [36] J. R. Tumbleston, B. A. Collins, L. Yang, A. C. Stuart, E. Gann, W. Ma, W. You, H. Ade, Nat. Photonics 2014, 8, 385. [37] X. Wang, H. Azimi, H.-G. Mack, M. Morana, H.-J. Egelhaaf, A. J. Meixner, D. Zhang, Small 2011, 7, 2793.
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Highly Crystalline Films of PCPDTBT with Branched Side Chains by Solvent Vapor Crystallization: Influence on Opto-Electronic Properties Florian S. U. Fischer, Daniel Trefz, Justus Back, Navaphun Kayunkid, Benjamin Tornow, Steve Albrecht, Kevin G. Yager, Gurpreet Singh, Alamgir Karim, Dieter Neher, Martin Brinkmann, and Sabine Ludwigs* Conjugated copolymers with alternating donor and acceptor units in the backbone have attracted great attention recently due to their superior charge transport properties and high solar cell efficiencies compared with conventional conjugated polymers such as poly(3-hexylthiophene) (P3HT). Understanding and controlling the semicrystalline morphology of these polymers remains a challenge. A highly interesting polymer class is based on the donor cyclopentadithiophene and the acceptor benzothiadiazol (CPDT and BT). While PCPDTBT (poly{[4,4bis-alkyl-cyclopenta-(2,1-b;3,4-b′)dithiophen]-2,6-diyl-alt-(2,1,3benzo-thiadiazole)-4,7-diyl}) based on linear C16-alkyl side chains shows strong crystallization and high mobilities[1–4] the crystallization tendency of PCPDTBT changes dramatically when the linear side chains are replaced with branched alkyl side groups. These systems are reported to be much less crystalline. PCPDTBT with ethylhexyl side chains was first introduced by Brabec and co-workers[5] in 2006. Bulk heterojunction solar cells of PCPDTBT have reached power conversion efficiencies of up to 4.5%–5.5%.[6–9] For high efficiencies the addition of high boiling point solvent additives such as diiodooctane (DIO) or alkanedithiols during solution processing F. S. U. Fischer, D. Trefz, J. Back, Prof. S. Ludwigs IPOC-Functional Polymers University of Stuttgart Pfaffenwaldring 55, 70569, Stuttgart, Germany E-mail: [email protected] Dr. N. Kayunkid,[+] Dr. M. Brinkmann Institut Charles Sadron (UPR22) CNRS, 23 Rue du Loess 67034, Strasbourg Cedex 2, France B. Tornow, S. Albrecht, Prof. D. Neher University of Potsdam Institute of Physics and Astronomy Karl-Liebknecht-Str. 24-25, 14476, Potsdam, Germany Dr. K. G. Yager Brookhaven National Laboratory Upton, New York, USA Dr. G. Singh,[†] Prof. A. Karim Department of Polymer Engineering University of Akron Ohio, USA [+]Present Address: College of Nanotechnology, King Mongkut’s Institute of Technology Ladkrabang (KMITL), Bangkok 10520, Thailand [†]Present Address: IBM Almaden Research Center, San Jose, California, USA
DOI: 10.1002/adma.201403475
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of the bulk heterojunction is reported to lead to an improved blend morphology as well as increased crystallinity of the polymer.[9–11] GIWAXS (grazing incidence wide angle x-ray scattering) experiments by Nelson and co-workers,[10] Bazan and co-workers,[12] and Russell co-workers[13] on the morphology of the homopolymer films display only very few reflections. In all cases, a broad reflection at a q value of 1.6 Å−1 corresponding to a distance of 3.8 Å is indexed as π-stacking (010) direction with an edge-on orientation. A strong out-of-plane reflection (OOP)-(100) at distances of 10.5 Å (q = 0.6 Å−1) without additive and 12 Å (q = 0.51 Å−1) with solvent additive were found by Nelson and co-workers,[10] and Russel and co-workers.[13] Bazan and co-workers[12] reported in both cases, i.e., with and without solvent additive, a mixture of the polymorph with (100) reflections at 11.4 Å (q = 0.55 Å−1) and edge-on orientation and an additional polymorph with (100) at 12.6 Å (q = 0.50 Å−1) and face-on orientation. Along the polymer chain axis Nelson and co-workers[10] reported a reflection at 11.8 Å (q = 0.53 Å−1). The available data, however, do not allow a detailed crystal structure analysis. Regarding surface morphology films deposited following the typical literature protocol, i.e., chlorobenzene (CB) solutions containing 2 w% DIO, can be described as randomly oriented aggregate structures in the size range of 35–45 nm, Figure 1E and Figure S1A (Supporting Information). This sample preparation will be abbreviated as CB/DIO-film in the following. Interestingly, adding of DIO to a number of different processing solvents leads to similar surface structures, compare Figure S2 (Supportingt Information). In a recent publication, we could show that the use of 1-chloronaphthalene as processing solvent leads to fiber-like patterns that also seem to be made of these 35–45 nm aggregates, Figure S1B (Supporting Information).[14] Though these aggregates appear to be crystalline, both GIWAXS and electron diffraction show very low crystallinity. Here, we show an approach based on solvent vapor annealing that allows crystallization of even a marginally crystallizable polymer such as PCPDTBT with bulky branched alkyl side chains into well-defined semicrystalline morphologies. A tentative crystal structure of PCPDTBT is proposed—that is to our knowledge the first determination of a crystal structure from highly crystalline films and electron diffraction of a p-type donor–acceptor copolymer in the literature. We compare our highly crystalline films with CB/DIO-films following the wellestablished literature protocol for solar cells. Both absorption
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Figure 1. Morphologies after different preparation methods of thin PCPDTBT films (annealing protocol see Figure S4, Supporting Information). A)–D) PCPDTBT films after CB-annealing. A) POM image of a sample with only spherulites (substrate: Si-wafer). B) POM image of a sample with spherulites and terrace-like structures, inset: magnification of a terrace-like structure (SiOx-wafer). C) SEM image of the dendritic structure of a spherulite (Au-substrate). D) AFM height image of a terrace-like structure (SiOx-wafer). E) AFM height image of CB/DIO-film, inset: chemical formula of PCPDTBT: F) AFM height image of a thin PCPDTBT film created by CS2-annealing at RT (Si-wafer).
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spectroscopy and structural data suggest that we have prepared a highly stable new crystal structure of PCPDTBT that does not give any evidence for classical π-stacking between chains and therefore differs from literature data. We were further able to compare the influence of different crystal structures on the performance of bilayer solar cells. Exposure of solvent-cast films to solvent vapor atmospheres has become a good alternative to conventional temperature annealing protocols both for conjugated polymers and blends of conjugated polymers with electron acceptors (e.g., PCBM) for bulk heterojunction solar cells.[15–17] Most research groups use saturated vapor, e.g., by putting the films in a closed vial filled with solvent. Some other groups, including ours have built setups that allow control of the atmosphere over a range of solvent vapor pressures.[17–19] These setups have proven to be highly successful for directing block copolymer self-assembly in thin films.[20–22] Recently, we used solvent vapor protocols for the controlled swelling of P3HT to erase any pre-history of the film and induce recrystallization into defined spherulitic structures.[19,23,24] In this manuscript, we used CB and carbon disulfide (CS2) for solvent vapor annealing of initially rather amorphous films spincoated from CHCl3, Figure S3 (Supporting Information).[14] In the following first CB-annealed films and then CS2-annealed films will be discussed. CB-Annealed Films: A typical polarized optical microscopy (POM) image of a film recrystallized at vapor pressures between 73%–70% after swelling to the solution like state at 80%–90%, is shown in Figure 1A (detailed protocol in Figure S4A, Supporting Information). These samples have a strong birefringence and the Maltese crosses indicate radial growth from specific nucleation points. SEM images, Figure 1C and Figure S5 (Supporting Information) suggest formation of dendritic spherulites. The center of the spherulite is made of parallel lamellar subunits, which further outside splay and branch radially.[25] Both temperature-dependent POM-measurements of spherulite samples, Figure S9A (Supporting Information), and DSCdata locate the melting transition around 280 °C, Figure S9B (Supporting Information). Varying the crystallization conditions allows further tuning of the nucleation density as is shown in Figure S6 (Supporting Information). A similar behavior was observed recently for P3HT spherulites.[19] In addition to the spherulites some of the films reveal deep blue almost circular regions in POM, Figure 1B. These features become especially prominent at low nucleation densities (compare for example Figure S6A bottom left corner, Supporting Information). Atomic force microscopy (AFM) suggests a terrace-like character with flat terraces and distinct height steps of 15 ± 3 nm, Figure 1D and corresponding height profile in Figure S7 (Supporting Information). It has to be pointed out that the crystallization behavior is substrate independent: the thin-film structures were found on silicon wafers with natural and 300 nm thick SiOx layers, Au- and transparent substrates such as ITO, glass, and quartz. The coexistence of mostly spherulites with lamellar substructure and fewer terrace-like features could also be visualized with transmission electron microscopy (TEM), Figure 2A. In particular the terrace-like features reveal very distinct reflections up to the sixth order and high symmetries, Figure 2B, which strongly suggests high crystallinity.[10,12,13] The sequence
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of reflections in Figure 2B along the equator can be indexed as (h 0 0) and the lengths are calculated to 9.7 Å (2 0 0), 4.8 Å (4 0 0), and 3.2 Å (6 0 0). Along the meridian only one reflection of 6.2 Å was found and indexed as (0 2 0). All other reflections are listed and indexed in Table S1 (Supporting Information). It should be noted that only the even reflections are present and that three selection rules for the crystal structure can be identified: h k 0: h + k = 2n, h 0 0: h = 2n, and 0 k 0: k = 2n. These selection rules support an orthorhombic unit cell with a Pn2n space group i.e., a high symmetry structure. Taking into account the defined height steps (≈15 nm) evident from the AFM images (Figure 1D) we suggest that the polymer chains are standing perpendicular on the substrate in these terracelike areas and that the a–b plane of the film (crystal structure) lies in the plane of the substrate, Figure 2E right sketch.
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Figure 2. TEM/ED studies of CB-annealed PCPDTBT thin films. A) TEM bright field image of a CB-annealed film, with terrace-like (red, dashed circle) and spherulite regions with lamellar substructure (green, solid circle). B) TEM/ED pattern of a terrace-like region as shown in A). C) and D) TEM/ED pattern of the sublamellar region of the spherulites (C) and after tilting the lamellar region around the fiber axis (020) by 45°–50° (D). E) Sketches of the proposed orientation of PCPDTBT chains with standing (terrace-like, red) and tilted (spherulites, green) chains with respect to the substrate. Unit cell parameters are indicated.
The electron diffraction (ED) pattern of the spherulites shows less distinct reflections, Figure 2C. The meridian reflection labeled as (0 2 0) in the ED points along the direction of the lamellar/fiber axis and is the same as for the terrace-like regions. Upon tilting of the sample relative to the electron beam around the (0 2 0) axis, the meridian reflection (0 2 0) remains unchanged, whereas additional peaks appear. Closer examination shows that the same ED pattern appears as for the terrace-like structures at a tilting angle of 45°–50°, compare Figure 2B and D. This means that after the tilting around the b-axis the a–b plane lies perpendicular to the e−-beam. The identical reflections prove that the spherulites crystallize in the same crystal structure as the terrace-like areas and all reflections visible in Figure 2C belong to the same polymorph. The polymer chains are just tilted by 45°–50° relative to the substrate as schematically shown in the left sketch in Figure 2E. GIWAXS measurements further support the findings from ED and are discussed in the supplementary information (Figures S10, 11, and Table S1, Supporting Information). CS2-Annealed Films: Furthermore, information on the crystal structure could be gained from the CS2-annealed samples that show fibers with widths in the 30–40 nm range, see Figure 1F. The fibers show no defined nucleation points and only weak birefringence is detected in POM, Figure S8 (Supporting Information), which strongly suggests that these samples are overall less crystalline. This is also supported by TEM/ED and GIWAXS showing a limited set of weak reflections (Figures S10C and S12, Supporting Information). However, since the reticular distances observed in these patterns are identical to those observed in the CB-annealed samples, it supports the same polymorph as the CB-annealed ones. One single reflection is detected OOP at q = 0.66 Å−1 (9.5 Å) that we index as (2 0 0) in accordance with the CB-annealed data, Figures S10C and S11 (Supporting Information). Additionally, we find in-plane a new reflection at q = 1.07 Å−1 (5.9 Å). This reflection can be indexed as (0 0 4) in the polymer chain direction and this fits also nicely with the data by Nelson and co-workers.[10] From polarized Raman experiments[37] of highly oriented CS2-annealed samples[24] we can identify an orientation of the polymer chains perpendicular to the fiber growth direction and parallel to the substrate, i.e., an in-plane orientation of the polymer chains on the substrate, see supplementary information (Figures S14–S17, Supporting Information). We note that while Russell and co-workers,[13] Bazan and co-workers,[12] and Nelson and co-workers[10,12,13] also report on an in-plane orientation of the PCPDTBT chains in CB/solvent additive films, their reflections do not coincide with the patterns obtained for the CB- and CS2-annealed films reported in this publication. Mainly, we do not see any evidence for efficient long-range π-stacking. This means that we have indeed prepared a new polymorph of PCPDTBT. Tentative Crystal Structure: Combining the diffraction data from CB- and CS2-annealed samples we can complement the unit cell parameters of this polymorph: a = 19.3 Å, b = 12.4 Å, c = 23.6 Å, and α = β = γ = 90° (Figure 2 and Figure S13C, Supporting Information). The ED pattern in Figure 2B is considered as a crystal with standing chains in accordance to AFM data. This pattern shows high symmetry and as such, is strikingly different from that observed for a polymer such as P3HT.[26] In particular, this pattern does not reveal a lattice
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characterized by a large lamellar period corresponding to layers of π-stacked chains separated by layers of side chains (usually in the range 15–25 Å), and a short 3.5–4.0 Å period corresponding to a regular π-stacking of the chains. In the new PCPDTBT polymorph, both the a and b cell parameters are large and difficult to reconcile with a short π-stacking distance. It is the fuzzy reflection indexed as (5 1 0) that could indicate a π-stacking period whereas in the present case its origin may as well be related to some characteristic stacking distance of alkyl side chains in the large unit cell of the new polymorph. Therefore, the structure of PCPDTBT cannot be analyzed similarly to the layered structure of P3HT and PBTTT.[26,27] The high symmetry of the observed ED pattern reminds that observed for much more rigid polymers such as polyfluorenes for which the chain packing is characterized by the absence of long-range π-stacking and the pairing of chains into dimer-like motives in the unit cell.[28] A tentative structural model of this new polymorph of PCPDTBT was therefore constructed along this line and is displayed in Figure S13 (Supporting Information). To obtain a reasonable crystal density of ≈1.3 g cm–3, four chains comprising two comonomers (cyclopentadithiophene–benzothiadiazole units) must be placed in the unit cell. Electron diffraction patterns corresponding to the [0 0 1] zone were calculated using a chain without the branched alkyl side chains. Figure S13 (Supporting Information) compares both calculated
and experimental ED patterns. Despite the absence of the branched side chains, one can reproduce the highly symmetric experimental ED pattern with the intense 2 0 0 and 0 2 0 reflections seen in Figure 2B. Additional structural studies are in progress to determine the role of the side chains on the packing of the PCPDTBT backbones in this polymorph and to obtain additional information on the exact backbone conformation. Optical and Opto-Electronic Properties: The different crystal structures that are obtained either from CB/DIO-films or from solvent-annealing (CB- and CS2-annealed) are also reflected by the different absorption behavior of the samples, Figure 3B and Figure S18 (Supporting Information). The rather prominent 800 nm band in the CB/DIO-films that is commonly attributed to π-stacking in the literature[10,13,29] is not obtained for the solvent-annealed samples. This is in accordance to the formation of a new polymorph. The films with the new polymorph even exhibit a new blue-shifted maximum at 680 nm, the origin of which is subject of ongoing studies. We made the interesting experimental observation that the 800 nm absorption in the film vanishes gradually when the CB/DIO-films are heated from room temperature up to 100 °C (Figure S18B, Supporting Information). This finding strongly supports the statement by Peet and co-workers[29] that these aggregates at 800 nm are not thermodynamically stable. GIWAXS experiments on temperature-annealed films further showed that upon heating a loss of
Figure 3. A) Schematic setup of the prepared bilayer solar cells and the in-plane chain orientation of the PCPDTBT films. B) Normalized absorption spectra of PCPDTBT thin films. C) and D) Bilayer solar cell results with polymer layers prepared by CS2-annealing and CB/DIO spincoating. C) J–V curves under AM 1.5 G illumination with 100 mW cm–2 (solid) are compared with J–V curves with pure polymer excitation using monochromatic LED light at 740 nm (dashed) with the intensity adjusted to give ca. the same Jsc as under AM 1.5G illumination. D) EQE spectra as function of applied bias ranging from –1.6 to 0.4 V (left scale) compared with the absorption spectra of the pure polymer layers (grey, right scale).
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ordering in the (010) occurred, which Russell and co-workers[13] attribute to a more equilibrium crystalline state of the PCPDTBT chains at higher temperatures. The solvent-annealed films on the other hand are much more stable and only show changes above 100 °C, Figures S18C and S18D (Supporting Information). Literature proposes that the morphology of conjugated polymers strongly affects the device performance in optoelectronic devices.[13,16,30] This is particularly true for solar cells, where the orientation and packing of the conjugated main chains was shown to have a huge influence on the efficiency of the generation and extraction of photogenerated charges. We have shown that PCPDTBT is a system in which we are able to fabricate layers of the very same material but with different structural parameters, tuned via solvent annealing. This provides us with the opportunity to study the performance of solar cells in relation to the polymer morphology. Here, a comparison of the CS2annealed and the CB/DIO-films is of great interest as both form homogenous films with an in-plane polymer chain orientation but different crystal structures, one without and one with classical π-stacking, respectively. In addition, layers with as-prepared low crystallinity CHCl3-films were included into the study. All studies were performed on bilayer solar cells (device design see Figure 3A), with evaporated C60 on top of the polymer films to prevent penetration of the fullerene into the polymer layers or changes in the morphology as could occur using solvent coating. The J–V curves of the bilayer solar cells are shown in Figure 3C (CS2-annealed and CB/DIO-film) and Figure S19 (CHCl3-film, Supporting Information) illuminated with either the entire solar spectrum (solid line) or by monochromatic 740 nm LED light that only excites the polymer (dashed line). Despite a large difference in the polymer morphology, the shape of the J–V characteristics is quite comparable. In particular, we see a similar field dependence of the photocurrent at reverse bias and comparable absolute photocurrents. The performance data are summarized in Table S2 (Supporting Information). The external quantum efficiency (EQE) spectra as function of applied bias are compared with the absorption spectra in Figure 3D and Figure S20 (Supporting Information). The high-energy region (400–540 nm) is dominated by the C60 absorption whereas the low-energy region (600–900 nm) mirrors the polymer absorption, with distinct changes of the onset and peak position depending on the polymer morphology (compare vertical dashed lines Figure 3D). This allows us to conclude that the polymer chains involved in the photogeneration of free charges exhibit the same packing and environment as the chains monitored by optical absorption spectroscopy and structural analysis. Importantly, EQEs in the polymer absorption region show almost no field dependence between –1.6 and 0.2 V for all morphologies tested. A stronger field dependence is, however, observed when exciting the fullerene layer, which we attribute to the field-induced split-up of CT excitations in pure fullerene domains.[31] Using time delayed collection field (TDCF) measurements, discrimination between the field dependence of charge generation and extraction is possible. In Figure S21 (Supporting Information) the J–V curve is compared with the EQE and the total extracted charge Qtot from the TDCF measurements with primary C60 (480 nm) or polymer (650 nm) excitation. Again, we find very similar properties, irrespective of the protocol to prepare
the polymer layer, with a very weak field dependence of generation under direct polymer excitation. This leads to the important conclusion that in these PCPDTBT/C60 bilayer devices, free charge generation is not controlled by the structural properties of the hole-transporting polymer phase. One might argue that this result is expected as the delocalization of electrons in pure and extended fullerene phases was proposed to lead to efficient split-up of CT excitons at the polymer: fullerene heterojunction. However, work by Köhler and co-workers[32,33] revealed a large variation in the field dependence of generation in polymer/C60 bilayer devices depending on the type of polymer used. These authors proposed the photocurrent field dependence to correlate mainly with the conjugation length (and with that hole delocalization) of the polymer backbone. Our work demonstrates that for bilayer solar cells with the very same donor polymer but different layer morphologies, the yield and field dependence of free charge generation is, surprisingly, the same. In summary, we have prepared unusual highly crystalline PCPDTBT films by controlled solvent vapor annealing protocols applied to precast thin films. Both CB- and CS2-annealing induce a new polymorph of PCPDTBT, only differing in the orientation of the polymer chains on the substrate. Our electron diffraction patterns and preliminary structural model of the new polymorph do not give evidence for any efficient π-stacking and differ from the reflections reported in literature for films deposited by the typical device protocol from CB in the presence of solvent additives (CB/DIO-films). In contrast to CB/DIO-films, our solvent-annealed films seem to be quasi-equilibrated and are highly stable up to the melting point that was determined to be at 280 °C. The difference in crystal structure is also reflected by differing absorption spectra between the solvent-annealed and the CB/DIO-films. Using the possibility to induce certain polymorphs we were able—to our knowledge for the first time—to prepare bilayer solar cells with evaporated C60 differing only in the polymer morphology. Our data suggest charge generation in such cells to be independent of the polymer morphology. This finding is intriguing as earlier theoretical and experimental work suggested that the CT split up is largely affected by the polymer conjugation length, the hole mobility, or the orientation of the polymer chains with respect to the fullerene phase.[32–36] Furthermore experiments with defined device architectures and full control of the polymer morphology are needed to arrive at a conclusive view of the parameters determining free charge generation across a polymer: fullerene heterojunction. Last but not the least, we want to highlight the potential of solvent vapor crystallization as an approach applicable to any semicrystalline polymer. The feasibility of using low boiling point solvents as well as high boiling solvents as demonstrated here offers significant control over the crystallization behavior. This will help the community to better understand and control crystallization of functional semicrystalline polymers and to establish structure function relationships.
Supporting Information Supporting Information is available from the Wiley Online Library or from the author.
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Acknowledgements Experimental work by M. Plachetta, the scientific discussions with C. Ruiz Delgado and B. Lotz and the assembly of the TDCF setup by J. Kurpiers are gratefully acknowledged. We thank the DFG for funding within the priority program “Elementary processes in organic solar cells” (SPP-1355), the Emmy Noether Program, the EU for funding within Smartonics, the BMBF for funding within the project PVCOMB (FKZ 03IS2151D) the European Community for funding via the Interreg IV-A program (C25 Rhin-Solar). Research carried out in part at the Center for the Functional Nanomaterials and the National Synchrotron Light Source, the Brookhaven National Laboratory, which are supported by the U.S. Department of Energy, Office of Basic Energy Sciences, under Contract No. DE-AC02–98CH10886. A. Karim and G. Singh acknowledge support by the U.S. Department of Energy, Division of Basic Energy Sciences under contract No. DE-FG02–10ER4779 for the research towards this work. Received: July 31, 2014 Revised: September 19, 2014 Published online:
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Adv. Mater. 2014, DOI: 10.1002/adma.201403475
