- Email: [email protected]
Crystallographic structure and morphology of bithiophene-fluorene polymer nanocrystals
Polymer 52 (2011) 3368e3373
Contents lists available at ScienceDirect
Polymer journal homepage: www.elsevier.com/locate/polymer
Crystallographic structure and morphology of bithiophene-fluorene polymer nanocrystals Oliver Werzer a, b, Roland Resel a, *, Boril Chernev c, Harald Plank c, d, Michael M. Rothmann e, Peter Strohriegl e, Gregor Trimmel f, g, Arnaldo Rapallo h, William Porzio h a
Institute of Solid State Physics, Graz University of Technology, Petersgasse 16, A-8010 Graz, Austria Centre for Organic Electronics, University of Newcastle, Australia Centre for Electron Microscopy, Graz, Austria d Institute for Electron Microscopy, Graz University of Technology, Austria e Institute of Macromolecular Chemistry, University of Bayreuth, Germany f Institute for Chemistry and Technology of Materials, Graz University of Technology, Austria g Christian Doppler Laboratory for Nanocomposite Solar Cells, Graz University of Technology, Austria h Istituto per lo Studio delle Macromolecole, Consiglio Nazionale delle Ricerche, Milano, Italy b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 26 January 2011 Received in revised form 28 April 2011 Accepted 29 April 2011 Available online 6 May 2011
Nanocrystals of the polymer poly(9,9-dioctylfluorenyl-co-bithiophene) (F8T2) with a molecular weight of 3.2 kg/mol are grown in a para-xylene solution. The typical morphology of the crystals is needle like with typical widths of 50 nm and lengths of about 200 nm. The crystal structure and morphology are stable up to a temperature of 353 K. The structure solution is obtained by x-ray powder diffraction (XRD) pattern with data modelling by a stochastic global optimization procedure which allows simultaneous indexing and molecular packing determination. Final Rietveld refinement was applied on the most promising crystal structure with a ¼ 1.376 nm, b ¼ 3.105 nm, c ¼ 2.690 nm and ß ¼ 109.5 within the space group C2/c choosing the polymer backbone parallel to the b-axis. The structural motifs of the molecular packing could be identified: aromatic units within a single polymer chain are slightly bent relative to the chain axis, octyl side chains are aligned along the polymer backbone and aromatic units of neighbouring molecules display a strong tendency to stack parallel to each other. XRD results of F8T2 with a molecular weight of 19 kg/mol reveal the same peak positions compared to the 3.2 kg/mol material, showing that both materials crystallise similarly and can be described by the same crystallographic unit cell. The smaller peak intensities together with the broader peak widths, however, show that the ability of crystal formation for the 19 kg/mol material is reduced. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: Crystal structure Crystal morphology F8T2
1. Introduction Recently poly(9,9-dioctylfluorenyl-co-bithiophene) (F8T2) becomes important for the realisation of electronic as well as for optoelectronic applications like thin film transistors, optical emitters, detectors and solar cells [1e6]. Despite the large variety of applicability, there is only small knowledge about the packing of these specific polymer backbones within the solid state. Generally, the arrangement of the conjugated molecules within the solid state has large influence on the electronic properties [7]. Conjugated polymers with branching side chains can arrange in many different ways relative to neighbouring polymer chains
* Corresponding author. Tel.: þ43 316 873 8476; fax: þ43 316 873 8466. E-mail address: [email protected] (R. Resel). 0032-3861/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2011.04.063
whereby the location and density of the alkyl chain attachment sites as well as the length of the alkyl chains has a decisive influence. The packing of the side chains leads to a variety of crystalline states ranging from amorphous to liquid crystalline or highly crystalline. F8T2 is a copolymer with an alternating sequence of a dioctylfluorene moiety and two thiophene rings along the polymer chain (see inset of Fig. 2). Besides F8T2, two further prominent conjugated polymers are poly(3-hexyl)thiophene (P3HT) and poly(dioctylfluorene) (PF8). The polymer P3HT e a polymer formed by thiophene units with regularly attached side chains - shows a crystalline character. The thiophene units form p-p stacking sheets that are separated by the hexyl chains from each other [8,9]. The lamella order results in crystals where even chain folding within stacks was observed [10]. The polymer PF8 e a polymer consisting of side chain substituted dioctyl-fluorene units - shows a variety of crystallographic phases like a-phase, b-phase as well as a nematic liquid
O. Werzer et al. / Polymer 52 (2011) 3368e3373
3369
Fig. 1. Atomic force microscopy of a dry drop casted film of the polymer F8T2 with molecular weight of 3.2 kg/mol. On large scales the height image is depicted (left), while a phase imaging micrograph is given at higher resolution (right).
crystalline phase [11e14]. All of these phases have in common that a close p-stacking of the phenyl rings is not observed due to a twist of adjacent fluorene units resulting in a helix conformation of the polymer backbone along the polymer axis. The polymer F8T2 shows a glass transition at 380 K due to melting of the octyl side chains, a thermal induced crystalline phase formed at a temperature of 393 K [15] and a thermal induced liquid crystalline phase at a temperature of 463 K [16,17]. While all these phases show only small molecular weight dependencies for Mw 40 kg/mol [18] the nematiceisotropic transition shows a strong molecular weight dependence; higher molecular weight F8T2 melts at higher temperature (e.g. 19 kg/ mol melts at 563 K [17] and 3.2 kg/mol 428 K [19]). Due to the chemical similarity of PF8 and F8T2 a similar crystalline arrangement is expected. However the introduction of two thiophene rings along the polymer backbone is expected to introduce additional restriction to the polymer e polymer confinement, similar to P3HT in which the sulphur group of adjacent chains are mirrored by 180B. This together with the additional open space, due to the lack of octyl-chains at the thiophene units, allows for different arrangements of the fluorene units. Within this work F8T2 nanocrystals are investigated by atomic force microscopy and x-ray diffraction. The molecular packing is
Fig. 2. Specular x-ray diffraction scans of as-prepared 3.2 kg/mol (red) and 19 kg/mol (black) F8T2 films. The green line represents the diffraction pattern of 3.2 kg/mol F8T2 after heat treatment at 353 K. The curves are shifted for clarity. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
determined by a combined experimental and theoretical approach revealing the arrangement of the polymer chains within the unit cell. 2. Materials and methods F8T2 with a molecular weight of 3.2 kg/mol (Mw/Mn ¼ 1.6) and 19 kg/mol were used for the experiment. The 3.2 kg/mol F8T2 was synthesised following the procedure given in [18]. The 19 kg/mol F8T2 polymer was purchased from American Dye Source Ltd. and used without further purification. Nanocrystals of F8T2 were obtained by dissolving F8T2 in para-xylene with a concentration of 25 mg/ml at a temperature of 353 K by subsequent cooling to room temperature. The initial clear solution of the 3.2 kg/mol F8T2 material turned turbid and non transparent after two days, revealing the formation of aggregates. Similarly, a turbid solution was obtained for the 19 kg/mol material, but after a time period of two weeks. These turbid solutions are stable up to 353 K, at which the solutions became clear again. On cooling, the solutions turned turbid again after the respective time period, revealing the reversibility of nanocrystal formation in para-xylene. Films of F8T2 nanocrystals were prepared by drop casting the turbid solution onto thermally oxidised silicon wafers. The bulk solvent was removed at a vacuum of 106 mbar at room temperature. Atomic force microscopy (AFM) measurements were performed with a Dimension3100 microscope equipped with a Hybrid closed loop scan head and a Nanoscope IVa controller (Digital Instruments, VEECO). All measurements were done in TappingModeÔ with different Olympus cantilevers (2e40 N/m) depending on the sample requirements. Fourier transform infrared absorption spectroscopy (FTIR) measurements were performed using a Bruker Equinox 55 FTIR spectrometer equipped with a Bruker Hyperion 3000 infrared microscope and a single - element MCT detector. Thin films were drop casted on BaF2 support and the transmission spectra were collected using a 15 Cassegrain objective with spectral resolution of 4 cm1 and averaging over 32 interferograms. The area for each measurement was approximately 100 100 mm2. X-ray diffraction (XRD) experiments were performed with a Bruker D8 Discover diffractometer equipped with a parallelizing polycapillar optic at the primary side. The radiation is provided by a copper sealed tube (l ¼ 0.154 nm) and is monochromatized by a secondary side graphite monochromator. Temperature dependent in-situ x-ray investigations were performed with the DHS 900 heating attachment [20] from Anton Paar GmbH (Austria) under a vacuum of 103 mbar to reduce sample degradation.
3370
O. Werzer et al. / Polymer 52 (2011) 3368e3373
The crystal structure determination of the F8T2 crystals was dealt with by means of the VARICELLA package [21]. The method allows for searching the cell parameters together with the polymer conformations, orientations and packing, so that a prior indexing of the spectrum before the structure solution is not required. The molecular model is completely flexible (i.e. no constraints are applied to the stiff molecular degrees of freedom, hence bond lengths and bending angles can vary during the search). The sampling of the conformational space available to the molecules is guided by two factors, a potential energy provided by the MM2 force field as implemented in TINKER [22] and a disagreement factor between the calculated and the experimental profiles. All the interactions among the different molecules within the primary cell and its periodic images are completely accounted for within a suitable cut-off, during the calculations. Molecular moves are obtained by means of the molecular dynamics, while the cell parameter values are changed in a random fashion. This hybrid iterative Monte Carlo procedure starts from any initial crystal structure and improves the solution according to a global optimization strategy. Input parameters to the method are the space group and the number of independent monomers/molecules. The method was enriched with different features with respect to its original formulation, including the possibility to take into account preferred orientations and procedures for better sampling of the octyl side chain conformations. The calculated structures yielding the best compromise between crystal energy and disagreement factor between the calculated and the experimental diffraction patterns were further optimised by a Rietveld refinement procedure. The optimization was performed by minimizing the weighted profile parameter (Rwp) which writes as:
Rwp
2 11=2 0P wi Y sim ð2qi Þ I exp ð2qi Þ þ Y back ð2qi Þ C B i C ¼ B 2 P exp A @ wi I ð2qi Þ i sim
where Y is the calculated diffraction pattern, Iexp the experimental one, Yback the background subtracted profile and wi an unitary weight applied at each point. The Rietveld optimization was performed with the MATSTUDIO package [23] which accounts for peak width and asymmetry, crystallite size and lattice strains, orientational parameter, main chain rotation around the b-axis and translations along the three cell axes. 3. Experimental results Fig. 1 (left) shows a typical 10 mm 10 mm AFM height scan of a drop casted film prepared from the turbid solution revealing a wavy morphology with maximum height variations of about 400 nm. This wide range height variation is attributed to the drop cast process. However, AFM phase images with larger magnification reveal needle like structures as shown in Fig. 1 (right). The needles show a typical width of 45 nm with lengths ranging from 100 to 300 nm 2D Fourier analyses of the height and phase images do not reveal a preferred lateral orientation suggesting a random alignment of the needles at the silica surface. Crystallographic studies on the F8T2 films were performed by specular XRD scans, i.e. the scattering vector was kept perpendicular to the substrate surface. The XRD pattern for the as-prepared 3.2 kg/mol and 19 kg/mol sample are depicted in Fig. 2 and reveal similar diffraction pattern of both samples. Firstly, two broad peaks with maxima at about 4.9 and 19.7 are noted which belong to the amorphous fraction of F8T2 without crystallographic long range order of the polymer chains. Secondly, numerous clear and sharp Bragg peaks are present. The diffraction patterns of the two
samples appear at the same peak positions, suggesting that the same crystalline phase has been formed for the different molecular weights. However, the higher peak intensities and the smaller peak width together with the less pronounced amorphous contribution of the 3.2 kg/mol sample compared to the 19 kg/mol indicates that the crystallinity in the 3.2 kg/mol sample is higher. The first and most intense peak is located at 6.88 which corresponds to a real space repeating distance (d-spacing) of 1.28 nm. In general such a distance is connected with the backboneebackbone distance separated by the alkyl side chains [24]. Please note that this distance is considerably smaller than the length of fully extended octyl units, which would result into a d-spacing of at least 2.1 nm, if interdigitation of the octyl side chains from neighbouring polymer backbones would be present. Next to the strong first peak there are several additional peaks present revealing a three - dimensional crystallographic long range order within the F8T2 films. This is in contrast to P3HT which shows three strong peaks connected to the lamella sheet separation and one weak peak connected to intrachain repeating units across the backbone [8]. As the samples were prepared from solutions, the question arises if xylene is incorporated within the crystalline structure. IR spectroscopy results (not shown) do not provide an evidence of incorporated solvent within the vacuum dried thin films. IR peaks belonging to xylene were only observed right after the drop casting process, but on vacuum treatment the amount of xylene (if any) dropped below the detection limit. For the crystal structure determination of the F8T2 nanocrystals the diffraction pattern of the 3.2 kg/mol sample was chosen. By considering the particular chemical species involved, suitable space groups were chosen and different searches were performed, where either one or two monomers in the asymmetric unit were included. Within the orthorhombic system, Pbca and P212121 were tested, while P2, P21/c,and C2/c were tried in the monoclinic setting. By far, C2/c emerged as the best choice, allowing for both a close packing of the polymer backbones and reasonable arrangements of the alkyl side chains within the cells. All the other groups enforced unrealistically bent molecular backbone conformations and too high crystal densities (up to 1.44 g/cm3), and were therefore discarded. Moreover the disagreement factor obtained from the initial calculations for the discarded groups was always much higher than for the C2/c case. The calculated pattern was further refined against the experimental diffraction pattern, the result is given in Fig. 3. A final value of Rwp ¼ 0.24 was obtained, which can be considered quite acceptable, given the evident high level of uncertainty contained in the experimental XRD data. Such a lack of details in the XRD pattern produces some too close contacts among adjacent macromolecules. In this sense the given structure must be considered as a likely structural motif of the polymeric crystal. The proposed structure contains 16 monomer units in the unit cell of a ¼ 1.376 nm, b ¼ 3.105 nm, c ¼ 2.690 nm and beta ¼ 109.5 , and has a crystal mass density of 1.355 g/cm3. The results reveal a significant preferred orientation (texture) along the (0 0 1) direction. Aside from the amorphous bump centred at 19 , which approximately contributes 50% to the diffraction signal, evident overlapping of more reflections takes place in the profile. As a result the majority of peak widths exceeds 0.5 . Hence the calculation of crystallite size along any crystallographic direction is affected by relevant peak deconvolution problem, except for the first peak belonging essentially to (0 0 2) and in small part to (1 1 1) reflections. For this direction, normal to the polymer backbone, a crystallite dimension of 15 nm has been derived using line profile analysis [25]. The obtained conformation of one polymer backbone is depicted in Fig. 4. The crystallographic unit cell has the length of two monomer units, and the repetition of the chain conformation appears due to the space group after two monomer
O. Werzer et al. / Polymer 52 (2011) 3368e3373
Fig. 3. Experimental (green line) and calculated diffraction pattern (red line) of crystalline F8T2. The scattering contribution of amorphous F8T2 chains (pink line) is evaluated by polynomial interpolation. The difference profile is also plotted as blue line. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
units. The low detailed diffraction pattern causes some level of uncertainty about each torsion value of the side octyl moieties. Moreover some bent thiophene rings are recognizable and the extent to which this is real or an artefact induced by the peculiarities of the experimental profile cannot be assessed with the available data. For these reasons great improvements towards a more detailed crystal model are possible provided that higher quality experimental information becomes available. The structural features of the molecular packing of the polymer F8T2 are: (i) along the polymer backbone neighbouring aromatic units of thiophene and the fluorene segments show rather large tilt angles so that locally the polymer chain is not fully stretched, (ii) the arrangement of the aromatic units from neighbouring polymer backbones is determined by the juxtaposition of octyl-chains, bent along the b-axis to allow for a compact packing, i.e. relatively high crystal density takes place. Fig. 5 show the temperature dependent evolution of the (002) diffraction peak for the 3.2 kg/mol film. At a temperature of 313 K the same diffraction pattern, as for the as-prepared film is obtained. However, a slight decrease of the peak intensity is observed suggesting that melting of the F8T2 nanocrystals starts. At a temperature of 353 K the peak intensity is completely lost revealing that the crystallites are melted due to heat treatment. A decrease in temperature after melting did not recover any crystallographic order (or Bragg peaks); only two broad features remain at around 3.5 nm1 and 14 nm1. These features are characteristic to the
Fig. 4. Conformation of one polymer chain within the crystal structure of F8T2. The repeating unit consists of two monomers.
3371
Fig. 5. In-situ temperature dependent x-ray diffraction scans of the 002 peak of a F8T2 film prepared from Mw ¼ 3.2 kg/mol.
disordered state of F8T2 [15], the complete XRD pattern is included in Fig. 2. The film morphology after heat treatment at 373 K reveals again a wavy surface character, but in comparison to the as-prepared film the height differences are considerably smaller (Fig. 6, left). In addition the surface structure on small length scales are smoother followed from roughness analysis in the sub-micrometer range that reveals a “fine” roughness reduction from w 15 nm to w 1 nm (RMS) due to heat treatment. According phase images give an explanation for this observation: the needle like structures of the as-prepared films (compare Fig. 1 right) are completely vanished as shown in Fig. 6 (right). 4. Discussion The performed XRD experiments show that the polymer F8T2 has a tendency to crystallise in the solvent para e xylene; a needle like morphology is observed for the 3.2 kg/mol polymer. The time period for the crystal formation is considerably faster for the polymer material with the molecular weight of 3.2 kg/mol compared to 19 kg/mol. The reason for this cannot be ascertained, but the increased chain length results in stronger entanglements of the polymer chains and thus in an increased viscosity of the bulk solution. This effect disfavour a fast crystallization of the F8T2 chains and is in accordance with the crystallinity - as observed in Fig. 2 e which is much higher for the 3.2 kg/mol material. Other xylene isomers were not tested for their ability to crystallise F8T2. However, their Hansen solubility parameters, especially their polarity, are distinct. Therefore, a slight variation in the crystal formation can be expected to occur [26]. The gradual change from the clear F8T2 - xylene solutions to a turbid appearance with time and their reversibility on heating and cooling are comparable to the changes of the crystallographic and morphology properties of the drop casted films of the F8T2 nanocrystals. The needle like crystallites are stable up to a temperature of 313 K. Above 353 K the crystallographic changes are accompanied by a morphological transition, i.e. the needle like structures are melted. However, the formation of the crystalline phase of F8T2 was only observed in xylene solutions, suggesting that the obtained crystallographic phase is solvent mediated. As the crystallographic phase does not recover in solvent free films it can be classified as metastable. Recently, an F8T2 oligomer consisting of three monomer units has been described as crystalline [27], but the crystallographic order is different to the present
3372
O. Werzer et al. / Polymer 52 (2011) 3368e3373
Fig. 6. Atomic force microscopy of the F8T2 (Mw ¼ 3.2 kg/mol) film after annealing at 373 K: height images (left) and phase images (right).
polymer investigation. However, this suggests, that the dioctylfluorene e bithiophene moieties have a strong tendency for crystallization. In the present case of the polymer with a high number of repeating units, crystallization is strongly effected by the increased number of possible polymer chain conformation as well as polymer backbone twisting along its axis. These entropical effects hinder the crystallization and a much less defined arrangement of the polymer chains compared to the oligomer results. The similarity of the x-ray diffraction pattern shows that both molecular weight F8T2 samples are representative for the
polymer limit; below the oligomer limit a different crystallographic phase is formed [27]. Within our structural motif, a polymer backbone with tilted and slightly bent aromatic units relative to each other together with directed octyl-chains parallel to the backbone itself, emerges from the present structural analysis. Two polymer backbones are always connected in single columns where the octyl-chains are in the centre of the columns and the aromatic units within the columns are closely packed together. This packing involves both thiophene and fluorene units. The overall packing is constituted by repetition
Fig. 7. Crystal packing as viewed along mainechain axis b (top) and perspective view (bottom), only two neighbouring chains are represented and hydrogen atoms were omitted for clarity.
O. Werzer et al. / Polymer 52 (2011) 3368e3373
of adjacent column couples of polymer chains developing along both a and c axes, separated by side octyl-substituents but also by parallel packing of the aromatic units (see Fig. 7). These main structural features can be compared to the structural features of the polymeric system PF8. Although a large variety of the polymer backbone confirmation is concluded for the different phases, neighbouring aromatic units are neither collinear nor coplanar [28]. Moreover in agreement with our results, the anti-parallel alignment of the octyl side-chains is proven for the b-phase of PF8 [13]. But the aromatic units of neighbouring chains stack parallel to each other, the main feature of the crystal structure P3HT. 5. Conclusion The formation of crytalline F8T2 in the solvent para e xylene is observed. XRD and AFM measurements of the 3.2 kg/mol F8T2 films show the presence of nanocrystallites with needle like morphology. On heat treatment the crystalline long range order and the needle like morphology is lost. The crystal structure of F8T2 is proposed based on x-ray diffraction studies combined with a recently developed crystal structure solution method for polymers. It is found that the crystal structure of F8T2 has structural elements comparable to poly(octly-fluorene) and poly(hexylthiophene), although the polymer chain conformation are neither colinear nor coplanar. But individual F8T2 polymer chains forms closely packed columns and the aromatic units from neighbouring chains are stacked parallel to each other. XRD measurements on the 3.2 kg/ mol and the 19 kg/mol F8T2 samples reveal the same peak positions and can be explained by the common crystallographic unit cell of a ¼ 1.376 nm, b ¼ 3.105 nm, c ¼ 2.690 nm and beta ¼ 109.5 . Variation of relative peak intensities are a result of different ability of crystallization which results in a different preferred orientations of the crystallites; the small number of repeating units in the 3.2 kg/ mol material allows for a faster adaptation of their conformation to neighbouring polymer chains which lead to a more defined crystallization. Acknowledgements Financial support by the Austrian Nano-Initiative (research project cluster ISOTEC), the Christian Doppler Research Association
3373
(CDG) and the Austrian Federal Ministry of Economy, Family and Youth (BMWFJ) is gratefully acknowledged.
Appendix. Supplementary data Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.polymer.2011.04.063.
References [1] Sirringhaus H, Wilson RJ, Friend RH, Inbasekaran M, Wu W, Woo EP, et al. Appl Phys Lett 2000;77:406. [2] Fujiwara T, Locklin J, Bao ZN. Appl Phys Lett 2007;90:232108. [3] Noh YY, Zhao N, Caironi M, Sirringhaus H. Nat Nanotech 2007;2:784. [4] Sagane K, Shakutsui M, Tsutsui T, Fujita K. J Photopolymer Sci Technol 2008; 21:193. [5] Levermore PA, Jin R, Wang XH, de Mello JC, Bradley DDC. Adv Funct Mater 2009;19:950. [6] Huang JH, Ho ZY, Kekuda D, Chang Y, Chu CW, Ho KC. Nanotechnology 2009; 20:025202. [7] Nabok D, Puschnig P, Ambrosch-Draxl C, Werzer O, Resel R, Smilgies DM. Phys Rev B 2007;76:235322. [8] Prosa TJ, Moulton J, Heeger AJ, Winokur MJ. Macromolecules 1999;32:4000. [9] Kline RJ, McGehee MD, Toney MF. Nat Mater 2006;5:222. [10] Brinkmann M, Rannou P. Macromolecules 2009;42:1125. [11] Kawana S, Durrell M, Lu J, Macdonald JE, Grell M, Bradley DDC, et al. Polymer 2002;43:1907. [12] Chen SH, Chou HL, Su AC, Chen SA. Macromolecules 2004;37:6833. [13] Arif M, Volz C, Guha S. Phys Rev Lett 2006;96:025503. [14] Knaapila M, Stepanyan R, Torkkeli M, Garamus VM, Galbrecht F, Nehls BS, et al. Phys Rev E 2008;77:051803. [15] Kinder L, Kanicki J, Petroff P. Synth Metals 2004;146:181. [16] Werzer O, Matoy K, Smilgies DM, Rothmann M, Strohriegl P, Resel R. J Appl Poly Sci 2008;107:1817. [17] Werzer O, Flesch HG, Smilgies DM, Resel R. J Poly Sci B 2009;47:1599. [18] Kinder L, Kanicki J, Swensen J, Petroff P. Proc SPIE 2003;5217:35. [19] Thiem H, Rothmann MM, Strohriegl P. Designed Monomers Polym 2005;8: 619. [20] Resel R, Tamas E, Sonderegger B, Hofbauer P, Keckes J. J Appl Cryst 2003;36: 80. [21] Rapallo AJ. Chem Phys 2009;131:44113. [22] Ren P, Ponder JW. J Phys Chem B 2003;107:5933. [23] MATSTUDIO modeling release 4.0. 9685 Scranton Rd. San Diego CA (USA): Accelrys Inc.; 2003. [24] Ballauff M, Schmidt GF. Mol Cryst Liq Cryst 1987;147:163. [25] Enzo S, Fagherazzi G, Benedetti A, Polizzi S. J Appl Cryst 1988;21:536. [26] Mulder MHV, Kruitz F, Smolders CA. J Membr Sci 1982;11:349. [27] Li N, Zhang X, Geng Y, Su Z. Polymer 2008;49:4279. [28] Chunwaschirasiri W, Tanto B, Huber DL, Winokur MJ. Phys Rev Lett 2005;94: 107402.
Contents lists available at ScienceDirect
Polymer journal homepage: www.elsevier.com/locate/polymer
Crystallographic structure and morphology of bithiophene-fluorene polymer nanocrystals Oliver Werzer a, b, Roland Resel a, *, Boril Chernev c, Harald Plank c, d, Michael M. Rothmann e, Peter Strohriegl e, Gregor Trimmel f, g, Arnaldo Rapallo h, William Porzio h a
Institute of Solid State Physics, Graz University of Technology, Petersgasse 16, A-8010 Graz, Austria Centre for Organic Electronics, University of Newcastle, Australia Centre for Electron Microscopy, Graz, Austria d Institute for Electron Microscopy, Graz University of Technology, Austria e Institute of Macromolecular Chemistry, University of Bayreuth, Germany f Institute for Chemistry and Technology of Materials, Graz University of Technology, Austria g Christian Doppler Laboratory for Nanocomposite Solar Cells, Graz University of Technology, Austria h Istituto per lo Studio delle Macromolecole, Consiglio Nazionale delle Ricerche, Milano, Italy b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 26 January 2011 Received in revised form 28 April 2011 Accepted 29 April 2011 Available online 6 May 2011
Nanocrystals of the polymer poly(9,9-dioctylfluorenyl-co-bithiophene) (F8T2) with a molecular weight of 3.2 kg/mol are grown in a para-xylene solution. The typical morphology of the crystals is needle like with typical widths of 50 nm and lengths of about 200 nm. The crystal structure and morphology are stable up to a temperature of 353 K. The structure solution is obtained by x-ray powder diffraction (XRD) pattern with data modelling by a stochastic global optimization procedure which allows simultaneous indexing and molecular packing determination. Final Rietveld refinement was applied on the most promising crystal structure with a ¼ 1.376 nm, b ¼ 3.105 nm, c ¼ 2.690 nm and ß ¼ 109.5 within the space group C2/c choosing the polymer backbone parallel to the b-axis. The structural motifs of the molecular packing could be identified: aromatic units within a single polymer chain are slightly bent relative to the chain axis, octyl side chains are aligned along the polymer backbone and aromatic units of neighbouring molecules display a strong tendency to stack parallel to each other. XRD results of F8T2 with a molecular weight of 19 kg/mol reveal the same peak positions compared to the 3.2 kg/mol material, showing that both materials crystallise similarly and can be described by the same crystallographic unit cell. The smaller peak intensities together with the broader peak widths, however, show that the ability of crystal formation for the 19 kg/mol material is reduced. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: Crystal structure Crystal morphology F8T2
1. Introduction Recently poly(9,9-dioctylfluorenyl-co-bithiophene) (F8T2) becomes important for the realisation of electronic as well as for optoelectronic applications like thin film transistors, optical emitters, detectors and solar cells [1e6]. Despite the large variety of applicability, there is only small knowledge about the packing of these specific polymer backbones within the solid state. Generally, the arrangement of the conjugated molecules within the solid state has large influence on the electronic properties [7]. Conjugated polymers with branching side chains can arrange in many different ways relative to neighbouring polymer chains
* Corresponding author. Tel.: þ43 316 873 8476; fax: þ43 316 873 8466. E-mail address: [email protected] (R. Resel). 0032-3861/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2011.04.063
whereby the location and density of the alkyl chain attachment sites as well as the length of the alkyl chains has a decisive influence. The packing of the side chains leads to a variety of crystalline states ranging from amorphous to liquid crystalline or highly crystalline. F8T2 is a copolymer with an alternating sequence of a dioctylfluorene moiety and two thiophene rings along the polymer chain (see inset of Fig. 2). Besides F8T2, two further prominent conjugated polymers are poly(3-hexyl)thiophene (P3HT) and poly(dioctylfluorene) (PF8). The polymer P3HT e a polymer formed by thiophene units with regularly attached side chains - shows a crystalline character. The thiophene units form p-p stacking sheets that are separated by the hexyl chains from each other [8,9]. The lamella order results in crystals where even chain folding within stacks was observed [10]. The polymer PF8 e a polymer consisting of side chain substituted dioctyl-fluorene units - shows a variety of crystallographic phases like a-phase, b-phase as well as a nematic liquid
O. Werzer et al. / Polymer 52 (2011) 3368e3373
3369
Fig. 1. Atomic force microscopy of a dry drop casted film of the polymer F8T2 with molecular weight of 3.2 kg/mol. On large scales the height image is depicted (left), while a phase imaging micrograph is given at higher resolution (right).
crystalline phase [11e14]. All of these phases have in common that a close p-stacking of the phenyl rings is not observed due to a twist of adjacent fluorene units resulting in a helix conformation of the polymer backbone along the polymer axis. The polymer F8T2 shows a glass transition at 380 K due to melting of the octyl side chains, a thermal induced crystalline phase formed at a temperature of 393 K [15] and a thermal induced liquid crystalline phase at a temperature of 463 K [16,17]. While all these phases show only small molecular weight dependencies for Mw 40 kg/mol [18] the nematiceisotropic transition shows a strong molecular weight dependence; higher molecular weight F8T2 melts at higher temperature (e.g. 19 kg/ mol melts at 563 K [17] and 3.2 kg/mol 428 K [19]). Due to the chemical similarity of PF8 and F8T2 a similar crystalline arrangement is expected. However the introduction of two thiophene rings along the polymer backbone is expected to introduce additional restriction to the polymer e polymer confinement, similar to P3HT in which the sulphur group of adjacent chains are mirrored by 180B. This together with the additional open space, due to the lack of octyl-chains at the thiophene units, allows for different arrangements of the fluorene units. Within this work F8T2 nanocrystals are investigated by atomic force microscopy and x-ray diffraction. The molecular packing is
Fig. 2. Specular x-ray diffraction scans of as-prepared 3.2 kg/mol (red) and 19 kg/mol (black) F8T2 films. The green line represents the diffraction pattern of 3.2 kg/mol F8T2 after heat treatment at 353 K. The curves are shifted for clarity. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
determined by a combined experimental and theoretical approach revealing the arrangement of the polymer chains within the unit cell. 2. Materials and methods F8T2 with a molecular weight of 3.2 kg/mol (Mw/Mn ¼ 1.6) and 19 kg/mol were used for the experiment. The 3.2 kg/mol F8T2 was synthesised following the procedure given in [18]. The 19 kg/mol F8T2 polymer was purchased from American Dye Source Ltd. and used without further purification. Nanocrystals of F8T2 were obtained by dissolving F8T2 in para-xylene with a concentration of 25 mg/ml at a temperature of 353 K by subsequent cooling to room temperature. The initial clear solution of the 3.2 kg/mol F8T2 material turned turbid and non transparent after two days, revealing the formation of aggregates. Similarly, a turbid solution was obtained for the 19 kg/mol material, but after a time period of two weeks. These turbid solutions are stable up to 353 K, at which the solutions became clear again. On cooling, the solutions turned turbid again after the respective time period, revealing the reversibility of nanocrystal formation in para-xylene. Films of F8T2 nanocrystals were prepared by drop casting the turbid solution onto thermally oxidised silicon wafers. The bulk solvent was removed at a vacuum of 106 mbar at room temperature. Atomic force microscopy (AFM) measurements were performed with a Dimension3100 microscope equipped with a Hybrid closed loop scan head and a Nanoscope IVa controller (Digital Instruments, VEECO). All measurements were done in TappingModeÔ with different Olympus cantilevers (2e40 N/m) depending on the sample requirements. Fourier transform infrared absorption spectroscopy (FTIR) measurements were performed using a Bruker Equinox 55 FTIR spectrometer equipped with a Bruker Hyperion 3000 infrared microscope and a single - element MCT detector. Thin films were drop casted on BaF2 support and the transmission spectra were collected using a 15 Cassegrain objective with spectral resolution of 4 cm1 and averaging over 32 interferograms. The area for each measurement was approximately 100 100 mm2. X-ray diffraction (XRD) experiments were performed with a Bruker D8 Discover diffractometer equipped with a parallelizing polycapillar optic at the primary side. The radiation is provided by a copper sealed tube (l ¼ 0.154 nm) and is monochromatized by a secondary side graphite monochromator. Temperature dependent in-situ x-ray investigations were performed with the DHS 900 heating attachment [20] from Anton Paar GmbH (Austria) under a vacuum of 103 mbar to reduce sample degradation.
3370
O. Werzer et al. / Polymer 52 (2011) 3368e3373
The crystal structure determination of the F8T2 crystals was dealt with by means of the VARICELLA package [21]. The method allows for searching the cell parameters together with the polymer conformations, orientations and packing, so that a prior indexing of the spectrum before the structure solution is not required. The molecular model is completely flexible (i.e. no constraints are applied to the stiff molecular degrees of freedom, hence bond lengths and bending angles can vary during the search). The sampling of the conformational space available to the molecules is guided by two factors, a potential energy provided by the MM2 force field as implemented in TINKER [22] and a disagreement factor between the calculated and the experimental profiles. All the interactions among the different molecules within the primary cell and its periodic images are completely accounted for within a suitable cut-off, during the calculations. Molecular moves are obtained by means of the molecular dynamics, while the cell parameter values are changed in a random fashion. This hybrid iterative Monte Carlo procedure starts from any initial crystal structure and improves the solution according to a global optimization strategy. Input parameters to the method are the space group and the number of independent monomers/molecules. The method was enriched with different features with respect to its original formulation, including the possibility to take into account preferred orientations and procedures for better sampling of the octyl side chain conformations. The calculated structures yielding the best compromise between crystal energy and disagreement factor between the calculated and the experimental diffraction patterns were further optimised by a Rietveld refinement procedure. The optimization was performed by minimizing the weighted profile parameter (Rwp) which writes as:
Rwp
2 11=2 0P wi Y sim ð2qi Þ I exp ð2qi Þ þ Y back ð2qi Þ C B i C ¼ B 2 P exp A @ wi I ð2qi Þ i sim
where Y is the calculated diffraction pattern, Iexp the experimental one, Yback the background subtracted profile and wi an unitary weight applied at each point. The Rietveld optimization was performed with the MATSTUDIO package [23] which accounts for peak width and asymmetry, crystallite size and lattice strains, orientational parameter, main chain rotation around the b-axis and translations along the three cell axes. 3. Experimental results Fig. 1 (left) shows a typical 10 mm 10 mm AFM height scan of a drop casted film prepared from the turbid solution revealing a wavy morphology with maximum height variations of about 400 nm. This wide range height variation is attributed to the drop cast process. However, AFM phase images with larger magnification reveal needle like structures as shown in Fig. 1 (right). The needles show a typical width of 45 nm with lengths ranging from 100 to 300 nm 2D Fourier analyses of the height and phase images do not reveal a preferred lateral orientation suggesting a random alignment of the needles at the silica surface. Crystallographic studies on the F8T2 films were performed by specular XRD scans, i.e. the scattering vector was kept perpendicular to the substrate surface. The XRD pattern for the as-prepared 3.2 kg/mol and 19 kg/mol sample are depicted in Fig. 2 and reveal similar diffraction pattern of both samples. Firstly, two broad peaks with maxima at about 4.9 and 19.7 are noted which belong to the amorphous fraction of F8T2 without crystallographic long range order of the polymer chains. Secondly, numerous clear and sharp Bragg peaks are present. The diffraction patterns of the two
samples appear at the same peak positions, suggesting that the same crystalline phase has been formed for the different molecular weights. However, the higher peak intensities and the smaller peak width together with the less pronounced amorphous contribution of the 3.2 kg/mol sample compared to the 19 kg/mol indicates that the crystallinity in the 3.2 kg/mol sample is higher. The first and most intense peak is located at 6.88 which corresponds to a real space repeating distance (d-spacing) of 1.28 nm. In general such a distance is connected with the backboneebackbone distance separated by the alkyl side chains [24]. Please note that this distance is considerably smaller than the length of fully extended octyl units, which would result into a d-spacing of at least 2.1 nm, if interdigitation of the octyl side chains from neighbouring polymer backbones would be present. Next to the strong first peak there are several additional peaks present revealing a three - dimensional crystallographic long range order within the F8T2 films. This is in contrast to P3HT which shows three strong peaks connected to the lamella sheet separation and one weak peak connected to intrachain repeating units across the backbone [8]. As the samples were prepared from solutions, the question arises if xylene is incorporated within the crystalline structure. IR spectroscopy results (not shown) do not provide an evidence of incorporated solvent within the vacuum dried thin films. IR peaks belonging to xylene were only observed right after the drop casting process, but on vacuum treatment the amount of xylene (if any) dropped below the detection limit. For the crystal structure determination of the F8T2 nanocrystals the diffraction pattern of the 3.2 kg/mol sample was chosen. By considering the particular chemical species involved, suitable space groups were chosen and different searches were performed, where either one or two monomers in the asymmetric unit were included. Within the orthorhombic system, Pbca and P212121 were tested, while P2, P21/c,and C2/c were tried in the monoclinic setting. By far, C2/c emerged as the best choice, allowing for both a close packing of the polymer backbones and reasonable arrangements of the alkyl side chains within the cells. All the other groups enforced unrealistically bent molecular backbone conformations and too high crystal densities (up to 1.44 g/cm3), and were therefore discarded. Moreover the disagreement factor obtained from the initial calculations for the discarded groups was always much higher than for the C2/c case. The calculated pattern was further refined against the experimental diffraction pattern, the result is given in Fig. 3. A final value of Rwp ¼ 0.24 was obtained, which can be considered quite acceptable, given the evident high level of uncertainty contained in the experimental XRD data. Such a lack of details in the XRD pattern produces some too close contacts among adjacent macromolecules. In this sense the given structure must be considered as a likely structural motif of the polymeric crystal. The proposed structure contains 16 monomer units in the unit cell of a ¼ 1.376 nm, b ¼ 3.105 nm, c ¼ 2.690 nm and beta ¼ 109.5 , and has a crystal mass density of 1.355 g/cm3. The results reveal a significant preferred orientation (texture) along the (0 0 1) direction. Aside from the amorphous bump centred at 19 , which approximately contributes 50% to the diffraction signal, evident overlapping of more reflections takes place in the profile. As a result the majority of peak widths exceeds 0.5 . Hence the calculation of crystallite size along any crystallographic direction is affected by relevant peak deconvolution problem, except for the first peak belonging essentially to (0 0 2) and in small part to (1 1 1) reflections. For this direction, normal to the polymer backbone, a crystallite dimension of 15 nm has been derived using line profile analysis [25]. The obtained conformation of one polymer backbone is depicted in Fig. 4. The crystallographic unit cell has the length of two monomer units, and the repetition of the chain conformation appears due to the space group after two monomer
O. Werzer et al. / Polymer 52 (2011) 3368e3373
Fig. 3. Experimental (green line) and calculated diffraction pattern (red line) of crystalline F8T2. The scattering contribution of amorphous F8T2 chains (pink line) is evaluated by polynomial interpolation. The difference profile is also plotted as blue line. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
units. The low detailed diffraction pattern causes some level of uncertainty about each torsion value of the side octyl moieties. Moreover some bent thiophene rings are recognizable and the extent to which this is real or an artefact induced by the peculiarities of the experimental profile cannot be assessed with the available data. For these reasons great improvements towards a more detailed crystal model are possible provided that higher quality experimental information becomes available. The structural features of the molecular packing of the polymer F8T2 are: (i) along the polymer backbone neighbouring aromatic units of thiophene and the fluorene segments show rather large tilt angles so that locally the polymer chain is not fully stretched, (ii) the arrangement of the aromatic units from neighbouring polymer backbones is determined by the juxtaposition of octyl-chains, bent along the b-axis to allow for a compact packing, i.e. relatively high crystal density takes place. Fig. 5 show the temperature dependent evolution of the (002) diffraction peak for the 3.2 kg/mol film. At a temperature of 313 K the same diffraction pattern, as for the as-prepared film is obtained. However, a slight decrease of the peak intensity is observed suggesting that melting of the F8T2 nanocrystals starts. At a temperature of 353 K the peak intensity is completely lost revealing that the crystallites are melted due to heat treatment. A decrease in temperature after melting did not recover any crystallographic order (or Bragg peaks); only two broad features remain at around 3.5 nm1 and 14 nm1. These features are characteristic to the
Fig. 4. Conformation of one polymer chain within the crystal structure of F8T2. The repeating unit consists of two monomers.
3371
Fig. 5. In-situ temperature dependent x-ray diffraction scans of the 002 peak of a F8T2 film prepared from Mw ¼ 3.2 kg/mol.
disordered state of F8T2 [15], the complete XRD pattern is included in Fig. 2. The film morphology after heat treatment at 373 K reveals again a wavy surface character, but in comparison to the as-prepared film the height differences are considerably smaller (Fig. 6, left). In addition the surface structure on small length scales are smoother followed from roughness analysis in the sub-micrometer range that reveals a “fine” roughness reduction from w 15 nm to w 1 nm (RMS) due to heat treatment. According phase images give an explanation for this observation: the needle like structures of the as-prepared films (compare Fig. 1 right) are completely vanished as shown in Fig. 6 (right). 4. Discussion The performed XRD experiments show that the polymer F8T2 has a tendency to crystallise in the solvent para e xylene; a needle like morphology is observed for the 3.2 kg/mol polymer. The time period for the crystal formation is considerably faster for the polymer material with the molecular weight of 3.2 kg/mol compared to 19 kg/mol. The reason for this cannot be ascertained, but the increased chain length results in stronger entanglements of the polymer chains and thus in an increased viscosity of the bulk solution. This effect disfavour a fast crystallization of the F8T2 chains and is in accordance with the crystallinity - as observed in Fig. 2 e which is much higher for the 3.2 kg/mol material. Other xylene isomers were not tested for their ability to crystallise F8T2. However, their Hansen solubility parameters, especially their polarity, are distinct. Therefore, a slight variation in the crystal formation can be expected to occur [26]. The gradual change from the clear F8T2 - xylene solutions to a turbid appearance with time and their reversibility on heating and cooling are comparable to the changes of the crystallographic and morphology properties of the drop casted films of the F8T2 nanocrystals. The needle like crystallites are stable up to a temperature of 313 K. Above 353 K the crystallographic changes are accompanied by a morphological transition, i.e. the needle like structures are melted. However, the formation of the crystalline phase of F8T2 was only observed in xylene solutions, suggesting that the obtained crystallographic phase is solvent mediated. As the crystallographic phase does not recover in solvent free films it can be classified as metastable. Recently, an F8T2 oligomer consisting of three monomer units has been described as crystalline [27], but the crystallographic order is different to the present
3372
O. Werzer et al. / Polymer 52 (2011) 3368e3373
Fig. 6. Atomic force microscopy of the F8T2 (Mw ¼ 3.2 kg/mol) film after annealing at 373 K: height images (left) and phase images (right).
polymer investigation. However, this suggests, that the dioctylfluorene e bithiophene moieties have a strong tendency for crystallization. In the present case of the polymer with a high number of repeating units, crystallization is strongly effected by the increased number of possible polymer chain conformation as well as polymer backbone twisting along its axis. These entropical effects hinder the crystallization and a much less defined arrangement of the polymer chains compared to the oligomer results. The similarity of the x-ray diffraction pattern shows that both molecular weight F8T2 samples are representative for the
polymer limit; below the oligomer limit a different crystallographic phase is formed [27]. Within our structural motif, a polymer backbone with tilted and slightly bent aromatic units relative to each other together with directed octyl-chains parallel to the backbone itself, emerges from the present structural analysis. Two polymer backbones are always connected in single columns where the octyl-chains are in the centre of the columns and the aromatic units within the columns are closely packed together. This packing involves both thiophene and fluorene units. The overall packing is constituted by repetition
Fig. 7. Crystal packing as viewed along mainechain axis b (top) and perspective view (bottom), only two neighbouring chains are represented and hydrogen atoms were omitted for clarity.
O. Werzer et al. / Polymer 52 (2011) 3368e3373
of adjacent column couples of polymer chains developing along both a and c axes, separated by side octyl-substituents but also by parallel packing of the aromatic units (see Fig. 7). These main structural features can be compared to the structural features of the polymeric system PF8. Although a large variety of the polymer backbone confirmation is concluded for the different phases, neighbouring aromatic units are neither collinear nor coplanar [28]. Moreover in agreement with our results, the anti-parallel alignment of the octyl side-chains is proven for the b-phase of PF8 [13]. But the aromatic units of neighbouring chains stack parallel to each other, the main feature of the crystal structure P3HT. 5. Conclusion The formation of crytalline F8T2 in the solvent para e xylene is observed. XRD and AFM measurements of the 3.2 kg/mol F8T2 films show the presence of nanocrystallites with needle like morphology. On heat treatment the crystalline long range order and the needle like morphology is lost. The crystal structure of F8T2 is proposed based on x-ray diffraction studies combined with a recently developed crystal structure solution method for polymers. It is found that the crystal structure of F8T2 has structural elements comparable to poly(octly-fluorene) and poly(hexylthiophene), although the polymer chain conformation are neither colinear nor coplanar. But individual F8T2 polymer chains forms closely packed columns and the aromatic units from neighbouring chains are stacked parallel to each other. XRD measurements on the 3.2 kg/ mol and the 19 kg/mol F8T2 samples reveal the same peak positions and can be explained by the common crystallographic unit cell of a ¼ 1.376 nm, b ¼ 3.105 nm, c ¼ 2.690 nm and beta ¼ 109.5 . Variation of relative peak intensities are a result of different ability of crystallization which results in a different preferred orientations of the crystallites; the small number of repeating units in the 3.2 kg/ mol material allows for a faster adaptation of their conformation to neighbouring polymer chains which lead to a more defined crystallization. Acknowledgements Financial support by the Austrian Nano-Initiative (research project cluster ISOTEC), the Christian Doppler Research Association
3373
(CDG) and the Austrian Federal Ministry of Economy, Family and Youth (BMWFJ) is gratefully acknowledged.
Appendix. Supplementary data Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.polymer.2011.04.063.
References [1] Sirringhaus H, Wilson RJ, Friend RH, Inbasekaran M, Wu W, Woo EP, et al. Appl Phys Lett 2000;77:406. [2] Fujiwara T, Locklin J, Bao ZN. Appl Phys Lett 2007;90:232108. [3] Noh YY, Zhao N, Caironi M, Sirringhaus H. Nat Nanotech 2007;2:784. [4] Sagane K, Shakutsui M, Tsutsui T, Fujita K. J Photopolymer Sci Technol 2008; 21:193. [5] Levermore PA, Jin R, Wang XH, de Mello JC, Bradley DDC. Adv Funct Mater 2009;19:950. [6] Huang JH, Ho ZY, Kekuda D, Chang Y, Chu CW, Ho KC. Nanotechnology 2009; 20:025202. [7] Nabok D, Puschnig P, Ambrosch-Draxl C, Werzer O, Resel R, Smilgies DM. Phys Rev B 2007;76:235322. [8] Prosa TJ, Moulton J, Heeger AJ, Winokur MJ. Macromolecules 1999;32:4000. [9] Kline RJ, McGehee MD, Toney MF. Nat Mater 2006;5:222. [10] Brinkmann M, Rannou P. Macromolecules 2009;42:1125. [11] Kawana S, Durrell M, Lu J, Macdonald JE, Grell M, Bradley DDC, et al. Polymer 2002;43:1907. [12] Chen SH, Chou HL, Su AC, Chen SA. Macromolecules 2004;37:6833. [13] Arif M, Volz C, Guha S. Phys Rev Lett 2006;96:025503. [14] Knaapila M, Stepanyan R, Torkkeli M, Garamus VM, Galbrecht F, Nehls BS, et al. Phys Rev E 2008;77:051803. [15] Kinder L, Kanicki J, Petroff P. Synth Metals 2004;146:181. [16] Werzer O, Matoy K, Smilgies DM, Rothmann M, Strohriegl P, Resel R. J Appl Poly Sci 2008;107:1817. [17] Werzer O, Flesch HG, Smilgies DM, Resel R. J Poly Sci B 2009;47:1599. [18] Kinder L, Kanicki J, Swensen J, Petroff P. Proc SPIE 2003;5217:35. [19] Thiem H, Rothmann MM, Strohriegl P. Designed Monomers Polym 2005;8: 619. [20] Resel R, Tamas E, Sonderegger B, Hofbauer P, Keckes J. J Appl Cryst 2003;36: 80. [21] Rapallo AJ. Chem Phys 2009;131:44113. [22] Ren P, Ponder JW. J Phys Chem B 2003;107:5933. [23] MATSTUDIO modeling release 4.0. 9685 Scranton Rd. San Diego CA (USA): Accelrys Inc.; 2003. [24] Ballauff M, Schmidt GF. Mol Cryst Liq Cryst 1987;147:163. [25] Enzo S, Fagherazzi G, Benedetti A, Polizzi S. J Appl Cryst 1988;21:536. [26] Mulder MHV, Kruitz F, Smolders CA. J Membr Sci 1982;11:349. [27] Li N, Zhang X, Geng Y, Su Z. Polymer 2008;49:4279. [28] Chunwaschirasiri W, Tanto B, Huber DL, Winokur MJ. Phys Rev Lett 2005;94: 107402.