Phase Separation in the Melt and Confined Crystallization as the Key

May 21, 2013 - Microphase-separated donor–acceptor block copolymers have been discussed as ideal systems for morphology control in organic ...
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Phase Separation in the Melt and Confined Crystallization as the Key to Well-Ordered Microphase Separated Donor−Acceptor Block Copolymers Ruth H. Lohwasser,† Gaurav Gupta,‡ Peter Kohn,‡ Michael Sommer,† Andreas S. Lang,† Thomas Thurn-Albrecht,‡,* and Mukundan Thelakkat†,* †

Applied Functional Polymers, Department of Macromolecular Chemistry I, University of Bayreuth, Universitätsstraße 30, 95444 Bayreuth, Germany ‡ Eperimental Polymer Physics Group, Martin-Luther-Universität Halle-Wittenberg, Von-Danckelmann-Platz 3, 06120 Halle, Germany S Supporting Information *

ABSTRACT: Microphase-separated donor−acceptor block copolymers have been discussed as ideal systems for morphology control in organic photovoltaics. Typical microphases as known from coil−coil systems were not observed in such systems due to crystallization dominating over microphase separation. We show how this problem can be overcome by the synthesis of high molecular weight block copolymers leading to a high enough χN parameter and microphase separation in the melt. A combination of copper-catalyzed azide-alkyne click reaction and nitroxide mediated radical polymerization (NMRP) was used for the synthesis of donor−acceptor poly(3-hexylthiophene)-block-poly perylene bisimide acrylate (P3HT-b-PPerAcr) block copolymers. With this synthetic strategy, high molecular weights are possible and no triblock copolymer byproducts are formed, as observed with former methods. Two different block copolymers with a high molecular weight P3HT block of 19.7 kg/mol and a PPerAcr content of 47 and 64 wt % were obtained. X-ray scattering measurements show that the diblock copolymers exhibit microphase separation in the melt state. Furthermore, upon cooling confined crystallization occurs inside the microphase separated domains without destroying the microphase order. The observed microstructures fit well to the respective volume fractions and the crystalline packing within the individual blocks is analogous to those in the respective homopolymers. For the first time, typical lamellar or cylindrical phase separated structures as known for amorphous coil−coil systems are realized for a crystalline−liquid crystalline, donor−acceptor block copolymer. A similar block copolymer synthesized with an earlier method exhibits a crystallization-induced microphase separation.

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INTRODUCTION Conjugated polymers have become an intense field of research because of their high potential for application in organic electronics. The devices based on organic materials benefit from solution processability, possible fabrication on flexible substrates and lightweight.1 In organic photovoltaics, excitation leads to the formation of a bound electron−hole pair, an exciton. Free charge carriers are only obtained if the exciton reaches an interface within its lifetime, which requires interfaces between a donor and an acceptor material within the exciton diffusion length.1 For organic photovoltaics, despite the steadily increasing efficiencies,2 improved morphological control and a higher long-term stability are desired. The morphological instability mainly originates from the nonequilibrium structures of blends. Thus, device performance strongly depends on preparation conditions and post-treatment or thermal history.3−6 This is crucial especially for large area roll to roll processing where thermal annealing/drying steps are involved in the processing conditions.7 Therefore, block copolymers © XXXX American Chemical Society

containing both donor (D) and acceptor (A) segments have become of great scientific interest. The idea of using electronically active block copolymers is to control the D−A morphology by microphase separation and to create interfaces in the range of the exciton diffusion length (10−15 nm). Several reviews have emphasized the advantage of such D−A block copolymers for photovoltaics.8,9 The earlier works of Hadziioannou et al.10 and our own group11 have demonstrated the advantages of charge transfer in D−A block copolymers compared to the corresponding blends. In ideally aligned block copolymer systems, continuous charge transport pathways to the electrodes have been theoretically predicted to favor improved device efficiency.12 Therefore, the fundamental challenges in D−A block copolymers are (1) tailor-made synthesis of these systems with high enough molecular weights Received: October 9, 2012 Revised: April 25, 2013

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the previously reported polymer with the in situ synthesized macroinitiator was also studied with the new samples for their structure formation and microphase separation, both in molten state as well as during cooling from melt by temperature dependent SAXS/WAXS measurements in bulk. X-ray scattering measurements demonstrate that in the block copolymer synthesized with the in situ macroinitiator, the microphase separation is induced by the simultaneous crystallization of both blocks, whereas in the other samples microphase separation occurs in the melt itself and confined crystallization of the respective blocks takes place without altering or destroying the microphase separated structure. Thus, lamellar and cylindrical microstructures are obtained in bulk for the new samples, which are consistent with the respective volume fractions as known from simple coil−coil systems. Wide angle X-ray analysis shows that these microphase separated domains consist of crystalline P3HT and liquid crystalline PPerAcr. The observed crystalline morphologies fit well to those of the respective homopolymers.

for good charge transport properties, (2) a deep understanding and control of their microphase separation, and (3) alignment of the nanoscale morphology. A variety of different block copolymers containing a conjugated poly(3-alkylthiophene), P3AT, block as donor segment and various electronically inactive amorphous blocks have been synthesized, and in some of these systems the question of microphase separation has been addressed in detail.13−17 But only very few of these crystalline−coil block copolymers showed the typical microphase-separated structures known from amorphous coil−coil systems.14,16 The classical microphase separation in amorphous systems depends on the incompatibility factor χN (product of the degree of polymerization, N and the Flory−Huggins interaction parameter, χ) and the volume fraction. In crystalline−amorphous block copolymers, it might also depend on additional parameters of interaction such as the Maier−Saupe parameter (μN) as shown for some conjugated block copolymers.18 The complexity of microphase separation in double-crystalline block copolymers as well as in crystalline−amorphous systems involving poly(ethylene) has been studied in detail by Register et al.19,20 Among the block copolymers containing poly(3-hexylthiophene), P3HT segments, only a few of them contained an electronically active second block: either carrying fullerenes21,22 or perylene bisimides23−26 as pendant moieties in acceptor block. Some of these block copolymers have also found application in solar cells. In the case of all-conjugated block copolymers,27−29 nanostructure formation was observed in a P3HT-b-poly(fluorene) system coincident with the crystallization of P3HT.29 Indications for complex self-assembly processes were also reported. Otherwise, up to now, D−A block polymers with crystalline blocks have not shown the typical microphase separated structures, most likely since crystallization of one of the blocks dominated over microphase separation. In consequence, elongated fibrils or crystalline lamellar structures of P3HT were observed in D−A block copolymers instead of well-ordered microphase separated structures as known from coil−coil systems. Additionally, the P3HT block length incorporated into such D−A block copolymers were usually short and thus high N values were never reached. High molecular weights on the other hand, are synthetically challenging and often limited by the decrease in solubility with increasing molecular weight of P3HT. Starting from these facts, we set some prerequisites for the synthesis of donor−acceptor block copolymers containing the most common semiconductor P3HT. The first requirement is a long P3HT block for good charge carrier mobility of holes and high N values. Second, a balanced charge transport of donor and acceptor blocks is necessary, which can be tuned by the block ratio. We have recently demonstrated that a content of about 60−70 wt % of PPerAcr is necessary for good performance in solar cells.30 Recently we demonstrated that a combination of coppercatalyzed azide−alkyne click reaction and nitroxide mediated radical polymerization (NMRP) is highly useful for the synthesis of P3HT block copolymers with a long P3HT block and a poly(4-vinylpyridine) segment.31 In the present work, we use this method to synthesize two novel high molecular weight P3HT-b-PPerAcr donor−acceptor block copolymers. With this approach no triblock copolymer byproducts were formed as in the earlier reported method of direct in situ synthesis of an alkoxyamine macroinitiator followed by polymerization of PerAcr.32−34 For comparison,

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EXPERIMENTAL SECTION

Materials and Characterization. 1H NMR spectra were recorded in chloroform on a Bruker Avance 250 spectrometer at 300 MHz. Coupling constants are given in hertz. The spectra were calibrated according to the solvent signal at 7.26 ppm. Size exclusion chromatography (Poly-SEC) measurements were performed utilizing a Waters 515-HPLC pump with stabilized THF as the eluent at a flow rate of 0.5 mL/min. A 20 μL volume of a solution with a concentration of approximately 1 mg/mL was injected into a column setup, which consists of a guard column (Varian, 50 × 0.75 cm, ResiPore, particle size 3 μm) and two separation columns (Varian, 300 × 0.75 cm, ResiPore, particle size 3 μm). The compounds were monitored with a Waters UV detector at 254 nm. Polystyrene was used as external standard and o-dichlorobenzene as an internal standard for calibration. Matrix assisted laser desorption ionization spectroscopy with time-offlight detection mass spectroscopy (MALDI-TOF MS) measurements were performed on a Bruker Reflex III using Dithranol as matrix and a mixture of 1000:1 (matrix:polymer). The laser intensity was set to around 70%. Differential scanning calorimetry (DSC) measurements were done on Perkin-Elmer DSC apparatuses (DSC 7 and Diamond) with a heating and cooling rate of 10 K/min. The first heating run after preparation was not considered. Background signals were subtracted, resulting in a measurement of the apparent heat capacity cp(T). Samples for transmission electron microscopy (TEM) were annealed for 1−3 h at 265 °C (above Tm) and slowly cooled to room temperature with a cooling rate of 1 °C/min. Subsequently the samples were cut with an Ultramicrotom, stained 5 min with RuO4 and measured on a Zeiss 9220mega. The monomer 2,5-dibromo3-hexylthiophene, the catalyst 1,3-bis(diphenylphosphino)propanenickel(II) chloride [Ni(dppp)Cl2], perylene bisimide acrylate PerAcr, 2,2,5-trimethyl-4-phenyl-3-azahexane-3-oxyl, and 2,2,5-trimethyl-3-(1′-p-azidomethylphenylethoxy)-4-phenyl-3-azahexane were synthesized according to the literature.35−38 The P3HT-macroinitiator was synthesized as recently published.31 All glass apparatus for polymerization were heated and cooled down under argon. Dry THF was distilled over calcium hydride and potassium. Dry odichlorobenzene, ethynylmagnesium chloride (0.6 M in THF/ tolouene), and t-BuMgCl (1.7 M in THF) were purchased from Acros and the Grignard reagents were titrated according to Krasovskiy and Knochel.39 LiCl puriss p.a. water free was purchased from Fluka and dried prior to use. Synthesis of P3HT-b-PPerAcr (5, 6). In a 5 mL Schlenk tube, the P3HT-macroinitiator 3 (300 mg, 0.023 mmol), perylene bisimide acrylate PerAcr 4 (1910 mg, 2.3 mmol) and free nitroxide 2,2,5trimethyl-4-phenyl-3-azahexane-3-oxyl (1.03 mg, 0.0046 mmol) were mixed with 1300 μL o-dichlorobenzene. After degassing with three freeze, pump, thaw cycles the mixture was stirred at 125 °C. After 28 h B

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Scheme 1. Synthesis of P3HT-Alkoxyamine Macroinitiator 3 from P3HT-Alkyne 1 and P3HT-b-PPerAcr 5, 6 Using a Combination of Copper-Catalyzed Azide-Alkyne Click Reaction and Nitroxide Mediated Radical Polymerization

Table 1. Molecular Weights, Polydispersities, and Thermal Properties of P3HT-Alkyne 1, P3HT-Alkoxyamine Macroinitiator 3, and P3HT-b-PPerAcr 5, 6 As Determined via SEC, MALDI-TOF MS, 1H NMR, and Differential Scanning Calorimetrya polymer P3HT-alkyne 1 P3HT-macroinitiator 3 P3HT-b-PPerAcr 5 P3HT-b-PPerAcr 6 P3HT- macroinitiator in situ P3HT-b-PPerAcr in situ 7

Mn (SEC)/ Mn (MALDI) [kg/mol]

Mp (SEC) [kg/mol]

19.7/12.4 19.7/− 28.4/− 35.5/− 17.0/−

22.5 23.0 31.4 36.4

29.5/−

PDI (SEC)

repeating units P3HT (MALDI)/ PPerAcr (1H NMR)

wt % PPerAcr (1H NMR)

Tm [°C]

Tc [°C]

ΔHm [J/g]

75/− 75/− 75/13 75/27

0 0 47 64 0

− 228/239 193/228/239 213/232/246 −

− 198 180 190/177 −

− 21.6 15.2 (P3HT 11.9) 16.4 (P3HT 8.2) −

55

−

−

−

1.17 1.15 1.13 1.19 1.12 1.15

a

Mn = number-average molecular weight, Mp = peak molecular weight. Additional data of an in situ synthesized macroinitiator and the respective P3HT-b-PPerAcr 7 from an earlier report are given for comparison.33

Figure 1. Size exclusion chromatography traces (a) of P3HT-alkyne 1 (black), P3HT-macroinitiator 3 (blue, dash) and the in situ functionalized macroinitiator (red) and of (b) P3HT-macroinitiator 3 (blue, dash), P3HT-b-PPerAcr 5 (dash), and 6 (solid). The red curve shows the P3HT-bPPerAcr block copolymer 7 synthesized with a former method using the in situ functionalized macroinitiator. δH (300 MHz; CDCl3): 8.65−7.32 (m × 8H, br, Har PPerAcr), 6.98 (n × 1H, s, Har P3HT), 5.21−4.45 (m × 1H, br, CH swallow tail PPerAcr), 4.35−3.70 (m × 4H, br, OCH2 and NCH2 PPerAcr), 2.80 (2 H, t, Jαβ 7.6, α-CH2 P3HT), 2.30−1.05 (m × 45H and n × 8H, br, CH and CH2 PPerAcr and P3HT), 1.00−0.70 (m × 6H and n × 3H, br, CH3 PPerAcr and P3HT); n = degree of polymerization of P3HT and m = degree of polymerization of PPerAcr.

(for sample 5) and 46 h (for sample 6), the reaction was stopped through rapid cooling in liquid nitrogen. The highly viscous product was dissolved in chloroform and precipitated in acetone. The purple powder was extracted in a Soxhlet apparatus with acetone and CH2Cl2 to remove unreacted monomer and P3HT macroinitiator. The pure polymer was obtained after dissolution in chloroform and precipitation in methanol. C

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X-ray Scattering of Samples 5 and 6. X-ray scattering experiments at small and intermediate angles were performed in transmission geometry with a laboratory setup consisting of a Rigaku rotating anode, a focusing X-ray optics device (Osmic confocal max flux), and a Bruker 2D-detector (HighStar). The optics also served as monochromator for Cu Kα radiation (λ = 0.154 nm). Aluminum disks with a central hole of 0.8 mm diameter were used as sample holders. Measurements taken on two sample stages with different sample-todetector distances were combined in order to access a larger q range. The largest scattering vector q accessible in these experiments was about q ≈ 6.5 nm−1. The samples were mounted on a Linkam hotstage for temperature control. Heat conducting paste was used to ensure good thermal contact. Wide angle X-ray scattering measurements were performed on an Empyrean diffractometer from PANalytical operating in Bragg−Brentano geometry using a programmable divergence slit to keep the illuminated area constant and a PIXEL detector. Here also Cu Kα radiation (λ = 0.154 nm) was used.

of the methylene protons in proximity to the acrylate unit and the imide unit of the perylene bisimide acrylate. The weight ratio of PPerAcr was determined as 47 wt % for P3HT-bPPerAcr 5 and 64 wt % for P3HT-b-PPerAcr 6. Using the number of P3HT repeating units as determined from MALDITOF MS, the number of repeating units of PPerAcr for 5 and 6 were calculated as 13 and 27 respectively. All the SEC and thermal data of macroinitiators, 5, 6, and the earlier reported P3HT-b-PPerAcr 7 are given in Table 1. In the earlier reported in situ method, we have used a smaller P3HT macroinitiator of 17 kg/mol (see SEC in Figure 1a).32−34 The SEC trace of the block copolymer sample 7 synthesized with the in situ functionalized macroinitiator shows a clear high molecular weight shoulder due to triblock impurities32 (see Figure 1b). The triblock copolymer in the reported sample 7 arises from the presence of bifunctional macroinitiator.32,40 The molecular weights of the respective block copolymer 7 is also given in Table 1 for a comparison. With 29.5 kg/mol the molecular weight of 7 seems similar or even slightly higher than that for P3HT-b-PPerAcr 5, however this is due to the triblock copolymer content, which shifts the number-average value to higher Mn values. According to 1H NMR the weight percent of PPerAcr in this sample was 55 wt %. Because of the triblock copolymer content, calculation of the respective repeating units was not possible. In contrast to the method of in situ end-capping with a Grignard functionalized alkoxyamine, the present synthetic approach using P3HTalkyne did not lead to a mixture of di- and triblocks, even when a high molecular weight macroinitiator was used. Apart from the avoidance of triblock copolymer contamination the new method also results in a small difference in the chemical connectivity of the two blocks. While in the former method the blocks are only separated by the NMRP starting unit (structure see Supporting Information Figure SI 1), the new block copolymers are separated by an additional triazole unit arising from the click reaction. Structural Analysis by X-ray Scattering and Transmission Electron Microscopy in Bulk. The two components of P3HT-b-PPerAcr were studied independently before; both blocks exhibit ordered structures at room temperature and melt in a similar temperature range around 200−250 °C.41,42 P3HT is a semicrystalline polymer. In comparison to many common semicrystalline polymers, it has a somewhat larger persistence length of about 3 nm,43 which on the other hand is still much smaller than the contour length of common molecular weights for this polymer and also for the P3HT block of the samples studied here. In contrast to many other conjugated polymers, P3HT should therefore not be considered as a rod polymer, but rather as a semiflexible polymer.43 Therefore, nematic phases are not expected and were so far also not observed for P3HT. PPerAcr also orders on a molecular scale, it forms a lamellacolumnar liquid crystalline phase.41 In consequence, these ordering processes of the two components (here simply called crystallization) are in competition to a classical microphase separation of the two blocks in the amorphous phase. A schematic overview over the resulting different possible scenarios for structure formation is presented in Figure 2. The crystallization of one or both blocks could either occur from an ordered, microphase separated melt, Figure 2(b), or directly from the disordered melt, as shown in Figure 2c. It was interesting to see how the here used block copolymers behave. Direct evidence about structure formation processes was obtained from temperature dependent X-ray scattering. Figure

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RESULTS AND DISCUSSION Synthesis. The P3HT-macroinitiator was synthesized as recently published.31 First an alkyne-functionalized P3HT (P3HT-alkyne 1) was synthesized via Kumada catalyst transfer polymerization followed by end-capping the active chain end by reacting with ethynylmagnesium chloride.40 Because of the unsaturated nature of the alkyne predominately monofunctionalized P3HT is obtained. The alkyne group was further reacted in a copper-catalyzed azide−alkyne reaction with 2,2,5trimethyl-3-(1′-p-azidomethylphenylethoxy)-4-phenyl-3-azahexane 2 to form the alkoxyamine functionalized P3HT macroinitiator (P3HT-macroinitiator 3) (Scheme 1). The number-average molecular weight of the P3HT-macroinitiator as determined by size exclusion chromatography (SEC) was 19.7 kg/mol and the polydispersity index (PDI) 1.15. (see Table 1 and Figure 1a). The small peak at higher molecular weight results from coupling of alkyne end groups.31 According to matrix assisted laser desorption ionization mass spectrometry (MALDI-TOF MS) this corresponds to a number-average molecular weight Mn of 12.4 kg/mol and about 75 repeating units. For the nitroxide mediated radical polymerization (NMRP) of perylene bisimide acrylate, the ratio of perylene bisimide acrylate monomer 4 to P3HT-macroinitiator 3 to free nitroxide was chosen as 99:1:0.2 in order to obtain a high percentage of PPerAcr in the block copolymer. The polymerization was performed at 125 °C in o-dichlorobenzene (oDCB). Because of the high viscosity of the polymerization mixture it is not possible to polymerize to conversions above 30−40%. Addition of more degassed o-DCB after 28 h of polymerization did not promote higher conversions, but limited the chain growth and led to a lower content of PPerAcr for P3HT-b-PPerAcr 5. P3HT-b-PPerAcr 6 was polymerized for 46 h until the polymerization mixture could no longer be stirred due to increased viscosity. The polymer was purified by dissolution in chloroform and reprecipitation in acetone. Traces of remaining monomer and most of the remaining macroinitiator were removed by Soxhlet extraction using acetone and dichloromethane. However a small amount of unreacted macroinitiator remained after extraction. The molecular weight distribution curves obtained by calibration with PS for all macroinitiators and block copolymers are given in Figure 1a and 1b. The block copolymers 5 and 6 had Mn values of 28.4 kg/mol and 35.5 kg/mol and low polydispersity indices (PDIs) of 1.13 and 1.19, respectively. The block ratios were determined with 1H NMR spectroscopy by comparing the integral of the α-CH2 proton signals of P3HT with the integral D

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signal, the position of the main peak remains unchanged indicating that crystallization of the components does not create a new microstructure. The change in intensity of the second order peak is most likely caused by additional roughness of the interface between the two components. Thus, the new synthetic approach, enabling higher molecular weight polymers and excluding triblock copolymer byproducts, leads to a segregated system with microphase separation in both crystalline and amorphous states (Figure 2b). While the scattering data for P3HT-b-PPerAcr 5 look similar, P3HT-bPPerAcr 7 exhibits the other scenario according to Figure 2c, i.e. crystallization directly from the disordered melt. This conclusion follows from the analysis of the shape of the peak in the small angle scattering (see Supporting Information and Figure SI 3). For high temperatures, this shape is well described by the structure factor calculated by Leibler for a disordered block copolymer melt. 44 The resulting value for the incompatibility χN indicates consistently a disordered state. Details are given in the Supporting Information. Additionally for P3HT-b-PPerAcr 7, the position of the small angle peak shifts upon crystallization of the components to smaller qvalues indicating that crystallization is not confined by a microphase structure. It is to be noted that this polymer 7 contains certain amounts of triblock copolymer. Generally, for given molecular weight and composition the disordered phase is more stable in triblock copolymers than in diblocks.45 To identify the microphase structure, TEM images were taken from samples which were annealed above the melting temperature and subsequently slowly cooled. The samples were stained with RuO4, which selectively stains the P3HT phase.25 Figure 4b shows a lamellar morphology with a periodicity of about 40 nm for P3HT-b-PPerAcr 5. A detailed analysis of the SAXS pattern shows that the same microphase structure exists in the melt state (Figure 4a). Two reflections are visible at qvalues of 0.181 and 0.362 nm−1. The ratio of 1:2 indicates a

Figure 2. Schematic illustration of possible ordering scenarios for block copolymers upon cooling (top to bottom): (a) A block copolymer with two noncrystallizable blocks forms a disordered melt at temperatures above the order−disorder temperature TODT (i), while for T < TODT (ii) in the ordered melt the well-known microphase separated morphologies develop. For BCPs with two crystallizable blocks, the crystallization of the individual blocks either occurs from an ordered melt (b) or directly from the disordered melt (c) depending on the relative locations of TODT and Tc. Here, for simplicity, the same melting temperature Tm and periodicity b was assumed for both crystalline materials.

3 shows the small and intermediate X-ray patterns of P3HT-bPPerAcr 6 measured at different temperatures during stepwise

Figure 3. X-ray diffraction pattern of P3HT-b-PPerAcr 6 recorded during cooling from the melt. Temperatures at which the block copolymer is in the molten state are marked in orange.

cooling as an example; data for the other samples is given in the Supporting Information. The obtained scattering curves can be divided into two temperature ranges. At high temperatures both components are in the molten state (orange lines), below Tc crystallization of the individual components occurs (black lines). Already in the molten state both polymers show clear indications for microphase separation, i.e., a peak at small angles whose width (fwhm = 0.03 nm−1) corresponds to the resolution of the instrument with higher orders at positions reflecting the symmetry of the microphase separated structure. TODT is not reached in the experimentally accessible temperature range. At the same time reflections of the crystalline phases at larger q are absent. Upon cooling, below the crystallization temperature, additional reflections show up, indicating crystallization of the individual blocks. In the SAXS

Figure 4. Lorentz corrected SAXS data at 270 °C of P3HT-b-PPerAcr 5 (a) and 6 (c) and transmission electron microscopy image of 5 (b) and 6 (d) in bulk showing a lamellar and cylindrical morphology, respectively. E

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Figure 5. Scattering pattern for the block copolymer P3HT-b-PPerAcr 6 at 30° for (a) intermediate and (b) high q range showing the reflections for π−π stacking, with the reflections from P3HT marked in blue and the reflections from PPerAcr marked in red.

lamellar phase. The long period calculated from q = 0.15 nm−1 at 30 °C is 42 nm consistent with result from TEM. For P3HT-b-PPerAcr 6, the TEM image shows hexagonally ordered dark P3HT cylinders in a bright matrix of PPerAcr (Figure 4d). Figure SI 7 (Supporting Information) shows an additional TEM image, in which both standing and lying cylinders can be observed. The reflections in the SAXS pattern are not as well resolved as in the case above, but the data are consistent with a hexagonal lattice for which reflections at positions corresponding to a ratio of 1:√3:2 are expected (cf. Figure 4c) The spacing between the (100) planes calculated from the position of the first reflection (q = 0.127 nm−1 at T = 30 °C) is 49.5 nm. On the basis of this value, the lattice parameter “a” of the hexagonal lattice is calculated to be 57 nm (inset Figure 4d). Detailed information about the crystalline structure of the individual blocks inside the microdomains can be obtained from the scattering pattern for q > 1 nm−1. Figure 5 shows the data of P3HT-b-PPerAcr 6 recorded at 30°, as a representative example for both block copolymers (data for P3HT-b-PPerAcr 5 can be found in Figure SI 6 in the Supporting Information). Since the q range for the temperature-dependent measurement apparatus is limited to q = 6.5 nm−1, the higher q ranges were measured on a different setup. They show reflections for the π−π stacking of both polymers (Figure 5b). The positions of the reflections fit to the structures known from the individual homopolymers,41,42 the P3HT reflections are marked in blue and the PPerAcr reflections in red. P3HT crystals can be described with the well-known orthorhombic unit cell with lattice parameters a = 1.79 nm and b = c = 0.76 nm. PPerAcr forms a lamello-columnar liquid crystalline structure with a two-dimensional oblique unit cell with lattice parameters a = 2.97 nm, b = 2.11 nm, c = 0.36 nm, and γ ∼ 81° which is in agreement with the recently published structure of the homopolymer.41 The crystallographic density of the unit cells is calculated as 1.06 g/cm2 for P3HT and 1.21 g/cm2 for PPerAcr. Using the weight fractions of the block copolymers as determined from 1H NMR spectroscopy, the volume fraction of PPerAcr were calculated as 44% for P3HT-b-PPerAcr 5 and 61% for 6. These volume fractions and the respective lamella and cylindrical structures are consistent with the phase diagram described for simple coil−coil systems. The TEM image of P3HT-b-PPerAcr 7 is given in the Supporting Information (Figure SI 4). In comparison to 5 and

6, crystallization from the disordered melt in 7 led to a less ordered structure. Thermal Characterization. To find out to what extent crystallization is influenced by confinement within the block copolymer microphase separated structure, we performed thermal characterization of P3HT-b-PPerAcr 5 and 6 by differential scanning calormetry (DSC) (Figure 6). Generally

Figure 6. Differential scanning calorimetry curves of P3HT-b-PPerAcr 5, 6.

for the samples under study the melting temperatures depend on molecular weight and the thermal behavior of the block copolymers changes accordingly with the chain length of each block.34 DSC measurements show that the P3HT-macroinitiator 3 has two melting points, a major one at 228 °C and a minor one at 239 °C (see Supporting Information and Table 1). This behavior could either be due to the existence of a layered phase with smectic symmetry at high temperatures or, more likely, be caused by recrystallization and a subsequent second melting during heating.42 The melting enthalpy is 21.6 J/g. Therefore, the sterically demanding alkoxyamine end group lowered the crystallinity only slightly compared to a nonfunctionalized homopolymer with similar molecular weight, which has a melting enthalpy of 22.4 J/g.42 For the P3HT-bPPerAcr 5 three melting points are observed, one at 193 °C belonging to the PPerAcr, and other two peaks at 228 and 240 °C belonging to P3HT. Only a single crystallization peak is visible at 180 °C indicating simultaneous crystallization of both blocks. While the reasons leading to shift of melting temperatures in a block copolymer might be complex and difficult to analyze from conventional DSC alone, the value of F

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ACKNOWLEDGMENTS We acknowledge financial support from DFG (SPP 1355) and EU (Grant 261936 LARGECELLS). We thank the European Synchrotron Radiation Facility for beamline time and facilities. We also thank M. Sztucki, ESRF for assistance at beamline ID2 and H. Schoberth, Bayreuth for help at the ESRF measurements. Furthermore, we thank S. Goerlitz and W. Lebek, Halle, and Melanie Förtsch, Bayreuth, for the TEM measurements.

the melting enthalpy is more reliable and allows a qualitative estimate of the crystallinity. Here, the total melting enthalpy of the block copolymer is 15.2 J/g from which 11.9 J/g can be attributed to the melting enthalpy of the P3HT block (two higher melting points). If we divide this value by the melting enthalpy of the pure P3HT block we get a ratio of about 0.55, which is similar to the weight fraction of P3HT in the block copolymer (53 wt %), indicating that the block copolymer structure does not suppress crystallization. The thermal properties of P3HT-b-PPerAcr 6 are similar, however the melting temperatures are slightly increased (PPerAcr block, 213 °C, and P3HT block, 232 and 246 °C). The temperature shift is most pronounced for the PPerAcr block which can be attributed to the higher molecular weight (27 repeating units in P3HT-b-PPerAcr 6 in comparison to 13 repeating units in 5) and the fact that PPerAcr forms a continuous phase. Again crystallization of both components occurs over a narrow temperature range. From the two melting peaks occurring at higher temperatures, which can be clearly attributed to the P3HT component, the melting enthalpy was determined as 8.2 J/g. This value again roughly corresponds to the weight fraction of P3HT in 6 indicating that crystallization in the block copolymer takes place basically without hindrance. For both block copolymers, the observed crystallization temperatures are consistent with the supercooling observed in the SAXS measurements.

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CONCLUSION We demonstrated that a combination of click chemistry and NMRP is a suitable route for the synthesis of high molecular weight P3HT-b-PPerAcr block copolymers with a low polydispersity. The increased molecular weight and the absence of triblock copolymer byproducts lead to microphase separation in the melt. Confined crystallization occurred within this microphase separated state without alteration of the morphology. Thus, for the first time lamella and cylindrical microstructures as known from simple coil−coil system were obtained for a crystalline−liquid crystalline donor−acceptor block copolymer. Additionally, the volume fractions at which these structures occur are analogous to that known from coil− coil block copolymers. The crystalline structures of the individual blocks remained unchanged in comparison to the respective homopolymers, which implies a high purity of the components in the respective domains. The results presented here demonstrate that block copolymer self-assembly, a process well understood, can be used to produce well-defined, adjustable donor−acceptor nanostructures, which is relevant for applications. ASSOCIATED CONTENT

S Supporting Information *

Experimental details, additional DSC curves, scattering patterns, and TEM images.This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: (M.T.) [email protected]; (T.T.A.) [email protected]. Notes

The authors declare no competing financial interest. G

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