Molecular Orientation and Phase Separation by Controlling Chain

Sep 2, 2016 - The partial melting of PPP blocks promoted the extension of P3HT ... processes, which could alter the movement ability of P3HT molecule...
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Molecular Orientation and Phase Separation by Controlling Chain Segment and Molecule Movement in P3HT/N2200 Blends Rui Zhang,†,‡ Hua Yang,§,∥ Ke Zhou,†,‡ Jidong Zhang,† Xinhong Yu,† Jiangang Liu,*,† and Yanchun Han*,† †

State Key Laboratory of Polymer Physics and Chemistry, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences, 5625 Renmin Street, Changchun 130022, China ‡ University of the Chinese Academy of Sciences, No. 19A Yuquan Road, Beijing 100049, China § Institute of High Energy Physics, Chinese Academy of Sciences (CAS), Beijing 100049, China ∥ Dongguan Neutron Science Center, No. 1, Zhongziyuan Road, Dalang, Dongguan 523803, China S Supporting Information *

ABSTRACT: The molecular orientation and phase separation of poly(3-hexylthiophene) (P3HT) and poly[[N,N-bis(2-octyldodecyl)napthalene-1,4,5,8-bis(dicarboximide)-2,6-diyl]-alt-5, 5′-(2,2′-bithiophene)] (N2200) blends are investigated by in situ temperatureresolved two-dimensional grazing incidence X-ray diffraction (2D GIXD) in step-by-step heating and cooling processes. When pristine amorphous films are thermally annealed below the melting temperature of P3HT, only the chain segment of P3HT moves, while P3HT molecule is crystallized, induced by N2200 with face-on orientation (thermodynamics unsteady state). The interpenetrating network forms during the crystallization process of the polymers. When thermally annealed above melting temperature of P3HT, the whole molecular chain of P3HT moves, the space restriction of N2200 could be broken. The P3HT crystallizes with edge-on orientation (thermodynamics steady state) while N2200 crystallizes with face-on orientation. The hierarchical morphology forms by nucleation and growth or liquid phase separation. The P3HT crystallization with face-on or edge-on orientation in different postannealing processes is explained in terms of P3HT chain segment and molecule movement confined or broken out by N2200 crystalline confinement.

1. INTRODUCTION Conjugated polymers have generated significant interests due to their electronic proprieties and solution process.1−7 Because of their low entropy of mixing, conjugated polymer blends tend to phase-separate into relatively pure phases of the two components which is negative to organic electronic devices.8 Fortunately, it is possible to achieve microstructures by controlling the kinetics and thermodynamics of film formation such as changing chain structure,9 choosing different solvents,10 adjusting polymer molecular weights and blend ratios,11 and so on. There are two phase separation mechanisms for conjugated polymer blend: one is nucleation and growth phase separation and the other is spinodal decomposition.12 When two components are well mixed on a molecular scale, some other nanoscale phase separation takes place. If the conjugated blend polymer phase separation can be specifically controlled, it will be an efficient way to assemble structures on the nanometer scale. Because of the anisotropy of charge transport properties, from the conjugated polymer chain, the π−π stacking direction changing to the lamellar alkyl stacking direction, the charge transport efficiency decreases step-by-step.13,14 Lin et al. had © XXXX American Chemical Society

expounded that optimizing the packing of conjugated polymer chains to increase the crystallization was very important for their optoelectronic properties.15,16 For different optical devices, they are different molecular orientation requirements. As to field-effect transistors, the favored molecular orientation is edge-on, in which the direction of π−π stacking and backbones is beneficial for carrier transport.17 On the other hand, for photovoltaic devices, the face-on molecular orientation is preferred for a better carrier transport in the vertical direction of active layers.18,19 In our previous work,20 we also found that the same face-on molecular orientation in the interface promoted the photophysical processes such as photoinduced charge separation and inhibition of recombination. Therefore, to control the charge transfer efficiency of these π-conjugated photovoltaic devices, an understanding of how to adjust the molecular orientation in polymer blend system is important. Received: July 15, 2016 Revised: August 19, 2016

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DOI: 10.1021/acs.macromol.6b01526 Macromolecules XXXX, XXX, XXX−XXX

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other hand, when thermal annealing temperature is above the melting temperature of P3HT molecule, the whole chain of P3HT molecule could move. The P3HT molecule crystallized with edge-on orientation (thermodynamics steady state) while the N2200 molecule crystallized with face-on orientation. That might demonstrate the nucleation and growth or liquid phase separation while the hierarchical morphology formed.

During the controlling of the nanoscale phase separation structure by thermal annealing,21 adjusting dynamics of spincoating and the composition of the two components,22 the modification of molecular orientation was paid great attention by more and more researchers.23−26 Therefore, a large number of techniques have been applied to control the molecular orientation in order to get the maximum charge carrier mobility. For example, using a rubbed polyimide surface as an orientation layer,27 directional crystallization,28,29 hightemperature rubbing,30 and zone-casting or blade-coating methods.31,32 Kilwon Cho et al.12 reported that they controlled the regioregular poly(3-hexylthiophene) (P3HT) chains in the monolayer films to adopt two different conformations (edge-on and face-on) with the aim of elucidating the surface-induced molecular ordering in P3HT monolayer films. Michael D. McGehee et al.33 demonstrated that the chain orientation in P3HT was changed by infiltrating it into anodic alumina films with pore diameters in the range of 20−120 nm. The size of the anodic alumina determined the orientation of P3HT molecule. Orientation-induced crystallization is another effective way of controlling in-plane anisotropy and molecular orientations. The tuning of alignment or orientation of molecules by epitaxy34 is often achieved through using specific patterning substrates. Another effective method to control the molecular orientation is thermal annealing. In our previous work,35 the orientation transition and self-epitaxial crystallization of allconjugated diblock copolymers were systematically investigated. In the heating process, P3HT started to crystallize during the heating process followed by PPP blocks melting with edgeon orientation. During cooling, PPP blocks also crystallized with edge-on orientation, which was induced by the formed edge-on P3HT crystals via self-epitaxial crystallization. Then the as-casted film was heated in the melting temperature region of PPP blocks. The partial melting of PPP blocks promoted the extension of P3HT blocks, making P3HT blocks crystallize in a face-on orientation under the epitaxy from face-on PPP blocks. Because of the steric limitation effect, PPP blocks also crystallized with a face-on orientation via the self-epitaxy from P3HT blocks. Thus, the PPP and P3HT blocks both crystallized with face-on orientation after this circle. Furthermore, the face-on orientation transformed to thermodynamically stable edge-on orientation in the melting annealing process. There were two necessary requirements for selfepitaxial crystallization: one was the covalent bonding between blocks, and the other was P3HT transition from amorphous to crystalline through annealing. Combined to the previous work about the block polymers, in this work, the molecular orientation and its phase separation process in poly(3-hexylthiophene) (P3HT) and poly[[N,Nbis(2-octyldodecyl)-napthalene-1,4,5,8 -bis(dicarboximide)-2, 6-diyl]-alt-5, 5′-(2,2′-bithiophene)] (N2200) blend system was investigated by controlling the heating and cooling processes, which could alter the movement ability of P3HT molecule. It was convenient to observe the spontaneous orientation-induced crystallization without heterogeneity. On one hand, when the thermal annealing temperature is below the melting temperature of P3HT, only the chain segment of P3HT could move. So the crystallization of N2200 molecule might have space restriction for the crystallization of the P3HT molecule. This space restriction induced P3HT molecule crystallized with thermodynamics unsteady state of face-on orientation similar to N2200 molecule. The same crystal orientation favored the rise of fibrous phase separation. On the

2. EXPERIMENTAL SECTION 2.1. Materials. P3HT with a weight-average molecular weight of 50 kDa (PDI = 2.2) was purchased from Solarmer Materials. N2200 with a weight-average molecular of 84 kDa (PDI = 3.1) was purchased from Polyera Corporation, respectively. The structures are shown in Scheme 1. The blends were prepared by solution-casting using

Scheme 1. Molecular Structures of P3HT and N2200 Used in the Blend System

chloroform (CF) as a solvent which was purchased from Beijing Chemical Reagent Co. Ltd., China. The solvent was used without a further purification. 2.2. Sample Preparation. The two conjugated polymers were dissolved in CF with a weight ratio of 1:1. The total concentration of the blend system was 10 mg mL−1. After the solvent dissolved completely, the samples were spin-coated on a glass wafer with dimensions of 1.5 × 1.5 cm. Prior to spin-coating, the wafers were obtained by sonication in acetone for 30 min, sonication in ethanol for 30 min, cleaning with a 70/30 v/v solution of 98% H2SO4/30% H2O2 (Piranha solution (Caution!)) between 90 and 110 °C for 30 min and sonication in deionized water for 30 min and thoroughly rinsed. The substrates were subsequently dried in a flow of nitrogen. During the spin-coating process, we kept the solvent at about 55 °C. At the same time, the background was irradiated by filament lamp to keep the wafer at about 55 °C too. The condition of spin-coating was 4000 rpm for 30 s in the air which kept the thickness of films to about 150 nm. Characterization. Grazing incidence X-ray diffraction (GIXD), two-dimensional GIXRD were performed to characterize the morphologies and the crystalline structures of the films. Twodimensional GIXRD images were conducted at the BL1W1A beamline of Beijing Synchrotron Radiation Facility (BSRF) (λ = 1.54 Å) and the BL14B1 beamline of Shanghai Synchrotron Radiation Facility (SSRF) (λ = 1.24 Å). The incidence angle was 0.16 deg, and the exposure time was 100 s. In situ temperature-resolved synchrotron GIXD experiments were performed at BL1W1A beamline of BSRF. A TA Instruments Q2000 differential scanning calorimeter (DSC) was used to characterize the thermal properties of P3HT and N2200 at a heating/cooling rate of +10/−10 °C/min under fluid nitrogen. The thermal behaviors of these two polymers were characterized by DSC. The crystallization temperatures of P3HT and N2200 (Tc,P3HT and Tc,N2200) were measured to be about 190 and 250 °C, respectively. The melting temperature of P3HT (Tm,P3HT) was measured to be about 216 °C. The in situ ultraviolet−visible (UV−vis) absorption spectra of blend thin films were recorded on a Lambda 750 spectrometer (PerkinElmer, Wellesley, MA). Thermal annealing was performed using a TMS-94 B

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Figure 1. TEM images of blend films for different temperature thermal annealing: (a−d) thermal annealing below the melting temperature of P3HT; (e) thermal annealing above the melting temperature of P3HT. All of the samples were placed at the corresponding temperatures for 10 min.

molecules started to occur under such conditions. At 190 °C, the fibers became the most obvious (Figure 1d). At last, an interpenetrating network formed in the blend. When heated up to 250 °C (above the melting temperature of P3HT) sequentially, the interpenetrating network was damaged and a hierarchical morphology formed in the blend (Figure 1e). They had a large length scale dark-white phase separation while some small length scale separation also existed in the dark and white regions. When the samples were heated up from room temperature to 250 °C, the phase separation morphology evolved from interpenetrating network to hierarchical morphology. A similar phase separation evolution when the samples were heated up from room temperature to 190 °C was also investigated by using atomic force microscopy (AFM) phase images (Figure S1). To investigate the phase separation behavior further, we took the photoluminescence (PL) spectroscopy measurement (Figure S2). We chose a 532 nm solid-state laser as the excitation light source. It was found that the as-cast film had a low PL intensity, suggesting there was a better degree of mixing and few phase separations. As the temperature rose, the PL intensity increased step-by-step, which suggested that the phase separation enlarged. The result was in keeping with our TEM (Figure 1) and AFM (Figure S1) data. Then the evolution of phase separation during cooling down from different temperatures was investigated in Figure 2 and Figure 3. As shown in Figure 2, when cooling down from 190 °C to room temperature, the interpenetrating network phase separation remained unchanged. At the same time, when cooling down from 250 °C to room temperature, the hierarchical morphology also remained unchanged (Figure 3). This phenomenon suggested that once the different phase separation structures formed, they would not be destroyed during the cooling down process. It was worth noting that when increasing the thermal annealing temperature from 190 °C, fibrous microphase separation disappeared and large size phase separation appeared. The domain coarsened continuously along with the cooling process from above the melting temperature of P3HT.

hot stage (Linkaman) which was connected to a TMS-94 temperature controller. Prior to annealing, the chamber of the hot stage was purged several times with nitrogen. During the thermal annealing treatment, the samples were heated to 90, 160, 190, and 250 °C at 10 °C/min under a flow of nitrogen, respectively. After being placed at the corresponding temperatures for 10 min, they were cooled to different temperatures at 10 °C/min. The morphology of the blend films were investigated by transmission electron microscopy (TEM). The TEM images were obtained by a JEOL JEM-1011 transmission electron microscope operated at 100 kV.

3. RESULTS AND DISCUSSION 3.1. Development of Phase Separation Behavior by Controlling the Molecular Movement Ability in P3HT/ N2200. For many polymer blend systems, they often show poor compatibility because of their high molecular weight, which reduces their effective entropy of mixing. To investigate the relationship between crystallization and phase separation of the P3HT/N2200 blend system, a method of increasing the speed of films forming was adopted to suppress the crystallization and phase separation for the two polymer components in a pristine state. In this method, the solvent was heated up to approximately 55 °C. During the spin-coating process, a filament lamp was used to irradiate the background simultaneously which could ensure the temperature of background kept about 55 °C. Thus, the obtained film was kept in an amorphous state. Then, the evolution process of film crystallization and phase separation was controlled by thermal annealing with different temperatures. The process of phase separation morphology evolution during the heating up process was investigated as shown in Figure 1a−e; there was little phase separation in the pristine blend film using the fast spin coating process (Figure 1a). When the temperature was 90 °C, there was little phase separation change. When it was heated up from room temperature to 160 °C, some obvious fibrous morphology started to form in the film. At the same, some dark and white regions were also observed at this temperature. We guess it was because the two-phase boundary of P3HT and N2200 C

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crystallization process, the amorphous polymer is preferentially removed from the crystal. Nucleation and growth phase separation might take place, respectively. 3.2. P3HT Induced to Face-on Orientation by N2200 Space Restriction. In the following, the relationship between the crystallization of orientation and phase separation during the thermal annealing process for the blend system was investigated. Before probing the blend film, the orientation behavior of pure P3HT and N2200 phase was observed. From Figure 4a−c, when the pure P3HT film was annealed at 190 and 250 °C, the major diffraction was confirmed at the vertical direction at qz = 3.6 nm−1, corresponding to a d-spacing of 17.6 Å, which was the diffraction peak of (100) indicating the packing of alkyl side chains of P3HT. Smaller diffraction was confirmed at the horizon direction at qxy = 3.6 nm−1. That showed that the P3HT molecule mainly presented an edge-on orientation that was the thermodynamics steady state of the molecule. On the contrary, when thermally annealing the pure N2200 film at 250 °C, we observed only one diffraction peak confirmed at the horizon direction at qxy = 2.6 nm−1, corresponding to a d-spacing of 24.2 Å, which was the diffraction peak of (100) indicating the packing of side chains of N2200. This phenomenon indicated that the major orientation was face-on after thermal annealing of the N2200 molecule. To explore the crystallization orientation behavior of a twocomponent blend film, we employed in situ temperatureresolved GIXD measurements. All samples were followed by stepwise heating and cooling at the rate of +10/−10 °C/min under vacuum conditions. During the stepwise heating and cooling process for P3HT/N2200 blend film, the horizontal and vertical axis indicate the magnitude of scatting vector q in the xy plane and z directions, respectively. First, we discussed the transformation of the two molecules orientation below the melting temperature of P3HT. In Figure 5, at room temperature, because of the fast spin-coating, the two components did not have time to crystallize, and not much crystallization was observed. When the sample was heated up to

Figure 2. TEM images of blend films during cooling down process from 190 °C to room temperature. All of the samples were placed at the corresponding temperatures for 10 min.

At last, hierarchical morphology formed in the blend. It is known that the phase separation for polymer blend system is driven by a low entropy of mixing. When thermal annealing at the crystallization temperature of P3HT (below melting temperature), only the chain segment could move. The crystallization process of polymers came into play for the formation of fibrous phase separation. While thermal annealing above the melting temperature of P3HT, the whole molecular chain could move. It would be more likely that droplet nucleation and the successive growth by secondary nucleation take place. However, it could also be a liquid phase separation, since thin film exhibited a much higher mobility. During the

Figure 3. TEM images of blend films during cooling down process from 250 °C to room temperature. All of the samples were placed at the corresponding temperatures for 10 min. D

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Figure 4. Temperature dependence of 2D GIXD patterns for neat P3HT and N2200 film: (a) neat P3HT film at room temperature; (b) neat P3HT film at 190 °C; (c) neat P3HT film at 250 °C; (d) neat N2200 film at room temperature; (e) neat N2200 film at 250 °C.

Figure 5. Temperature dependence of 2D GIXD patterns for different temperatures below the melting temperature of P3HT.

90 °C, the intensity of P3HT(100) in the horizon direction at qxy = 3.6 nm−1 and N2200(100) in the horizon direction at qxy = 2.6 nm−1 started to appear. As the temperature increased, the intensity of diffraction peak of the two components at the horizon increased as seen in Figure 5, parts c and d. As the

cooling process went on, we observed that the diffraction peak intensity becomes increasingly strong, and there was only faceon orientation in the blend film. This meant that, in the blend film, a large proportion of P3HT molecules adopted a face-on orientation compared to those in the pure P3HT film, which E

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Figure 6. Temperature dependence of 2D GIXD patterns for different ratios of P3HT/N2200 at room temperature and 190 °C.

Figure 7. Temperature dependence of 2D GIXD patterns for different temperatures above the melting temperature of P3HT.

both pure and blend films. These results implied that, in blend film, the P3HT molecule crystallized with a face-on orientation

adopted an edge-on orientation (thermodynamics steady state), while the N2200 molecule adopted a face-on orientation in F

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Figure 8. Temperature dependence of 2D GIXD patterns for different ratios of P3HT/N2200 at room temperature and 250 °C.

appeared. Although there was no diffraction peak in P3HT:N2200 = 8:2 blend film at room temperature, the diffraction peak of N2200(100) in the horizon direction at qxy = 2.6 nm−1 and P3HT(100) in the horizon direction at qxy = 3.6 nm−1 appeared simultaneously. This had the same results compared to P3HT:N2200 = 5:5 after the heating and cooling stages. This phenomenon showed that a small number of N2200 molecules (20 wt %) could induce P3HT to face-on orientation. All of the results suggested that the crystallization of N2200 has the role of space restriction with the crystallization of P3HT molecule during the heating and cooling stage no matter how they blend with different proportions. The space restriction made a result that the P3HT molecule was induced to face-on orientation by the N2200 molecule. 3.3. P3HT Broke Restriction of N2200 to Edge-On Orientation. As mentioned above, if there was some space restriction to P3HT orientation for N2200, it would be possible to break this confine with enough external energy. Here, we attempted to increase the temperature to 250 °C (above the melting temperature of the P3HT molecule) during the heating up process, at which level the P3HT molecule was molten and the N2200 molecule will still crystallize. From Figure 7e, we could find that the diffraction peak of N2200(100) in the horizon direction at qxy = 2.6 nm−1 still existed, while the diffraction peak of P3HT(100) in the horizon direction at qxy = 3.6 nm−1 disappeared. That means at this temperature, the P3HT molecule was amorphous and the N2200 was crystallized with a face-on orientation. At meltdown process from this temperature (Figure 7f−i), the diffraction peak of P3HT in the vertical direction at qz = 3.6 nm−1 indicated a d-spacing of 17.6 Å, which was the (100) peak of the P3HT molecule, indicating that the P3HT molecule crystallized with an edge-on orientation. However, for the N2200 molecule, it crystallized with face-on orientation during the whole process, although above the melting temperature of P3HT molecule. The previous result suggested that, during the heating up process,

induced by the face-on N2200 crystallized grown during thermal annealing. As we all know, if two blend components have a good degree of mixing and similar interplanar crystal spacing, one material could crystallize along with another one which has a better orientation and crystallization degree. It is always a phenomenon of surface-induced crystallization. In the P3HT/ N2200 blend system, the d-spacing of the P3HT main chain was about 3.8 Å and the d-spacing of the N2200 main chain was about 14.4 Å. We realized that c(N2200) ≈ 4c(P3HT). The miss ratio of 4c(P3HT) and c(N2200) was about 0.05, which was less than 0.15. During the thermal annealing, the N2200 molecules crystallized with a face-on orientation. The crystallization of the P3HT molecule (mixing well with N2200 molecule) might be space restricted by N2200. It is possible that about four P3HT molecules were crystallized with face-on orientation induced by one N2200 molecule. In the following, the P3HT molecules could also crystallize along with their own π−π stacking direction. The crystallization of P3HT molecule might be space restricted by the crystallization of N2200 during the heating and cooling stage. There was a faceon orientation for P3HT like N2200, which was in an unsteady thermodynamics state. Then, the orientation of different P3HT/N2200 proportional blend systems was investigated. The total concentration of the blend solvent was 10 mg/mL, and the proportions of 2/8, 5/5, and 8/2 (P3HT/N2200) were selected. From Figure 6 ,we could find that when P3HT:N2200 = 2:8 at room temperature, the diffraction peak of N2200(100) in the horizon direction at qxy = 2.6 nm−1 appeared in the initial film while few diffraction peaks of P3HT(100) were observed. That was mainly because the proportion of N2200 was increased, higher than the point where much more N2200 molecule could crystallize during fast spin-coating. When the sample was heated to 190 °C, the intensity of the diffraction peak of N2200(100) in the horizon direction at qxy = 2.6 nm−1 increased while the diffraction peak of P3HT(100) in the horizon direction at qxy = 3.6 nm−1 G

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Macromolecules the movement ability of P3HT was better than N2200. After N2200 molecule crystallized, the crystallization of P3HT molecule would be limited in space by the crystallization of the N2200 molecule. So the P3HT molecule would crystallize with a face-on orientation similar to the N2200 molecule. However, when the temperature heated up above the melting temperature of the P3HT molecule, further external energy was provided for the movement of the P3HT molecule. In this condition, the P3HT molecule would move with more freedom for the whole molecular main chain. The space restriction of crystallized N2200 molecule became a weakness and the selforganized ability was increased. So the P3HT molecule tended to crystallize with an edge-on orientation, which was the stable structure for thermal annealing. That is to say, the orientation of P3HT molecule was not limited by the N2200 molecule any more. Then, we also investigated the orientation of different P3HT/N2200 proportions blend systems when thermal annealing above the melting temperature of P3HT. We could find a complete different result compared to thermal annealing below the melting temperature of P3HT as shown in Figure 8. During the thermal annealing process, the crystallization of P3HT molecule changed. The diffraction peak of P3HT in the vertical direction at qz = 3.6 nm−1 indicating a d-spacing of 17.6 Å which was the (100) peak of P3HT molecule appeared in three different blend ratios, indicating that the P3HT molecule crystallized with an edge-on orientation in this process. Under the circumstances, the P3HT molecule had enough space and opportunity for the chain arrangement which decreasing the influence of N2200 molecule crystallized with face-on orientation. In heating up process, the P3HT molecule crystallized with edge-on orientation which was the thermodynamic stability. Even though the N2200 molecule crystallized with face-on orientation all the time, it could not induce the P3HT molecule to crystallize with face-on orientation because of the increasing movement ability of the P3HT molecule above the melting temperature. Combined with the morphology evolution process previously, when thermal annealing took place below the melting temperature of P3HT molecule, the blend film started to have fibrous phase separation and the fibers were obviously compared to the pristine blend film during the cooling down stage (Figure 2). That might be because the crystallization of P3HT molecules played an important role in forming the interpenetrating network phase separation. Under such condition, both of the P3HT molecule and N2200 molecule crystallized with face-on orientation (Figure 5). When thermal annealing above the melting temperature of P3HT molecule took place, there was a hierarchical morphology formed with nucleation and growth or liquid phase separation (Figure 3). Under such conditions, the two components had different orientations, with which the P3HT molecule crystallized with an edge-on orientation and the N2200 molecule crystallized with a face-on orientation Figure 6i. In summary, we concluded that the same face-on orientations of P3HT and N2200 molecule played an important role in forming fibrous phase separation in the blend film. 3.4. P3HT Orientation Change Due to Its Better Molecular Diffusion Ability during Heating and Cooling Processes. The in situ temperature-resolved UV−vis spectrum was used to investigate the dynamics of the two molecule in blend film during thermal annealing. As shown in Figure 9, the absorption peaks of the N2200 molecule were 400 and 700 nm,

Figure 9. In situ ultraviolet−visible (UV−vis) absorption spectra of blend thin films during the heating and cooling process.

and the main absorption peak of the P3HT molecule was about 520 nm. The absorption peak at 400 nm for N2200 represented the delocalization of the π−π conjugate plane while the absorption peak at about 520 nm represented the main chain for the P3HT molecule. At room temperature, because of the fast spin-coating, the blend film had not enough time to crystallize, so it was amorphous in pristine film. The absorption peak of P3HT was one peak at about 520 nm. When heating up the temperature from room temperature to 190 °C, the absorption peak of P3HT at about 520 nm started to blue-shift gradually. At a temperature of 190 °C, the peak blue-shifted to 500 nm which was to the fullest extent. As the cooling process went on, this absorption peak of P3HT molecule then started to red-shift; when cooling to room temperature, the absorption peak of P3HT red-shifted to 520 nm again. At the same time, during the process, the triplet peak of P3HT appeared and became obvious. However, there were not any changes to the absorption peak of N2200 at 400 and 710 nm, no matter the heating or cooling stage. It was noteworthy that during the whole thermal annealing process, the P3HT molecule had a stronger ability of movement compared to N2200 molecule. According to the above results, we guessed that during this stage, it was the N2200 molecule that started to crystallize with a face-on orientation. When providing external energy, the P3HT molecule had a stronger ability of movement compared with the N2200 molecule. As the temperature heated up, the degree of crystallization for N2200 increased. So the movement of P3HT molecule would be limited to self-organization with freedom in space by the N2200 molecule with a face-on orientation. It could not crystallize with an edge-on orientation, which was the thermodynamics steady state, however.

4. CONCLUSION In conclusion, the phase separation evolution process and molecular orientation in P3HT/N2200 blend system was investigated by controlling the heating and cooling process. In fast spin-coating film, both molecules were amorphous. When the thermal annealing temperature is below the melting temperature of the P3HT molecule, only the chain segment of the P3HT molecule could move. The P3HT molecule was induced to crystallize with face-on orientation (thermodynamics unsteady state) which was space restricted by the faceon N2200 molecule. That might be because the crystallization of P3HT molecules played an important role in forming the interpenetrating network phase separation. When thermal H

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annealing took place above the melting temperature of P3HT, the whole molecular chain of P3HT could move, and the space restriction from the N2200 molecule could be broken. The P3HT molecule crystallized with edge-on orientation (thermodynamics steady state) while the N2200 molecule crystallized with face-on orientation. That might be a result of the nucleation and growth or liquid phase separation and the hierarchical morphology formed.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.6b01526. AFM phase images and PL spectra of blend films (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (J.L.). *E-mail: [email protected] (Y.H.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the National Natural Science Foundation of China (21334006, 51577138), the National Basic Research Program of China (973 Program 2014CB643505) and the Strategic Priority Research Program of the Chinese Academy of Sciences (Grant No. XDB12020300). We also thank Beijing Synchrotron Radiation Facility (BSRF) 1W1A and Shanghai Synchrotron Radiation Facility (SSRF) BL14U1 for 2D GIXD measurements.

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DOI: 10.1021/acs.macromol.6b01526 Macromolecules XXXX, XXX, XXX−XXX

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DOI: 10.1021/acs.macromol.6b01526 Macromolecules XXXX, XXX, XXX−XXX