Exploiting π–π Stacking for Stretchable Semiconducting Polymers

Article Cite This: Macromolecules 2018, 51, 2572−2579

Exploiting π−π Stacking for Stretchable Semiconducting Polymers Sung Yun Son,† Jae-Han Kim,‡ Eunjoo Song,† Kyoungwon Choi,† Joohyeon Lee,† Kilwon Cho,*,† Taek-Soo Kim,*,‡ and Taiho Park*,† †

Department of Chemical Engineering, Pohang University of Science and Technology (POSTECH), Pohang, Gyeongbuk 37673, Korea ‡ Department of Mechanical Engineering, Korea Advanced Institute of Science and Technology (KAIST), Daejeon 34141, Korea S Supporting Information *

ABSTRACT: Although intermolecular charge transport is known to occur via π−π stacking, the influence of π−π stacking on the mechanical properties of polymers has received little attention compared with other dynamic noncovalent interactions. Herein, we demonstrate a method to enhance stretchability via lowering crystallinity and increasing π−π stacking of thiophene-based random copolymer chains, which causes π−π stacking-induced polymer networks to form within the fully conjugated semiconducting polymer matrix. The polymer networks contain coiled amorphous chains that aid energy dissipation when the polymer film is subjected to strain; furthermore, the π−π stacking prevents the chains from irreversible sliding out of place due to the applied strain and provides interchain charge transport. Consequently, we are able to improve the polymer’s mechanical properties such as elongation at break, tensile strength, and toughness along with charge mobility. Additionally, our polymer shows great tolerance to a 40% strain without a decrease in mobility while maintaining a stable electrical performance even after 5000 stretching cycles at 30% strain.

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INTRODUCTION Noncovalent interactions have been used extensively to control the supramolecular assembly or self-assembly of materials.1−4 Among such interactions, hydrogen bonding and metal−ligand coordination are the most widely used in the field of supramolecular chemistry because they are sufficiently strong interactions and enable well-defined supramolecular architectures.5−7 Both interactions are also very important in polymer science because they are important in determining polymer microstructure and consequently polymer properties, especially mechanical properties.8−10 Moreover, owing to the dynamic nature, these interactions can give polymers new properties, such as self-healing, stimulus responsiveness, and reversibility.11,12 The introduction of noncovalent interactions into semiconducting polymers is reported to enhance stretchability and allow self-healing because any applied stress could be effectively dissipated by breaking the interactions.13,14 Thus, even cracks could heal through the recovery of the interactions. Like hydrogen bonding and metal−ligand coordination, π−π stacking between aromatic rings is another dynamic, spontaneously forming noncovalent interaction (Figure 1a).15 However, exploiting the dynamic nature of π−π stacking to induce supramolecular assembly remains relatively underdeveloped compared with hydrogen bonding and metal−ligand coordination. Instead, the aim of introducing π−π stacking has been mainly to enhance the charge mobility in semiconducting polymers because intermolecular charge transport readily occurs via π−π stacking.16−19 Given that π−π stacking is © 2018 American Chemical Society

dynamic and facilitates charge transport, one can postulate that π−π stacking may be used to simultaneously enhance charge mobility and tune the mechanical properties of semiconducting polymers. The ability to tune the mechanical properties of polymers has recently become important because it can lead to the realization of intrinsically stretchable semiconducting polymers.20−25 However, manipulating the mechanical properties of polymers without negatively affecting their charge mobility remains very challenging.26,27 For example, lowering the crystallinity of poly(3-hexylthiophene) (P3HT) generally leads to a decrease in the tensile modulus but an increase in elongation at break because, in the amorphous domains, chains can more easily stretch under tensile stress than in the crystalline domains (Figure S1a).28 However, lowering the crystallinity also results in a significant decrease in charge mobility because intermolecular charge transport is considerably restricted in amorphous domains by the random orientation of the polymer chains.28 Meanwhile, cross-linking of polymer networks can effectively enhance both the tensile modulus and elongation at break (Figure S1b).29 Note that as a result of noncovalent interactions such as hydrogen bonding, the tensile modulus of semiconducting polymers decreases in some cases upon Received: January 15, 2018 Revised: March 14, 2018 Published: March 22, 2018 2572

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compared the mechanical stability of P3HT and RP33 as a function of strain. Finally, we investigated the strain-dependent microstructures and strain-dependent electrical properties by using P3HT and RP33 films coated on flexible polydimethylsiloxane (PDMS) substrates.

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EXPERIMENTAL SECTION

Material Synthesis. Synthesis of P3HT, RP17, RP25, and RP33 was performed using the Grignard metathesis method according to the literature.32 The number-average molecular weights (polydispersity indices) were 16.4 (1.06), 16.2 (1.07), 16.1 (1.10), and 16.1 (1.15) kg/mol for P3HT, RP17, RP25, and RP33, respectively. UV−Vis Absorption Measurement. For temperature-dependent solution UV−Vis absorption measurement, 3 mL samples of toluene solutions (0.1 mg/mL) were characterized using a temperaturecontroller-equipped UV−vis spectrophotometer. Since the absorption spectra of all solutions exhibited one broad peak without any shoulder peaks at higher wavelength regions at 65 °C, which indicates complete dissolution of the polymers, 65 °C was chosen as the starting temperature for the experiments. Cooling rates varied depending on the temperature of the chiller. For thin-film absorption measurement, films were spin-coated from 1 mg/mL chlorobenzene solutions on a glass substrate at 2000 rpm for 60 s. AFM Measurement. All films were spin-coated on a Si substrate from chlorobenzene solutions (1 mg/mL) at 2000 rpm for 60 s. A series of solutions were maintained at room temperature before spincoating. The other was maintained at 60 °C for 30 min and then at −10 °C for 30 min before warming to room temperature and spincoating. The surface morphology of the films was investigated in the tapping mode using Si tips. GIWAXS and XRD Measurements. All films were spin-coated on a poly(styrenesulfonate) PSS/glass substrate from chlorobenzene solution at 2000 rpm for 60 s. GIWAXS experiment was conducted with a sample-to-detector distance of 212 mm, an X-ray radiation beam energy of 10.26 keV, and an incidence angle of 0.12°. All films were analyzed in a vacuum chamber. XRD experiments were performed with an X-ray radiation beam energy of 11.57 keV and an incidence angle of 0.12°. The measurement angle was gradually increased from 2° to 22° in 0.05° steps in the out-of-plane and inplane directions. Tensile Testing of the Thin Films. The semiconducting polymers were spin-coated onto a PSS/glass substrate. The watersoluble PSS was used as a sacrificial layer to form a free-standing semiconducting polymer film on the water surface. Bone-shaped tensile specimens were patterned using a femtosecond laser patterning method. The specimen was delaminated from the glass substrate by dissolving the PSS layer in water, allowing the specimen to float on the water surface. To grip the specimens, PDMS-coated Al grips were attached to the specimen using van der Waals adhesion. The tensile test was performed by applying tensile strain using a linear actuator, and stress and strain data were obtained by a load cell and a digital image correlation device, respectively. Five specimens of each polymer were analyzed. Microscopic Observation of Thin Films on a Stretchable Substrate. Stretching test of P3HT and RP33 was performed on a PDMS substrate. The film/PDMS structure was prepared by stamping the PDMS substrate to a floating thin film. The stretching behavior of the thin films under increasing strain was observed using an optical microscope. Strain-Dependent GIWAXS and Field-Effect Mobility Measurement. For strain-dependent GIWAXS measurements, polymer solutions were spin-coated onto a UV/ozone-treated (30 min) PDMS substrate with a thickness of 500 μm. The prepared films were stretched (ε ∼ 30%) using a stretching jig and released prior to measurement. For strain-dependent field-effect mobility measurement, all films were spin-coated on a PDMS substrate from chlorobenzene solution. The stretched polymer films at different strains were contacted to heavily doped n-type Si wafers with 300 nm thick SiO2, onto which an ODTS self-assembled monolayer had been grown

Figure 1. (a) π−π stacking between thiophene oligomers. (b) Chemical structures of P3HT, RP17, RP25, and RP33. For RP17, RP25, and RP33, y can take the value of 1, 2, 3, 4, or more.

introducing chemical moieties for cross-linking because the introduced moieties can reduce the rigidity of the backbone.13 Although cross-linking of polymer chains is effective for enhancing the mechanical properties, it decreases the charge mobility of semiconducting polymers because the cross-linking can prevent the desired chain packing for intermolecular charge transport.30,31 In addition, for cross-linking via noncovalent interactions, the intrinsic charge mobility of semiconducting polymers decreases with increasing fraction of chemical moieties because of the insulating properties of such moieties and the rupture they cause in conjugation.13 Recently, we reported that incorporating thiophene units into the backbone of poly(3-hexylthiophene) (P3HT) improves charge mobility via random copolymerization of 3hexylthiophene and thiophene monomers.32 We used thiophene contents of 17, 25, and 33 mol % to obtain RP17, RP25, and RP33, respectively (Figure 1b). The improved charge mobility is attributed to enhanced π−π stacking in amorphous domains, in turn resulting from the random incorporation of thiophene units, which also facilitates close chain packing (Table S1). Meanwhile, the incorporated thiophene units also reduce the structural regularity and thus the crystallinity. Therefore, we reasoned that if the π−π stacking between chains in amorphous domains in RP17, RP25, and RP33 can not only act as junctions for intermolecular charge transport but also cross-link the chains to form a polymer network, the tensile modulus should increase along with the charge mobility in spite of the decrease in crystallinity (Figure S1c). Moreover, if π−π stacking in the polymer network helps dissipate tensile energy (as is the case for other noncovalent interactions13,14), elongation at break should also increase. Thus, in the present report, we discuss how π−π stacking affects the mechanical properties of semiconducting polymers. For this study, we first accurately measured the mechanical properties of P3HT, RP17, RP25, and RP33 thin films by applying direct tensile tests on the surface of water. Next, we 2573

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Figure 2. Normalized temperature-dependent UV−vis absorption spectra of (a) P3HT, (b) RP17, (c) RP25, and (d) RP33 in toluene upon cooling from 65 to −10 °C. (e) Tapping mode AFM images of the as-cast spin-coated films prepared from solutions kept at 60 °C for 30 min, then −10 °C for 30 min, and finally warmed to room temperature before spin-coating. Insets are tapping mode AFM images of as-cast spin-coated films prepared from solutions kept at room temperature before spin-coating. and source and drain electrodes (Au) had been deposited. The channel length (L) and width (W) were 50 and 1000 μm, respectively. The prepared OFET devices were characterized using Keithley 4200 units under ambient conditions. The transfer curves (ID−VG) were obtained in the saturation regime. At least five devices were measured at each strain. In the repeat-strain tests, the device performance was measured after releasing the films back to 0% strain.

another 30 min prior to warming to room temperature and spin-coating because this process provides thermodynamically suitable conditions for polymer chains to pack in a face-to-face manner.33,34 The resulting morphologies of the films are very different (Figure 2e). Unlike P3HT and RP17, which have smooth surfaces, RP25 has a small number of short nanofibrils embedded in its microstructure. Interestingly, RP33 forms very long and thick nanofibrils. Meanwhile, all the films show similar surface morphologies when solutions were maintained at room temperature before spin-coating (insets in Figure 2e). This result clearly confirms the increased π−π stacking propensity upon increasing the number of incorporated thiophene units. Microstructures and Crystallinity. The changes in semicrystalline microstructures of thin films caused by the random incorporation of thiophene units were systematically investigated using X-ray diffraction (XRD) and grazing incidence wide-angle X-ray scattering (GIWAXS). The GIWAXS pattern of P3HT presents distinct out-ofplane (h00) peaks indicating long-range, well-ordered lamellar structures (Figure 3a). In contrast, the (200) and (300) peaks are hardly observed in the GIWAXS patterns of RP17, RP25, and RP33, indicating that the incorporation of thiophene units significantly lowers their structural regularity and consequently their regular lamellar packing. This tendency is also clearly observed in the out-of-plane XRD spectra (Figure 3b). Since the crystallinity of a polymer thin film is commonly in line with its degree of lamellar packing, these results suggest that P3HT has a high degree of thin film crystallinity, whereas RP17, RP25, and RP33 have very low thin film crystallinity. The out-of-plane (100) peaks are observed at 3.75°, 4°, 4.4°, and 4.5°, corresponding to the dspacings of 16.38, 15.35, 13.96, and 13.65 Å for P3HT, RP17, RP25, and RP33. The diminution of the d-spacing indicates that the reduced side-chain bulkiness upon the incorporation of thiophene units promotes edge-to-edge chain accessibility. In the in-plane XRD spectra, only P3HT shows distinct (100) and (010) peaks representing regular in-plane lamellar packing and π−π stacking, respectively (Figure 3c). This also confirms the

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RESULTS AND DISCUSSION Evaluation of Degree of π−π Stacking. We used UV−vis absorption spectroscopy to monitor the π−π stacking propensity of the polymers in solution by lowering the solution temperature because upon lowering the solution temperature the chains pack more closely in a face-to-face manner.33,34 At approximately λ = 450 nm, the P3HT and RP17 solutions exhibit an absorption maximum associated with intrachain π−π* transitions,35,36 and no distinct shoulders appear at higher wavelengths (Figure 2a,b). This result indicates that the polymers dissolve well over the entire temperature range investigated and do not form aggregates.37,38 The slight redshift upon cooling from 65 to −10 °C is attributed to reduced backbone torsion at a lower temperature.39 In the absorption spectra of RP25 and RP33, distinct shoulders appear at approximately 600 nm at −10 and 50 °C, respectively (Figure 2c,d), which is due to enhanced interchain π−π* transitions in the solution state.39,40 This is particularly the case for RP33 because the reduced torsional rotation of thiophene rings with decreasing temperature induces planarization of the backbone, leading to enhanced π-orbital overlap between neighboring backbones. Furthermore, the absorption spectra of the as-cast films also present increasing shoulder peak intensity with increasing thiophene content (Figure S2). These results clearly reflect the increasing degree of π−π stacking with increasing incorporation of thiophene units. To confirm the π−π stacking propensity of the polymers, a nanofibril formation experiment was conducted since nanofibrils are generally formed by π−π stacking of extended polymer backbones along the nanofibril axis.41−43 Polymer solutions were kept at 60 °C for 30 min and then at −10 °C for 2574

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RP33 exhibits the highest degree of π−π stacking but much lower crystallinity than P3HT. This indicates that π-orbital overlap between polymer chains predominantly occurs in the amorphous domains in RP33 thin films but in the crystalline domains in P3HT thin films. This is because the face-to-face packing between backbones in the amorphous domains in P3HT is hindered by its side chains, whereas significant sections of the backbone in RP33 can orient face-to-face in the amorphous domains due to the reduced side-chain bulkiness Mechanical Properties. Mechanical properties of P3HT, RP17, RP25, and RP33 were measured to investigate the correlation between the microstructures and the mechanical properties. Because measuring the mechanical properties of semiconducting polymers by conventional methods is extremely difficult, we applied direct tensile tests to freestanding films floating on the surface of water to investigate the effect of π−π stacking on the mechanical properties of P3HT, RP17, RP25, and RP33 (Figure S3). This technique allows facile and accurate measurement of the mechanical properties of semiconducting polymer thin films.28 The stress−strain curves and measured mechanical properties appear in Figure S4 and Table S2, respectively. Figure 4a shows that the tensile moduli increase with increasing thiophene content. Since tensile moduli generally decrease with decreasing crystallinity,27,28,44 this result implies that upon incorporation of the thiophene units polymer networks are formed via π−π stacking, which prevents the chains from sliding out of place and thereby increases the resistance to mechanical deformation. Consequently, the robustness of the polymer network improves with increasing thiophene content. In addition, the elongation at break of the thin films is also gradually enhanced from P3HT to RP33 (Figure 4b). In general, elongation at break is known to increase as the crystallinity of a thin film decreases because an applied stress is dissipated by stretching the amorphous chains.27,28 Furthermore, regular chain arrangement in crystalline regions maximizes intermolecular interactions therein, leading to cracking at the boundaries of such domains.28,45 Thus, P3HT should exhibit the lowest elongation at break. However, the elongation at break still increases from RP17 to RP33, although the intensities of the (100) peaks remain similar (Figure 3b), which

Figure 3. (a) GIWAXS patterns; (b) out-of-plane and (c) in-plane XRD spectra of P3HT, RP17, RP25, and RP33 thin films.

highest thin film crystallinity of P3HT among the polymer studied in this study. Interestingly, we found that P3HT exhibits the lowest degree of π−π stacking among the polymers in this study as demonstrated by the temperature-dependent UV−vis absorption spectroscopy and nanofibril formation experiments, but the highest degree of π−π stacking in its crystalline domains owing to the regular chain arrangement therein. On the other hand,

Figure 4. (a) Tensile modulus, (b) elongation at break, (c) tensile strength, and (d) toughness of P3HT, RP17, RP25, and RP33 thin films. (e) Schematic diagrams illustrating mechanisms for energy dissipation during strain and for recovery of microstructure after release of stretching in a polymer network featuring π−π stacking. 2575

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Figure 5. Optical microscopy images of P3HT and RP33 thin films under strain from 0% to 30% and 60% for P3HT and RP33, respectively.

means that the crystallinity of the thin films remains similar. Therefore, these results suggest that energy dissipation becomes more efficient with increasing thiophene content. Because the incorporation of thiophene units decreases the crystallinity and increases π−π stacking, the mechanism for energy dissipation can be explained by two steps (Figure 4e). At the onset of stretching, the energy is dissipated by stretching coiled chains in amorphous domains. Next, as more tensile stress is applied, energy is dissipated by the gradual disruption of π−π stacking (applied stress effectively dissipates upon breaking reversible noncovalent interactions13,14). These results indicate that the polymer thin films become more robust but more ductile with increasing content of thiophene units. Consequently and surprisingly, both tensile strength and toughness, which heretofore have been considered mutually exclusive properties,46 increase with increasing thiophene content, indicating that the mechanical stability of RP33 is significantly enhanced compared with that of P3HT (Figure 4c,d). Given the large difference in the measured mechanical properties of P3HT and RP33, we used optical microscopy to monitor the surface of thin films on a PDMS substrate under strain to compare their mechanical stability (Figure 5). At 10% strain, the optical microscopy images of P3HT thin films already reveal many cracks, which propagate further as the strain increases to 30%. Conversely, no cracks appear in RP33 thin films, even at 30% strain. Tiny cracks begin to appear at 40% strain and gradually propagate with increasing strain. This confirms that energy dissipation is effective in RP33, which is attributed both to the stretching of coiled amorphous chains and to the disruption of π−π stacking. Strain-Dependent Microstructures and Field-Effect Mobility. Furthermore, to understand the semicrystalline microstructure upon stretching, we investigated the straindependent microstructural features of P3HT and RP33 thin films as a function of the number of stretches. The procedure for this experiment is depicted in Figure S5. After spin-coating polymer solutions on flexible PDMS substrates, the films were stretched to 30% strain by using a stretching jig. Next, the microstructures of the films were characterized by using GIWAXS. For easy comparison, the out-of-plane GIWAXS spectra were extracted from the GIWAXS patterns (see Figure 6a,b and Figure S6 for the out-of-plane spectra of P3HT, RP33, and their

Figure 6. Out-of-plane GIWAXS spectra of (a) P3HT and (b) RP33 thin films as a function of number of stretches (0, 5, 10, 50, and 100 at 30% strain).

GIWAXS patterns, respectively). For P3HT films, the intensity of the (100) peak rapidly decreases as the films are stretched five times (Figure 6a). This result indicates that the crystalline regions are destroyed or transformed into amorphous regions by stretching. Furthermore, the intensity of the (100) peak decreases further with increasing number of stretches. Conversely, the intensities of the (100) peak in RP33 remains constant, even after 100 stretches, which indicates that the semicrystalline microstructure does not change as a result of repeated stretching (Figure 6b). This result suggests that the two-step energy dissipation by coiled amorphous chains and π−π stacking effectively prevents permanent microstructural deformation. Given the superior robustness of the RP33 microstructure, we expect that the electrical characteristics of RP33 are also 2576

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CONCLUSIONS We demonstrate that the mechanical properties of P3HT, RP17, RP25, and RP33, and especially their stretchability, may be enhanced by exploiting π−π stacking. This work shows that the charge mobility and the mechanical properties can be simultaneously enhanced when the π−π stacking between chains in amorphous domains can not only act as junctions for intermolecular charge transport but also cross-link the chains to form a polymer network. We expect that this approach may also be applied to other high-mobility semiconducting polymers, such as diketopyrrolopyrrole- or isoindigo-based polymers.

more robust against tensile strain than is the case for P3HT. Therefore, the mobility of the stretched P3HT and RP33 films along the strain direction in a field-effect-transistor (FET) configuration was evaluated by varying the strain (Figure 7). The measured strain-dependent transfer curves for the P3HTand RP33-FETs are shown in Figure S7.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.8b00093. Experimental details, Figures S1−S8, and Tables S1−S4 (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected], phone +82-54-279-5295, fax +82-54-279-8293 (T.P.). *E-mail [email protected] (T.-S.K.). *E-mail [email protected] (K.C.). ORCID

Kilwon Cho: 0000-0003-0321-3629 Taek-Soo Kim: 0000-0002-2825-7778 Taiho Park: 0000-0002-5867-4679

Figure 7. (a) FET mobility along the strain direction in P3HT- and RP33-FETs as a function of strains from 0% to 40% and 60% for P3HT and RP33, respectively. (b) Normalized FET mobility along the strain direction as a function of number of stretching cycles (ε = 0% → 30% → 0%). μ0 is the initial FET mobility before stretching.

Present Address

J.-H.K.: Korea Atomic Energy Research Institute, Daejeon, 34057, Korea. Author Contributions

S.Y.S., J.-H.K., and E.S. contributed equally to this work.

When the P3HT films are stretched to 10% strain, the FET mobility significantly decreases and finally reaches ∼10−8 cm2 V−1 s−1 at 40% strain (Figure 7a and Table S3). This result agrees with the optical microscopy observations of severe cracks in P3HT at 10% strain that propagate further upon increasing tensile strain. Conversely, the RP33 film tolerates up to 40% strain with a negligible decrease in FET mobility. Furthermore, under repetitive stretching, the FET characteristics of RP33 are more stable than those of P3HT (Figure 7b and Table S4). The ratio μ/μ0 (where μ0 is the initial mobility) of the P3HT FETs rapidly decreases with increasing number of stretching cycles, and the device ceases to function after 300 stretches. This result is consistent with the observations of the semicrystalline microstructure failure in P3HT upon repetitive stretching. Conversely, RP33-FETs retain normal FET characteristics even after 5000 stretches. Since the semicrystalline microstructure does not change and no cracks appear in the films after repetitive stretching (Figure 6b and Figure S8), the slight decrease in the mobility indicates that some π−π stacking that is disrupted upon stretching is not recovered once the strain is released. This result shows clearly the excellent capacity of RP33 to dissipate strain energy, which imbues the polymer with excellent mechanical and electrical robustness against external deformation.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS GIWAXS and XRD measurements were performed at a synchrotron radiation on the beamline 3C and 5A, respectively, at the Pohang Accelerator Laboratory (PAL), Korea. This work was supported by the Technology Development Program to Solve Climate Changes of the National Research Foundation of Korea (NRF) (2016M1A2A2940912), Center for Advanced Soft Electronics under the Global Frontier Research Program (Code No. NRF-2012M3A6A5055225 and Code No. 20110031628), and Wearable platform materials technology center (2016R1A5A1009926) funded by NRF under the Ministry of Science, ICT and Future Planning.

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DOI: 10.1021/acs.macromol.8b00093 Macromolecules 2018, 51, 2572−2579