Structure–Property Relationships of Semiconducting Polymers for

Nov 1, 2017 - Department of Science Education, Ewha Womans University, Seoul 03760, ... Department of Chemistry, Kunsan National University, Kunsan-si...
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Structure-Property Relationships of Semiconducting Polymers for Flexible and Durable Polymer Field-Effect Transistors Min Je Kim, A-Ra Jung, Myeongjae Lee, Dongjin Kim, Suhee Ro, SeonMi Jin, Hieu Dinh Nguyen, Jee Hye Yang, Kyung-Koo Lee, Eunji Lee, Moon Sung Kang, Hyunjung Kim, Jong-Ho Choi, BongSoo Kim, and Jeong Ho Cho ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b12435 • Publication Date (Web): 01 Nov 2017 Downloaded from http://pubs.acs.org on November 4, 2017

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ACS Applied Materials & Interfaces

Structure-Property Relationships of Semiconducting Polymers for Flexible and Durable Polymer Field-Effect Transistors Min Je Kim,1,† A-Ra Jung,2,† Myeongjae Lee,3,† Dongjin Kim,4 Suhee Ro,2 Seon-Mi Jin,5Hieu Dinh Nguyen,6 Jeehye Yang,7 Kyung-Koo Lee,6 Eunji Lee,5 Moon Sung Kang,7 Hyunjung Kim,4 Jong-Ho Choi,3 BongSoo Kim,2,* and Jeong Ho Cho1,* 1

SKKU Advanced Institute of Nanotechnology (SAINT), School of Chemical Engineering,

Sungkyunkwan University, Suwon 16419, Republic of Korea 2

Department of Science Education, Ewha Womans University, Seoul 03760, Republic of

Korea M. Lee, Prof. J.-H.Choi 3

Department of Chemistry, Korea University, Seoul 02841, Republic of Korea

4

Department of Physics, Sogang University, Seoul 121-742, Republic of Korea

5

Graduate School of Analytical Science and Technology, Chungnam National University,

Daejeon 34134, Republic of Korea 6

Department of Chemistry, Kunsan National University, Kunsan-si 54150, Republic of Korea

7

Department of Chemical Engineering, Soongsil University, Seoul 06978, Republic of Korea

†

M. J. Kim, A.-R. Jung, and M. Lee contributed equally.

Keywords: Flexible field effect transistors, Structure-property relationship, Organic semiconductors, Carrier mobility, Mechanical stability

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Abstract We report high-performance top-gate bottom-contact flexible field-effect polymer transistors (FETs) fabricated by flow-coating diketopyrrolopyrrole (DPP)-based and naphthalene diimide (NDI)-based polymers (P(DPP2DT-T2), P(DPP2DT-TT), P(DPP2DTDTT), P(NDI2OD-T2), P(NDI2OD-F2T2), and P(NDI2OD-Se2)) as semiconducting channel materials. All of the polymers displayed good FET characteristics with on/off current ratios exceeding 107. The highest hole mobility of 1.51 cm2V-1s-1 and the highest electron mobility of 0.85 cm2V-1s-1 were obtained from the P(DPP2DT-T2) and P(NDI2OD-Se2) polymer FETs, respectively. The impacts of the polymer structures on the FET performance are well explained by the interplay between the crystallinity, the tendency of the polymer backbone to adopt an edge-on orientation, and the interconnectivity of polymer fibrils in the film state. Additionally, we demonstrated that all of the polymer-based flexible FETs were highly resistant to tensile stress, with negligible changes in their carrier mobilities and on/off ratios after a bending test. Conclusively, these high-performance, flexible, and durable FETs demonstrate the potential of semiconducting conjugated polymers for use in flexible electronic applications.

1. Introduction A number of organic semiconductors have been developed for field-effect transistor (FET) applications such as foldable displays,1 sensors,2,

3

radio-frequency identification

(RFID) tags,4and complementary circuit elements.5-7 Organic semiconductors can meet the requirement of large-area electronics by solution processing. They can also have high mechanical stability under bending and stretching.6,8,9 Semiconducting small molecule- and polymer-based FETs have already demonstrated high carrier mobilities (> 10 cm2V-1s-1), exceeding those of amorphous silicon-based FETs (0.5‒1.0 cm2V-1s-1).8,10-13 Recently developed high-performance semiconducting polymers typically have donor-acceptor type aromatic backbones.2,14 p-Type semiconducting polymers are commonly based on diketopyrrolopyrrole (DPP) 8,10,11,15 or benzothiadiazole (BT)11,16-18 acceptor moieties, while n-type semiconducting polymers are commonly based on naphthalene diimide (NDI) and perylene diimide (PDI) acceptor moieties.7, 19 In addition to material developments, novel processing methods have also significantly improved FET performance. For instance, bar-coating, 7,16 shear-coating,13,20 and flow-coating methods21 have produced polymer FET devices with largely improved carrier mobilities. However, these improvements have typically been obtained only when hard silicon (Si) wafers or glass substrates are used. These substrates provide flat surfaces and enable 2 ACS Paragon Plus Environment

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high-temperature treatment of the semiconducting polymer films above 200 °C. With the aid of these features, polymer semiconductors can be assembled in a well-controlled manner. However, when the polymers are deposited on flexible substrates through solution-based processes, the resulting device performances are less impressive. A record carrier mobility obtained with a flexible substrate is just around 1 cm2V-1s-1, and only a few related results have been reported.5-7 The lower mobilities of polymer semiconductors on flexible substrates can be attributable to the high surface roughness and poor thermal stability of plastic substrates. For instance, the 125-µm-thick polyethylene naphthalate (PEN) flexible substrate (Teijin DuPont Films, Teonex Q65HA) that have been often used5 has an rms roughness of 4.54 nm (Figure S1) and is only thermally stable up to 150 °C (glass transition temperature (Tg) = 120 °C). In order to develop further-advanced flexible devices, it is essential to develop polymer semiconductors demonstrating high-performance even under such limitations. Moreover, the polymer-structure-electrical property relationship should be established particularly for flexible devices processed at low temperatures. Here, we report high-performance polymer FETs prepared by flow-coating a series of three DPP-based semiconductors and a series of three NDI-based semiconductors onto plastic substrates. A comparative study on these polymer films, after they have been treated at a moderate temperature (150 °C), reveals how the chemical structure, film crystallinity, and morphology affect the electrical properties of polymer FETs. Note that we achieved a maximum hole mobility of 1.51 cm2V-1s-1 and a maximum electron mobility of 0.85 cm2V-1s-1 from the highly edge-on oriented polymers of P(DPP2DT-T2) and P(NDI2OD-Se2), respectively. Additionally, all of the polymer FETs demonstrated excellent mechanical stability and high bias stability.

2. Results and Discussion Figure 1 shows the schematic device structure of the top-gate bottom-contact (TGBC) flexible polymer FETs as well as the chemical structures of the materials employed in this study. Three DPP-based semiconductors (P(DPP2DT-T2), P(DPP2DT-TT), and P(DPP2DTDTT)) and three NDI-based semiconductors (P(NDI2OD-T2), P(NDI2OD-F2T2), and P(NDI2OD-Se2)) were utilized as semiconducting channel materials. Details of the polymer syntheses are provided in the Supporting Information (Figures S2-S19). These semiconducting polymer layers were formed on PEN substrates (with pre-patterned Cr/Au source and drain electrodes) by flow-coating of polymer solutions.21 The as-coated polymer films were annealed at 100 and 150 °C. We note that although thermal annealing of the 3 ACS Paragon Plus Environment

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polymer films was applied to improve the FET performance,15,22 annealing temperature was not increased beyond 150 °C as it would cause thermal damage to the PEN substrate. The poly(methyl methacrylate) (PMMA) gate dielectric layer, with a thickness of around 500 nm, was then spin-coated onto the semiconducting layers. Finally, 40-nm-thick Al was deposited thermally through a shadow mask onto the channel region to form the gate electrode. Figure 2a shows the representative transfer characteristics of the FETs made with the DPP-based and NDI-based semiconductors annealed at three different temperatures: 25 °C (as-coated), 100 °C, and 150 °C. Table 1 and Table 2 summarize the carrier mobilities, threshold voltages, and on/off current ratios of the devices. V-shaped transfer curves, indicating ambipolar charge transport, were observed for all of the polymers. The DPP-based polymers exhibited hole-dominant transport, whereas the NDI-based polymers showed electron-dominant transport. Hysteresis during the forward and reverse sweeps was negligible in the transfer curves. The carrier mobility (µ) of the FETs was evaluated in the respective saturation regimes according to the equation ID = CS·µ·W·(VG – VTH)2/2L, where CS is the specific capacitance of the gate dielectric layer, W is the channel width, VG is the gate voltage, VTH is the threshold voltage, and L is the channel length.23 For the DPP-based polymers, the as-coated P(DPP2DT-DTT) film showed the average hole mobilities of 0.013 cm2V-1s-1. The as-coated P(DPP2DT-TT) film exhibited an average hole mobility that was twice as high (i.e., 0.026 cm2V-1s-1). The top-performing polymer among the DPP series was P(DPP2DT-T2), which exhibited an average hole mobility of 0.089 cm2V-1s-1. The electron mobilities of the DPP-based FETs were approximately 5-100 times lower than the hole mobilities. This strong asymmetry between hole and electron mobilities was attributed mainly to the lower hole injection barrier from the Au electrode (with a high work function of ~5.1 eV) to the semiconductor;24, 25 this will be further explained below. After thermal annealing at 100 and 150 °C, the electrical properties, including the carrier mobility and on-off current ratio (Ion/Ioff), were improved significantly. The 150 °C-annealed P(DPP2DT-T2) films exhibited a maximum hole mobility of 1.51 cm2V-1s-1 (with an average hole mobility of 1.43 ± 0.13 cm2V-1s-1), which is the highest hole mobility ever reported for a flexible polymer FET device. The NDI-based FETs exhibited n-channel-dominant characteristics with weak hole transport, which might be caused by the long hole-hopping distance between oligothiophene bridges in the twisted polymer backbone; this will be discussed in greater detail below. The as-coated P(NDI2OD-T2) and P(NDI2OD-F2T2) films showed the average electron mobilities of 0.020 and 0.027 cm2V-1s-1, respectively. The P(NDI2OD-Se2) film yielded the highest average electron mobility (0.050 cm2V-1s-1) among the NDI series. Similar to the 4 ACS Paragon Plus Environment

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DPP-based FETs, thermal annealing improved the carrier mobilities considerably. The maximum electron mobility of 0.85 cm2V-1s-1 (with an average electron mobility = 0.61±0.20 cm2V-1s-1) was obtained from the 150 °C-annealed P(NDI2OD-Se2), which is the highest electron mobility ever reported for a flexible polymer FET device. The slight hysteresis in ntype NDI semiconductors was originated from the fact that electrons were easily trapped in the trap sites generated by hydrogenated oxygen.26 Figure 2b and 2c summarize the maximum carrier mobilities of the 150 °C-annealed DPP and NDI samples, respectively. All devices exhibited the extremely low leakage current level of 0.1 nA (Figure S20). The output characteristics of the 150 °C-annealed P(DPP2DT-T2) and P(NDI2OD-Se2) FETs (Figure S21) showed a transition from non-Ohmic behavior at a low VG to saturation behavior at a high VG. The electrical properties of the FETs were strongly affected by the energy level alignment of the frontier orbitals relative to the electrode Fermi level, conformation of the polymer backbone, film crystallinity, and morphology. First, the backbone geometries and electronic structures of the semiconducting polymers were investigated by performing quantum mechanics calculations using the Gaussian 09 package at a density functional theory (DFT) level. To reduce calculation costs, long alkyl chains of 2-decyltetradecyl or 2octyldodecyl groups were replaced with methyl groups, and the three monomeric structures of the polymers were used. Becke’s three-parameter, gradient-corrected function (B3LYP) with a 6-311G(d) at ground state was used as the basis sets for full geometry optimization and energy calculations. Figure 3 displays the DFT-optimized molecular geometries. All of the DPP-based trimers were planar. The dihedral angles between DPP moieties and the neighboring thienyl groups were below 1.2°. Slightly twisted conformations occurred at the connecting bonds between the thienyl groups and oligothiophene bridges (< 15°). The highest dihedral angle of 20.2° was observed in the bithiophene units in the (DPP2DT-T2)3. Note that, although there were some differences in the dihedral angles of the DPP-based trimer series, all of the angles were lower than 20°; these differences would not be a critical factor.27 Alternatively, the NDI-based trimers exhibited highly twisted backbones. The main twists occurred between NDI moieties and neighboring thienyl groups where the dihedral angles were over 40°, which was presumably due to the steric hindrance between the C-H bond in the thienyl group and the C=O bond in the DPP unit. The slightly higher dihedral angles (~48°) between NDIs and neighboring selenophenes might be caused by the stronger repulsion between the lone-pair electrons of the electron-rich, large-sized selenophenes and C-H bonds in the NDI structures. Moreover, different dihedral angles were observed 5 ACS Paragon Plus Environment

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depending on the bridging groups: 2,2’-bithiophene (T2), 3,3’-difluoro-2,2’-bithiophene (F2T2), and 2,2’-biselenophenes (Se2). The insertion of fluorine atoms into the T2 units led to the more planar geometry (dihedral angle: 3.6° vs. 25° for F2T2 and T2 structures) because of F···S interactions.16, 28-31 Replacing sulfur with selenium resulted in intermediate planarity (dihedral angles of 26°, 2.5°, and 28° for Se2 structures). These dihedral angle characteristics were directly reflected in the frontier orbital overlapping. The surface plots of the HOMO and LUMO orbitals are shown in Figure 4, and the corresponding energy levels are summarized in Figure 5a and Table 3. The planar DPP-based trimers displayed good electronic overlapping in both the HOMO and LUMO orbitals along the molecular backbone. In contrast, the highly twisted NDI-based trimers revealed rather localized electronic structures; the electron-rich thiophene or selenophene bridge units contributed mainly to the HOMO orbitals, while the electron-deficient NDI structures contributed to the LUMO orbitals. Both the intrinsic electron richness of the constituting chemical moieties and the degree of electronic overlap led to a bandgap of ~1.7 eV for the DPP-based trimers and a bandgap of ~2.0 eV for the NDI-based trimers. In addition, the energy levels and bandgaps were experimentally measured using cyclic voltammetry (CV) and UV-visible absorption spectroscopy. Figure S22 shows the CV curves for the six semiconducting polymers. The DPP-based and NDI-based polymers exhibited only clear oxidation and reduction, respectively. Thus, the HOMO and LUMO levels were well-resolved for DPP-based and NDI-based polymers, respectively; unresolved energy levels were estimated by considering the optical bandgaps (Eg,Opts) that were determined from the UV-visible absorption edges of the polymer films. The HOMO and LUMO levels of the polymers are summarized in Figure 5a. Note that, because the optical bandgaps can be smaller than the true electronic bandgaps due to the exciton binding energy (~0.3 eV),32 the true LUMO levels of the DPP-based polymers lie slightly above the estimated values. Likewise, the true HOMO levels of the NDI-based polymers would lie slightly below the estimated values. Thus, all of the DPP-based polymers would be suitable for hole transport, while all of the NDI-based polymers would be preferable for electron transport. This conclusion coincides with the electrical properties that were obtained from the polymer FETs. In addition to being useful for characterizing absorption features, UV-visible absorption spectroscopy can also provide information about interchain interactions between polymers. Figure 5b shows the UV-visible absorption spectra of the solution and film states of DPP-based and NDI-based polymers. All three planar DPP-based polymers in the solution 6 ACS Paragon Plus Environment

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state (in chloroform) showed strong absorption in the range of 550-900 nm, with maximum peaks at 799, 818, and 812 nm and weak shoulders around 740 nm for P(DPP2DT-T2), P(DPP2DT-TT), and P(DPP2DT-DTT), respectively. In the solid states, the main absorption peaks remained the same, but the shoulder peaks increased in size. These findings suggest that the DPP-based polymers might exist in the J-aggregate form in solution, while H-aggregates might have developed upon solidification.33,34 For instance, the P(DPP2DT-T2) polymer film displayed a larger shoulder at 717 nm and a 10-nm-blue-shifted absorption peak at 789 nm compared to its solution-phase spectrum. In the cases of the P(DPP2DT-T2) and P(DPP2DTTT) polymers, thermal annealing promoted even more H-aggregation (Figure S23). The NDIbased polymers showed different behavior. The P(NDI2OD-T2) and P(NDI2OD-F2T2) polymers in the solution and film states were similar, which was likely due to both polymers already being in the aggregated form in the solution state.35 In contrast, the P(NDI2OD-Se2) polymer showed a large red-shift in the film state relative to the solution state, implying that it was fully dissolved in the chloroform solution state (unlike the other two NDI polymers) and formed J-aggregates during solidification. The film crystallinity, surface morphology, and thermal annealing effects were investigated by two-dimensional grazing incidence X-ray diffraction (2D GIXD) and tappingmode atomic force microscopy (AFM). Figure 6a shows the 2D GIXD patterns of P(DPP2DT-T2) and P(NDI2OD-Se2) films annealed at three different temperatures (25, 100, and 150 °C). The as-coated P(DPP2DT-T2) exhibited no obvious diffraction peaks. However, upon thermal annealing at 100 and 150 °C, intense edge-on orientation along the qz direction was developed. This aspect indicated that the alkyl chains were aligned normal to the substrate, and that the backbone was aligned parallel to the substrate. This molecular orientation is beneficial for lateral charge transport in the FET channel. More detailed information was obtained by extracting the one-dimensional (1D) profiles along the out-ofplane and in-plane directions from the GIXD patterns. The out-of-plane profiles were extracted along the qz direction at qxy = 0.00 Å-1, and the in-plane profiles were extracted along the qxy direction at qz = 0.03 Å-1. Figures 6b and 6c show the out-of-plane and in-plane profiles of the P(DPP2DT-T2) films, respectively; these peaks were fitted with a Lorentzian function. Figure S24 and Tables S1 and S2 show the peak-fitting results, and Table 4 summarizes the important parameters for the (200) and (010) peaks in the qz direction. Upon thermal annealing, the (100) peak (with higher–order peaks) became well-defined and more pronounced in both the out-of-plane direction and the in-plane direction. The (200) peaks in the out-of-plane direction appeared at qz = 0.567 and 0.571 Å-1 for the P(DPP2DT-T2) films 7 ACS Paragon Plus Environment

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annealed 100 and 150 °C, respectively, corresponding to lamellar spacings of 22.2 and 22.0 Å and crystal correlation lengths of 15.4 and 15.8 nm. The tighter packing and increased longrange ordering explain the enhanced carrier mobilities that occur upon thermal annealing. As the annealing temperature increased from 25 to 150 °C, the out-of-plane (l00) peak intensities increased significantly, and the in-plane (010) peak became more distinct at qz = 1.690 Å-1, which corresponds to a π-π interchain stacking distance of 3.72 Å. Moreover, to analyze this preferential molecular orientation quantitatively, the (100) intensity (IOUT(100)) in the qz direction was compared with both the (010) intensity (IOUT(010)) in the qz direction and the (100) intensity (IIN(100)) in the qxy direction (Figure 6d). Both IOUT(100)/IOUT(010) and IOUT(100)/IIN(100) ratios increased considerably with increasing annealing temperature, indicating that annealing at higher temperatures promoted a transition from the nearly amorphous state of the as-coated P(DPP2DT-T2) film to a highly edge-on-oriented crystalline state. The orientation distribution of the crystallites with respect to the substrate was confirmed by analyzing the intensity pole figure plot for the (100) peak reflection (Figure 6e). This thermally-induced edge-on orientation was also observed in the NDI-based semiconductor of P(NDI2OD-Se2) (Figures 6f, 6g, and 6h). The as-coated P(NDI2OD-Se2) showed an intense (100) peak in the qxy direction and a weak (100) peak in the qz direction, indicating that most P(NDI2OD-Se2) adopted the face-on orientation. Annealing at 150 °C enhanced the (h00) peaks significantly, with a more pronounced increase in the higher-order peaks in the qz direction. Both IOUT(100)/IOUT(010) and IOUT(100)/IIN(100) ratios increased considerably as the annealing temperature increased (Figure 6h), indicating that thermal annealing facilitated the formation of edge-on-oriented crystalline states while increasing crystallinity in all directions. The orientation distribution of the crystallites with respect to the substrate was confirmed by analyzing the intensity pole figure plot for the (100) peak reflection (Figure 6i). The increased crystallinity and improved edge-on orientation achieved by thermal annealing can facilitate carrier transport in the FET channel, accounting for the significant carrier mobility improvement in P(NDI2OD-Se2) FETs. Figure S25 shows the AFM images of the P(DPP2DT-T2) and P(NDI2OD-Se2) films annealed at three different temperatures (25, 100, and 150 °C). The as-coated P(DPP2DT-T2) films displayed finely aggregated surfaces, while the thermally-annealed films formed densely packed interconnected nanofibrillar structures. Likewise, the as-coated P(NDI2OD-Se2) films exhibited smooth surfaces with some nanofibrillar aggregates; the number of fibrillar aggregates increased with increasing annealing temperature. The development of crystalline nanostructures in both films was accompanied by increased surface roughness. The 150 °C8 ACS Paragon Plus Environment

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annealed P(DPP2DT-T2) and 150 °C-annealed P(NDI2OD-Se2) films revealed relatively high root-mean-square (rms) values of 1.22 and 0.99 nm, respectively. Next, the influence of the oligothiophene bridging units (connected to DPP or NDI units) on the crystalline microstructures was investigated. Figure 7a shows the 2D GIXD patterns of the DPP- and NDI-based polymers annealed at 150 °C. All of the DPP-based polymers exhibited multiple, well-defined (h00) diffraction patterns up to the fourth order in the out-of-plane direction (Figure 7b); this indicates that all of the films have highly-ordered lamellar structure. The observed lamellar spacing values were close but slightly different from one another: 22.74, 22.55, and 22.00 Å for P(DPP2DT-DTT), P(DPP2DT-TT), and P(DPP2DT-T2), respectively. The crystal correlation lengths were estimated to be 10.96, 11.19, and 15.81 nm for P(DPP2DT-DTT), P(DPP2DT-TT), and P(DPP2DT-T2), respectively. Furthermore, weak (010) peaks were observed in the qz direction for both P(DPP2DT-DTT) and P(DPP2DT-TT), while the (010) peak of P(DPP2DT-T2) was much weaker. The most distinct (010) peak and weakest (100) peak were observed in the in-plane profile of P(DPP2DT-T2), as shown in Figure 7c. The π-π interchain stacking distances were 3.731, 3.726, and 3.717 Å for P(DPP2DT-DTT), P(DPP2DT-TT), and P(DPP2DT-T2), respectively. All of the important estimated parameters indicate that the film crystallinity increased in the order of P(DPP2DT-DTT)