Crystallinity-Controlled Naphthalene-alt-diketopyrrolopyrrole

Article pubs.acs.org/JPCC

Crystallinity-Controlled Naphthalene-alt-diketopyrrolopyrrole Copolymers for High-Performance Ambipolar Field Effect Transistors Hyo-Sang Lee,†,‡,△ Joong Suk Lee,§,△ Sanghyeok Cho,¶ Hyunjung Kim,¶,△ Kyung-Won Kwak,∥ Youngwoon Yoon,† Seon Kyoung Son,†,§ Honggon Kim,† Min Jae Ko,† Doh-Kwon Lee,† Jin Young Kim,† Sungnam Park,‡ Dong Hoon Choi,‡ Se Young Oh,# Jeong Ho Cho,*,⊥ and BongSoo Kim*,† †

Photo-electronic Hybrids Research Center, Korea Institute of Science and Technology (KIST), Seoul 136-791, Republic of Korea Department of Chemistry, Korea University, Seoul 136-713, Republic of Korea § Department of Organic Materials and Fiber Engineering, Soongsil University, Seoul 156-743, Republic of Korea ¶ Department of Physics, Sogang University, Seoul 121-742, Republic of Korea ∥ Department of Chemistry, Chung-Ang University, Seoul 156-756, Republic of Korea ⊥ SKKU Advanced Institute of Nanotechnology (SAINT) and Center for Human Interface Nano Technology (HINT), School of Chemical Engineering, Sungkyunkwan University, Suwon 440-746, Republic of Korea # Department of Chemical and Biomolecular Engineering, Sogang University, Seoul 121-742, Republic of Korea ‡

S Supporting Information *

ABSTRACT: We report high-performance of ambipolar organic fieldeffect transistors (FETs) based on the low band gap copolymers of pDPPT2NAP-HD and pDPPT2NAP-OD. The polymers are composed of electron-rich 2,6-di(thienyl)naphthalene (T2NAP) and electrondeficient diketopyrrolopyrrole (DPP) units with branched alkyl chains of 2-hexyldecyl (HD) or 2-octyldodecyl (OD). The polymers were polymerized via Suzuki coupling, yielding optical band gaps of ∼1.4 eV. In the transistor performance test, we observed good ambipolar transport behavior in both polymer films, and pDPPT2NAP-OD displayed hole and electron mobilities 1 order of magnitude higher than the corresponding properties of pDPPT2NAP-HD. Thermal annealing of the polymer films increased the carrier mobilities. Annealing at 150 °C provided optimal conditions yielding saturated film crystallinity and maximized carrier mobility. The highest hole and electron mobilities achieved in these polymers were 1.3 and 0.1 cm2/(V s), respectively, obtained from pDPPT2NAP-OD. The polymer structure and thermal annealing affected the carrier mobility, and this effect was investigated by fully characterizing the polymer films by grazing incidence X-ray diffraction (GIXD), atomic force microscopy (AFM), and transmission electron microscopy (TEM) experiments. The GIXD data revealed that both polymers formed highly crystalline films with edge-on orientation. pDPPT2NAP-OD, which included longer alkyl chains, showed a higher tendency to form long-range order among the polymer chains. Thermal annealing up to 150 °C improved the polymer film crystallinity and promoted the formation of longer-range lamellar structures. AFM and TEM images of the films were consistent with the GI-XD data. Theoretical calculations of the polymer structures provided a rationale for the relationship between the torsional angle between aromatic rings and the carrier mobility. From the intensive electrical measurements and full characterizations, we find that the chemical structure of polymer backbone and side alkyl chain has a profound effect on film crystallinity, morphology, and transistor properties. thienyl)-2,2′-bithiophene)] (PQT),13 and poly(2,5-bis(3-alkylthiophene-2-yl)thieno[3,2-b]thiophenes) (pBTTTs), 14 which exhibit hole mobilities of 0.1−0.6 cm2/(V s). These polymer films display highly edge-on oriented closely packed crystalline lamellar structures, a characteristic that is essential for charge transport through the π−π stacks between polymer chains.

1. INTRODUCTION The development of organic field-effect transistors (OFETs) is of paramount importance for the future of large-area low-cost printed electronic device applications, such as flexible displays, radiofrequency identification (RF-ID) tags, memory devices, and sensors.1−9 Conjugated polymers are highly attractive transistor channel materials in that they are solutionprocessable and display higher field-effect carrier mobilities than amorphous silicon semiconductors. Previous studies have focused on thiophene-based polymers,5,10 such as poly(3hexylthiophene) (P3HT),2,11,12 poly[5,5′-bis(3-dodecyl-2© 2012 American Chemical Society

Received: September 17, 2012 Revised: November 24, 2012 Published: November 26, 2012 26204

dx.doi.org/10.1021/jp309213h | J. Phys. Chem. C 2012, 116, 26204−26213

The Journal of Physical Chemistry C

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Figure 1. Representative donor−acceptor type polymer structures.

transistor behavior by exploring new conjugated polymer backbones with the fine-tuning of side alkyl chains. In this article, we present the FET performances of 2,6di(thienyl)naphthalene-alt-diketopyrrolopyrrole polymers (pDPPT2NAPs) with branched alkyl chains of 2-hexyldecyl (HD) and 2-octyldodecyl (OD) groups. The naphthalene units are highly conjugated and planar, and they provided closely stacked crystalline lamella-like microstructures. Such units introduced sufficient spacing between alkyl chains to facilitate interdigitation. Incorporation of the naphthalene and DPP units into a polymer backbone, as in pDPPT2NAPs, provided good transistor performance. The charge transport behavior in this system was elucidated by fully characterizing the polymer films using UV−visible spectrometry, cyclic voltammetry, grazingincidence X-ray diffraction, atomic force microscopy, transmission electron microscopy, and density functional theory calculations. Overall, this work describes the good correlation between chemical structure, film crystallinity, and carrier mobility. It is also highlighted that planar connections among the conjugated backbone aromatic rings are essential for achieving a high carrier mobility.

Recently, we and others achieved even higher carrier mobilities of 0.6−3.3 cm2/(V s) from donor−acceptor-type low band gap polymers.15−17 Those polymers consisted of electron-rich donor and electron-deficient acceptor moieties with side alkyl chains. Figure 1 shows representative donor− acceptor-type polymer structures. Donor moieties commonly comprise linearly connected or fused thiophenes and selenophenes;18−22 acceptor moieties comprise isoindigos,23−25 diketopyrrolopyrroles,26−34 benzothiadiazoles,35−38 and naphthalenedicarboximide.1,39,40 These high carrier mobilities have been accomplished by employing an elegant combination of side alkyl chains, donors, and acceptors in the polymer structure. First, side alkyl chains not only endow polymers with solution processability, but they also promote the self-organization of polymer chains in the film state. Optimal alkyl chain lengths and alkyl chain spacing distances yield polymers that are easily soluble and form lamellar structures with well-interdigitated alkyl chains.27,32 Second, the chemical structures of donor and acceptor moieties govern the backbone planarity and energy level alignment. Planar polymer backbones facilitate close interchain π−π stacking.10,27,41 Moreover, planar backbone structures promote orbital overlap between the donor and acceptor units, which profoundly affects the positions of HOMO and LUMO levels of the polymers. Ultimately, the HOMO and LUMO level alignment dictates the main carrier type in the semiconducting polymer channel. Besides the improvement of carrier mobility that is a central issue in the field of organic transistors, there is also considerable interest to fabricate complementary metal oxide semiconductor (CMOS)-like inverters using semiconducting polymers.3,8,42−44 For this purpose, it would be more profitable to use singlecomponent ambipolar transistors than a combination of p- and n-channel transistors. Low band gap polymers have great potential for ambipolar transistors since both their HOMO and LUMO levels lie close to the Fermi levels of contact metals. However, low band gap polymers have still rarely exhibited ambipolar behavior.26,27,30,39,45 Therefore, for the further improvement in both carrier mobility and hole/electron transport balance in ambipolar transistors, it is still strongly needed to understand the relationship of chemical structure to

2. EXPERIMENTAL SECTION 2.1. Materials and Synthesis. 2,6-Dibromonaphthalene, n-BuLi in hexane, 2-isopropoxy-4,4,5,5-tetramethyl-1,3,2-dioxaborolane, Pd2dba3, P(o-tolyl)3, K3PO4, and Aliquat 336 were purchased from Sigma-Aldrich, Acros, and TCI. Common organic solvents were purchased from Daejung and J. T. Baker. Tetrahydrofuran (THF) was dried over sodium and benzophenone prior to use. All other reagents were used as received without further purification. The synthetic routes and chemical structures of the polymers used in this study are shown in Scheme 1. Monomers 1 and 2 were synthesized according to literature procedures.22,26,27,32,46 2.1.1. Synthesis of 2,6-Bis(4,4,5,5-tetramethyl-1,3,2-dioxaborolan-2-yl)naphthalene. n-BuLi (2.5 M solution in hexane, 6.15 mL, 15.4 mmol) was added dropwise to a solution of 2,6dibromonaphthalene (2 g, 6.99 mmol) in dry THF (20 mL) at −78 °C. The mixture was stirred at this temperature for 2 h and then 2-isopropoxy-4,4,5,5-tetramethyl-1,3,2-dioxaborolane (5.14 mL, 25.2 mmol) was added. After stirring for 30 min, the reaction was allowed to warm to room temperature and 26205



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dependence on the transistor performance.26 Therefore, our discussion focused on the impact of the alkyl chain length, i.e., HD versus OD, primarily.) 2.1.3. Synthesis of pDPPT2NAP-OD. pDPPT2NAP-OD was synthesized as described for pDPPT2NAP-HD, using 2,6bis(4,4,5,5-tetramethyl-1,3,2-dioxaborolan-2-yl)naphthalene (0.290 g, 0.551 mmol) and monomer 2 (0.562 g, 0.551 mmol). The resulting polymer was obtained in 0.470 g (86.3%). 1H NMR (CDCl3, 400 MHz) δ (ppm) 9.28−8.27 (b, 2H), 7.94− 6.18 (b, 8H), 4.46−3.01 (b, 4H), 2.37−0.25 (b, 78H); GPC Mn = 23 611 g/mol, PDI = 3.11. 2.2. Material Characterization. 1H NMR spectra were recorded on a Bruker advance 400 spectrometer (400 MHz). The molecular weights of the polymers were measured by gel permeation chromatography (GPC) using chloroform as an eluent and polystyrene as a standard. Differential scanning calorimetry (DSC) was recorded on a Perkin-Elmer Pyris 1 DSC instrument at a heating rate of 10 °C/min between 20 and 300 °C under nitrogen. UV−visible absorption spectra were collected on a Perkin-Elmer Lamb 9 UV−vis spectrophotometer. Cyclic voltammetry (CV) was performed using a CH Instruments electrochemical analyzer, and the solutions were prepared using a degassed acetonitrile solution containing 0.1 M tetrabutylammonium hexafluorophosphate (TBAPF6) as the electrolyte. The scan rate was 50 mV/s. A Pt wire electrode coated with a thin film of the polymer was used as the working electrode, a Pt wire was used as the counter electrode, Ag/Ag+ was the reference electrode, and ferrocene was used as the internal standard. The crystalline structures of the ambipolar polymer films were characterized by grazing incidence X-ray diffraction (GIXD) studies at the 8-ID-E beamline of the Advanced Photon Source (APS) at Argonne National Laboratory (energy 7.35 keV; incidence angle 0.2°). The surface morphology of the samples was investigated by tapping mode atomic force microscopy (AFM) (D3100 Nanoscope V, Veeco) and transmission electron microscopy (TEM) (Philips CM-30). Transistor current−voltage characteristics were measured using Keithley 2400 and 236 source/measure units at room temperature under vacuum conditions (10−5 Torr) in a dark environment. 2.3. Device Fabrication. PFETs were fabricated on a highly doped n-type Si wafer with a thermally grown 300 nm thick oxide layer as the substrate. The wafer served as the gate electrode and a thermally grown 300 nm thick SiO2 layer acted as a gate insulator. Prior to treating the silicon oxide surface, the wafer was cleaned in piranha solution for 30 min at 100 °C and washed with copious amounts of distilled water. A SiO2 layer was modified using octadecyltrichlorosilane (ODTS, Gelest, Inc.) to reduce electron trapping by the silanol groups on SiO2. A 50 nm thick layer of the ambipolar semiconducting polymer was spin-coated from a 0.5 wt % chloroform solution onto the ODTS-treated substrates. After spin-coating, the samples were dried in a vacuum chamber for 24 h. The ambipolar polymer films were thermally annealed for 30 min in a vacuum chamber

Scheme 1. Chemical Structures and Synthetic Routes of Di(thienyl)naphthalene-alt-diketopyrrolopyrrole Polymers (pDPPT2NAP-HD and pDPPT2NAP-OD)

stirred for 4 h. The reaction was quenched with water and extracted with ether, then dried over MgSO4. After filtration organic solvent was concentrated in vacuo. The residue was chromatographed on SiO2 to afford the title compound (1.5 g, 56%). 1H NMR (CDCl3, 400 MHz) δ (ppm) 8.35 (s, 2H), 7.84 (q, 4H), 1.39 (s, 24H). 2.1.2. Synthesis of pDPPT2NAP-HD. To a degassed 5 mL toluene solution of 2,6-bis(4,4,5,5-tetramethyl-1,3,2-dioxaborolan-2-yl)naphthalene (0.209 g, 0.551 mmol), monomer 1 (0.500 g, 0.551 mmol), K3PO4 (0.585 g, 2.76 mmol), and three drops of Aliquat 336 and 1 mL of degassed demineralized water were added. Pd2dba3 (0.0202 mg, 0.022 mmol) and tri(otolyl)phosphine (0.0268 mg, 0.088 mmol) were then added to the reaction mixture. The reaction solution was stirred for 1 h at 60 °C under an argon atmosphere. The reaction mixture was poured into a methanol/H2O (4/1 v/v%) solution. The precipitated polymer was redissolved in chloroform and reprecipitated in MeOH. The collected polymer was further purified by Soxhlet extraction using methanol, hexane, acetone, and chloroform. The chloroform fraction was collected, reprecipitated in methanol, and then filtered. The polymer was dried under vacuum, yielding 0.350 g (72.5%). 1H NMR (CDCl3, 400 MHz) δ (ppm) 9.24−8.55 (b, 2H), 8.00−6.90 (b, 8H), 4.32−3.20 (b, 4H), 2.10−0.40 (b, 62H); GPC Mn = 6856 g/mol, PDI = 3.37. (Note that the relatively low molecular weight of this polymer is mainly due to its lower solubility compared to pDPPT2NAP-OD. However, recently Bljleveld et al. showed that PDPP3T, which has close structural similarity to this polymer, exhibited very weak molecular weight

Table 1. Redox Potentials and Energy Levels of pDPPT2NAP-HD and pDPPT2NAP-OD pDPPT2NAP-HD pDPPT2NAP-OD

Eonset,ox (V)

Eonset,red (V)

HOMOa (eV)

LUMOb (eV)

Eg,ECc (eV)

Eg,optc (eV)

0.5 0.5

−1.4 −1.4

−5.3 −5.3

−3.4 −3.4

1.9 1.9

1.4 1.4

HOMO = Eonset,ox + 4.8 eV. bLUMO = Eonset,red + 4.8 eV. cEg,EC and Eg,opt were determined from (Eonset,ox − Eonset,red) and the absorption edge of the UV−visible spectra, respectively.

a

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Figure 2. UV−visible absorption spectra of (a) pDPPT2NAP-HD and (b) pDPPT2NAP-OD films annealed at various temperatures: 25, 100, 150, and 200 °C. (c) Cyclic voltammograms of pDPPT2NAP-HD and pDPPT2NAP-OD films. The insets in (a) and (b) show the red shift in the maximum light absorption of polymer films upon annealing.

Figure 3. Transfer characteristics at a fixed VD of −80 V for PFETs based on (a) pDPPT2NAP-HD and (b) pDPPT2NAP-OD films annealed at various temperatures: 25, 100, 150, and 200 °C. (c) Hole and electron mobilities of the pDPPT2NAP-HD and pDPPT2NAP-OD PFETs as a function of the annealing temperature. Electron mobilities are shown in the inset.

at various temperatures, 25, 100, 150, and 200 °C. The 50 nm source/drain electrodes (Au, Al, LiF/Al, and Ag) were vacuumdeposited through a shadow mask on the polymer film, with channel lengths and widths of 50 and 800 μm, respectively.

(see the insets of Figure 2a,b). This trend indicated that the thermal energy promoted self-assembly of the polymer chains via strong chain-to-chain interactions,19,47 consistent with the GIXD, TEM, and AFM results (discussed below). Figure 2c shows cyclic voltammograms of the polymers. The HOMO and LUMO levels of both polymer films were estimated to be −5.3 and −3.4 eV, respectively, while the levels were independent of the branched alkyl chains. No clear thermal transitions were observed in the differential scanning calorimetry (DSC) curves over the range 30−300 °C (Figure S1 in the Supporting Information). Bottom-gate top-contact polymer field-effect transistors (PFETs) were used to evaluate the electrical properties of the pDPPT2NAP polymers as a function of the branched alkyl chain length and thermal annealing temperature. Figure 3a,b depicts the drain current (ID)−gate voltage (VG) characteristics at a fixed drain voltage (VD) of −80 V for PFET devices based, respectively, on pDPPT2NAP-HD and pDPPT2NAP-OD films annealed at various temperatures, and with Au source/drain contacts. The transistors exhibited typical ambipolar behavior in hole enhancement (VD = −80 V) and electron enhancement (VD = +80 V) mode operations (Figure 3 and Figure S2 in the Supporting Information, respectively). The carrier mobilities of each PFET were calculated in each respective saturation regime

3. RESULTS AND DISCUSSION Low band gap 2,6-di(thienyl)naphthalene-alt-diketopyrrolopyrrole polymers with different lengths of branched alkyl chains (pDPPT2NAP-HD and pDPPT2NAP-OD) were synthesized as shown in Scheme 1. The electronic properties of pDPPT2NAPs were measured by UV−visible absorption spectroscopy and cyclic voltammetry (CV). The UV−visible absorption and CV results are summarized in Table 1. Figure 2a,b, shows the film states of either polymer absorbed photons across a wide spectral range that extended to ∼900 nm (1.4 eV). The main peak of pDPPT2NAP-OD was red-shifted relative to the peak of pDPPT2NAP-HD by ∼15 nm, and its absorption coefficient was higher than the coefficients of the shoulder peaks, indicating that the pDPPT2NAP-OD polymer chains carrying longer branched alkyl chains were more closely packed. Film annealing at 25, 100, 150, or 200 °C for 30 min showed that the absorption peaks at 723 and 738 nm, for pDPPT2NAP-HD and pDPPT2NAP-OD, respectively, became increasingly red-shifted as the annealing temperature increased 26207



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Table 2. Field-Effect Mobilities of PFETs Based on pDPPT2NAP-HD and pDPPT2NAP-OD Annealed at Various Temperatures with Au Source/Drain Contacts pDPPT2NAP-HD pDPPT2NAP-OD

hole electron hole electron

RT

100 °C

150 °C

200 °C

0.037 (±0.011) 0.023 (±0.008) 0.28 (±0.051) 0.037 (±0.014)

0.063 (±0.012) 0.040 (±0.010) 0.38 (±0.087) 0.040 (±0.021)

0.11 (±0.024) 0.081 (±0.015) 1.1 (±0.21) 0.07 (±0.017)

0.084 (±0.014) 0.036 (±0.012) 0.71 (±0.18) 0.040 (±0.010)

Table 3. Field-Effect Mobilities of PFETs Based on the 150 °C Annealed pDPPT2NAP-HD and pDPPT2NAP-OD with Au, Al, LiF/Al, and Ag Source/Drain Contacts Au pDPPT2NAP-HD pDPPT2NAP-OD

hole electron hole electron

Al

0.11 (±0.024) 0.081 (±0.015) 1.1 (±0.21) 0.070 (±0.017)

0.023 0.007 0.061 0.030

(±0.009) (±0.002) (±0.010) (±0.011)

LiF/Al 0.040 0.007 0.077 0.017

(±0.012) (±0.003) (±0.023) (±0.007)

Ag 0.044 (±0.018) 0.010 (±0.004) 0.17 (±0.060) 0.014 (±0.004)

Figure 4. 2D GIXD patterns of (a) pDPPT2NAP-HD and (b) pDPPT2NAP-OD films annealed at various temperatures: 25, 100, 150, and 200 °C. The top (sky blue) and middle (orange) panels of the inset show the 1D out-of-plane and in-plane X-ray diffraction profiles extracted along the qz direction at qy = 0.00 Å−1 and the qy direction at qz = 0.03 Å−1, respectively. The bottom (green) panels of the inset show the azimuthal X-ray diffraction profile (qr) at the qz = (200) peak.

using the relationship ID = CiμW(VG − Vth)2/2L, where W and L are the channel width and length, respectively, Ci is the specific capacitance of the gate dielectric (11 nF/cm2), Vth is the threshold voltage, and μ is the carrier mobility. PFETs based on as-cast pDPPT2NAP-HD showed hole and electron mobilities of 0.037 and 0.023 cm2/(V s), respectively. In contrast, PFETs based on as-cast pDPPT2NAP-OD with longer branched alkyl chains yielded higher hole and electron mobilities of 0.28 and 0.037 cm2/(V s), respectively. Thermal annealing of the polymer films at 100 or 150 °C resulted in a significant increase in carrier mobility, as shown in Figure 3c. The average hole mobilities of the PFETs were 0.11 and 1.1 cm2/(V s) for the 150 °C annealed pDPPT2NAP-HD and pDPPT2NAP-OD, respectively. The PFETs based on pDPPT2NAP-OD carrying longer alkyl side chains were highly responsive to the annealing temperature. Interestingly, annealing at a higher temperature of 200 °C led to a slight decrease in the carrier mobilities in both polymers. The

electrical performances of the devices were optimized by fabricating PFETs with Al, LiF/Al, or Ag source/drain contacts, as shown in Figure S3 in the Supporting Information; however, the device performances of PFETs with these lower work function metal contacts were not better than the performance obtained using Au contacts, possibly due to the poor interfacial contacts induced by the oxidation of the metals. The electrical properties of PFETs based on either polymer are summarized in Tables 2 and 3. In addition, Figure S4 in the Supporting Information shows the ID−VD curves at 11 different VGs in PFETs based on pDPPT2NAP-HD and pDPPT2NAP-OD annealed at 100 °C. A decrease in VG from −100 to −20 V reduced the hole accumulation in the semiconducting channel, which then decreased ID. An increase in VG from 0 to 100 V was accompanied by a gradual increase in ID due to electron accumulation. The maximum ID for the pDPPT2NAP-OD PFETs in the hole saturation regime (VD = −100 V) was much higher than that of the pDPPT2NAP-HD PFETs. All of these 26208



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Figure 5. 1D out-of-plane and in-plane X-ray profiles extracted along the qz direction at qy = 0.00 Å−1 and the qy direction at qz = 0.03 Å−1 from the 2D GIXD patterns of (a) pDPPT2NAP-HD and (b) pDPPT2NAP-OD films annealed at various temperatures: 25, 100, 150, and 200 °C.

electrical characteristics of PFETs as a function of alkyl chain length and annealing temperature were strongly associated with the evolution of the crystalline and morphological structures of the semiconducting polymers, as confirmed by GIXD, TEM, and AFM studies (see below). Crystalline structures of pDPPT2NAP-HD and pDPPT2NAP-OD were characterized using synchrotron 2D GIXD measurements. Parts a and b of Figure 4 show the 2D GIXD patterns of the pDPPT2NAP-HD and pDPPT2NAPOD films, respectively. The x coordinate indicates the probing beam direction, the y coordinate indicates the direction parallel to the substrate (in-plane), and the z coordinate indicates the direction normal to the substrate (out of plane). The figures in the first columns of Figure 4a,b show the 2D GIXD patterns of the as-cast pDPPT2NAP-HD and pDPPT2NAP-OD films, from which the 1D out-of-plane and in-plane X-ray diffraction profiles in the top (sky blue) and middle (orange) insets were extracted along the qz direction at qy = 0.00 Å−1 and the qy direction at qz = 0.03 Å−1, respectively. The orientations of the ordered domains with respect to the substrate were similar in both polymers; specifically, both polymers showed intense (100) reflections with higher order peaks along the qz direction and (010) reflections along the qy direction. The (010) diffraction peaks were almost undetectable along the qz direction for both films. This observation indicates that the alkyl side chains were aligned along the normal to the substrate, and the polymer backbones were essentially parallel to the substrate. This molecular orientation was particularly favorable for charge transport through the π−π stacks in the PFETs, such that the drain current flowed along the semiconducting channel parallel to the substrate. The out-of-plane spacings of the ordered pDPPT2NAP-HD and pDPPT2NAP-OD phases were 18.5 and 20.0 Å, respectively, estimated from the spacing between the (h00) peaks. A comparison of these spacings to the fully extended polymer widths, i.e., two alkyl side chain lengths plus the aromatic backbone width, suggested that the side chains of the polymer were interdigitated and/or tilted out of the polymer backbone plane. Moreover, the out-of-plane spacings of pDPPT2NAP-HD differed from those of pDPPT2NAP-OD by 1.5 Å, suggesting that pDPPT2NAPOD polymer chains were stacked more tightly than pDPPT2NAP-HD polymer chains. By contrast, the π−π

interchain stacking spacing for pDPPT2NAP-HD (∼3.92 Å) was almost identical to the spacing of pDPPT2NAP-OD. Higher-order diffraction peaks and much weaker diffuse scattering intensities around the (h00) diffraction peaks along the Debye rings were observed for the pDPPT2NAP-OD films, although these features were less prominent for the pDPPT2NAP-HD films. The full width half-maximum (fwhm) of the azimuthal-angle intensity of the (h00) peak provides information about the broadness of the crystalline plane orientations with respect to the substrate surfaces. Better alignment of the (h00) planes with respect to the substrate is important for charge transport through the π−π stacking structures among polymer chains. The bottom insets (green) in Figure 4a,b show the azimuthal X-ray diffraction profile (qr) at the qz = (200) peak (corresponding qx = 0.0639 Å−1 for pDPPT2NAP-HD and qx = 0.0513 Å−1 for pDPPT2NAP-OD). Higher order diffraction peaks and a narrow orientation distribution in the pDPPT2NAP-OD were clearly observed. This observation indicated that the pDPPT2NAP-OD polymer with longer alkyl chains produced longer-range upright crystalline plane orientations, and such long alkyl chains were critical to the control over film crystallinity and orientation. Figure 4a,b also illustrates the effects of thermal annealing on film crystalline structure. As the annealing temperature increased, the film crystallinity increased and more ordered crystalline domains developed. A more in-depth analysis of the 2D GIXD patterns was conducted by extracting the 1D X-ray profiles along both the out-of-plane and in-plane directions, as shown in Figure 5. Figure 5a presents the 1D out-of-plane and in-plane X-ray diffraction profiles extracted along the qz direction at qy = 0.00 Å−1 and the qy direction at qz = 0.03 Å−1 from the 2D GIXD patterns of the pDPPT2NAP-HD films. In the out-of-plane profile, the (h00) diffraction peak intensity increased with the annealing temperature up to 150 °C and then flattened out around 200 °C. In the in-plane profile, the (010) diffraction peak intensity increased with the annealing temperature up to 150 °C. The most pronounced (010) peak was observed at 150 °C. The increase in the GIXD peak intensity upon thermal annealing arose from a combination of reorganization among the semiconducting polymers during heating and the growth of existing crystalline regions. The crystallinity enhancement, observed in both the out-of-plane (h00) and in-plane (010) peaks with thermal annealing at 26209



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Figure 6. TEM images of (a) pDPPT2NAP-HD and (b) pDPPT2NAP-OD films annealed at various temperatures: 25, 100, 150, and 200 °C.

temperatures between 25 and 150 °C, suggested that the thermal energy promoted the development of crystalline lamellar microstructures and shifted the orientations of these microstructures to favor charge transport. Annealing at 200 °C reduced the (010) peak, while the (100) peak increased in the in-plane profile. These observations imply that a fraction of the polymer chains in the film crossed over from an edge-on orientation to a face-on orientation in the conjugated polymer backbone. The π−π stacking distance also increased slightly from 3.92 to 3.95 Å with thermal annealing at 200 °C. Both effects resulted in a slight drop in the carrier mobility for polymers annealed at 200 °C. A similar trend was observed in the pDPPT2NAP-OD films, as shown in Figure 5b, while the intensity of the (h00) diffraction peak was more pronounced and sharper for pDPPT2NAP-OD. This sensitive dependence of the film crystallinity on the annealing temperature in conjugated polymers with long branched alkyl chains is consistent with the temperature-dependent carrier mobilities (discussed above). In addition, the evolution of film morphology upon thermal annealing at 25, 100, 150, or 200 °C was examined by TEM and AFM experiments. Figure 6 and Figure S5 in the Supporting Information show TEM and AFM images, respectively, of the pDPPT2NAP-HD and pDPPT2NAP-OD films. The pDPPT2NAP-HD films in Figure S5a displayed aggregated granular domains, whereas the pDPPT2NAP-OD films in Figure S5b exhibited more crystalline structures. The TEM images clearly revealed the film morphological changes, in detail, upon thermal annealing. Figure 6a shows that, in the pDPPT2NAP-HD films, fibril lamellar structures developed with increasing annealing temperature. Fibril structures in an as-cast film are narrow and often disconnected. These structures become thicker and better-connected in films annealed at 150 and 200 °C. By contrast, wavy fibril structures were observed in the pDPPT2NAP-OD films (Figure 6b). Upon annealing, the fibril structures reorganized in a fashion similar to the reorganization observed among pDPPT2NAPHD, but the change was much more dramatic. Highly fibrous morphological features are consistent with an increase in the crystallinity, as revealed in the GIXD analysis. The GIXD, TEM, and AFM experiments conclusively showed that the

enhancement in carrier mobility upon thermal annealing resulted directly from both the favorable orientations of the π−π stacked polymer chains and the formation of wellconnected crystalline fibril structures in the films. Complementary inverter devices were successfully fabricated by connecting two identical ambipolar PFETs based on pDPPT2NAP-HD and pDPPT2NAP-OD. Figure 7 shows the

Figure 7. Output voltage versus input voltage plots of inverters based on pDPPT2NAP-HD and pDPPT2NAP-OD films annealed at 150 °C at a constant supply voltage.

output voltage (VOUT) as a function of the input voltage (VIN) at a constant supply voltage (VDD). The circuit diagram of an inverter is shown in the inset of Figure 7. A reasonable inverter action was observed as the VIN sweep. Signal inversion was obtained at both positive (+80 V) and negative VDD values (−80 V) due to the ambipolar nature of the constituent transistors. Note that better inverter action was observed in the pDPPT2NAP-HD PFETs because the hole and electron currents were more balanced than in the pDPPT2NAP-OD PFETs, as shown in Figure 3 and Table 2. (μh/μe = 1.4 and 15.7 for pDPPT2NAP-HD and pDPPT2NAP-HD PFETs, respectively). The signal inverter gains, defined as the absolute 26210



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Figure 8. Energy levels and surface plots of frontier orbitals (top) and torsional angles (bottom) for model molecules representing polymer repeating units. Energy unit is eV.

pDPPT2TT-OD, and pDPPT3-OD, respectively.27 (Note that these out-of-plane spacings were estimated not from (010) peaks but from the spacing between the (h00) peaks) Longer distances between alkyl chains, as were observed in pDPPT2NAP-OD, provide more room for interdigitation/ packing of the branched alkyl side chains and more efficient space filling in the qz direction, which facilitated the crystallization of the polymers. (4) The hole mobilities of PFETs showed a strong dependence on the donor structure type. Polymers annealed at 150 °C displayed average hole mobilities of 1.1, 1.6, and 1.9 cm 2 /(V s), for pDPPT2NAP-OD, pDPPT3-OD, and pDPPT2TT-OD, respectively. The critical factors for obtaining these high mobilities were interchain packing, lamellar orientation, and film roughness. Another crucial factor was identified here: the torsional angle between backbone aromatic rings, as demonstrated by conducting density function theory (DFT) calculations. The simplified backbone structures of TDPP-T-NAP-T, T-DPP-T-T-T, and T-DPP-T-TT-T (representing pDPPT2NAP-OD, pDPPT2TT-OD, and pDPPT3OD, respectively) were geometrically optimized to an energy minimum using Gaussian 09 at the DFT B3LYP level with the 6-31+G(d,p) basis set. To expedite the calculation, alkyl chains were shortened to methyl groups on the DPP nitrogen. Figure 8 shows the energy levels of the frontier orbitals, their surface plots, and geometric models of their structures. The DFT results offered an explanation for why the band gap increased in the order of pDPPT3-OD, pDPPT2TT-OD, and pDPPT2NAP-OD, consistent with the UV−visible spectra. More importantly, the geometrical information from the DFT calculations supported the correlation between the carrier mobility and the π−π stacking distance. The torsional angles of φ1 and φ2 were similar for T-DPP-T-T-T and T-DPP-T-TT-T (φ1 ≈ 12° and φ2 ≈ 14°), whereas the angles were higher in TDPP-T-NAP-T (φ1 ≈ 16° and φ2 ≈ 26°). The π−π stacking distances in pDPPT2NAP and pDPPT3, based on the GIXD experiments, were 3.95 and 3.65 Å, respectively. Considering the DFT and GIXD data together, we postulate that the reduced torsional angles in the pDPPT3 and pDPPT2TT

value of dVOUT/dVIN, were 5.9 or 3.0 for the PFETs based, respectively, on pDPPT2NAP-HD or pDPPT2NAP-OD at VDD = −80 V. The value of the signal inverter gain demonstrates the importance of the balanced hole and electron mobilities. Lastly, we discuss our observations on the crystallinity, morphology, and transistor performances of a series of polymers that we have synthesized recently. The polymers were constructed using three types of electron-rich donor moieties: (i) thiophene−2,6-naphthalene−thiophene (T-NAPT), (ii) thiophene−thiophene−thiophene (T-T-T), and (iii) thiophene−thieno[3,2-b]thiophene−thiophene (T-TT-T). Diketopyrrolopyrrole (DPP) moieties carrying two branched alkyl chains of 2-hexyldecyl (HD) and 2-octyldodecyl (OD) acted as acceptor moieties and were paired with the donor moieties. The corresponding polymers were pDPPT2NAP, pDPPT3,27 and pDPPT2TT,27 respectively. The chemical structures of pDPPT3 and pDPPT2TT are shown in Figure 1. The series of donors and solubilizing groups tested here elucidated the relationship between the film structure and the electronic properties. Several important conclusions may be drawn: (1) Longer alkyl chains facilitate film crystallization. The OD-attached polymers showed a higher degree of stacking among the polymer chains than in the HD-attached polymers. For instance, pDPPT2NAP-OD showed strong crystalline diffraction peaks with a narrower distribution of crystal orientations. As a result, the carrier mobility was higher for the OD-attached polymers than the HD-attached polymers. (2) Polymers with long alkyl chains respond more rapidly to thermal annealing than polymers with short alkyl chains. This characteristic was attributed to the fact that thermal energy facilitates rearrangement and self-assembly among longer alkyl side chains, thereby increasing the van der Waals interactions more rapidly than in shorter chains.27,48,49 (3) The distance between alkyl chains in the polymer backbone plays a crucial role in forming interdigitated lamellar structures. For instance, for a given OD alkyl chain, the out-ofplane spacings of the ordered polymer lamella were calculated to be 20.0, 21.1, and 23.6 Å for pDPPT2NAP-OD, 26211



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polymers afforded close π−π stacking among the polymer chains. Of course, the close polymer chain stacking is beneficial for the carrier hopping and carrier mobility.

PAL through the abroad beamtime program of Synchrotron Radiation Facility Project under MEST.

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4. CONCLUSIONS Here, we describe the performances of ambipolar FETs based on the polymers pDPPT2NAP-HD and pDPPT2NAP-OD. Both polymers have low band gaps around 1.4 eV. The FET performance measurements revealed good ambipolar transport behavior in both polymer films. The hole and electron mobilities increased with the alkyl chain length and thermal annealing. The highest hole and electron mobilities were 1.3 and 0.1 cm 2 /(V s), respectively, obtained from the pDPPT2NAP-OD film annealed at 150 °C. GIXD, AFM, and TEM experiments revealed that the transistor results were strongly correlated with the crystalline features and edge-on orientations. That is, the features of longer range ordering among the polymer chains and a thicker fibrous crystalline structure were observed in the pDPPT2NAP-OD than in the pDPPT2NAP-HD. In addition, DFT calculations of the polymer backbone structure provided deeper insight into the relationship between the backbone torsional angles along a polymer chain and the carrier mobility and interchain packing. The collective data analysis supports our conclusion that planar conjugated polymer backbones favor high carrier mobilities. For a given polymer backbone structure, the crystallinity and carrier mobility could be controlled by introducing solubilizing groups of a suitable length and applying thermal treatment. Importantly, this work highlights the important role of the torsional angles between aromatic rings in the backbone for π−π interchain packing and carrier hopping transport.

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ASSOCIATED CONTENT

S Supporting Information *

DSC curves, transfer and output characteristics of PFETs, and AFM images of polymer films (Figures S1−S5). This material is available free of charge via the Internet at http://pubs.acs.org.

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REFERENCES

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (J.H.C.); [email protected] (B.K.). Author Contributions △

These authors equally contributed to this work.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by New and Renewable Energy Program of the Korea Institute of Energy Technology Evaluation and Planning (KETEP) grant funded by the Ministry of Knowledge Economy (MKE) (20113030010060 and 20113010010030) and by Korea Research Council of Fundamental Science and Technology (KRCF) and Korea Institute of Science and Technology (KIST) for “NAP National Agenda Project Program” and Basic Science Research Program (2009-0083540 and 2010-0026294) through the National Research Foundation of Korea funded by the Ministry of Education, Science and Technology. Use of the Advanced Photon Source was supported by the US Department of Energy, Office of Basic Energy Science, under Contract No. DE-AC02-06CH11357. H.K. and S.C. thank the support by 26212



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Crystallinity-Controlled Naphthalene-alt-diketopyrrolopyrrole

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