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A Highly Planar Fluorinated Benzothiadiazole-Based Conjugated Polymer for High-Performance Organic Thin-Film Transistors Benjamin Nketia-Yawson, Hyo-Sang Lee, Dongkyun Seo, Youngwoon Yoon, Won-Tae Park, Kyungwon Kwak, Hae Jung Son, BongSoo Kim,* and Yong-Young Noh* Development of high-performance conjugated polymers is of primary importance for the fabrication of organic-based devices — such as organic thin film transistors (OTFTs), organic photovoltaics, organic memories, and sensors — on flexible plastic substrates at low temperature by cost-effective graphic art printing processes.[1–5] Recently, the performance of conjugated-polymer-based OTFTs has been enhanced significantly; this improvement has occurred primarily through developing donor (D)–acceptor (A) type conjugated polymers.[6–9] The stateof-the-art high-performance D–A polymers consist of alternating electron-rich donor moieties, such as oligothiophenes, and electron-deficient acceptor moieties, such as diketopyrrolopyrrole (DPP), isoindigo (IIG), and naphthalenedicarboximide; they have yielded very high carrier mobilities exceeding 14 and 7 cm2 V−1 s−1 for holes and electrons, respectively, in OFETs.[10,11] Among those acceptor moieties, it is well known that the highly electron-deficient DPP and IIG moieties having coplanar conjugated structures promote high charge delocalization over the polymer backbone and strong cofacial π–π interactions between polymer chains. On the other hand, benzothiadiazole (BT) based D–A-type polymers were developed earlier; their performance remained relatively lower (hole mobilities ≈1 × 10−3 to 3.6 cm2 V−1 s−1),[12–23] though the BT polymers can form a long-range ordered structures as well.
B. Nketia-Yawson, W.-T. Park, Prof. Y.-Y. Noh Department of Energy and Materials Engineering Dongguk University, 26 Pil-dong, 3 ga Jung-gu, Seoul 100-715, South Korea E-mail: [email protected] H.-S. Lee, Y. Yoon, Dr. H. J. Son Photo-electronic Hybrids Research Center Korea Institute of Science and Technology (KIST) Hwarangno14-gil 5, Seongbuk-gu, Seoul 136-791, Republic of Korea H.-S. Lee, Dr. H. J. Son Green School (School of Energy and Environment) Korea University 1 Anam-dong 5-ga, Seongbuk-gu, Seoul 136-701, Republic of Korea D. Seo, Prof. K. Kwak Department of Chemistry Chung-Ang University 84 Heukseok-ro, Dongjak-gu, Seoul 156-756, Republic of Korea Prof. B. Kim Department of Science Education Ewha Womans University 52 Ewhayeodae-gil, Seodaemun-gu, Seoul 120-750, Republic of Korea E-mail: [email protected]
DOI: 10.1002/adma.201500233
Adv. Mater. 2015, DOI: 10.1002/adma.201500233
In order to improve the carrier mobility of BT-based polymers, we explored two synthesizing strategies. First, fluorine atoms were introduced to the BT units. It is known that the replacement of hydrogens with fluorines can induce a planar geometry when fluorines are geometrically close to the neighboring thiophene moieties via F•••S non-covalent attractive interaction.[24–31] Additionally, C H(aromatic)•••F and C F•••πF interchain interactions become available.[27,28,32] These interactions can promote electrostatic and dispersion interactions between aromatic chains; these enhance the planarity of polymer backbones further, resulting in quite close (3.4–3.8 Å) π−π polymer chain stacking.[32] Furthermore, the insertion of fluorine atoms directly into the backbone of conjugated polymers may engender high thermal and oxidative stability, as well as high resistance to degradation, by lowering both the highest occupied molecular orbital (HOMO) and the lowest unoccupied molecular orbital (LUMO) energy levels of the conjugated polymer through the electron-withdrawing effect of the fluorine atoms.[32–34] After this first step, dithienosilole moieties were incorporated to construct a polymer backbone instead of the commonly used cyclopentadiene moieties.[12,35] Dithienosilole moieties are interesting to use due to their efficient charge transport through tight interchain packing of conjugated polymers.[36] Silole-containing organic materials have shown higher charge transport[16,37] and better solar cell properties[19,37,38] compared to their carbon analogues.[36,37] A key reason for these advantages is that the Si–C bond length (1.89 Å) in the dithienosilole unit is longer than the C–C bond length (1.53 Å) in the cyclopentadiene unit, allowing efficient interchain packing and strong π–π interactions.[36] With these favorable features in mind, we attempted to combine difluorobenzothiadiazole and dithienosilole moieties to synthesize a semiconducting polymer while expecting high coplanarity of the resulting polymer backbone and strong π−π interactions. In this work, we report the synthesis and electrical characterization of a highly planar poly(4-(4,4-bis(2-ethylhexyl)4H-silolo[3,2-b:4,5-b′]dithiophen-2-yl)-7-(4,4-bis(2-ethylhexyl)6-(thiophen-2-yl)-4 H -silolo[3,2- b :4,5- b ′ ]dithiophen-2-yl)5,6-difluorobenzo[c][1,2,5]thiadiazole) (PDFDT) (Figure 1) based on the dithienosilole and difluorobenzothiadiazole moieties. Excellent p-type characteristics with highest hole mobilities of 2.6, 2.8, and surprisingly 9.0 cm2 V−1 s−1 were obtained from top gate/bottom contact (TG/BC) OTFTs incorporating PDFDT as the active material (Figure 1a) and poly(methylmethacrylate) (PMMA), high-k poly(vinylidenefluoride–trifluoroethylene– chlorotrifluoroethylene) (P(VDF–TrFE–CTFE), and poly(vinylidenefluoride–trifluoroethylene) P(VDF–TrFE) as dielectric
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The optical properties of PDFDT were characterized using UV–vis absorption spectroscopy. The UV–vis absorption spectrum of the PDFDT polymer solution in chlorobenzene (CB) taken at 25 °C (Figure 1c) showed three distinct absorption bands of around 470, 645, and 693 nm. These bands correspond to π–π* transitions between frontier molecular orbitals in both dissolved and aggregated polymer chains (see the time-dependent density functional theory (TD–DFT) calculations below for the detailed transition assignments).[39] The existence of aggregations of polymers in the CB solution was confirmed by use of gel-permeation chromatograms. Unlike the gel-permeation chromatography taken at 150 °C in TCB, a gel-permeation chromatogram obtained at 80 °C in o-dichlorobenzene (DCB) showed a dual molecular weight distribution (Figure S3, Supporting Information). The higher molecular weight fraction was attributed to polymer Figure 1. a) Chemical structure of PDFDT polymer, b) top gate/bottom contact OTFT device structure, and c,d) UV–vis absorption spectra of PDFDT polymer in 0.008 mg mL−1 CB solution aggregates present in the DCB solution at (c), and in film state (d) at various temperatures. 80 °C due to the strong interchain interactions and the lower molecular weight (Mn = 18 kDa with a PDI of 1.76), relative to fully dissolved insulators, respectively. These mobilities, to the best of our knowledge, are the highest hole mobility values among the polymer chains. In addition, temperature-dependent UV–vis conjugated polymers employed in OTFTs except DPP and IIGabsorption behavior observed in both CB and DCB solutions based polymers. proved that the solutions contained some aggregated polymer The PDFDT polymer structure was carefully designed with chains. Both UV–vis absorption spectra from CB and DCB soluthe aid of density-functional-theory (DFT) calculations. First, tions (Figure 1c and Figure S4a, Supporting Information) cona simple monomeric structure of 4,7-bis(4,4-bis(methyl)-4Hsistently showed that with increasing solution temperature, the silolo[3,2-b:4,5-b′]dithiophen-2-yl)-5,6-difluorobenzo[c][1,2,5] maximum of the second band blueshifted slightly from 654 nm at 25 °C to 642 nm at 120 °C; in addition, the third band associthiadiazole in three conjugation motifs was calculated to check ated with polymer aggregates became weak. In solid states, the the potential effects of the different bond connections on the second and third bands displayed a very slight redshift of 4 nm. electronic and conformational structures. The three chemical The third band became relatively more pronounced compared structures are shown in Figure S1, Supporting Information. to the second band and the thermal annealing promoted the The orientations of 4,4-bis(methyl)-4H-silolo[3,2-b:4,5-b′] polymer aggregation (Figure 1d and Figure S4b, Supporting dithiophen-2-yl units with respect to the 5,6-difluorobenzo[c][1,2,5] Information). All these observations indicated that the PDFDT thiadiazole unit were varied. All the optimized structures of the polymer contained strong interactions between polymer chains HOMO/LUMO orbitals appeared to be nearly independent of even in solution state. It is also instructive to compare the UV– the bond connection type in terms of energy levels and elecvis absorption feature of PDFDT film with that of a structural tronic conjugation. Importantly, all the chemical structures were analog polymer with the non-fluorinated BT acceptors (poly(4extremely flat, with dihedral angles of 0°; this implied that this (4,4-bis(2-ethylhexyl)-4H-silolo[3,2-b:4,5-b′]dithiophen-2-yl)-7monomeric system would provide a flat expanded π-conjugation (4,4-bis(2-ethylhexyl)-6-(thiophen-2-yl)-4H-silolo[3,2-b:4,5-b′] system, and thus promote the interchain electronic interacdithiophen-2-yl)benzo[c][1,2,5]thiadiazole (PDHDT).[40] Much tions. This monomeric structure, substituted 2-ethylhexyl groups in place of methyl groups, was polymerized by Stille stronger aggregate-type UV–vis absorption was observed in the polymerization using 4,7-bis(6-bromo-4,4-bis(2-ethylhexyl)-4HPDFDT than in the PDHDT. This comparison reflects that the silolo[3,2-b:4,5-b′]dithiophen-2-yl)-5,6-difluorobenzo[c][1,2,5] interchain interaction in the PDFDT polymer was strengthened by the replacement of hydrogens with fluorines in the BT units, thiadiazole and 2,5-bis(trimethylstannyl)thiophene. After the as was intended.[24,41–43] polymerization was completed, palladium catalysts and oligomers were removed by successive Soxhlet extractions using The electrochemical properties of PDFDT were measured methanol, acetone, hexane, and chloroform. The polymer by cyclic voltammetry. A cyclic voltammogram of PDFDT structure was confirmed by H-NMR and elemental analysis. (Figure S5, Supporting Information) displays the oxidationThe number average molecular weight (Mn) was determined onset potential (Eox) of 0.44 V and the reduction onset potential to be 14 kDa with a polydispersity index (PDI) (Mw/Mn) of 2.4, (Ered) of −1.43 V with respect to ferrocene/ferrocenium (Fc/Fc+) measured in trichlorobenzene (TCB) solution at 150 °C (see the oxidation potential (assumed to be −4.8 eV below vacuum Supporting Information for more details). level), which corresponds to the HOMO and LUMO levels of
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−5.24 and −3.37 eV, respectively. The electrochemical HOMO– LUMO gap (Eg,cv) was 1.87 eV. The HOMO level was lowered by 0.14 eV due to the introduction of the electron-withdrawing fluorine atoms, compared to similarly structured polymers containing non-fluorinated BT and dithienosilole moieties.[19,40,44] Note also that keeping the HOMO level low is important for improving the stability of OTFTs.[37,45,46] In addition, the thermal properties of the PDFDT polymer were examined by using differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA). The DSC curve displayed a cold crystallization transition (Tc) of 90 °C without melting in the temperature range of 30–340 °C (Figure S6a, Supporting Information). The TGA curve showed a 5% weight loss temperature of 410 °C (Figure S6b, Supporting Information). To understand the electronic structure, electronic transitions, and molecular geometry of PDFDT polymer chains, density functional theory and time-dependent DFT calculations were conducted on a model molecule of DFDT3 consisting of three polymer repeating units. DFT calculations used the B3LYP functional with the 6-311G(d,p) basis set in Gaussian 09, with the assumption that the model molecule is dissolved in the CB solvent. Figure 2 presents the energy levels and surface plots of
frontier orbitals, a simulated UV–vis absorption spectrum, and optimized geometries of the DFDT3 molecule. The HOMO and LUMO orbitals displayed well-delocalized π-conjugation across the molecular backbone. The sub-HOMO and sub-LUMO orbitals showed a more localized π-conjugation; in particular, the sub-LUMO orbitals were governed by the electron-deficient BT moieties (Figure 2a). TD–DFT calculations predicted two excitation bands with maximum peaks at 424 and 609 nm (Figure 2b). Only primary excitations involved in each wavelength are indicated here and a full description is provided in Table S1 in the Supporting Information. The overall features agreed well with the UV–vis absorption spectrum of the fully dissolved PDFDT polymer. Thus, the first band around 470 nm can be attributed to the π–π* transitions, mainly HOMO − 3 → LUMO + 2 and HOMO → LUMO + 3; the second band around 645 nm is largely associated with the HOMO → LUMO excitation, as explained above. Figure 2c displays energy minimum geometries. The molecular backbone appeared to be very flat, with dihedral angles around single bonds less than 15°; this characteristic allows for the effective overlap of π orbitals.[47,48] Molecular widths were determined to be 12.25 and 16.01 Å in the parallel and perpendicular directions to the aromatic backbone plane, respectively.
Figure 2. a) Surface plots and energy levels of frontier orbitals of the DFDT3 molecule. Note that molecular energy levels are simply even-spaced for clarity. b) UV–vis absorption spectrum of the DFDT3 molecule predicted by TD–DFT method (solid line) with oscillator strengths (circles). c) Energy minimum geometries of the DFDT3 molecule in three directional views with molecular widths and dihedral angles.
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Figure 3. a) GIXD image of 50 nm thick PDFDT film. b) Line-cut profiles in the out-of-plane (qz) and in-plain (qy) directions. The inset shows the same line-cut profiles in the linear-scaled intensity to compare the dramatic difference in the (100) peak intensities.
To gain insight into the crystallinity and microstructure of PDFDT polymer film, grazing-incidence X-ray diffraction (GIXD) experiments were conducted.[49] Figure 3a displays the GIXD image of PDFDT film. A strong crystalline (100) peak at qz = 0.359 Å−1 was observed in the out-of-plane (qz) direction, while the much weaker (100) peak was shown in the in-plane (qy) direction. Similarly, a radial diffuse π–π stacking peak at 1.638 Å−1 was also detected, with higher intensity in the outof-plane direction than in the in-plane direction. Another broad radially distributed peak appeared at ≈1.3 Å−1. For detailed analysis, the line-cut profiles in the out-of-plane and in-plane directions were extracted along the qz direction at qy = 0.00 Å−1 and along the qy direction at qz = 0.05 Å−1, respectively (Figure 3b). The inset of Figure 3b shows a strong dependence of (100) peak intensities on the directions. The (100) peak corresponds to a d-spacing of 17.5 Å, which matches closely with the sum of the DFT-calculated molecular width of 16.01 Å and the van der Waals radii of atoms. A correlation length (Lc) was estimated based on Scherrer's equation[35] to be 12.7 nm (approximately seven stacks), which is comparable to values found in high performance DPP and IIG-based polymer films[10,50–53] In addition, a clear π–π stacking peak was observed at 1.630 Å−1 (3.85 Å) with an Lc of 3.50 nm (approximately nine stacks) and another weak π–π stacking peak at 1.839 Å−1 (3.42 Å) with an Lc of 2.87 nm (approximately eight stacks). Note that the very 4
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close (3.42 Å) π–π stacking has not been previously observed in the D–A-type conjugated polymers. Its presence reflects that polymer backbones can be put in contact with each other very intimately, as the strong tendency of polymer aggregate formation in solution and in film states was observed in UV–vis absorption experiments. Importantly, the highly planar backbone geometries of PDFDT polymer would be responsible for the strong interchain interactions between polymer chains. It should be also noted that the high intensities of (100) peak and π–π stacking peaks appeared in the same (i.e., out-of-plane) direction. This behavior is different from typical high-performance polythiophenes, DPP-based polymers, and long alkyl chained BT-contained polymers displaying strong (h00) peaks in the out-of-plane direction and π–π stacking (010) peaks inplane (qy) direction.[50,53–56] To explain this unusual characteristic, the polymer geometries were considered. The optimized model compound of DFDT3 displayed the 2-ethylhexyl chains lying perpendicular to the molecular backbone. The molecular width of the optimized model compound of DFDT3 in the side view was dictated by the stretched-out 2-ethylhexyl chains, and coincided with the GIXD lamellar spacing. This suggested that the polymer backbones lie parallel to the substrate with face-on orientation. However, if the polymer chains stacked in a simple fashion of one chain by one chain, such close π–π stacking cannot be observable. Thus, we propose a double-cabled polymer structure — as suggested previously by the Müllen group — with poly(2,6-(4,4-bis(hexadecyl)-4H-cyclopenta[2,1b:3,4-b′]-dithiophene)-alt-4,7-(2,1,3-benzothiadiazole).[12] That is, two wavy-shaped polymer chains stack in a fashion such that one planar linear chain stretches out; the upper planar linear chain lies on its top after a 180° rotation with respect to the polymer backbone axis (rotated stacking mode). A tight stacking of polymer chains is also possible when the polymer chains are slipped toward each other with a certain distance (slipped stacking mode). These polymer interchain packing motifs are illustrated in Figure S7 in Supporting Information. Basically, both stacking modes would exist in the polymer film to reduce the steric hindrance between bulky 2-ethylhexyl chains and promote the close π–π stacking between polymer backbones. With the proposed interchain packing model, the fact that both the lamellar spacing and the π–π stacking are pronounced in the same direction can be explained. Note that the PDFDT polymer structure is still under investigation to confirm our proposed structure and identify the details of interchain stacking modes further. The surface morphological microstructures of PDFDT thin-films were investigated using tapping mode atomic force microscopy (AFM). As shown in Figure S8 in the Supporting Information, the AFM height images of PDFDT thin-films spin-coated at different spin speeds and then annealed at various temperatures show tightly packed nanofibers that establish well-interconnected polymer chain networks to enhance chargecarrier transport. This morphological feature was confirmed by transmission electron microscopy imaging (Figure S9, Supporting Information). The 150 °C-annealed polymer films exhibited more uniform and more numerous grown nanofibers of networked polymer chains compared to the unannealed polymer films at room temperature (RT), and consistent with the UV–vis absorption spectra of the annealed PDFDT films.
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The electrical properties of the PDFDT polymer were measured by using staggered TG/BC OTFT device geometry with various polymer gate dielectrics. The PDFDT active layer (thickness, t ≈ 8.2 nm) was spin-coated on Au source and drain electrode patterned glass substrates, from 2 mg mL−1 concentration in CB at 2000 rpm for 60 s. Details of these fabrication processes are described in the Experimental section. The PDFDT– OTFTs exhibited typical p-channel output and transfer characteristics, as shown in Figure 4. The PDFDT–OTFTs showed various characteristics with annealing temperatures of semiconducting layers, dimensions of channel width (W) and length (L), and used dielectrics. The optimized charge carrier mobility of 2.6 cm2 V−1 s−1 for PMMA devices was obtained at 150 °C annealing, with W and L of 1000 and 10 µm, respectively; this is a marked improvement compared to the carrier mobility of 3 × 10−3 cm2 V−1 s−1 from previously reported BT-based polymer (TS6T1-BTD) using bottom-gate/bottom-contact organic field-effect transistor (OFET) structure on SiO2 dielectric with hexamethyldisilazane surface treatment.[40] Interestingly, the output curves of PDFDT–OTFTs with L = 10 µm exhibited non-saturating behavior due to shortchannel effect.[57] However, with longer channel length; output curves of PDFDT–OTFTs exhibited very good saturation behavior with PMMA dielectric resulting from increase channel resistance (Figure S11d and S13, Supporting Information). The highly improved mobility by thermal annealing can be attributed to the tightly packed, fibrously well-connected nature of the polymer chains in thin films when subjected to thermal annealing.
Adv. Mater. 2015, DOI: 10.1002/adma.201500233
The charge carrier mobility is dramatically improved even further by applying high-k ferroelectric polymers, P(VDF– TrFE–CTFE) (ε ≈ 20.9) and P(VDF–TrFE) (ε ≈ 10.4), as gate dielectrics for PDFDT OTFTs. Interestingly, PDFDT OTFTs with P(VDF–TrFE–CTFE) exhibited improved mobility of 2.8 cm2 V−1 s−1; more impressively, a highest mobility of 9.0 cm2 V−1 s−1 (average 6.6 cm2 V−1 s−1) was recorded with P(VDF–TrFE) dielectrics. The improvement in mobility is due to the induced band bending emanating from the large dipole polarization at the interface of the ferroelectric gate insulators and PDFDT polymer. This improves the accumulation of positive charges carriers, leading to a remarkable boost in hole mobility with the high-k P(VDF–TrFE) gate insulator.[10] In addition, the chlorine atom incorporated in the bulky CTFE units distorts the crystal sequence length, reducing the degree of crystallization in the P(VDF–TrFE–CTFE) backbone.[58] These features manifest in the negligible hysteresis observed in the transfer plots with P(VDF–TrFE–CTFE) compared to P(VDF–TrFE), which exhibit larger hysteresis owing to the strong polarization behavior that results from crystalline β-phase formation in the P(VDF–TrFE).[10] Another benefit of using those high-k ferroelectric polymer gate dielectrics is a reduction in operating voltage. The P(VDF–TrFE–CTFE) and P(VDF–TrFE) devices showed relatively low threshold voltage (VTh) values of −1.93 ± 0.80 V and −20.9 ± 0.62 V, respectively, compared with PMMA devices (VTh = −51.8 ± 1.6 V) due to the high capacitance of gate dielectric layers (capacitance: P(VDF– TrFE–CTFE) ≈ 68.5 nF cm−2, P(VDF–TrFE) ≈ 36.8 nF cm−2 and
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www.MaterialsViews.com Table 1. Summary of the electrical parameters of PDFDT OTFTs with PMMA, P(VDF–TrFE), and P(VDF–TrFE–CTFE) dielectrics annealed at various temperatures. Conditiona) Gate dielectric (Thickness) PMMA (≈500 nm)
T [°C]
µh, max (µh,avg) [cm2 V−1 s−1]
VTh [V]
S.S. [V dec−1]
Rc.W [kΩ cm]
Spin speedb) [rpm]
µh,max (µh,avg) [cm2 V−1 s−1]
VTh [V]
RT
0.24 (0.15)
−33.1 ±1.2
−11.2 ±1.6
29.2–101
1000
0.67 (0.41)
−38.8 ±2.5
100
0.31 (0.23)
−35.4 ±1.9
−13.3 ±0.8
18.6–77.2
1500
1.92 (1.42)
−48.9 ±1.6
150
2.63 (1.87)
−51.8 ±1.6
−15.5 ±0.6
1.50–21.2
2000
2.63 (1.87)
−51.8 ±1.6
200
2.22 (1.26)
−48.4 ±3.1
−14.9 ±0.7
3.06–65.1
3000
0.79 (0.61)
−42.7 ±1.2
P(VDF–TrFE) (≈250 nm)
150
9.05 (6.56)
−20.9 ±0.6
−3.92 ±0.3
7.18–17.4
N.A.
P(DVF–TrFE–CTFE) (≈270 nm)
150
2.82 (2.01)
−1.93 ±0.8
−2.37 ±0.2
0.62–2.37
N.A.
a) The TFT performance of more than 15 devices with channel length and width of 10 and 1000 µm, respectively, measured in a N2-filled glove box; b)All the devices were annealed at 150 °C.
PMMA ≈ 6.2 nF cm−2). All the basic transistor parameters are summarized in Table 1. In order to achieve optimum device performance, different thicknesses of PDFDT-layer OTFTs were fabricated by varying the spin speed from 1000 rpm (t ≈ 14.2 nm) to 3000 rpm (t ≈ 7.6 nm) for 60 s. Increasing the spin speed improves device performance due to a reduction in the contact resistance (Rc) of PDFDT OTFTs (Figure S10, Supporting Information). A further increase at 3000 rpm resulted in high Rc due to a PDFDT active layer that was too thin. Clearly, PDFDT OTFT devices showed resistance dependency with gate voltage and channel length using the transfer line method (TLM) (Figure S11, Supporting Information). A width-normalized Rc range of 3.15–10.66 kΩ cm was observed for gate voltages from −55 to −80 V, respectively, by TLM, which is comparable to the 1.50–21.2 kΩ cm range extracted using the Y-function method for 150 °C-annealed devices. Sequential estimation of contact resistance using that Y-function method has been detailed in our previous report and by other research.[59,60] In addition, we conducted a bias-stress stability test of PDFDT OTFTs with PMMA gate insulator in a nitrogen atmosphere and in air (Figure S12, Supporting Information). The devices showed good stability in both nitrogen and air owing to the lower HOMO level (−5.24 eV) of the PDFDT polymer. However, a steady decrease in on-current was observed in air, resulting from the unintended oxygen doping; this caused a drastic decrease in the off-current by the trapping of electrons in the presence of oxygen and moisture. In conclusion, we have designed and synthesized a new fluorinated BT-based conjugated polymer, PDFDT, as an active material for OTFT devices. PDFDT OTFTs showed high charge carrier mobilities of 2.6, 2.8, and 9.0 cm2 V−1 s−1 in TG/BC OTFTs using common PMMA, and high-k P(VDF–TrFE–CTFE) and P(VDF–TrFE) dielectrics, respectively. The remarkable boost in hole mobility results from molecular energy modulation through dipoles of the fluorinated ferroelectric dielectrics and PDFDT, which decrease the gap between HOMO and Fermi level by pull down the Fermi level with respect to the source-drain electrodes. Moreover, high bias-stress stability was obtained for PDFDT devices presumably because of the lower HOMO energy level. This work highlights that fluorinated BT can be an important class of acceptor moiety to pose polymer HOMO level in a zone of high chemical stability and produce
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strong interchain interactions. It can also induce a highly planar conjugated-polymer backbone. These features allow us to achieve high charge-carrier mobility and high device-operational stability.
Experimental Section OTFT Fabrication and Characterization: Top-gate/bottom-contact geometry OTFT devices were fabricated. For a charge-injection electrode, Au was deposited on 3 nm Ni adhesive layer on Corning Eagle XG glass substrates. The source and drain electrodes (Au/Ni = 13/3 nm) were patterned using a photolithography process on the glass substrates. The channel lengths were 10, 20, 30, and 50 µm, and the width was 1000 µm. A 2 mg mL−1 solution of the PDFDT polymer in CB was spin-coated at 1000–3000 rpm for 60 s and thermally annealed at temperatures ranging from 100 to 200 °C for 30 min and then allowed to cool down slowly in a N2-purged glove box. After thermal annealing, PMMA (Sigma–Aldrich, Mn: 120 kDa) film as a gate dielectric layer was spin-coated on the active layer from 80 mg mL−1 solution in n-butyl acetate at 2000 rpm for 60 s and the coated PMMA film was then baked at 80 °C for 2 h in a N2 purged glove box. In the case of ferroelectric dielectric insulators, P(VDF–TrFE–CTFE) (Solvay, 63:28:9 mol% random terpolymer, Mw: 290 kDa) and P(VDF–TrFE) (Solvay, 70:30 mol% random copolymer) were dissolved in methyl ethyl ketone at the same concentration of 30 mg mL−1. Thereafter, each solution was spin-coated on the active layer at 2000 rpm for 60 s and then baked at 80 °C for 2 h. The OTFTs were completed by depositing 50 nm top-gate aluminum electrodes via thermal evaporation using a metal shadow mask. OTFT devices were characterized using a Keithley 4200-SCS semiconductor parameter analyzer connected to an N2purged glove box probe station.
Supporting Information Supporting Information is available from the Wiley Online Library or from the author.
Acknowledgements B.N.-Y. and H.-S.L. contributed equally to this work. This work was supported by Nano·Material Technology Development Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT and Future Planning (NRF-2014M3A7B4051749), the Center for Advanced Soft-Electronics funded by the Ministry
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Received: January 16, 2015 Revised: March 4, 2015 Published online:
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of Science, ICT and Future Planning as Global Frontier Project (2013M3A6A5073183), a New and Renewable Energy Program of the Korea Institute of Energy Technology Evaluation and Planning (KETEP) grant funded by the Korea Government Ministry of Trade, Industry and Energy (MTIE) (20133030000130), and the National Research Foundation of Korea Grant funded by the Korean Government (MSIP) (2013, University-Institute corporation program). The authors thank Prof. Hyunjung Kim in Department of Physics at Sogang University for help and discussion on the analysis of GIXD data.
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A Highly Planar Fluorinated Benzothiadiazole-Based Conjugated Polymer for High-Performance Organic Thin-Film Transistors Benjamin Nketia-Yawson, Hyo-Sang Lee, Dongkyun Seo, Youngwoon Yoon, Won-Tae Park, Kyungwon Kwak, Hae Jung Son, BongSoo Kim,* and Yong-Young Noh* Development of high-performance conjugated polymers is of primary importance for the fabrication of organic-based devices — such as organic thin film transistors (OTFTs), organic photovoltaics, organic memories, and sensors — on flexible plastic substrates at low temperature by cost-effective graphic art printing processes.[1–5] Recently, the performance of conjugated-polymer-based OTFTs has been enhanced significantly; this improvement has occurred primarily through developing donor (D)–acceptor (A) type conjugated polymers.[6–9] The stateof-the-art high-performance D–A polymers consist of alternating electron-rich donor moieties, such as oligothiophenes, and electron-deficient acceptor moieties, such as diketopyrrolopyrrole (DPP), isoindigo (IIG), and naphthalenedicarboximide; they have yielded very high carrier mobilities exceeding 14 and 7 cm2 V−1 s−1 for holes and electrons, respectively, in OFETs.[10,11] Among those acceptor moieties, it is well known that the highly electron-deficient DPP and IIG moieties having coplanar conjugated structures promote high charge delocalization over the polymer backbone and strong cofacial π–π interactions between polymer chains. On the other hand, benzothiadiazole (BT) based D–A-type polymers were developed earlier; their performance remained relatively lower (hole mobilities ≈1 × 10−3 to 3.6 cm2 V−1 s−1),[12–23] though the BT polymers can form a long-range ordered structures as well.
B. Nketia-Yawson, W.-T. Park, Prof. Y.-Y. Noh Department of Energy and Materials Engineering Dongguk University, 26 Pil-dong, 3 ga Jung-gu, Seoul 100-715, South Korea E-mail: [email protected] H.-S. Lee, Y. Yoon, Dr. H. J. Son Photo-electronic Hybrids Research Center Korea Institute of Science and Technology (KIST) Hwarangno14-gil 5, Seongbuk-gu, Seoul 136-791, Republic of Korea H.-S. Lee, Dr. H. J. Son Green School (School of Energy and Environment) Korea University 1 Anam-dong 5-ga, Seongbuk-gu, Seoul 136-701, Republic of Korea D. Seo, Prof. K. Kwak Department of Chemistry Chung-Ang University 84 Heukseok-ro, Dongjak-gu, Seoul 156-756, Republic of Korea Prof. B. Kim Department of Science Education Ewha Womans University 52 Ewhayeodae-gil, Seodaemun-gu, Seoul 120-750, Republic of Korea E-mail: [email protected]
DOI: 10.1002/adma.201500233
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In order to improve the carrier mobility of BT-based polymers, we explored two synthesizing strategies. First, fluorine atoms were introduced to the BT units. It is known that the replacement of hydrogens with fluorines can induce a planar geometry when fluorines are geometrically close to the neighboring thiophene moieties via F•••S non-covalent attractive interaction.[24–31] Additionally, C H(aromatic)•••F and C F•••πF interchain interactions become available.[27,28,32] These interactions can promote electrostatic and dispersion interactions between aromatic chains; these enhance the planarity of polymer backbones further, resulting in quite close (3.4–3.8 Å) π−π polymer chain stacking.[32] Furthermore, the insertion of fluorine atoms directly into the backbone of conjugated polymers may engender high thermal and oxidative stability, as well as high resistance to degradation, by lowering both the highest occupied molecular orbital (HOMO) and the lowest unoccupied molecular orbital (LUMO) energy levels of the conjugated polymer through the electron-withdrawing effect of the fluorine atoms.[32–34] After this first step, dithienosilole moieties were incorporated to construct a polymer backbone instead of the commonly used cyclopentadiene moieties.[12,35] Dithienosilole moieties are interesting to use due to their efficient charge transport through tight interchain packing of conjugated polymers.[36] Silole-containing organic materials have shown higher charge transport[16,37] and better solar cell properties[19,37,38] compared to their carbon analogues.[36,37] A key reason for these advantages is that the Si–C bond length (1.89 Å) in the dithienosilole unit is longer than the C–C bond length (1.53 Å) in the cyclopentadiene unit, allowing efficient interchain packing and strong π–π interactions.[36] With these favorable features in mind, we attempted to combine difluorobenzothiadiazole and dithienosilole moieties to synthesize a semiconducting polymer while expecting high coplanarity of the resulting polymer backbone and strong π−π interactions. In this work, we report the synthesis and electrical characterization of a highly planar poly(4-(4,4-bis(2-ethylhexyl)4H-silolo[3,2-b:4,5-b′]dithiophen-2-yl)-7-(4,4-bis(2-ethylhexyl)6-(thiophen-2-yl)-4 H -silolo[3,2- b :4,5- b ′ ]dithiophen-2-yl)5,6-difluorobenzo[c][1,2,5]thiadiazole) (PDFDT) (Figure 1) based on the dithienosilole and difluorobenzothiadiazole moieties. Excellent p-type characteristics with highest hole mobilities of 2.6, 2.8, and surprisingly 9.0 cm2 V−1 s−1 were obtained from top gate/bottom contact (TG/BC) OTFTs incorporating PDFDT as the active material (Figure 1a) and poly(methylmethacrylate) (PMMA), high-k poly(vinylidenefluoride–trifluoroethylene– chlorotrifluoroethylene) (P(VDF–TrFE–CTFE), and poly(vinylidenefluoride–trifluoroethylene) P(VDF–TrFE) as dielectric
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The optical properties of PDFDT were characterized using UV–vis absorption spectroscopy. The UV–vis absorption spectrum of the PDFDT polymer solution in chlorobenzene (CB) taken at 25 °C (Figure 1c) showed three distinct absorption bands of around 470, 645, and 693 nm. These bands correspond to π–π* transitions between frontier molecular orbitals in both dissolved and aggregated polymer chains (see the time-dependent density functional theory (TD–DFT) calculations below for the detailed transition assignments).[39] The existence of aggregations of polymers in the CB solution was confirmed by use of gel-permeation chromatograms. Unlike the gel-permeation chromatography taken at 150 °C in TCB, a gel-permeation chromatogram obtained at 80 °C in o-dichlorobenzene (DCB) showed a dual molecular weight distribution (Figure S3, Supporting Information). The higher molecular weight fraction was attributed to polymer Figure 1. a) Chemical structure of PDFDT polymer, b) top gate/bottom contact OTFT device structure, and c,d) UV–vis absorption spectra of PDFDT polymer in 0.008 mg mL−1 CB solution aggregates present in the DCB solution at (c), and in film state (d) at various temperatures. 80 °C due to the strong interchain interactions and the lower molecular weight (Mn = 18 kDa with a PDI of 1.76), relative to fully dissolved insulators, respectively. These mobilities, to the best of our knowledge, are the highest hole mobility values among the polymer chains. In addition, temperature-dependent UV–vis conjugated polymers employed in OTFTs except DPP and IIGabsorption behavior observed in both CB and DCB solutions based polymers. proved that the solutions contained some aggregated polymer The PDFDT polymer structure was carefully designed with chains. Both UV–vis absorption spectra from CB and DCB soluthe aid of density-functional-theory (DFT) calculations. First, tions (Figure 1c and Figure S4a, Supporting Information) cona simple monomeric structure of 4,7-bis(4,4-bis(methyl)-4Hsistently showed that with increasing solution temperature, the silolo[3,2-b:4,5-b′]dithiophen-2-yl)-5,6-difluorobenzo[c][1,2,5] maximum of the second band blueshifted slightly from 654 nm at 25 °C to 642 nm at 120 °C; in addition, the third band associthiadiazole in three conjugation motifs was calculated to check ated with polymer aggregates became weak. In solid states, the the potential effects of the different bond connections on the second and third bands displayed a very slight redshift of 4 nm. electronic and conformational structures. The three chemical The third band became relatively more pronounced compared structures are shown in Figure S1, Supporting Information. to the second band and the thermal annealing promoted the The orientations of 4,4-bis(methyl)-4H-silolo[3,2-b:4,5-b′] polymer aggregation (Figure 1d and Figure S4b, Supporting dithiophen-2-yl units with respect to the 5,6-difluorobenzo[c][1,2,5] Information). All these observations indicated that the PDFDT thiadiazole unit were varied. All the optimized structures of the polymer contained strong interactions between polymer chains HOMO/LUMO orbitals appeared to be nearly independent of even in solution state. It is also instructive to compare the UV– the bond connection type in terms of energy levels and elecvis absorption feature of PDFDT film with that of a structural tronic conjugation. Importantly, all the chemical structures were analog polymer with the non-fluorinated BT acceptors (poly(4extremely flat, with dihedral angles of 0°; this implied that this (4,4-bis(2-ethylhexyl)-4H-silolo[3,2-b:4,5-b′]dithiophen-2-yl)-7monomeric system would provide a flat expanded π-conjugation (4,4-bis(2-ethylhexyl)-6-(thiophen-2-yl)-4H-silolo[3,2-b:4,5-b′] system, and thus promote the interchain electronic interacdithiophen-2-yl)benzo[c][1,2,5]thiadiazole (PDHDT).[40] Much tions. This monomeric structure, substituted 2-ethylhexyl groups in place of methyl groups, was polymerized by Stille stronger aggregate-type UV–vis absorption was observed in the polymerization using 4,7-bis(6-bromo-4,4-bis(2-ethylhexyl)-4HPDFDT than in the PDHDT. This comparison reflects that the silolo[3,2-b:4,5-b′]dithiophen-2-yl)-5,6-difluorobenzo[c][1,2,5] interchain interaction in the PDFDT polymer was strengthened by the replacement of hydrogens with fluorines in the BT units, thiadiazole and 2,5-bis(trimethylstannyl)thiophene. After the as was intended.[24,41–43] polymerization was completed, palladium catalysts and oligomers were removed by successive Soxhlet extractions using The electrochemical properties of PDFDT were measured methanol, acetone, hexane, and chloroform. The polymer by cyclic voltammetry. A cyclic voltammogram of PDFDT structure was confirmed by H-NMR and elemental analysis. (Figure S5, Supporting Information) displays the oxidationThe number average molecular weight (Mn) was determined onset potential (Eox) of 0.44 V and the reduction onset potential to be 14 kDa with a polydispersity index (PDI) (Mw/Mn) of 2.4, (Ered) of −1.43 V with respect to ferrocene/ferrocenium (Fc/Fc+) measured in trichlorobenzene (TCB) solution at 150 °C (see the oxidation potential (assumed to be −4.8 eV below vacuum Supporting Information for more details). level), which corresponds to the HOMO and LUMO levels of
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−5.24 and −3.37 eV, respectively. The electrochemical HOMO– LUMO gap (Eg,cv) was 1.87 eV. The HOMO level was lowered by 0.14 eV due to the introduction of the electron-withdrawing fluorine atoms, compared to similarly structured polymers containing non-fluorinated BT and dithienosilole moieties.[19,40,44] Note also that keeping the HOMO level low is important for improving the stability of OTFTs.[37,45,46] In addition, the thermal properties of the PDFDT polymer were examined by using differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA). The DSC curve displayed a cold crystallization transition (Tc) of 90 °C without melting in the temperature range of 30–340 °C (Figure S6a, Supporting Information). The TGA curve showed a 5% weight loss temperature of 410 °C (Figure S6b, Supporting Information). To understand the electronic structure, electronic transitions, and molecular geometry of PDFDT polymer chains, density functional theory and time-dependent DFT calculations were conducted on a model molecule of DFDT3 consisting of three polymer repeating units. DFT calculations used the B3LYP functional with the 6-311G(d,p) basis set in Gaussian 09, with the assumption that the model molecule is dissolved in the CB solvent. Figure 2 presents the energy levels and surface plots of
frontier orbitals, a simulated UV–vis absorption spectrum, and optimized geometries of the DFDT3 molecule. The HOMO and LUMO orbitals displayed well-delocalized π-conjugation across the molecular backbone. The sub-HOMO and sub-LUMO orbitals showed a more localized π-conjugation; in particular, the sub-LUMO orbitals were governed by the electron-deficient BT moieties (Figure 2a). TD–DFT calculations predicted two excitation bands with maximum peaks at 424 and 609 nm (Figure 2b). Only primary excitations involved in each wavelength are indicated here and a full description is provided in Table S1 in the Supporting Information. The overall features agreed well with the UV–vis absorption spectrum of the fully dissolved PDFDT polymer. Thus, the first band around 470 nm can be attributed to the π–π* transitions, mainly HOMO − 3 → LUMO + 2 and HOMO → LUMO + 3; the second band around 645 nm is largely associated with the HOMO → LUMO excitation, as explained above. Figure 2c displays energy minimum geometries. The molecular backbone appeared to be very flat, with dihedral angles around single bonds less than 15°; this characteristic allows for the effective overlap of π orbitals.[47,48] Molecular widths were determined to be 12.25 and 16.01 Å in the parallel and perpendicular directions to the aromatic backbone plane, respectively.
Figure 2. a) Surface plots and energy levels of frontier orbitals of the DFDT3 molecule. Note that molecular energy levels are simply even-spaced for clarity. b) UV–vis absorption spectrum of the DFDT3 molecule predicted by TD–DFT method (solid line) with oscillator strengths (circles). c) Energy minimum geometries of the DFDT3 molecule in three directional views with molecular widths and dihedral angles.
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Figure 3. a) GIXD image of 50 nm thick PDFDT film. b) Line-cut profiles in the out-of-plane (qz) and in-plain (qy) directions. The inset shows the same line-cut profiles in the linear-scaled intensity to compare the dramatic difference in the (100) peak intensities.
To gain insight into the crystallinity and microstructure of PDFDT polymer film, grazing-incidence X-ray diffraction (GIXD) experiments were conducted.[49] Figure 3a displays the GIXD image of PDFDT film. A strong crystalline (100) peak at qz = 0.359 Å−1 was observed in the out-of-plane (qz) direction, while the much weaker (100) peak was shown in the in-plane (qy) direction. Similarly, a radial diffuse π–π stacking peak at 1.638 Å−1 was also detected, with higher intensity in the outof-plane direction than in the in-plane direction. Another broad radially distributed peak appeared at ≈1.3 Å−1. For detailed analysis, the line-cut profiles in the out-of-plane and in-plane directions were extracted along the qz direction at qy = 0.00 Å−1 and along the qy direction at qz = 0.05 Å−1, respectively (Figure 3b). The inset of Figure 3b shows a strong dependence of (100) peak intensities on the directions. The (100) peak corresponds to a d-spacing of 17.5 Å, which matches closely with the sum of the DFT-calculated molecular width of 16.01 Å and the van der Waals radii of atoms. A correlation length (Lc) was estimated based on Scherrer's equation[35] to be 12.7 nm (approximately seven stacks), which is comparable to values found in high performance DPP and IIG-based polymer films[10,50–53] In addition, a clear π–π stacking peak was observed at 1.630 Å−1 (3.85 Å) with an Lc of 3.50 nm (approximately nine stacks) and another weak π–π stacking peak at 1.839 Å−1 (3.42 Å) with an Lc of 2.87 nm (approximately eight stacks). Note that the very 4
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close (3.42 Å) π–π stacking has not been previously observed in the D–A-type conjugated polymers. Its presence reflects that polymer backbones can be put in contact with each other very intimately, as the strong tendency of polymer aggregate formation in solution and in film states was observed in UV–vis absorption experiments. Importantly, the highly planar backbone geometries of PDFDT polymer would be responsible for the strong interchain interactions between polymer chains. It should be also noted that the high intensities of (100) peak and π–π stacking peaks appeared in the same (i.e., out-of-plane) direction. This behavior is different from typical high-performance polythiophenes, DPP-based polymers, and long alkyl chained BT-contained polymers displaying strong (h00) peaks in the out-of-plane direction and π–π stacking (010) peaks inplane (qy) direction.[50,53–56] To explain this unusual characteristic, the polymer geometries were considered. The optimized model compound of DFDT3 displayed the 2-ethylhexyl chains lying perpendicular to the molecular backbone. The molecular width of the optimized model compound of DFDT3 in the side view was dictated by the stretched-out 2-ethylhexyl chains, and coincided with the GIXD lamellar spacing. This suggested that the polymer backbones lie parallel to the substrate with face-on orientation. However, if the polymer chains stacked in a simple fashion of one chain by one chain, such close π–π stacking cannot be observable. Thus, we propose a double-cabled polymer structure — as suggested previously by the Müllen group — with poly(2,6-(4,4-bis(hexadecyl)-4H-cyclopenta[2,1b:3,4-b′]-dithiophene)-alt-4,7-(2,1,3-benzothiadiazole).[12] That is, two wavy-shaped polymer chains stack in a fashion such that one planar linear chain stretches out; the upper planar linear chain lies on its top after a 180° rotation with respect to the polymer backbone axis (rotated stacking mode). A tight stacking of polymer chains is also possible when the polymer chains are slipped toward each other with a certain distance (slipped stacking mode). These polymer interchain packing motifs are illustrated in Figure S7 in Supporting Information. Basically, both stacking modes would exist in the polymer film to reduce the steric hindrance between bulky 2-ethylhexyl chains and promote the close π–π stacking between polymer backbones. With the proposed interchain packing model, the fact that both the lamellar spacing and the π–π stacking are pronounced in the same direction can be explained. Note that the PDFDT polymer structure is still under investigation to confirm our proposed structure and identify the details of interchain stacking modes further. The surface morphological microstructures of PDFDT thin-films were investigated using tapping mode atomic force microscopy (AFM). As shown in Figure S8 in the Supporting Information, the AFM height images of PDFDT thin-films spin-coated at different spin speeds and then annealed at various temperatures show tightly packed nanofibers that establish well-interconnected polymer chain networks to enhance chargecarrier transport. This morphological feature was confirmed by transmission electron microscopy imaging (Figure S9, Supporting Information). The 150 °C-annealed polymer films exhibited more uniform and more numerous grown nanofibers of networked polymer chains compared to the unannealed polymer films at room temperature (RT), and consistent with the UV–vis absorption spectra of the annealed PDFDT films.
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COMMUNICATION Figure 4. The output and transfer characteristics of PDFDT–OTFTs with PMMA as gate dielectric are shown in (a) and (d), respectively; the output and transfer characteristics of PDFDT–OTFTs with P(VDF–TrFE–CTFE) as gate dielectric are shown in (b) and (e), respectively; the output and transfer characteristics of PDFDT–OTFTs with P(VDF–TrFE) as gate dielectric are shown in (c) and (f), respectively. The insets of (a)–(c) show chemical structure of gate dielectric materials employed in the OTFTs.
The electrical properties of the PDFDT polymer were measured by using staggered TG/BC OTFT device geometry with various polymer gate dielectrics. The PDFDT active layer (thickness, t ≈ 8.2 nm) was spin-coated on Au source and drain electrode patterned glass substrates, from 2 mg mL−1 concentration in CB at 2000 rpm for 60 s. Details of these fabrication processes are described in the Experimental section. The PDFDT– OTFTs exhibited typical p-channel output and transfer characteristics, as shown in Figure 4. The PDFDT–OTFTs showed various characteristics with annealing temperatures of semiconducting layers, dimensions of channel width (W) and length (L), and used dielectrics. The optimized charge carrier mobility of 2.6 cm2 V−1 s−1 for PMMA devices was obtained at 150 °C annealing, with W and L of 1000 and 10 µm, respectively; this is a marked improvement compared to the carrier mobility of 3 × 10−3 cm2 V−1 s−1 from previously reported BT-based polymer (TS6T1-BTD) using bottom-gate/bottom-contact organic field-effect transistor (OFET) structure on SiO2 dielectric with hexamethyldisilazane surface treatment.[40] Interestingly, the output curves of PDFDT–OTFTs with L = 10 µm exhibited non-saturating behavior due to shortchannel effect.[57] However, with longer channel length; output curves of PDFDT–OTFTs exhibited very good saturation behavior with PMMA dielectric resulting from increase channel resistance (Figure S11d and S13, Supporting Information). The highly improved mobility by thermal annealing can be attributed to the tightly packed, fibrously well-connected nature of the polymer chains in thin films when subjected to thermal annealing.
Adv. Mater. 2015, DOI: 10.1002/adma.201500233
The charge carrier mobility is dramatically improved even further by applying high-k ferroelectric polymers, P(VDF– TrFE–CTFE) (ε ≈ 20.9) and P(VDF–TrFE) (ε ≈ 10.4), as gate dielectrics for PDFDT OTFTs. Interestingly, PDFDT OTFTs with P(VDF–TrFE–CTFE) exhibited improved mobility of 2.8 cm2 V−1 s−1; more impressively, a highest mobility of 9.0 cm2 V−1 s−1 (average 6.6 cm2 V−1 s−1) was recorded with P(VDF–TrFE) dielectrics. The improvement in mobility is due to the induced band bending emanating from the large dipole polarization at the interface of the ferroelectric gate insulators and PDFDT polymer. This improves the accumulation of positive charges carriers, leading to a remarkable boost in hole mobility with the high-k P(VDF–TrFE) gate insulator.[10] In addition, the chlorine atom incorporated in the bulky CTFE units distorts the crystal sequence length, reducing the degree of crystallization in the P(VDF–TrFE–CTFE) backbone.[58] These features manifest in the negligible hysteresis observed in the transfer plots with P(VDF–TrFE–CTFE) compared to P(VDF–TrFE), which exhibit larger hysteresis owing to the strong polarization behavior that results from crystalline β-phase formation in the P(VDF–TrFE).[10] Another benefit of using those high-k ferroelectric polymer gate dielectrics is a reduction in operating voltage. The P(VDF–TrFE–CTFE) and P(VDF–TrFE) devices showed relatively low threshold voltage (VTh) values of −1.93 ± 0.80 V and −20.9 ± 0.62 V, respectively, compared with PMMA devices (VTh = −51.8 ± 1.6 V) due to the high capacitance of gate dielectric layers (capacitance: P(VDF– TrFE–CTFE) ≈ 68.5 nF cm−2, P(VDF–TrFE) ≈ 36.8 nF cm−2 and
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www.MaterialsViews.com Table 1. Summary of the electrical parameters of PDFDT OTFTs with PMMA, P(VDF–TrFE), and P(VDF–TrFE–CTFE) dielectrics annealed at various temperatures. Conditiona) Gate dielectric (Thickness) PMMA (≈500 nm)
T [°C]
µh, max (µh,avg) [cm2 V−1 s−1]
VTh [V]
S.S. [V dec−1]
Rc.W [kΩ cm]
Spin speedb) [rpm]
µh,max (µh,avg) [cm2 V−1 s−1]
VTh [V]
RT
0.24 (0.15)
−33.1 ±1.2
−11.2 ±1.6
29.2–101
1000
0.67 (0.41)
−38.8 ±2.5
100
0.31 (0.23)
−35.4 ±1.9
−13.3 ±0.8
18.6–77.2
1500
1.92 (1.42)
−48.9 ±1.6
150
2.63 (1.87)
−51.8 ±1.6
−15.5 ±0.6
1.50–21.2
2000
2.63 (1.87)
−51.8 ±1.6
200
2.22 (1.26)
−48.4 ±3.1
−14.9 ±0.7
3.06–65.1
3000
0.79 (0.61)
−42.7 ±1.2
P(VDF–TrFE) (≈250 nm)
150
9.05 (6.56)
−20.9 ±0.6
−3.92 ±0.3
7.18–17.4
N.A.
P(DVF–TrFE–CTFE) (≈270 nm)
150
2.82 (2.01)
−1.93 ±0.8
−2.37 ±0.2
0.62–2.37
N.A.
a) The TFT performance of more than 15 devices with channel length and width of 10 and 1000 µm, respectively, measured in a N2-filled glove box; b)All the devices were annealed at 150 °C.
PMMA ≈ 6.2 nF cm−2). All the basic transistor parameters are summarized in Table 1. In order to achieve optimum device performance, different thicknesses of PDFDT-layer OTFTs were fabricated by varying the spin speed from 1000 rpm (t ≈ 14.2 nm) to 3000 rpm (t ≈ 7.6 nm) for 60 s. Increasing the spin speed improves device performance due to a reduction in the contact resistance (Rc) of PDFDT OTFTs (Figure S10, Supporting Information). A further increase at 3000 rpm resulted in high Rc due to a PDFDT active layer that was too thin. Clearly, PDFDT OTFT devices showed resistance dependency with gate voltage and channel length using the transfer line method (TLM) (Figure S11, Supporting Information). A width-normalized Rc range of 3.15–10.66 kΩ cm was observed for gate voltages from −55 to −80 V, respectively, by TLM, which is comparable to the 1.50–21.2 kΩ cm range extracted using the Y-function method for 150 °C-annealed devices. Sequential estimation of contact resistance using that Y-function method has been detailed in our previous report and by other research.[59,60] In addition, we conducted a bias-stress stability test of PDFDT OTFTs with PMMA gate insulator in a nitrogen atmosphere and in air (Figure S12, Supporting Information). The devices showed good stability in both nitrogen and air owing to the lower HOMO level (−5.24 eV) of the PDFDT polymer. However, a steady decrease in on-current was observed in air, resulting from the unintended oxygen doping; this caused a drastic decrease in the off-current by the trapping of electrons in the presence of oxygen and moisture. In conclusion, we have designed and synthesized a new fluorinated BT-based conjugated polymer, PDFDT, as an active material for OTFT devices. PDFDT OTFTs showed high charge carrier mobilities of 2.6, 2.8, and 9.0 cm2 V−1 s−1 in TG/BC OTFTs using common PMMA, and high-k P(VDF–TrFE–CTFE) and P(VDF–TrFE) dielectrics, respectively. The remarkable boost in hole mobility results from molecular energy modulation through dipoles of the fluorinated ferroelectric dielectrics and PDFDT, which decrease the gap between HOMO and Fermi level by pull down the Fermi level with respect to the source-drain electrodes. Moreover, high bias-stress stability was obtained for PDFDT devices presumably because of the lower HOMO energy level. This work highlights that fluorinated BT can be an important class of acceptor moiety to pose polymer HOMO level in a zone of high chemical stability and produce
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strong interchain interactions. It can also induce a highly planar conjugated-polymer backbone. These features allow us to achieve high charge-carrier mobility and high device-operational stability.
Experimental Section OTFT Fabrication and Characterization: Top-gate/bottom-contact geometry OTFT devices were fabricated. For a charge-injection electrode, Au was deposited on 3 nm Ni adhesive layer on Corning Eagle XG glass substrates. The source and drain electrodes (Au/Ni = 13/3 nm) were patterned using a photolithography process on the glass substrates. The channel lengths were 10, 20, 30, and 50 µm, and the width was 1000 µm. A 2 mg mL−1 solution of the PDFDT polymer in CB was spin-coated at 1000–3000 rpm for 60 s and thermally annealed at temperatures ranging from 100 to 200 °C for 30 min and then allowed to cool down slowly in a N2-purged glove box. After thermal annealing, PMMA (Sigma–Aldrich, Mn: 120 kDa) film as a gate dielectric layer was spin-coated on the active layer from 80 mg mL−1 solution in n-butyl acetate at 2000 rpm for 60 s and the coated PMMA film was then baked at 80 °C for 2 h in a N2 purged glove box. In the case of ferroelectric dielectric insulators, P(VDF–TrFE–CTFE) (Solvay, 63:28:9 mol% random terpolymer, Mw: 290 kDa) and P(VDF–TrFE) (Solvay, 70:30 mol% random copolymer) were dissolved in methyl ethyl ketone at the same concentration of 30 mg mL−1. Thereafter, each solution was spin-coated on the active layer at 2000 rpm for 60 s and then baked at 80 °C for 2 h. The OTFTs were completed by depositing 50 nm top-gate aluminum electrodes via thermal evaporation using a metal shadow mask. OTFT devices were characterized using a Keithley 4200-SCS semiconductor parameter analyzer connected to an N2purged glove box probe station.
Supporting Information Supporting Information is available from the Wiley Online Library or from the author.
Acknowledgements B.N.-Y. and H.-S.L. contributed equally to this work. This work was supported by Nano·Material Technology Development Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT and Future Planning (NRF-2014M3A7B4051749), the Center for Advanced Soft-Electronics funded by the Ministry
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Adv. Mater. 2015, DOI: 10.1002/adma.201500233
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Received: January 16, 2015 Revised: March 4, 2015 Published online:
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of Science, ICT and Future Planning as Global Frontier Project (2013M3A6A5073183), a New and Renewable Energy Program of the Korea Institute of Energy Technology Evaluation and Planning (KETEP) grant funded by the Korea Government Ministry of Trade, Industry and Energy (MTIE) (20133030000130), and the National Research Foundation of Korea Grant funded by the Korean Government (MSIP) (2013, University-Institute corporation program). The authors thank Prof. Hyunjung Kim in Department of Physics at Sogang University for help and discussion on the analysis of GIXD data.
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