Anisotropy of charge transport in a uniaxially aligned fused electron deficient polymer by solution shear coating
Mingfei Xiao[1], Boseok Kang[1][2][3], Seon Baek Lee[2], Luís M.A. Perdigão[4], Alex Luci[4], Satyaprasad P. Senanayak [1], Mark Nikolka[1], Martin Statz[1], Yutian Wu[1], Aditya Sadhanala[1], Sam Schott[1], Remington Carey[1], Qijing Wang[1], Mijung Lee[1][5], Chaewon Kim[5][6], Ada Onwubiko[7], Cameron Jellett[7], Hailiang Liao[8], Wan Yue[7][8], Kilwon Cho[2], Giovanni Costantini[4], Iain McCulloch[7][9], Henning Sirringhaus[1] 

[1] Mingfei Xiao, Prof. Boseok Kang, Prof. Satyaprasad P. Senanayak, Dr. Mark Nikolka, Martin Statz, Yutian Wu, Dr. Aditya Sadhanala, Dr. Sam Schott, Remington Carey, Dr. Qijing Wang, Prof. Mijung Lee, Prof. Henning Sirringhaus.  
Optoelectronics Group, Cavendish Laboratory, JJ Thomson Avenue, Cambridge, CB3 0HE, UK
E-mail: hs220@cam.ac.uk
[2] Prof. Boseok Kang, Seon Baek Lee, Prof. Kilwon Cho 
Department of Chemical Engineering, Pohang University of Science and Technology, Pohang 790-784, Korea.
[3] Prof. Boseok Kang
SKKU Advanced Institute of Nanotechnology and Department of Nano Engineering, Sungkyunkwan University, Suwon, 16419 South Korea
[4] Dr Luís M.A. Perdigão, Alex Luci, Prof. Giovanni Costantini
Department of Chemistry, University of Warwick, Coventry CV4 7AL, UK.
[5] Prof. Mijung Lee, Dr. Chaewon Kim
School of Advanced Materials Engineering, Kookmin University, Seoul 02707, Republic of Korea.
[6] Dr. Chaewon Kim
Institute of Functional Nano and Soft Materials, Soochow University, Suzhou 215123, China
[7] Ada Onwubiko, Cameron Jellett, Prof. Wan Yue, Prof. Iain McCulloch 
Department of Chemistry, Imperial College London, South Kensington, SW7 2AZ, UK.
[8] Hailiang Liao, Prof. Wan Yue
Key Laboratory for Polymeric Composite and Functional Materials of Ministry of Education, School of Materials Science and Engineering, Sun Yat-Sen University, Guangzhou 510275, China.
[9] Prof. Iain McCulloch
KSC, King Abdullah University of Science and Technology (KAUST), Thuwal, 23955-6900, Saudi Arabia.

 

Abstract

Precise control of the microstructure in organic semiconductors (OSC) is essential for developing high-performance organic electronic devices. Here we report a comprehensive charge transport characterization of two recently reported rigid rod conjugated polymers that do not contain single bonds in the main chain. We demonstrate that the molecular design of the polymer makes it possible to achieve an extended linear backbone structure which could be directly visualized by high-resolution scanning tunneling microscopy (STM). The rigid structure of the polymers allows the formation of thin films with uniaxially aligned polymer chains by using a simple one-step, solution-shear / bar coating technique. These aligned films show a high optical anisotropy with a dichroic ratio of up to a factor of 6. Transport measurements performed using top-gate bottom-contact field-effect transistors exhibit a high saturation electron mobility of 0.2 cm2 V1s1 along the alignment direction, which is more than six times higher than the value reported in our previous work. This work demonstrates that this new class of polymers is able to achieve mobility values comparable to state-of-the-art n-type polymers and identifies an effective processing strategy for this class of rigid-rod polymer system to optimize their charge transport properties.





Conjugated semiconducting polymers demonstrate unique electronic and optoelectronic characteristics significantly different from inorganic semiconductors[1][2][3][4][5][6][7]. For instance, their van der Waals bonded structure makes them appealing candidates for large-area, low-cost and highly uniform printable and flexible electronics, hitherto not achievable with traditional inorganic semiconductor technology[8]. An important application for conjugated polymers is the active layer in organic field-effect transistors (OFETs)[9]. Similar to conventional field-effect transistors, OFETs act as electronic switches and form the basic building block for organic electronics[10]. Moreover, OFETs are also being widely used as an effective platform for the study of the charge transport mechanisms in organic semiconductors[11]. Based on extensive research performed with a range of polythiophene derivatives, it has been generally observed that semi-crystalline conjugated polymers like poly(3-hexylthiophene) (P3HT)[12] and poly (2,5-bis (3-alkylthiophen-2-yl) thieno [3,2-b]-thiophene) (pBTTT)[13] facilitate charge delocalization through close  stacking, leading to moderately high carrier mobilities > 0.1 cm2 V1s1. In addition, strategies based on intelligent chemical design have resulted in donor-acceptor copolymers with highly planar backbone design[14][15][16][17], low energetic disorder[14][18], high microcrystalline order[1], high degree of in-plane alignment of the backbone[19] and long persistence length[3]. These donor-acceptor copolymers have been shown to exhibit field-effect mobilities (µFET) > 1 cm2 V1s1 in several cases, though these polymers tend to be less crystalline than pBTTT. It has been argued that a high degree of crystallinity is not required in these systems as long as there exist at least small crystalline domains, aggregate regions or close crossing points in which interchain charge transfer can occur and the energetic disorder associated with conformational variations along the backbone is sufficiently small such that intrachain transport in between these close crossing points is sufficiently fast. [10] 
    In this work we present a detailed charge transport study of a recently developed and unusual family of conjugated polymers that differs substantially from the design motives followed by conventional polythiophene or donor-acceptor copolymers. These rigid-rod polymers are produced by a catalyst-free aldol condensation reaction between bis-isatin and bis-oxindole[20]. One of their unique features is that they do not contain any single bonds along the polymer backbone, but only fused ring structures linked together by double bonds. This should minimize torsional defects along the polymer backbone and favor a highly conjugated, planar backbone structure. The other unique feature is that they are not donor-acceptor systems, but comprise only electron accepting groups along the polymer backbone. This induces a relatively deep electron affinity and makes them exhibit stable electron transport with very good environmental stability of up to 300 hours[20]. This is remarkable since the development of n-channel conjugated polymers with long term air stability still lags behind their p-type counterparts[21], due to electron transport being prone to redox reactions with water and oxygen[22][23][24]. However, the highest electron mobilities demonstrated in these systems so far have only reached somewhat disappointing low values on the order of 0.03 cm2 V1s1[20], which is significantly lower than what is achievable with other state-of-the-art n-type polymers[21]. In this work, we visualize the unique fused backbone structure of these promising rigid-rod polymers with monomeric precision by state-of-the-art STM imaging[25] and demonstrate a strategy to optimize their electron transport properties by uniaxial chain alignment. Our aim is to investigate whether these rigid-rod polymers are capable of exhibiting similarly high electron mobilities as donor-acceptor systems. 
Several techniques have been proved to be effective to induce uniaxial chain alignment and charge transport anisotropy, such as solution shear[26], brush printing[27], and bar coating[28], etc. To deposit the films we use solution shear coating, a simple technique for controlled film deposition that is able to induced uniaxial chain alignment: The unidirectional motion of the solution meniscus relative to the substrate produces a shear force on the fluid that promotes polymer chain alignment during the film drying process.[26] Intuitively, it can be expected that this technique would be particularly favorable if the polymer has a rigid backbone structure allowing the realization of microstructures comparable to aligned CNT networks [27]. Our fused polymer system should exhibit the ideal rigid backbone structure needed for polymer chain alignment and achieving a high degree of charge transport anisotropy that is expected to improve carrier mobilities along the chain alignment direction.
    The polymers used in this study are shown in Figure1 a&b. The two polymers comprise two core naphthalene units linked together, we therefore refer to them here as NN1 and NN2 (in Ref. 20 these polymers are simply refereed to as P3 and P4, respectively). They have the same fused backbone structure but vary in the alkyl side chains attached. In order to visualize their backbone structure directly, we performed scanning tunneling microscopy (STM) measurements of polymer sub-monolayer films fabricated by vacuum electrospray deposition (ESD) on atomically clean and flat single crystal surfaces[25]. While the solubility of NN2 was too low for a successful electrospray deposition, the extra solubility imparted by the fully branched side-chains of NN1 was enough for this latter polymer to be deposited by ESD on Au(111) and Ag(111) substrates and for its backbone structure to be imaged by STM (Figure 1 c&d). The surface appears to be covered by long strands displaying a brighter central core (the polymer backbone) flanked on each side by a slightly dimmer region (the side-chains). The flat-laying individual polymers are randomly oriented and dispersed on the surface and only occasionally arrange in small locally ordered patches where a few of them lay parallel to each other. There was no noticeable difference between the results on Au(111) and Ag(111), implying that the herringbone reconstruction of Au(111)[28] has only a negligible tempting effect on NN1. A remarkable difference with other polymers analyzed so far by ESD-STM[25][29] is the absence of interdigitation of the solubilizing side-chains of NN1 which is probably prevented by their high attachment density[30]. The NN1 chains exhibit remarkable straight conformations which extend for several tens of nanometers occasionally interrupted by a few and rather sharp bends (Figure1 c&d). These kinks are probably the result of cis-defects in the double bond linkage between monomers. The overall straight-chain conformation observed for NN1 is believed to reflect the intrinsic rigid-rod backbone nature of both polymers (since they share the same backbone structure) originating from their unique double-bonded fused backbones. 
The persistence length lp of a polymer chain is a measure of how far a chain extends in a straight direction and reflects the stiffness of the polymer.[31] Recently, it has been noted that large lp values of conjugated polymers might be beneficial for their applications in optoelectronic devices, and several high performance conjugated polymers used in OFETs and photovoltaics demonstrate significantly higher persistence lengths compared with the widely investigated P3HT[32]. Stiffer chains are likely to make better use of the fast charge transport along the polymer backbone, facilitate on-chain charge delocalization, and provide short and straight paths for charge transport within polymer film[33]. While the persistence length is typically evaluated in three dimensions (3D), an effective “2-dimensional (2D) persistence length (lp2D)” can be determined based on our STM measurements of the NN1 polymer, resulting in a value of lp2D  19 nm (see the Supplementary Information, SI, for more details on obtaining this value). A systematic study of the 2D persistence length is currently missing for other surface-adsorbed high-mobility conjugated polymers; however, preliminary evaluations of this parameter we did for P3HT and PBTTT seem to indicate significantly smaller values of lp2D, closer to 1-5 nm. Since backbone deflection and torsion angles influence the polymer’s persistence length[34], these results can be attributed to the fused backbones of the NN1 and NN2 polymers and can be traced back to the larger energetic barrier associated with 180° rotation of double bonds (~25 kcal/mol) compared with single bonds (~ 2-8 kcal/mol)[20]. When extrapolated from 2D to 3D, our STM-based results suggest that a long persistence length may also be achievable in the bulk of thin films fabricated with NN1 and NN2. As a consequence, they motivated us to explore the usage of alignment techniques to fully exploit the rigid polymer backbone structure, with the aim of enhancing charge transport along the polymer chain direction.
In order to align the polymer films, we used a solution shearing technique with a PTFE rod working as the coating blade. After a series of optimization trials, we adopted a process where a temperature-controlled glass substrate (140 °C) was moved at a fixed speed of 146 μm/s underneath a PTFE coating blade. By using this technique, it was possible to restrict the solvent evaporation to the edge of the meniscus (schematically shown in Figure S2) and hopefully align polymer chains uniaxially along the shearing direction. It is worth noting that only 10 μL of the solution are required to fully cover a 16mm × 16mm substrate by solution shearing, compared to 100 μL solution required to obtain uniform film coverage in the spin coating process; this indicates the efficacy of the solution shearing technique for large-area electronics applications. Detailed microstructure characterization was performed on the solution-sheared and spin-coated films of both NN1 and NN2 polymers. Atomic Force Microscope (AFM) measurements performed on spin-coated NN2 films indicated a disordered surface morphology with small grains of a few nm in size (Figure 2a). In contrast, upon solution shearing of the NN2 polymer, it was possible to observe an anisotropic microstructure with well defined, aligned fibrils each with a width on the order of 100 nm. (Figure 2b, Figure S3). In addition, the root mean square roughness of the film increased to 16.9 nm compared with 5.3 nm for spin-coated films. However, under similar conditions, we were unable to observe comparable aligned surface microstructure in the AFM images of solution-sheared films of NN1 (Figure S4), and the AFM images of solution-sheared films of NN1 highly resembles their spin-coated counterparts (Figure S4). This behavior can possibly be attributed to the bulkier structure and the higher attachment density of the side-chains of NN1, which is likely to result in a relatively weaker interchain interaction when compared with NN2, due to a sterically hindered π-π stacking and a low degree of side-chain interdigitation (as observed in the STM images). This is reflected in a lower degree of crystallinity of spin-coated films of NN1 compared with spin-coated films of NN2 from previous grazing incidence wide-angle X-ray scattering (GIWAXS) measurement results[20], and also in a much lower OFET mobility of such NN1 films (around 0.001 cm2 V1s1) when compared with spin-coated films of NN2 (around 0.03 cm2 V1s1)[20].  
One of the simplest ways to quantify the degree of alignment of the films is to determine the dichroic ratio from optical absorption measurements acquired in directions parallel and perpendicular to the polymer alignment direction. In NN2 films it was possible to observe a dichroic ratio close to 6 for the main absorption peak around 2.3 eV[20]. It is important to note that the optical transition dipole moment of NN2 is likely to have a non-zero component perpendicular to the polymer backbone direction, which implies that the dichroic ratio can only provide a lower bound for the degree of alignment. In contrast, for NN1 no absorption anisotropy could be observed, coherently with the absence of any alignment feature in the AFM surface morphology of these films. In the following we therefore focus the discussion only on aligned films of NN2.  
The band-tail region of the absorption spectrum reveals important information, on the amount of energetic disorder. Energetic disorder which creates trap states in the bandgap would broaden the absorption onset and create an exponential sub-bandgap tail (Urbach tail)[35], the characteristic width of which is defined as the Urbach energy (EU) and can be interpreted as a measure of the width of the joint density of states (DOS) [36]. We performed photothermal deflection spectroscopy (PDS) measurements on the aligned NN2 samples to obtain the Urbach energies parallel and perpendicular to the chain alignment direction. For the aligned film, it is possible to extract a value of EU = 43 meV for measurement parallel to the aligned direction and of EU = 51 meV for measurement perpendicular to the alignment direction (Figure 2c), indicating a decrease in the energetic disorder along the alignment direction. This presumably implies that the chains or chain segments that are aligned along the shearing direction within the aligned fibrils exhibit lower conformational/energetic disorder than those chain segments or chains that are in regions of the film that are poorly aligned, such as grain boundaries or more disordered regions in between the aligned fibrils. The values of EU are not particularly low compared to some donor-acceptor copolymer systems, in which values less than 30 meV can be obtained[14]. This may at first seem surprising given the rigid rod nature of the NN2 polymer, which should be less susceptible to variations in backbone torsion angles than polymers that comprise single bonds along the backbone. However, it is also important to consider that NN2 is not a donor-acceptor copolymer; adjacent conjugated units along the backbone interact with each other more strongly than in a donor-acceptor system. This is expected to result in a larger intrachain bandwidth, which also tends to induce broader band tails than in a relatively narrow band donor-acceptor system[37].  
To further characterize the effect of alignment we performed detailed structural analysis of the NN2 polymers through GIWAXS measurements on both spin-coated and sheared samples. Lattice spacing parameters that are extracted from these measurements are summarized in Table S1. The π–π stacking distance and lamella stacking distance are calculated from the related peak spacing qhkl using the equation dhkl = 2π/qhkl, with qhkl expressed in nm-1, while the associated coherence lengths were estimated from the radial full width at half maximum (FWHM), ΔFWHM, via Scherrer’s equation[38][39]. For the spin-coated sample, a bimodal morphology with both face-on and edge-on oriented crystallites is observed. A π–π stacking peak corresponding to the face-on crystallites is measured in the out of plane direction, with a π–π stacking distance of 0.38 nm, which is typical of polymer systems [10]. For the edge-on crystallites, two orders of alkyl-stacking (h00) peaks are observed along the out-of-plane direction. In contrast, aligned samples show preferential edge-on orientation, with significant scattering anisotropy between the incident X-ray beam being along the direction parallel and perpendicular to the chain alignment direction. Along the chain alignment direction, the (h00) peaks as well as the weak π–π stacking diffraction intensity exhibit an arc texture indicating a distribution of grain orientations around the surface normal, while, perpendicular to the chain alignment direction, the arc disappears. Considering the alignment process during which lamella crystals form, the arc appearance could be related to the perturbation of the growth of the lamella, resulting in a crystalline disorder which is frozen-in after film drying. Based on this interpretation, the arc signal is expected to be more intense when the diffraction pattern is measured along the alignment direction, as reported here. For both directions, alkyl stacking (h00) peaks up to three orders could be observed, implying increased crystallinity compared with spin-coated samples. The lamella stacking distance in the out-of-plane direction is reduced compared with spin-coated films, showing that the shearing process impacts the molecular packing within the unit cell. The π–π stacking distance stays similar upon alignment, indicating there is no significant lattice strain along this direction or other non-equilibrium effects that were observed for sheared TIPS-pentacene[43]. In conclusion, the analysis of GIWAXS patterns further confirms the structural order improvement and strong structural anisotropy induced by this solution shearing process.
Transport measurements on the aligned NN2 polymer films were performed using a top-gate, bottom-contact (TGBC) device architecture (details in the methods section). To distinctively measure the influence of the polymer alignment on the transistor properties, we carefully fabricated two transistor channels oriented at 90 degrees with respect to each other on the same substrate. This allows us to compare transistor characteristics for current flow parallel and perpendicular to the chain alignment direction. The devices exhibit ambipolar transport characteristics with a current modulation of > 104 and evidence for hole injection and hole transport when VD=50V and VG<10V (Figure 4a,b). However, higher currents are obtained in the electron regime and we focus on discussion of the electron transport regime here. Output curves with well-defined linear and saturation regimes and some evidence for ambipolar transport at VG = 0V (Figure S5) were obtained for transistors fabricated with polymer chains aligned parallel and perpendicular to the channel/current flow direction. The output curves indicate the existence of a finite contact resistance at low gate voltage VG (Figure S5). This contact resistance possibly originates from the mismatch of energy levels between the gold electrode and the conjugated polymer film, the injection barrier caused by insulating alkyl chains of edge-on crystallites in direct contact to gold electrodes, or from the roughness of the semiconductor thin film, thus indicating scope for further improvement in device architecture and also the shearing procedure to further enhance the transport properties. A maximum channel current of 80 µA at drain voltage VD = 50 V and VG = 60 V was obtained for devices with polymer chains aligned parallel to the channel direction. In comparison, devices with polymer chains aligned perpendicular to the channel exhibited a much lower current of 18 µA under the same conditions. Correspondingly, saturation mobility as high as 0.2 cm2 V1s1 was obtained for the device with the channel along the alignment direction, which decreases to 0.03 cm2 V1s1 for the device with polymer chains aligned perpendicular to the channel direction. The mobility anisotropy was thus obtained to be around 6. 
In order to understand the transport mechanism of OFETs with aligned NN2 films, temperature-dependent transfer characteristics were measured over a range of 200 K to 350 K. In order reduce the influence of contact resistance effects, we estimated µFET in the saturation regime of operation with standard 20 µm channel length transistors (details of device architecture can be found in the experimental section). It is important to note that for the whole temperature range of the measurements, the magnitude of µFET along the chain alignment direction is higher than the value of µFET evaluated for devices with channels perpendicular to the chain alignment direction (Figure S6), proving the effectiveness of our shearing technique to improve charge carrier mobility along the alignment direction. Furthermore, the temperature-dependent mobility comparison reveals different charge transport mechanisms parallel and perpendicular to the chain alignment direction. OFETs fabricated with the polymer alignment direction perpendicular to the channel direction, exhibit thermally activated charge transport within the whole temperature range measured. The activation energy was extracted to be 81 meV for charge transport along this direction (Figure 4d). However, OFETs fabricated with the current flow in the channel parallel to the polymer chain alignment direction exhibit three distinct temperature regions (Figure 4d): in the range from 200 K to 240 K, charge transport is thermally activated with an activation energy of 69 meV; for the temperature region of 250-300 K, the activation energy increases to 110 meV. A very similar behavior has been reported for aligned DPP-BTz OFETs along the chain alignment direction[29]. However, when the temperature is increased to the 300-350 K range, the transfer characteristics overlap with each other (Figure S7), indicating that µFET is independent of temperature. In this regime, the transfer curves are highly reproducible, which eliminates device degradation as a possible explanation of this temperature independence. At the same time, taking into consideration the high activation energy in the 250-300 K temperature regime, the relatively high Urbach energy of 43 meV, and the mobility value below 1 cm2 V1s1, it is highly unlikely that a disorder-free transport regime has been achieved in the 300-350 K range. However, a possible explanation could be that in this high-temperature regime the mobility for transport along the alignment direction gets a significant boost from charges being able to access delocalized electronic states along the polymer backbone. The density of states of conjugated polymer is expected to comprise localized states near the band edges and states that are more delocalized along the backbone towards the center of the band[40]. If these delocalized states become thermally accessible, they are likely to benefit strongly the charge transport, particularly in devices with the channel aligned along the chain alignment direction. In this way one might qualitatively explain also the observed increase in the activation energy with increasing temperature and the observation of a temperature independent regime above room temperature. This interesting charge transport regime warrants further detailed study.  
In conclusion, we have performed detailed structural and microscopic characterization on a new class of rigid, fused-ring aldol condensation polymers without single bonds along the polymer backbone. In STM it was possible to directly visualize single polymer chains with high 2D persistence lengths and long straight chain segments. By taking advantage of this exceptional rigidity of the backbone, it was possible to uniaxially align one of the polymers by using a simple solution shear / bar coating technique and demonstrate significantly improved structural order as observed by GIWAXS. Consequently, upon alignment, the charge mobility increases by six times compared to the spin-coated films, reaching a value of 0.2 cm2 V1s1. These mobility values are within a factor of 5 compared to some state-of-the-art n-type polymers[21], which are encouraging and this class of polymers warrants further study. We see significant potential for further improvements in mobility by optimizing polymer chain structure or alignment technique[44], for example, in donor-acceptor copolymers synthesized by aldol condensation, as these would retain the rigid rod nature with no single bonds along the backbone but are likely to exhibit smaller bandwidths and lower energetic disorder. Combining future better alignment technique and improved molecular design, even higher charge transport property with further improved anisotropy would be possible to achieve with these newly emerged, completely fused conjugated polymers, approaching the anisotropy and performance demonstrated previously by aligned P (NDI2OD-T2) system[28].   

Acknowledgement
The authors gratefully acknowledge support from the European Research Council (ERC, Synergy Grant 610115) and the Engineering and Physical Sciences Research Council (EPSRC, programme grant EP/M005143/1). 


Figures:
 



Figure 1 Chemical structure of (a) NN1 and (b) NN2 polymers. STM images showing of the NN1 polymer adsorbed on (c) Ag(111) and (d) Au(111). (c) Overview image acquired with tunneling parameters Vbias = 1.5 V and I = 30 pA. (d) Higher magnification image demonstrating that the polymers are mostly straight with occasional sharp bends; a few examples of these bends are indicated by the orange segments superposed on the STM image. Image obtained with tunneling parameters Vbias = 2.0 V and I = 30 pA.
 

 
Figure 2 (a) AFM image of spin-coated NN2 polymer film. (b) AFM image of aligned NN2 polymer film. (c) Polarized PDS spectrum of the aligned NN2 film with the incident light polarized along the directions parallel and perpendicular to the chain alignment direction.



Figure 3 (a) 2D GIWAXS patterns of solution sheared and spin-coated NN2 films. For the solution sheared sample the incident X-ray beam was aligned parallel and perpendicular to the chain alignment direction (left and center panels). (b) Corresponding GIWAXS line profiles (out-of-plane direction). (c) GIWAXS line profiles (in-plane direction).





Figure 4 Transfer curves measured on top-gate, bottom-contact FETs (L = 20 µm, W = 1mm) fabricated from the aligned NN2 film with charge transport (a) parallel to the chain alignment direction, and (b) perpendicular to the chain alignment direction. (c) Comparison of the saturation mobility between devices with charge transport direction parallel and perpendicular to chain alignment direction. (d) Temperature-dependent saturation mobility for charge transport along with the two directions.





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