The Influence of Backbone Fluorination on the Dielectric Constant of Conjugated Polythiophenes

The ability to modify or enhance the dielectric constant of semiconducting polymers can prove valuable for a range of optoelectronic and microelectronic applications. In the case of organic photovoltaics, increasing the dielectric constant of the active layer has often been suggested as a method to control charge generation, recombination dynamics, and ultimately, the power conversion efficiencies. In this contribution, the impact that the degree and pattern of fluorination has on the dielectric constant of poly(3‐octylthiophene) (P3OT), a more soluble analogue of the widely studied conjugated material poly(3‐hexylthiophene), is explored. P3OT and its backbone‐fluorinated analogue, F‐P3OT, are compared along with a block and alternating copolymer version of these materials. It is found that the dielectric constant of the polymer thin films increases as the degree of backbone fluorination increases, in a trend consistent with density functional theory calculations of the dipole moment.


Introduction
Manipulation of the dielectric constant of conjugated materials through design is an attractive prospect for microelectronic and optoelectronic applications. In the microelectronic arena, sev eral purposes are readily apparent. Enhancing the permittivity of organic dielectric layers in capacitors can lead to enhanced different trend. Indeed, manipulation of monomer polarity and size, along with chemical crosslinking can lead to very high ε r values up to 70, as is the case for P(VDFTrFECTFE). [18,19] Introducing highly polar groups to increase ε r in conjugated polymers has been an approach favored by some research groups. [7,10,13,[20][21][22] Such efforts have mainly focused on the introduction of the polar groups within the side chains to avoid disturbing the conjugated backbone, and although ε r increased, the subsequent performance in an OPV was variable, quite likely due to the influence of such polar groups on blend micro structure. An alternative approach has been the blending of higher dielectric additives with the active layer. A range of such additives have been explored. For example, Torabi et al. investi gated bilayer devices in which the donor was blended with cam phoric anhydride, resulting in an enhancement of permittivity from 4.5 to 10.8. [13,23] Hyperbranching of conjugated backbones has also been suggested to increase the dielectric constant through longrange polaron delocalisation. [24][25][26] Fluorination in conjugated polymer backbones has been explored as a route to increase dielectric constant. For example, difluorination of a partially conjugated thermoplastic polymer led to a 10% increase in the dielectric constant. [27] Similarly, two separate studies have shown that partial backbone fluorination of quinoxalinebased conjugated polymers led to substantial increases in permittivity, and in both cases, fluorination drove increased OPV performance. [4,11] In addition to these efforts, backbone fluorination of conjugated polymers is also moti vated by factors other than dielectric constant improvements. Stabilization of both the highestoccupied molecular orbital and the lowestunoccupied molecular orbital through the addition of electron withdrawing fluorine atoms is a common approach to increase the V oc . [28][29][30] Reductions in recombination (and the subsequent increase in J sc ) upon fluorination have also been widely reported, these, however, are often explained by more favorable microstructure and domain compositions of the active layers, or increased polarization of the exciton, which reduces bimolecular and geminate recombination respectively. [28,[31][32][33] With the high level of interest in backbone fluorination for organic semiconductors and conjugated polymers, we decided to explore the effects of such fluorination on the dielectric constant of the benchmark materials that are poly(3alkylth iophenes) (P3AT). We recently reported the synthesis and characterization of poly(3alkyl4fluorothiophenes) (FP3ATs) and observed that FP3HT, the fluorinated analogue of P3HT exhibits very low solubility due to aggregation, making its pro cessing difficult. [34] We, therefore, opted to investigate the effect of fluorination on the permittivity of the more soluble octyl analogue, FP3OT. In this contribution, we probe the effect of backbone fluorination by comparing P3OT with FP3OT, as well as possible effects of asymmetric fluorination through alternating and block copolymers of these thiophene deriva tives. We observe that fluorination of the P3OT polymer leads to an increased ε r from 2.7 to 4.8, and that the substitution pat tern has some impact on the extent of this increase.

Synthesis
For the preparation of P3OTaltFP3OT, we decided to focus upon the Grignardmetathesis (GRIM) [35] polymerization of the dimer 6 for better control of the regioregularity, and to keep residual catalyst and impurities as consistent as possible across the samples investigated. We note that a related material, P3HT altFP3HT, has recently been reported by a similar polymeriza tion. [36] Exploiting the exclusive selectivity for the magnesium/ bromine exchange of 3, [34] monobrominated compound 4 was produced, which was coupled with 2 to yield the headtotail bithiophene 5 in moderate yield after purification by reverse phase chromatography (Scheme 1). Subsequent bromination afforded the monomer 6, which was polymerized using GRIM conditions, similar to the other polymers (Scheme 2). P3OTalt FP3OT was purified by Soxhlet extraction, washing with meth anol, acetone, hexane, and finally chloroform, before extracting with chlorobenzene. The resulting polymer only exhibited a single resonance in the 19 F NMR, suggestive of very high back bone regioregularity (see Supporting Information). No peak attributable to a tailtotail defect was observed in our case. [36] The homopolymers P3OT and FP3OT were synthesized via GRIM polymerization from the activated monomers 8 and 9, as previously reported (Scheme 2). [34] The polymers were purified by Soxhlet extraction, washing sequentially with methanol, ace tone, and hexane (and chloroform for FP3OT). P3OT was then extracted using chloroform, and FP3OT with chlorobenzene. P3OTbFP3OT was also synthesized by GRIM polymerization, using the method we recently reported, with the more soluble P3OT block grown first from the activated monomer 8 fol lowed by the addition of activated monomer 9 to the P3OT mac roinitiator. [37] The polymer was purified by Soxhlet extraction as with P3OT. However, a further wash with dichloromethane was performed before extraction with chloroform in order to remove traces of P3OT homopolymer. The presence of P3OT homopoly mer within P3OTbFP3OT is typically characterized by melt and crystallization peaks (at 198 and 155 °C, respectively) observed in the differential scanning calorimetry thermograms (see Figure S1, Supporting Information). [37] After purification, 1 H NMR analysis showed that the two blocks were of approxi mately equal length. 19 F NMR only afforded a single peak, again suggestive of very high degrees of backbone regioregularity. In agreement with our earlier studies, [38] all fluorinated poly mers were less soluble than their nonfluorinated analogues and required processing from hot chlorinated solvents like chlo robenzene or 1,2dichlorobenzene. The molecular weights, as measured by gel permeation chromatography (GPC) in either hot chlorobenzene (P3OTbFP3OT and P3OTaltFP3OT or hot 1,2,4trichlorobenzene (FP3OT), of all fluorinated poly mers were similar ( Table 1).

Thin-Film Properties
In preparation for the subsequent device studies, the film forming properties and basic characterization from all four materials were examined. Figure 1a shows the absorption spectra, with the absorption coefficients extracted from trans mission spectra using the Beer-Lambert expression. All spectra are qualitatively similar, exhibiting a main absorption peak around 530-550 nm, with a clear vibronic progression apparent toward lower energies-a behavior often ascribed to aggrega tion of the conjugated backbone. The most striking conse quence of backbone fluorination is the resulting blue shift of both absorption [34] and emission spectra shown in Figure 1b.
Regarding the absorption coefficient, the P3OT thin film shows the highest value (≈8.3 × 10 4 cm −1 ). While in compar ison, the peak absorption for the fully fluorinated counterpart (FP3OT) is almost half of this value with the alternating and block copolymers taking on intermediate values. For latter two cases, the P3OTaltFP3OT and P3OTbFP3OT differ predom inantly in the strength of the vibronic structure, with the two main vibronic peaks taking almost equal value in the instance of the alternating compound (P3OTaltFP3OT), suggestive of increased interchain coupling and greater backbone planaritytransitioning from H to Jlike aggregation. [39] Interestingly, the peak positions in the absorption profile of P3OTbFP3OT appear dominated by FP3OT contributions despite exhibiting the lower absorption coefficient. This is apparent when com paring the differences between the normalized absorbance spectra of P3OTbFP3OT with that obtained from a linear combination of the two constituent polymers, P3OT and FP3OT ( Figure S3, Supporting Information). In common to other diblock polymers, such as thiopheneselenophenebased systems, there is a reasonable match between the spectra and the linear combination of the two constituent polymers, [40,41] although similar to the 3:1 diblock [37] polymer it appears that the enhanced backbone planarity and elongated chain of the fluorinated block is partially inhibiting the ordering of the non fluorinated block in this 1:1 copolymer.
The decreasing trend in peak absorption value with increasing fluorination (Figure 1) may be expected to have a similar impact on the refractive index properties, following the usual Kramers-Kronig arguments. [42] As a preliminary inves tigation, we have modeled the measured specular reflection recorded from the thinfilm samples (Figure 2), concentrating on the spectral region beyond the main absorption band (i.e., >700 nm). Employing a Cauchytype description for the refrac tive indices of each material (Table S1, Supporting Information) does indeed suggest that the refractive index reduces upon increasing fluorination (inset of Figure 2).
Further influence to the extent of backbone fluorination on optoelectronic properties is evident from the measured work functions from polymer thinfilm samples, listed in Table 2. Consistent with our previous report on the homopolymers, a stabilization of around 0.3 eV was observed upon fluorina tion. [34] P3OTaltFP3OT exhibits an intermediate work func tion, close to halfway between the two homopolymers. Similar behavior was reported for an alternating thiophene-thiazole copolymer and an allthiophene system where 3hexylthiophene and 3(2ethylhexyl)thiophene were randomly copolymer ized. [44,45] Conversely, P3OTbFP3OT has a work function closer to that of the readily ionized P3OT. The phenomena of block copolymers displaying the work function of the most easily ionized block have previously been reported for polythio phene derivatives. [46,47] We note, however, that in the case of P3OTbFP3OT, the work function is still 0.1 eV higher than P3OT.  Measured by GPC against polystyrene standards in chlorobenzene at 80 °C except for F-P3OT, which was measured in 1,2,4-trichlorobenzene at 140 °C.

Density Functional Theory (DFT) Calculations
DFT calculations on hexameric thiophene analogues were employed to predict the effect of fluorination on the dipole moment of the repeat units and, therefore, potentially on the dielectric constant of the polymer. Our calculations are based on the favored transoid conformation of the thiophene-thio phene bond since recent studies demonstrate that this is lower in energy compared with the cisoid. [34,36] We observe that fluorination of the backbone leads to a large increase in the dipole moment (Figure 3; Figure S4, Supporting Information). Indeed, P3OT has a small dipole moment of 0.37 D along the backbone, but upon fluorina tion, this is increased to 7.72 D and remains largely directed along the backbone. Alternating the fluorination, while also increasing the dipole moment to 4.55 D, leads to a signifi cant proportion of the dipole pointing away from the back bone. Translating these calculated values into a macroscopic property such as the dielectric constant is awkward since a number of factors come into play, such as the orientation of the repeat unit dipoles, their nanoscale ordering, and the projected alignment to the electric field. Nevertheless, a clear trend emerges that suggests increasing the backbone fluorination should lead to an increased dielectric constant, on the understanding that the materials all have a similar microstructure.

Dielectric Constant
Capacitance-voltage (C-V) characteristics of metal-insulatorsemiconductor diodes were measured to extract the lowfre quency dielectric constant of the materials. [43,[48][49][50] The devices, based on a ptype semiconductor (Si ++ ), include the thinfilm polymer layers as part of the insulating region in addition to a native SiO 2 layer. In the usual nomenclature, we refer to the structure as a metal-oxidesemiconductor (MOS) capacitor structure (Figure 4; Figure S5, Supporting Information). The extracted dielectric constants (ε r ) of the polymers show that, compared with the reference material P3OT, increasing back bone fluorination increases ε r for all cases (Figure 4b). In particular, full fluorination leads to a near doubling of the die lectric constant, i.e., from 2.71 to 4.82 for P3OT and FP3OT, respectively. We note the value of ε r for P3OT, obtained from our MOS structures, was in good agreement with a value of 2.84 obtained from a metal-insulator-metal (MIM) architecture   ( Figure S6, Supporting Information). Both, however, are some what lower than the value of 3.24 previously reported for P3OT by ellipsometry. [50] The increase in ε r upon fluorination reflects the trend observed in quinoxalinebased conjugated polymers and further suggests the dielectric constant of conjugated poly mers respond differently to fluorination when compared with nonconjugated poly mers. [4,11] At present, it is unclear from our results whether this difference can be assigned solely to the increased dipole moment, as predicted by DFT calculations, or if nanoscale ordering, and other microstructure effects (e.g., crystallite dipoles) are at play, as found for P(VDFTrFECFE) and related high dielectric polymers. [19] In the case of these polyethylenebased fluorinated polymers, the anisotropy in crystalline domains resulting from the introduction of different comonomers is exploited. FP3OT has indeed been reported to frustrate crystallization and has shorter coherence lengths in both the lamellar and π-π stacking directions than P3OT, which may indicate some differences in crystalline domains. [34] The block copolymer, P3OTbFP3OT, exhibits a dielec tric constant that is close to the average of ε P3OT and ε FP3OT , at 3.97, whereas the alternating copolymer, P3OTaltFP3OT, recorded a dielectric constant approaching the fully fluorinated homopolymer. The latter case appears in line with previous reports, involving polyphenylene vinylene (PPV) copolymers with alkyl and ethylene glycol sidechains, where the alternating copolymers take on intermediate ε values of the two homopoly mers. [20] The contrasting results obtained from the alternating and block copolymers confirm that the pattern of fluorination along a polymer chain has an influence. Despite the earlier per ceived difficulties (and reservations) with directly translating DFT results, it is gratifying to note that of the two copolymer materials, the predicted dipole moment for the alternating copolymer was expected to be larger than the block copolymer, which is indeed the case with the dielectric constant. In fact, the overall trend in predicted dipole moment faithfully mirrors the trend observed with the measured dielectric constant ( Table 2).

Conclusions
The ability to modify the electrical properties of semiconducting polymers, such as the dielectric constant, through chemical modification offers great promise and opportunities in many application areas. In this contribution, we have reported the synthesis of two novel copolymers based on partially fluorinated P3OT to probe the effect fluorination has on the dielectric con stant, as well as the influence the pattern of fluorination can have on such properties. Through measurements of the lowfre quency dielectric constant (ε) in MOS devices, we find that fluor ination of only half of the thiophene units leads to a significant increase in the dielectric constant, i.e., from 2.7 for P3OT to 3.97 in a block copolymer configuration, and 4.3 for an alternating configuration. Full fluorination leads to an even larger increase to 4.8. The fact that both the alternating and block copolymers show substantial increases in ε compared with P3OT confirm that backbone fluorination is a useful tool for increasing the dielectric constant of semiconducting polymers. Deconvolving the factors that lie behind such an increase is difficult-never theless, we find a strong qualitative link between the calculated dipole moments and the measured dielectric constant-which correctly reproduce the increasing order. Inevitably the details and differences of microstructure undoubtedly play a significant  role, as indeed highlighted from the absorption spectra and the degree of intrachain coupling between the two copolymers. Fur ther studies, perhaps specifically targeting the microstructure of these poly mers and the corresponding blends, would prove useful to elucidate the origin of the increase in ε r . Finally, of par ticular interest would be the subsequent performance of these polymers in photo voltaic devices. [51]

Experimental Section
Materials: Reagents and chemicals were purchased from commercial sources such as Sigma-Aldrich and Acros etc. unless otherwise stated. 2,5-dibromo-3-fluoro-4-octylthiophene (3) was synthesized according to previously reported procedures. [34] All reactions were carried out under argon using solvents and reagents as commercially supplied, unless otherwise stated.
Synthesis of 2-Tributylstannyl-4-Octylthiophene (2): In a dry 2-neck flask, 3-octylthiophene (1.80 g, 9.17 mmol) was dissolved in dry tetrahydrofuran (THF, 20 mL) THF (20 mL) and cooled to −78 °C. Lithium diisopropylamide (5.05 mL, 2.0 m in THF/heptane/ethylbenzene) was then added dropwise, and the resulting solution was stirred at −78 °C for 2 h, after which tributyltin chloride (4.48 g, 13.76 mmol) was slowly added. The reaction mixture was further stirred at −78 °C for 2 h, before being warmed to room temperature, and finally poured into water (100 mL) and extracted with hexane (3 × 50 mL). The organic extracts were washed with water (3 × 100 mL), and acetonitrile (4 × 50 mL), and dried over sodium sulfite. The solvent was removed in vacuo to yield the crude product (2) as pale yellow oil, which was used for the next step without further purification. 1

Synthesis of Poly(3-Octylthiophene-2,5-Diyl) (P3OT):
Grignard solution freshly prepared from 2,5-dibromo-3-octylthiophene (29.1 mL, 0.18 m solution in THF) was added via cannula to a dry 100 mL 2-neck flask under Ar containing dichloro(1,3-bis(diphenylphosphino)propane)nickel (28.9 mg, 1 mol%) and fitted with a reflux condenser. The reaction was then refluxed for 12 h, before being poured into methanol (200 mL) acidified with a few drops of concentrated HCl. The precipitate was filtered through a cellulose thimble, and the solid purified by Soxhlet extraction with methanol, acetone, hexane (in each case until the extracting solvent was colorless), and finally extracted with chloroform, before being precipitated into methanol and filtered. The solid was dried under vacuum to give P3OT (740 mg, 72%). M n 26 kDa, M w 33 kDa. 1  In a dry 20 mL microwave vial, 2,5-dibromo-3-fluoro-4-octylthiophene (1.00 g, 2.69 mmol) was dissolved in dry THF (17 mL), and to the stirred solution, isopropylmagnesium chloride lithium chloride complex (1.97 mL, 1.3 m in THF) was added dropwise. After 30 min, a suspension of dichloro(1,3bis(diphenylphosphino)propane)nickel (7.3 mg, 0.5 mol%) in dry THF (0.7 mL) was added to the Grignard monomer solution, and the resulting solution was stirred at 70 °C for 12 h, before being poured into methanol (200 mL) acidified with a few drops of concentrated HCl. The precipitate was filtered through a glass thimble, and the solid purified by Soxhlet extraction with methanol, acetone, hexane, chloroform (in each case until the extracting solvent was colorless), and finally extracted with chlorobenzene, before being precipitated into methanol and filtered. The solid was dried under vacuum to give F-P3OT (420 mg, 73%).

Synthesis of Poly(3-Octylthiophene-Block-3-Fluoro-4-Octylthiophene) (P3OT-b-F-P3OT):
In a dry 2-5 mL microwave vial charged with dichloro(1,3-bis(diphenylphosphino)propane)nickel (2.7 mg, 0.5 mol%) was added Grignard solution freshly prepared from 2,5-dibromo-3octylthiophene (2.67 mL, 0.28 m in THF), and the reaction mixture was stirred at 40 °C for 1 h, after which Grignard solution freshly prepared from 2,5-dibromo-3-fluoro-4-octylthiophene (0.89 mL, 0.28 m) was added, and the reaction heated to 70 °C for 12 h before being poured into methanol (200 mL) acidified with a few drops of concentrated HCl. The precipitate was filtered through a cellulose thimble, and the solid purified by Soxhlet extraction with methanol, acetone, and hexane (in each case until the extracting solvent was colorless). In order to determine if substantial amounts of P3OT homopolymer still remained, a differential scanning calorimetry (DSC) was run on a sample, and after confirmation that this was indeed the case, the solid was further washed with dichloromethane and finally extracted with chloroform, before being precipitated into methanol and filtered. The solid was dried under vacuum to give P3OT-b-F-P3OT (50 mg, 25%). M n 49 kDa, M w 78 kDa. 1 -F-P3OT): In a dry 20 mL microwave vial, 6 (750 g, 1.33 mmol) was dissolved in dry THF (17 mL), and to the stirred solution isopropylmagnesium chloride lithium chloride complex (1.00 mL, 1.3 m in THF) was added dropwise. After 1 h, a suspension of dichloro(1,3bis(diphenylphosphino)propane)nickel (10.8 mg, 1.5 mol%) in dry THF (1 mL) was added to the Grignard monomer solution, and the resulting solution was stirred at 70 °C for 5 h, before being poured into methanol (200 mL) acidified with a few drops of concentrated HCl. The precipitate was filtered through a cellulose thimble, and the solid purified by Soxhlet extraction with methanol, acetone, hexane, and chloroform (in each case until the extracting solvent was colorless). The residual polymer was dissolved in chlorobenzene, precipitated into methanol, and filtered. The solid was dried under vacuum to give P3OT-alt-F-P3OT (313 mg, 58%). M n 64 kDa, M w 88 kDa. 1 19 F, and 13 C NMR spectra were recorded on a Bruker AV-400 (400 MHz), using the residual solvent resonance of chloroform-d or 1,1,2,2-tetrachloroethane-d [2] and are given in ppm. Microwave experiments were performed in a Biotage initiator V 2.3. Number-average (M n ) and weight-average (M w ) molecular weights were determined by using an Agilent Technologies 1200 series GPC running in chlorobenzene at 80 °C, using two PL mixed B columns in series, and calibrated against narrow-polydispersity polystyrene standards. Due to poor solubility, M n and M w of F-P3OT was measured in 1,2,4-trichlorobenzene at 140 °C using a Polymer Laboratories PL-220 high-temperature GPC instrument calibrated against polystyrene standards. Electrospray mass spectrometry was performed with a Thermo Electron Corp. DSQII mass spectrometer. DSC measurements, using ≈3 mg of material, were conducted under nitrogen at scan rate of 10 °C min −1 with a TA DSC-Q20 instrument.
DFT Calculations: These were carried out using the B3LYP hybrid functional and the 6-31 g(d) basis set in the GAUSSIAN09 software package. [52] Alkyl chains were replaced with a methyl group to simplify calculations and reduce computational time.
Optical Characterization: Transmission and specular reflection spectra of polymer films on quartz substrates were acquired in the UV-vis and near infrared regions (220-1400 nm) using a Shimadzu UV-2600 spectrometer fitted with ISR-2600Plus integrating sphere option.
Scanning Kelvin Probe (SKP) measurements were performed at room temperature under ambient conditions by using an SKP5050 system from KP Technology Ltd. Samples were prepared by spin coating on ITO glass substrates, and work function values were calculated by averaging the last 500 measurement points of each data set (2000 points in total). For reference, a highly oriented pyrolytic graphite (HOPG) sample, with a nominal work function of 4.480 eV was used. The measurement uncertainty was calculated as the square of the quadratic sum of errors (the standard deviation) of both HOPG and sample data sets.
Device Fabrication: For the MOS device configuration, thin films of the polymers were spin coated from hot (≈150 °C) solution in 1,2,4-trichlorobenzene (5 mg mL −1 ) at 3000 rpm for 2 min onto hot Si ++ substrates with a 100 nm thick native SiO 2 layer (except P3OT, which was spun on room temperature substrates due to issues with dewetting). Metal top contacts (Au, thickness 40 nm, areas between 0.01 and 0.16 cm 2 ) were subsequently deposited by thermal evaporation (rate 0.1 nm s −1 ) using shadow masks. Additional MIM devices, comprising P3OT as the insulator, were prepared by spin coating a film (thickness 141 ± 5 nm) onto a glass substrate where bottom Au electrodes were previously deposited. Top Au contacts were then evaporated on the P3OT film, thus defining metal/polymer/metal structures. All polymer thickness values were measured using a Dektak 150 surface profilometer by taking several measurements for each film and calculating the average and std. deviation.
Dielectric Characterization: A Schlumberger SI 1260 impedance/ gain phase analyzer was used for the determination of the dielectric constants of the polymers by impedance spectroscopy. Capacitancefrequency (C-f) and Capacitance-Voltage (C-V) measurements were performed in the 10-10 6 Hz frequency range and with applied bias varied between −15 and +15 V. The capacitance of the MOS structure depends on the bias voltage applied, which influence the charge accumulation or depletion at the interface between the polymer and the oxide. The maximum capacitance, obtained in the accumulation regime at high negative bias for a p-type semiconductor, is determined exclusively by the oxide layer max S iO 0 SiO where A is the device area, ε 0 the vacuum permittivity, SiO 2 ε the dielectric constant of the insulator and SiO 2 d its thickness. The minimum capacitance is obtained when the polymer film is fully depleted at positive bias. The depleted polymer layer acts as a capacitor in series with the oxide layer, hence the total capacitance is The capacitance for the polymer, C polymer, can therefore be extracted as to obtain an estimate for the dielectric constant, ε polymer , of the polymer. [34][35][36]

Supporting Information
Supporting Information is available from the Wiley Online Library or from the author.