Synthesis of High-Crystallinity DPP Polymers with Balanced Electron

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On the synthesis of high crystallinity DPP polymers with balanced electron and hole mobility. Riccardo Di Pietro, Tim Erdmann, Joshua H. Carpenter, Naixiang Wang, Rishi Ramdas Shivhare, Petr Formanek, Cornelia Heintze, Brigitte Voit, Dieter Neher, Harald Ade, and Anton Kiriy Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.7b04423 • Publication Date (Web): 13 Nov 2017 Downloaded from http://pubs.acs.org on November 15, 2017

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Chemistry of Materials

On the synthesis of high crystallinity DPP polymers with balanced electron and hole mobility. Riccardo Di Pietro1*, Tim Erdmann2,3, Joshua H. Carpenter4, Naixiang Wang5, Rishi Ramdas Shivhare3, Petr Formanek2, Cornelia Heintze6, Brigitte Voit2,3, Dieter Neher5, Harald Ade4, Anton Kiriy2,3* 1

Hitachi Cambridge Laboratory, J. J. Thomson Avenue, CB3 0HE Cambridge, United Kingdom Leibniz-Institut für Polymerforschung Dresden e.V. (IPF), Hohe Straße 6, 01069 Dresden, Germany. 3 Technische Universität Dresden, Center for Advancing Electronics Dresden (cfaed), Dresden, Germany 2

4

North Carolina State University, 2401 Stinson Drive, Raleigh, NC 27695, United States of America Institute of Physics and Astronomy, University of Potsdam, Karl-Liebknecht-Str. 24-25, 14476 Potsdam 6 Helmholtz-Zentrum Dresden-Rossendorf, Institute of Resource Ecology, Bautzner Landstraße 400, 01328 Dresden, Germany 5

Abstract We

review

the

Stille

coupling

synthesis

of

P(DPP2OD-T)

(Poly[[2,5-di(2-

octyldodecyl)pyrrolo[3,4-c]pyrrole-1,4(2H,5H)-dione-3,6-diyl]-alt-[2,2’:5’,2’’-terthiophene5,5’’-diyl]]) and show that high quality, high molecular weight polymer chains are already obtained after as little as 15 minutes of reaction time. The results of UV-Vis spectroscopy, grazing incidence wide angle X-ray scattering (GIWAXS) and atomic force microscopy show that longer reaction times are unnecessary and do not produce any improvement in film quality. We achieve the best charge transport properties with polymers batches obtained from short reaction times and demonstrate that the catalyst washing step is responsible for the introduction of charge trapping sites for both holes and electrons. These trap sites decrease the charge injection efficiency strongly reducing the measured currents. The careful tuning of the synthesis allows us to reduce the reaction time by more than 100 times, achieving a more environmentally friendly, less costly process that leads to high and balanced hole and electron transport, the latter being the best reported for an isotropic, spin coated DPP polymer.

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Introduction

Dithienyl-diketopyrrolopyrrole (DPP) based copolymers are a class of semiconducting macromolecules which have been widely studied since their first introduction in 2007, due to the strong aggregating and electron donating properties of the DPP moieties.1–3 The initial reports on DPP polymers obtained by copolymerizing DPP with one or two thiophene rings showed good solubility and promising ambipolar transport.4,5 The combination of a very versatile chemistry and the strong π-π stacking capability of the dithienyl-DPP unit, with π-π stacking distances (around 3.6 - 3.9 Å) comparable to other high mobility semiconductors such as poly(2,5-bis(3alkylthiophen-2-yl)thieno[3,2- b]thiophenes)6 has since led to an unprecedented synthetic effort to find the right chemical modifications to improve device performance. The conventional AA+BB Stille coupling polycondensation protocol, first reported by Bao et al. in 1993, implies reaction of aryl dihalides and aryl distannanes usually in high boiling point solvents (e.g., toluene, chlorobenzene) in the presence of Pd catalysts (usually ligated by tri-aryl phosphines) at temperatures above 100°C usually over several days.7 Attractiveness of this approach originates from high functional group tolerance of Stille cross-coupling and its universality (that is a wide variety of structurally different polymers can be produced by the same synthetic protocol without optimizations of reaction conditions). With few exceptions,8,9 the original procedure reported by Bao underwent little modifications over 20 years and it is still used for preparation of most donor-acceptor copolymers. Exploiting its versatility, several modifications have been introduced in the chemical structure. Fused aromatic rings have been inserted in the polymer backbone to further improve π-π stacking, backbone planarity and crystalline packing.10–13 Solubilizing side chains also proved to impact both π-π stacking distance and size of crystalline domains.14–18 A careful tuning of the HOMO – LUMO structure by copolymerizing the DPP unit with electron donating units of different strength has been undertaken to produce ambipolar polymers with high and balanced hole and electron mobility,19–25 correlating the donor acceptor character to the improvement in the charge transport properties.26 Even the effect of additional electron withdrawing groups on the HOMO – LUMO structure,27–30 and their role in controlling texture and crystallinity through diffused non-covalent interactions has been studied.31,32 Heteroatom substitution allowed a reduction of the HOMO – LUMO gap and a further planarization of the polymer backbone driven the larger Van der Waals radius of selenium atoms.33–36 Overall, a very wide parameter 2 ACS Paragon Plus Environment

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space has been explored, causing a dramatic improvement of the charge transport properties of the synthesized polymers, with a correlation between the chemical and structural properties of the resulting polymers.37 However, while reported mobilities soared beyond 10 cm2 V-1 s-1 the measured current densities have not increased accordingly and remained very similar among the different publications, being limited by severe trapping effects. These traps have been shown to strongly impact charge carrier injection crippling the accuracy of the extracted mobility values.27,35 With the exception of a few limiting cases where clean injection and transistor behavior could be measured even at very low operating voltages, there is limited knowledge on the nature of these trapping sites and on how to reduce and possibly remove them.4,20,38 Here, we focus our analysis on one of the simplest possible molecular structures for a dithienylDPP polymer, P(DPP2OD-T), where the DPP acceptor has conventional C20 branched solubilizers and is paired to an unsubstituted thiophene as the donor unit. This structure was already reported in 2009 by Bijleveld et al. who found a relatively balanced hole and electron mobility of 0.04 and 0.01 cm2 V-1 s-1.5 Further optimization by Liu et al. and Zhang et al. led to a promising hole mobility in 0.3-1 cm2 V-1 s-1 range but with strongly hindered electron mobility.39,40 We review the standard Stille polycondensation method and show that high quality P(DPP2OD-T) is formed rapidly (within 15 minutes) and prolonged heating does not lead to a further increase in molecular weight P(DPP2OD-T). We further show that the catalyst washing step causes a significant decrease in both hole and electron current, an effect which is mitigated by air exposure but exclusively for hole transport.41 Through a careful scrutiny of the synthetic process we demonstrate hole and electron mobility in higher than 1 cm2/Vs in ambipolar OFETs of P(DPP2OD-T) over a broad voltage range, with the highest normalized electron current density for a spin-coated DPP polymer.20 The results presented here imply that high and balanced charge transport may be obtainable for a variety of other DPP based polymers using more environmentally friendly synthetic procedures. The critical aspect that needs to be addressed however is the presence of electron and hole trap sites in the polymer film, whose origin is shown here to be partially caused by specific synthetic steps. We enable high hole OR electron mobility by passivation of the respective trapping sites, but only eliminating such trapping sites altogether will allow to obtain true high ambipolar performance.



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2 2.1

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Experimental methods Synthesis of P(DPP2OD-T)

A Schlenk tube was heated and flushed with nitrogen before 408 mg (0.40 mmol, 1 eq.) 3,6-bis(5-bromothiophenyl)-2,5-di(2-octyldodecyl)pyrrolo[3,4-c]pyrrole-1,4(2H,5H)-dione

(1)

were

added, evacuated for 3 hours and dissolved with 20 ml dry and degassed chlorobenzene. In a glovebox solutions of 164 mg (0.04 mmol, 1 eq.) 2,5-bis(trimethylstannyl)thiophene (2), 9.8 mg (0.032 mmol, 0.08 eq.) tri(o-tolyl)phosphine and 7.4 mg (0.008 mmol, 0.02 eq.) tris(dibenzylideneacetone)dipalladium(0) were prepared using 20 ml of dry and degassed chlorobenzene and added to the Schlenk tube placed in an oil bath at 130 °C afterwards. In case of monitoring the polymerization progress, samples were withdrawn from the polymerization, quenched in 5 ml of HClc/MeOH (1:12.5), extracted by CHCl3, dried over MgSO4, dried under vacuum (40°C) and investigated by GPC. After stirring at 130 °C for the desired polymerization time, the cooled crude polymer was precipitated in HClc/MeOH (16 ml/200 ml), stirred overnight and the collected solids were purified by Soxhlet extraction using MeOH, acetone, and hexane (24 h each). The final polymer material was washed from the Soxhlet thimble by DCM and diluted with CHCl3. The polymer was isolated as dark solid by precipitation in methanol, filtration and drying under vacuum at 40 °C (346 mg, 92 %). Three different polymerization runs have been performed: P1, following the standard recipe reported in ref 42, was polymerized for 72 hours and used as a reference for our study. Two further batches P2-A and P2-B were polymerized for 15 minutes to assess the batch to batch reproducibility of our results. Independently of the specific synthetic protocol, all P(DPP2OD-T) batches were obtained as dark solids in high yields (91-95 %). The effect of the post-polymerization washing step was also studied by subjecting P1 and P2-B to different washing treatments. The not washed batches (NW) were not subjected to this step, the washed with water batches (WW) were washed with 40 ml deionised water at 60 °C for 2 h and the washed with ligand solution batches (WL) were washed with 40 ml aqueous sodium N,Ndiethyldithiocarbamate (SDEDTC) solution at 60 °C for 2 h and afterwards three times with water at room temperature.


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2.2

UV-Vis spectroscopy

UV-Vis spectra have been measured using a Varian Cary 5000 spectrometer, using a 1 cm quartz cuvette for the solution spectra (one cuvette containing only the solvent was placed in the reference beam) and borosilicate glass substrate coated with the polymer film for the thin film ones. 2.3

Grazing incidence wide angle X-ray scattering

2D GIWAXS patterns were acquired at beamline 7.3.3 at the Advanced Light Source at Lawrence Berkeley National Lab operating at 10 keV using a Pilatus 2M detector.43 For each sample, a 2D pattern was collected using a grazing angle within 0.01° above the critical angle (~0.13°), as well as a pattern with a grazing angle of 0.2°. To quantify d-spacings and coherence lengths in the π-π stacking direction, first in- and out-of-plane 1D profiles were obtained by azimuthally averaging +/-2° cake slices of the first 2D pattern about the given direction and plotting versus q. Peak positions were converted from reciprocal space to real space to obtain dspacings and the FWHM of the peaks were converted to coherence lengths using the Scherrer equation with a shape factor of 1 after correcting the measured widths for sample length and instrumental broadening. For relative degree of crystallinity (DoC) calculations, first pole figures for the π-π stacking peaks from the patterns recorded near the critical angle were calculated. This was done by integrating the 2D data about the peak in q such that all observable peak intensity was included, subtracting an appropriate background, and plotting vs. azimuthal angle. Pole figure intensities were corrected for illuminated volume based on thickness measurements made with a profilometer and sample length and also normalized to the relative peak intensity of the pattern recorded at a grazing angle of 0.2°, for which intensity scales more linearly with illuminated volume. Pole figures were Lorentz corrected assuming a 2D powder distribution and then integrated over azimuthal angle. The results were normalized to the most crystalline sample to give relative degree of crystallinity. All analysis was done in Igor Pro with a modified version of the NIKA package.44 2.4

Atomic force microscopy

Morphological surface analyses were carried out using a Bruker Dimension Icon microscope in tapping mode equipped with silicon-SPM-sensors (BudgetSensors, Bulgaria) with spring constant of about 40N/m and a resonance frequency of about 300 kHz. The tip radius was lower than 10 nm. 5 ACS Paragon Plus Environment


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2.5

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Field-effect transistors

P(DPP2OD-T) samples were characterized by using bottom contact top-gate (BCTG) OFETs. Chromium (1.5 nm)/gold (30 nm) source-drain electrodes were deposited on the surface of borosilicate glass substrates through shadow mask evaporation. Before deposition of the semiconductor, the substrates with gold electrodes were cleaned in ultrasonic bath by using chloroform, acetone and isopropanol, each for 5 min. P(DPP2OD-T) was dissolved under nitrogen (glovebox) either in chloroform (5 g l-1 concentration) or 1-methylnaphthalene (1-MN, 8 g l-1 concentration) and was spin coated on the cleaned substrates at 2000 rpm for 60 s at room temperature, yielding a semiconductor layer with thickness of 50 nm. To further tune the transistor performance the devices were then either dried in a vacuum oven (as cast) or annealed at 200 °C for 20 min under nitrogen. PMMA from Sigma Aldrich (Mn = 120 kg mol-1, 45 g l-1 in butyl acetate) was spin coated on top of the annealed films at 1500 rpm for 90 s, yielding a dielectric layer with thickness of 350 nm, and then heated at 80 °C for 30 min to remove the remaining solvent from the film. 20 nm thick aluminium gate electrodes were deposited on top of the PMMA layer via shadow mask evaporation. All the devices have the same channel length (L=20 µm) and channel width (W=1 mm). All the devices have been measured in a nitrogenfilled glovebox using an Agilent 4155 Semiconductor parameter analyser. 2.6

Transmission Electron Microscopy (TEM)

TEM was conducted with a Libra120 (Zeiss, Germany) operated at 120 kV. High-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) and energy dispersive Xray mappings (EDX) were conducted with a FEI Talos F200X (FEI, USA) operated at 200 kV and equipped with Super-X-EDX detector (Bruker, USA). The thin films for TEM were prepared by spin coating of the polymer solution in chloroform at concentration 5 mg/ml onto a PEDOT:PSS - coated glass. The glass was immersed in de-ionised water so that the polymer film was afloat on the water surface and collected onto a lacey-type carbon/copper TEM grids.

3 3.1 The

Results Synthesis DPP-based

monomer

3,6-bis-(5-bromothiophenyl)-2,5-di(2-octyldodecyl)pyrrolo[3,4-

c]pyrrole-1,4(2H,5H)-dione (1) was synthesized in 3 steps with an overall yield of 22 % 6 ACS Paragon Plus Environment

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according to modified literature procedures.45,46 Poly[[2,5-di(2-octyldodecyl)pyrrolo[3,4c]pyrrole-1,4(2H,5H)-dione-3,6-diyl]-alt-[2,2’:5’,2’’-terthiophene-5,5’’-diyl]]

(P(DPP2OD-T))

was prepared from dibromide (1) and commercially available 2,5-bis(trimethylstannyl)thiophene (2) by using a Stille coupling polycondensation protocol previously applied in the synthesis of a very high-molecular weight DPP-based polymer with a thieno[3,2-b]thiophene moiety.42 According to this procedure, the polymerization was conducted in chlorobenzene by using 2 mol-% of the palladium catalyst tris(dibenzylideneacetone)dipalladium(0) and 8 mol-% of tri(otolyl)phosphine at 130 °C for 3 days and led to high molecular weight P(DPP2OD-T) batch P1 with weight average molecular weight MW=214 kg·mol-1 and PDI=4.1 (GPC measurements in CHCl3 at 40 °C) in 90-92% yield (Table 1).

Scheme 1. Applied Pd-catalyzed Stille polycondensation.

Although the polymerization was conducted for a long time, we noticed that the significant color change from orange-red to dark green occurred within first few minutes indicating early formation of long enough conjugated chains and suggesting a high polymerization rate. To clarify whether long reaction times are necessary for obtaining high molecular weight P(DPP2OD-T), we investigated samples by gel permeation chromatography (GPC) taken at different reaction times. Figure 1 shows GPC chromatograms of reaction mixtures which correspond to different reaction times ranging from 3 minutes to 3 days and the development of Mw, Mn and Mw/Mn over the polymerization course. Oligomers with Mw of 5.2 kg mol-1 (Mw/Mn = 2.4) are formed within the first 3 min although substantial amount of unreacted monomer (28 %) is still present in the reaction mixture at this point. After 6 min, the monomer is fully consumed and the Mw of P(DPP2OD-T) reaches ~390 kg·mol-1. and only fluctuates around that value. The variation is much larger for Mw compared to Mn which rapidly saturates around ~390 kg·mol-1 and does not fluctuate much anymore the different behavior could be caused by different degrees of aggregation of the polymer chains in the different mixtures.3,45,47 We then 7 ACS Paragon Plus Environment


prepared the two polymer batches P2-A and P2-B to investigate the effect of such shorter polymerization time on the charge transport properties.

monomer 1

600000

Mw 500000

400000 8 300000

7

200000

6

100000

5

0

6

8

10

12

14

16

10

Mn Mw/Mn 9

0

18

Mw/Mn

3 min 6 min 12 min 24 min 60 min 4h 10 h 1d 2d 3d

Molecular weight (g mol-1)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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30

60

4 90 120 150 180 210 240 270 300

Reaction Time (min)

Elution time [min]

Figure 1. Synthesis of P(DPP2OD-T): collections of GPC chromatograms and development of Mw, Mw and Mw/Mn depending on the reaction time.

To observe the role played by aggregation we compared GPC in chloroform (CHCl3) at 40˚C and in trichlorobenzene (TCB) at 150˚C. The obtained values are reported in Table 1. All three batches show similar values (Mw ~ 220 kg mol-1 and Mn ~ 50 kg mol-1) and a broad molecular weight distribution when measured at low temperature. The same measurements at high temperature show however a radically different picture, with all the different polymer batches displaying much lower values of Mw in the range 40 - 60 kg mol-1 and Mn ~ 20 kg mol-1 (Table 1), confirming the strong aggregation in the polymer solution already observed, among others, in other polymers containing the DPP unit or in naphthalenediimide based polymers.45,47 Table 1. Molecular weight and polydispersity index of the different polymer batches extracted from GPC measurements at low and high temperature. Values have been confirmed by repeating the measurement. CHCl3 40 ˚C batch P1 P2-A P2-B

MN 52.7 53.9 49.9

MW 214.0 236.3 217.8

TCB 150 ˚C PDI 4.1 4.4 4.4

MN 27.4 17.7 23.3

MW 67.3 45.2 60.2

PDI 2.5 2.6 2.6


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Further characterization by high-temperature 1H-NMR spectroscopy, TGA, DSC and CV (Figure S1-S4) did not show any significant difference between the three different polymer batches.

3.2

UV-Vis absorption spectroscopy

We collected the UV-Vis absorption spectra in solution at a concentration of 10 µg ml-1 in different solvents at room temperature (Figure 2a). All the different spectra are characterized in solution by a strong J-aggregate behavior, dominated by a strong low energy absorption feature peaked at 1.51 eV with a lower intensity shoulder around 1.64 eV. We observe a markedly different behavior in CHCl3 compared to the other chlorinated aromatic solvents, possibly induced by a different solvatochromism in such solvent. In aromatic solvents the absorption peaks are instead narrower and better separated and it is possible to observe a slight increase in the relative intensity of the 1.64 eV absorption feature going from chlorobenzene (CB) to 1methylnaphthalene (1-MN), while no difference is observed for P1. This is consistent with what has already been reported for many different DPP copolymers and is a further evidence of a strong pre-aggregation in solution.48 In case of J-aggregates, the increase in the higher energy transition intensity can be in fact explained with a reduced conjugation length that decreases the strength of intra-chain coupling.49 In the inset of figure 5a the absorption edge is shown in a logarithmic scale: the absorption onset shifts from 1.19 eV to 1.21 eV going from chloroform to 1-MN, although the slope of the absorption edge is similar for all different preparation conditions.



b 100

1.51 eV

10-1

1.0

1.19 eV

1.64 eV

10-2

1.21 eV

10-3 0.5

1.45 eV 1.48 eV

1.20 eV

CHCl3

1.21 eV

1.2 DCB

CB

1.4

Norm. absorbance

a

Norm. absorbance

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1.0

1.61 eV

1.61 eV

0.5 as cast

1-MN

0.0 1.2

1.4

1.6

ann.

1.8

2.0

2.2

0.0 1.2

1.4

Energy (eV)

1.6 1.8 Energy (eV)

2.0

2.2

Figure 2. Panel a shows the solution UV-Vis absorption spectra of polymer P2-A dissolved in different solvents. The absorption onset is shown in the inset, together with the obtained onset values. Panel b shows the thin film absorption spectra of polymer P2-A coated from 1-MN as cast (dried in vacuum) and annealed at 200˚C.

The solid-state spectra (Figure 2b) are red-shifted and broadened compared to the solution ones, but show a very similar ratio between the low and high energy absorption peaks. The main difference between films annealed at 200˚C and film dried in a vacuum oven at room temperature is a red shift of the low energy absorption peak from 1.48 eV to 1.45 eV upon annealing. Furthermore, the peak ratio between the low and high energy absorption peaks increases significantly upon annealing, suggesting an increase in conjugation length along the polymer backbone.


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Figure 3. a-d) AFM topography maps of P2-A in CHCl3 and 1-MN: a) P2-A dissolved in CHCl3 as cast, b) P2-A dissolved in 1-MN as cast, c) P2-A dissolved in CHCl3 annealed at 200˚C, d) P2-A dissolved in 1-MN annealed at 200˚C. e-h) GIWAXS 2D plots of films prepared in the same conditions: e) P2-A dissolved in CHCl3 as cast, f) P2-A dissolved in 1-MN as cast, g) P2-A dissolved in CHCl3 annealed at 200˚C, h) P2-A dissolved in 1-MN annealed at 200˚C.

3.3

Structural characterization

We characterized the texture and morphology of the thin films coated from the different polymer batches by measuring GIWAXS and AFM topography. In Figure 3 the results for polymer P2-A



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are shown. When coated from chloroform solution, the films shows small needle like features even when spin coated from chloroform, both without and with thermal annealing (figure 3a-c). The presence of such structures even without annealing and the similar morphology upon annealing is a further indication of the strong aggregation tendency of the polymer in such solvent. When spin coating the polymer from 1-MN solution, the as cast sample shows a rather featureless surface (Figure 3b) while upon annealing extremely large needle like features are observed (Figure 3d). Higher boiling point aromatic solvents reduce the density of nucleation points in the film, allowing for features to grow larger upon annealing, while low boiling point solvent (and the different solvation observed from the solution UV-Vis spectra) are responsible for the very fast formation of smaller features. More information on the formation of crystalline domains is obtained from GIWAXS measurements: films coated from 1-MN solution and vacuum dried exhibit weak features relative to background and diffuse scattering with π-π stacking out of plane, indicating a face-on orientation of the crystalline domains (figure 3f). In chloroform, where aggregation is stronger, a more structured diffraction can be observed with both π-π stacking and lamellar stacking taking place in- and out-of- plane, indicating the presence of both edge-on and face-on crystallites (figure 3e). Upon annealing the amorphous halo is largely reduced and a more unimodal orientation distribution is visible in both chloroform (face-on orientation, figure 3g) and 1-MN (edge-on orientation, figure 3h) although some contribution from a face-on orientation is still visible in the latter. It has been recently shown by Mueller et al. and Kim et al. how the diffusive non covalent interaction between polymer chains can be used to control the orientation of the polymer chains.31,32 Here we show that orientation control is possible also through the choice of solvents, similar to what has already been observed for P(NDI2OD-T2).50 All three polymer batches showed similar results. From the GIWAXS plots we extracted the d-spacing along the π-π stacking direction (the direction most relevant for charge transport), the disorder induced coherence length (LC) along that direction and its relative degree of crystallinity (DoC) for all the different samples (Table 2). We measured a d-spacing between 3.73 - 3.81 Å for all the different preparation conditions regardless of polymer batch, chosen solvent or annealing condition, with the exception of P2-A coated from 1-MN as cast (no annealing), for which we found a slightly larger d-spacing along


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the π-π stacking direction of 3.86 Å. More complex is the behavior of the coherence length LC. It increases with annealing from 6 to 7.2 nm for polymer P2-A coated from 1-MN and from 4.4 to 6 nm for P2-B and P1. The trend is opposite when P2-A is coated from CHCl3 with LC decreasing from 6.9 to 4.3 nm upon annealing, highlighting the different behavior of such solution as already observed in the UV-Vis spectra and AFM characterization. We observe an overall higher coherence length for polymer P2-A compared to all other polymer batches. Overall it can be seen how shorter polymerization times lead to the best film properties, with P2A having the longest coherence length of the whole set (7.2 nm) and the large supramolecular aggregates in solid state. These conclusions are also supported by the calculation of the relative degree of crystallinity, which shows similar trends as we discussed for the coherence length, and the highest degree of crystallinity again for polymer P2-A. No improvement is observed when extending polymerization from 0.25 to 72 h. Table 2. Coherence length, d spacing and degree of crystallinity along the dominant (in- or out-of-plane) π-π stacking direction for the different films. Also reported is the average maximum mobility extracted from thin film transistor characterization. Mobilities are quoted only for sample for which no significant trapping is observed. All films are coated from 1-methylnaphthalene (1-MN) except where explicitly stated. Also indicated is whether the film has been annealed at 200˚C (Ann.) or dried in vacuum at room temperature (as cast). Sample

Annealing

Texture

d (Å) (±0.01)

LC (nm) (±0.2)

# of spacings

Rel. DoC

µhole (cm2/Vs)

µelectrons (cm2/Vs)

P1 P1 P2-A (CHCl3) P2-A (CHCl3) P2-A P2-A P2-B P2-B

RT (vac.) 200˚C RT (vac.) 200˚C RT (vac.) 200˚C RT (vac.) 200˚C

Face-on Edge-on Face-On Face-On Face-on Mixed Face-on Edge-on

3.81 3.76 3.73 3.74 3.86 3.81 3.80 3.73

4.5 6.0 6.9 4.3 6.0 7.2 4.3 5.9

11.8 16.1 18.4 11.4 15.5 18.8 11.5 15.7

0.29 0.72 0.95 0.24 0.31 1.00 0.32 0.72

0.7±0.3 2.0±0.2 0.079±0.015 0.088±0.005 0.54±0.07 2.4±1.2 1.0±0.4 2.0±0.2

0.54±0.09 2.1±0.5 0.05±0.03 0.16±0.014 0.477±0.005 6.3±1.5 0.60±0.14 1.6±0.3

4

Charge transport

Charge transport was quantified using field-effect transistors in the bottom contact top gate geometry (BCTG), following the procedure optimized by Chen et al.25



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OFET characterization and mobility extraction

The mobility extraction method and the impact of the atmosphere in which the measurements are carried out require clarification before a discussion of the impact of morphology and synthesis on the charge transport properties can be undertaken. In Figure 4a output curves (both holes and electrons) of P2-A coated from 1-MN and annealed are shown (the device was baked for 10h at 80˚C before being measured in nitrogen). Currents are in line with what has been already reported in literature for P(DPP2OD-T), with higher hole current compared to the electron one (only the output curves for VG=±20 V, VG=±30 V, VG=±40 V for holes and electrons respectively are shown). From the shape of the output curve at low drain voltages a high injection barrier for both holes and electrons can be inferred. This can be expected since we used solvent cleaned gold which has a slightly lower work function compared to the oxygen plasma cleaned counterpart. Transfer curves (Figure 4b) have been acquired both with fixed drain voltage VD=±60 V (in black) and with VD= VG (in green).


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d

c

with air exp.

VG=-40V

0.10 0.05

0V

VG=0 V

-10V

30V -30V

10V

20V

-20V

0.00 -40 -30 -20 -10 0 10 20 Drain voltage (V)

30

VD=-60V

40

e

10-1 VD=+40V VD=-40V 10-2 10-3 10-4 VD=VG 10-5 VD=VG 10-6 10-7 10-8 -40 -30 -20 -10 0 10 20 30 Gate voltage (V)

40

f

VD=+60V

100 10-1 10-2 VD=VG VD=VG 10-3 -4 10 10-5 -40 -30 -20 -10 0 10 20 Gate Voltage (V)

Abs. Idrain (mA)

0.15

VG=40V

Abs. Idrain (mA)

Abs. Idrain (mA)

b

output

Mobility (cm2/Vs)

Abs. Idrain (mA)

a

Mobility (cm2/Vs)

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30

40

0.15

w/o air exp.

0.10 0.05 0.00 -40 -30 -20 -10 0 10 20 Drain voltage (V)

30

40

30

40

30

40

10-1 10-2 10-3 10-4 10-5 10-6 10-7 -40 -30 -20 -10 0 10 20 Gate voltage (V) 100 10-1 10-2 10-3 10-4 10-5 -40 -30 -20 -10 0 10 20 Gate Voltage (V)

Figure 4. Output curves at VG=0V, ±20V, ±40V, ±60V (a) and transfer characteristics (b) of a field-effect transistor prepared with a 5 g/l solution of P2-A in 1-methylnaphthalene. The hole and electron mobility (c) are calculated using the method reported in ref. 51. Output curves (d) and transfer characteristics (e) of a thin film transistor prepared using polymer P2-B, with the measurements performed without air exposure in black and the ones taken after storing the device in air in the dark for two days reported in red (the output curve for electrons after air exposure is too low to be seen in the linear scale used in panel d). The gate voltage dependent mobility is reported in panel f and is calculated from the curves in panel e.

By applying the same voltage to drain and gate we observe unipolar transport throughout the voltage range, while by keeping the gate voltage fixed the typical v shape of an ambipolar transistor is observed (the two measurements overlap perfectly at high gate voltages, while they differ at low gate voltages). From an operational point of view, it is relevant to measure the ambipolar transfer (fixed VG) since it shows the device behavior in an electronic circuit,52 but for material characterization it is more useful to measure the unipolar saturation transfer curves 15 ACS Paragon Plus Environment


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(VD= VG) for both holes and electron. This method (which is the one used in all the following analysis) allows us to measure hole and electron transport in ambipolar devices down to very low gate voltages and it also avoids some of the possible artifacts that have been recently discussed in literature, such as the increase in current by arbitrarily applying an artificially high drain voltage during the measurement, by reflecting the actual saturation condition used in the gradual channel approximation(VD= VG-Vth).42 The mobility is then calculated according to the procedure outlined in ref.

51

, leading to the gate voltage dependent mobility curves reported in Figure 4c,

where it is also possible to observe how the mobility extracted from the two methods perfectly overlap in the range where both transfer curve are dominated by the same carrier type. Measuring the unipolar transfer curves, it is possible to measure hole and electron transport all the way to the onset voltage (VON=-20 V for holes and VON=+20 V for electrons for this specific device, which can be measured directly). It is then possible to observe a dramatic drop in the extracted mobility at low gate voltages. Such artificial drop in mobility by many order of magnitude has been shown to be caused by charge trapping taking place in the bulk of the semiconductor and causing a strong increase in contact resistance, as we already detailed in the case of hole only devices based on P(DPP2OD-TT).45 In the same work, we also showed how these trapping sites could be passivated by exposing the device to air. The output characteristics (Figure 4d) unipolar transfer characteristics (Figure 4e) and mobility (Figure 4f) of a P2-B transistor coated from 1-MN and annealed is shown. Without air exposure (black trace) we obtain rather similar hole and electron currents, with very little hysteresis and a rather symmetric behavior (VON=±5V for both hole and electrons). Both hole and electron mobility are strongly gate voltage dependent and increase from 10-3 to about 2 cm2/Vs. Upon air exposure, the device becomes practically a hole-only device. Hole mobility is increased by two orders of magnitude at VG=-10V, while being unaffected at high gate voltages. This is in complete agreement with our recent work on P(DPP2OD-TT), and points to the presence of electron donating traps in the bulk of the polymer film strongly hindering hole injection which can be passivated through air exposure.45 For an ambipolar device however air exposure has a dramatic impact also on electron transport, due to electrons reacting with water and oxygen complexes.53,54 Through the very same mechanism water and oxygen can be responsible for both the passivation of the electron donating trap site and the trapping of electrons injected in the LUMO of the polymer film. A similar trap passivating effect can be 16 ACS Paragon Plus Environment

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obtained for holes by evaporating F4TCNQ, which leads to a moderate doping of the film. A recent work by Nikolka et al. demonstrated this to be a rather general effect which can be obtained also with additives that do not lead to doping of the polymer film.41 However the approach reported by Nikolka et al. cannot be followed here if we want to characterize both hole and electron transport since it passivates hole transport only while at the same time blocking electron transport. We therefore opted to measure electron transport before air exposure and hole transport after air exposure. This ensures that the morphology of the films is comparable to what we measured in the previous sections while at the same time it minimizes the influence of charge trapping. 4.2

Impact of morphology and synthesis

a 10-1 Abs. Idrain (mA)

10-2 10-3 10-4 10-5

10-6 -40 -30 -20 -10 0 10 20 Gate voltage (V)

b

103 Mobility (cm2/Vs)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


102

1-MN ann. 1-MN as cast

30

40

CHCl3 ann. CHCl3 as cast

101 100 10-1 10-2 10-3 10-4 -40 -30 -20 -10 0 10 20 Gate voltage (V)

30

40

Figure 5. a) Transfer characteristics of transistors fabricated with polymer P2-A and the same preparation conditions used in figure 3. b) Calculated mobility of the two different P(DPP2OD-T) polymer batches. All the transistors have 1 mm channel width and 20 µm channel length. Error bars on mobility estimation are shown as grey areas around the respective trace. Electron characterization is



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performed before air exposure, the devices are then exposed to air in the dark for 3 days and hole transport is measured.

In Figure 5a the transfer characteristics of transistors prepared using P2-A with the same preparation conditions used in Figure 3 are shown, together with the calculated mobilities in b. We observe high and balanced hole and electron mobilities. Despite the presence of clear injection issues (as already discussed) which limit the accuracy of the extraction especially at low gate voltages, it is still possible to draw some important correlations between film morphology and charge transport. The as cast samples prepared from chloroform solution show a mobility of 4·10-2 cm2 V-1 s-1 for both electrons and holes at VG= ±40V, which increases to about 1-2·10-1 cm2 V-1 s-1 upon annealing (although in the case of electron transport the comparison is hindered by a large decrease in injection efficiency). On the other hand, spin coating polymer P2-A from 1-MN leads to much better performance, increasing the calculated mobility to 6·10-1 cm2 V-1 s-1 for the as cast sample and around 3-5 cm2 V-1 s-1 for the annealed sample. Care should be taken when quoting these values due to the non-ideal injection efficiency, however it is important to notice that such increase in mobility is always accompanied by a similar increase in current, clear evidence of an improvement in charge transport. For our best performing devices mobility rises above 1 cm2 V-1 s-1 throughout the range 20 < |VG| < 40 V for both holes and electrons, in conjunction with a current that increases from 7·10-5 A for the as cast sample to 4·10-4 A for the annealed one. It is also interesting to notice that we always observe better injection (higher currents at low gate voltages) for both holes and electrons in the as cast samples, compared to the annealed ones. This effect has been observed also for other DPP polymers, and might provide further clues on the nature of possible trapping sites involved in the process. We obtain similar values for all three polymer batches (P1, P2-A and P2-B) which enable us to conclusively state that shortening the reaction time by more than two orders of magnitude, from 72h to 0.25h, has no impact on film formation. 5

Charge trapping induced by catalyst removal

It was shown in previous works that the presence of residual transition metal catalysts in semiconducting polymers may adversely affect their optoelectronic properties, by degrading performance and stability in photovoltaic devices.55 One of the most efficient way for removal of the catalysts implies an extraction of polymer solutions by aqueous solutions of specially 18 ACS Paragon Plus Environment

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designed ligands, such as N,N-diethylphenylazothioformamide, having high complexing ability with transition metals.56 Since this process is carried using water-based solutions, it is also important to analyze whether the washing step itself (i.e., treatment with pure water) can have a detrimental impact on charge transport. To systematically address these issues, we compared the solution and thin film properties of the same polymer batch subjected to different washing treatments: the first portion was left untreated (P2-B NW and P1 NW); the second portion was treated with deionized water (P2-B WW and P1 WW) whereas the third portion was washed with a ligand solution (P2-B WL and P1 WL). High-temperature

1

H-NMR spectroscopy, TGA, DSC, GPC and CV (Figure S1-S4)

measurement conducted on samples of the same polymer batch (all three of them were tested) subjected to the different washing steps showed no significant difference (details are reported in the supplementary information) as did the UV-Vis spectroscopy and AFM characterization. Only the coherence length extracted from the GIWAXS plot shows a slight decrease for washed polymer batches, more evident for longer polymerization times. While P1 annealed shows a coherence length of 6 nm NW, similar to P2-B annealed (5.9 nm) and only slightly lower than P2-A annealed (7.2 nm), while for P1 WW and WL coherence lengths decreases to 4.7 nm. Films prepared from the short polymerization batch P2-B WW and WL show only a minor decrease in coherence length to 5.6 and 5.2 nm respectively. We estimated the content of residual Pd in the polymer and correlated it with the semiconducting properties of the polymer, by investigating the differently-treated polymer samples by TEM. We found Pd nanoparticles in the polymer batches which were not treated with the speciallydesigned ligand SDEDTC for the Pd removal (both non-treated and water-washed), 3 nm to 8 nm in size, randomly distributed in the film (Figure 2a). In contrast, no Pd nanoparticles were found in the SDEDTC-washed samples (Figure 2b). Interestingly, the largest amount of the nanoparticles was observed in the samples polymerized over short time whereas Pd nanoparticles are rare in the long-time polymerized samples (see SI). We quantified the Pd content in the highest containing sample P2-B NW. The particle size varied between 3.6 nm and 7.2 nm with ensemble mean and standard deviation 5.4 ± 1.0 nm (evaluated from 70 nanoparticles) and 4.1 ± 2.1 nanoparticles in each cluster. Clusters with more than 8 nanoparticles occasionally occur (evaluated from 120 clusters, see SI). The volume fraction of the nanoparticles with respect to



the volume to the polymer is ~8x10-5 (assuming film thickness of ~50 nm). The density of cluster per area is approx. 12 µm-2, the volume concentration of clusters is ~2.5 x 1014 cm-3. c

a Drain current (A)

10-4 10-5 10-6 10-7 10-8 10

-9

-40 d

b

P2-A NW P2-A WW P2-A WL -20 0 0 20 Gate voltage (V)

40

101 100 Mobility (cm2/Vs)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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10

-1

10-2 10-3 10

-4

10-5 10-6 -40

w/o air exp. with air exp. -20

0 20 Gate voltage (V)

40

Figure 6. TEM images of ~50 nm films of pristine P2-B NW (a) and ligand-treated sample P2-B WL (b). In c the transfer characteristics of transistors based on polymer P2-B with different washing conditions are shown. In d the mobility calculated from the measured transfer characteristics in c is plotted.

The impact of the catalyst washing step on the charge transport properties is shown in Figure 5c (transfer characteristics) and d (mobility) for polymer P2-B treated with the three different washing steps. Both hole and electron transport are shown with (dotted line) and without (straight line) air exposure. The first and most important conclusion we obtain is that despite no significant difference is observed from the chemical characterization and only a moderate decrease in coherence length, the washing step significantly impacts the charge transport properties. No washing step should be undertaken to obtain the best charge transport properties. Washing with DI-water has limited impact on hole transport but reduces the maximum measured electron current by 30%. When a water based ligand solution is used, hole current decreases by at least 60% while the electron current by 75%. The effect is even stronger at lower gate voltages. Upon air exposure hole transport is recovered and a much smaller difference between 20 ACS Paragon Plus Environment

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devices with different washing steps is observed, mainly at low gate voltages. The mobility curves extracted from these devices highlight the bulk trapping nature of the reduction in transport. It is in fact very difficult to observe differences between the mobility of the device with the different washing step at high gate voltages: beyond |VG|= 20 V all the extracted mobilities are very similar and between 1 – 3 cm2 V-1 s-1, confirming the similar morphology and molecular packing of the different polymer films. However major changes are observed at low gate voltage with the differences becoming significantly larger for the ligand based solution. The density (ca. 1014 cm-3), size (ca. 5 nm) and spacing (ca. 100 nm) of the Pd nanoparticle rules out the possibility of transport being shorted through the Pd nanoparticles and also their possible role in passvating already present bulk traps, which to affect injection should be in the 1016 cm-3 range. Therefore we conclude that hole and electron traps are directly introduced by water and N,N-diethylphenylazothioformamide, and the washing step should be avoided in order to improve charge transport. The presence of Pd nanoparticles in the film, on the other hand, does not affect device performance. This point is clear if we compare the performance of P1 and P2B. The two polymer batches have very similar performance despite P1 showing at least one order of magnitude lower Pd nanoparticle density (see supplementary material). 6

Discussion and conclusions

The presented results clearly show that the chosen synthetic pathway has a dramatic influence on morphology and charge transport properties of the polymer film, despite the result of the synthesis being nominally the same when characterized with standard and commonly employed techniques. Even a seemingly harmless step such as the catalyst removal, which has no measurable impact on the final product of the synthesis, proves to introduce bulk trapping mechanisms which harm hole transport and kill electron transport. These trap states have usually densities that are far below what is possible to detect with standard bulk characterization techniques and even with more sophisticated ones such as charge accumulation spectroscopy.53 In Figure 7a, we plotted the measured coherence length (in orange) and the relative DoC versus the calculated maximum mobility. The values are also reported in the disorder induced coherence length (LC) along that direction and its relative degree of crystallinity (DoC) for all the different samples (Table 2). We measured a d-spacing between 3.73 - 3.81 Å for all the different preparation conditions regardless of polymer batch, chosen solvent or annealing condition, with the exception of P2-A coated from 1-MN as cast (no annealing), for which we found a slightly 21 ACS Paragon Plus Environment


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larger d-spacing along the π-π stacking direction of 3.86 Å. More complex is the behavior of the coherence length LC. It increases with annealing from 6 to 7.2 nm for polymer P2-A coated from 1-MN and from 4.4 to 6 nm for P2-B and P1. The trend is opposite when P2-A is coated from CHCl3 with LC decreasing from 6.9 to 4.3 nm upon annealing, highlighting the different behavior of such solution as already observed in the UV-Vis spectra and AFM characterization. We observe an overall higher coherence length for polymer P2-A compared to all other polymer batches. Overall it can be seen how shorter polymerization times lead to the best film properties, with P2A having the longest coherence length of the whole set (7.2 nm) and the large supramolecular aggregates in solid state. These conclusions are also supported by the calculation of the relative degree of crystallinity, which shows similar trends as we discussed for the coherence length, and the highest degree of crystallinity again for polymer P2-A. No improvement is observed when extending polymerization from 0.25 to 72 h. Table 2 Although as we already pointed out these values are not extremely accurate due to the

reported injection issues, the same correlation would be obtained if current instead of mobility was used. There are some anomalous values, most notably the high coherence length and DoC of the CHCl3 as cast sample which are extremely high, which could be caused by the bulk values (measured by GIWAXS) being different from the surface ones probed by FET measurements. This effect could be stronger for fast drying solvents. In any case, the plot highlights a correlation between charge transport and morphology: mobility is low for samples coated from low boiling point solvents and samples without thermal annealing show rather low mobility (consistently with their rather poor morphology), and increases for samples annealed from high boiling point solvents. For the latter in particular a correlation between performance and crystallinity can be observed: the increase in crystallinity proceeds through the enlargement of the crystalline domains, which are then responsible for the observed increase in current and charge carrier mobility. This correlation has been recently discussed in detail for P(NDI2OD-T2) and P3HT, for which the better injection properties allowed a more accurate correlation, but the general effect can be observed also for P(DPP2OD-T).57 It is important to stress that the highest crystallinity values have been obtained with a short polymerization time polymer batch (P2-A).


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Quantifying how well the device is performing is rather straightforward when all samples are prepared the same way, however it becomes a much more complex task when devices with very different geometries are compared and charge injection issues are sever. To provide a meaningful comparison with the literature results on DPP polymers we compared the maximum currents that have been reported. To do this however we need to normalize the current density for all the geometric parameters of the device, which change significantly from study to study. A way to do this is to take the maximum current value reported for a specific device and calculate the “ideal mobility” that would be required for an ideal transistor with the same device geometry, constant mobility and 0V threshold voltage to output the measured current value, according to the following formula: ‫ܫ‬஽௦௔௧ =

ܹ 2‫ܮ‬ ߤ‫ܥ‬௜ ሺܸீ ሻଶ → ߤ = ‫ ܫ‬௦௔௧ 2‫ܮ‬ ܹ‫ܥ‬௜ ሺܸீ ሻଶ ஽

We can then compare the ideal mobility we obtain across multiple studies, since it directly represents the current corrected for device geometry and charge carrier density (a similar approach has been used in 58). b 7.0

6.0 0.6

5.5

1-MN ann. (edge-on)

5.0

0.4

4.5 0.2

4.0

face-on

3.5

0

1 2 mobility (cm2/Vs)

3

Ideal FET mobility (cm2/Vs)

0.8

Rel. Degree of Crystallinity

1.0

6.5

c 64 16 20(aligned) 38 65 33 63 26 9 39 19 34 46 20 10 59 60 21 27 24 13 66 4 11 35 25 14 62 27

1

0.1

5

0.01

29

13

DPP-T P2-A DPP-TT DPP-BTZ 0

2

4

6

8 12

10 -2

Hole density (10 cm )

Ideal FET mobility (cm2/Vs)

a

Coherence length (nm)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


1

0.1

20(aligned) 29 N2200 65 27 60 61 19 24 35

0.01

0.001

25 21 13

33

63 62

13 4 27

5 14

0 2 4 6 8 10 12 14 16 Electron density (1012 cm-2)

Figure 7. Comparison of DPP based OFET performance. Panel a shows the correlation between coherence length, degree of crystallinity and the calculated maximum hole mobility. In panel b and c the ideal mobility values (holes and electrons respectively) are obtained as the mobility of an ideal transistor which shows the same maximum current as that measured in the actual device. This method allows comparison of the actual measured currents across several studies, corrected for gate voltage, dielectric capacitance and device geometry, rather than relying on the reported mobility. DPP-T stands for



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P(DPP2OD-T) and DPP-TT for P(NDI2OD-T2). The numbers next to the data points indicate the reference from which the data was taken.59–66

The comparison is reported in Figure 7b and c, where these ideal mobility values are reported against charge density. All the P(DPP2OD-T) values are reported in red, with the results from our best performing batch P2-A shown with a red star. An interesting consideration is that the spread in performance for a specific molecular structure is larger than the difference between the different molecular structures. We also point out that we did never calculate an ideal mobility higher than 2 cm2/Vs, despite some of these references claiming mobility above 10 cm2/Vs. Although it would not seem we are reporting an improvement in device performance compared to earlier reports, the polymer batches we used in our work consistently end in the top half of both graphs. Extremely high hole performance has been reported for P(DPP2OD-T) already in ref.

38,39

, where however pure p-type transport was observed. With our careful optimization of

the synthesis we demonstrate one of the highest electron current density for a spin coated DPP transistor, close to the what has been reported for an aligned DPP-BTZ polymer.20 This is an extremely relevant result as it clearly demonstrates that DPP polymers are intrinsically able to allow very high and balanced hole and electron transport regardless of the specific choice of substituents, which can then be chosen to tailor other electronic and mechanical properties of the material. The crucial step that needs to be avoided to observe electron transport is the catalyst washing step, which is directly responsible for the introduction of electron trapping sites. In conclusion, we have carefully reviewed the synthetic route to produce P(DPP2OD-T), starting from the optimization of reaction time and obtained molecular weight. We show that the reaction is completed within as little as 15 minutes, shortening the reaction time by a factor of 290, with highly significant consequences in terms of scalability cost and energy savings. By optimizing the synthesis, we demonstrate that highly crystalline films of DPP polymers achieve balanced hole and electron mobility above 1 cm2 V-1 s-1 over a very broad voltage range, evidence that electron transport does not need to be promoted through the presence of specific electron accepting units such as nitrile or pyridine groups in the polymer backbone,24,27,62 but is indeed accessible throughout the DPP family with the necessary precautions during synthesis. Such precautions involve even seemingly harmless steps such as the catalyst scavenger treatment which has a dramatic negative impact on electron transport. Future development will need to


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focus on the identification and quantitative characterization of the bulk trapping sites we have observed in P(DPP2OD-T), requiring an even higher control over the quality of the synthesized polymer. This will require at the same time the development of techniques that can quantitatively characterize such defect states, which are clearly invisible with the standard suite of chemical / structural characterisation methods. Supporting Information. Additional Instrumentation description. 1H-NMR Spectroscopy of compounds P2-B and P1. Thermogravimetric analysis of compounds P1 and P2-B. GPC of all compounds subjected to different washing steps. DSC of compounds P1 and P2-B. Cyclic voltammetry of compounds P1 and P2-B. Uv-Vis and GIWAXS analysis of P1 and P2-B. TEM of compounds P1 and P2-B subjected to different washing steps. Acknowledgements R. D. P. would like to acknowledge the European Union ERC Synergy Grant SC2 (No. 610115) for funding and Mr Martin Statz for useful discussions regarding the fabrication process. This work was partly supported by DFG within the Cluster of Excellence “Center for Advancing Electronics Dresden” (cfaed). The use of HZDR Ion Beam Center TEM facilities and the support by its staff is gratefully acknowledged. In particular, we acknowledge the funding of TEM Talos by the German Federal Ministry of Education and Research (BMBF), Grant No. 03SF0451 in the framework of HECMP.

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