Control of Polymer-Packing Orientation in Thin Films through Synthetic

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Control of Polymer-Packing Orientation in Thin Films through Synthetic Tailoring of Backbone Coplanarity Mark S. Chen, Jeremy R Niskala, David A. Unruh, Crystal K. Chu, Olivia P. Lee, and Jean M. J. Frechet Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/cm402489a • Publication Date (Web): 18 Sep 2013 Downloaded from http://pubs.acs.org on October 1, 2013

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Chemistry of Materials

Control of Polymer-Packing Orientation in Thin Films through Synthetic Tailoring of Backbone Coplanarity Mark S. Chen,*,†,‡ Jeremy R. Niskala,†,‡ David A. Unruh,†,‡ Crystal K. Chu,† Olivia P. Lee,†,‡ and Jean M. J. Fréchet*,†,‡,§ †

Department of Chemistry and Department of Chemical Engineering, University of California, Berkeley, California 94720-1460, United States ‡

Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, California, 94720, United States

§

King Abdullah University of Science and Technology, Thuwal, Saudi Arabia 23955-6900 Supporting Information Placeholder ABSTRACT: Controlling solid-state order of π-conjugated polymers through macromolecular design is essential for achieving high electronic device performance; yet it remains a challenge, especially with respect to polymer-packing orientation. Our work investigates the influence of backbone coplanarity on a polymer’s preference to pack face-on or edgeon relative to the substrate. Isoindigo-based polymers were synthesized with increasing planarity by systematically substituting thiophenes for phenyl rings in the acceptor comonomer. This increasing backbone coplanarity, supported by density functional theory (DFT) calculations of representative trimers, leads to the narrowing of polymer band gaps as characterized by UV-vis-NIR spectroscopy and cyclic voltammetry. Among the polymers studied, regiosymmetric II- and TIIpolymers exhibited the highest hole mobilities in organic field-effect transistors (OFETs), while in organic photovoltaics (OPVs), TBII-polymers that display intermediate levels of planarity provided the highest power conversion efficiencies. Upon thin-film analysis by atomic force microscropy (AFM) and grazing incidence X-ray diffraction (GIXD), we discovered that polymer-packing orientation could be controlled by tuning polymer planarity and solubility. Highly soluble, planar polymers favor face-on orientation in thin films while the less soluble, non-planar polymers favor an edge-on orientation. This study advances our fundamental understanding of how polymer structure influences nanostructural order and reveals a new synthetic strategy for the design of semiconducting materials with rationally engineered solid-state properties. KEYWORDS: conjugated polymers, packing orientation, backbone coplanarity, isoindigo, organic field-effect transistors, organic photovoltaics INTRODUCTION Organic electronics is a rapidly growing field driven by the discovery of new carbon-based semiconducting materials.1 Innovative polymer design has led to remarkably high charge mobilities in organic field-effect transistors (OFETs)2,3 and power conversion efficiencies (PCEs) with organic photovoltaics (OPVs).4,5 Extensive studies of these materials have enabled chemists to devise numerous strategies for engineering optoelectronic properties.6 Yet for polymers to demonstrate state-of-the-art device performance, it is also essential that they be designed with specific nanoscale, solid-state properties.7 Toward this goal, optimization of side-chain length and bulk has emerged as an effective approach for designing materials that pack with tight π-π spacing and long-range order.8 Despite these advances, there are currently no general synthetic strategies for controlling the orientation of polymer packing in thin films.9 Most semiconducting organic materials display anisotropic charge transport, where transport occurs most efficiently in the direction of π-stacking.10 Thus, optimal device performance is often achieved when the π-stacking direction in a film aligns with the direction of intended

charge transport.11,12 For example, due to the horizontal arrangement of OFET electrodes, hole mobility of poly(3hexylthiophene) films is two orders of magnitude higher when the polymer chains reside edge-on to the substrate (with in-plane π-stacking), when compared to films composed of face-on aligned polymers.13 Conversely, since OPVs have vertically arranged electrodes, materials that favor face-on packing and out-of-plane π-stacking often demonstrate high PCEs.14 Although a preference for πstacking direction is observed with many polymers, the fundamental interactions that govern polymer alignment in thin films are not well understood.15,16 Since interpolymer and polymer-substrate interactions likely govern deposition mechanisms, we hypothesized that polymerpacking orientation might be influenced by synthetic tailoring of macromolecular structure, specifically backbone coplanarity and solubility.17 A high level of backbone coplanarity is a central design principle for developing top-performing semiconducting polymers.18 Structural planarity extends effective conjugation length and charge delocalization and, in the case of OPV operation, leads to a narrow band gap and improved light absorption.19,20 It also promotes favorable solid-state

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packing through strong π-π interactions that often result in high degrees of crystallinity.21,22 For these aforementioned reasons, most top-performing polymers are designed to display π-conjugated backbones with near complete coplanarity. Curiously, isoindigo-based polymers have recently been reported to demonstrate relatively high OFET23 and OPV24 performance, despite possessing non-coplanar, twisted backbones. Isoindigo-based materials, first reported by Reynolds,25 have quickly become popular for organic electronics due to their ease of synthesis. X-ray crystal analysis of isoindigo reveals that unfavorable steric interactions between the two indolinones lead to significant twist around the central C=C bond.26 Additionally, the steric effects of phenyl protons induce a torsional strain between isoindigo and adjacent π-systems.27 We hypothesized that replacing the fused phenyl rings of isoindigo with thiophenes would diminish the twist within the acceptor chromophore, while simultaneously reducing torsion with adjacent donor monomers.28 Herein, we report the synthesis of new thienyl derivatives of isoindigo, the resultant polymers, and the effects of increased polymer planarity on optoelectronics, nanostructural order, and device performance. Notably, we discovered that polymerpacking orientation can be controlled by tailoring backbone coplanarity and solubility to generate thin films that favor edge-on (in-plane π-stacking) or face-on (out-ofplane π-stacking) packing relative to the substrate.

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(3) was constructed similarly through the condensation of 6 with thieno-oxypyrrole (7), followed by bromination. We found this route for assembling 3 to be more reliable than the reductive dimerization of 6, which has been the strategy employed in previous synthetic reports.30 Intermediates 6 and 7 were derived from 8 by either cyclization with oxalyl chloride to provide 6, or acylation with 2chloro-acetylchloride, followed by Pd-catalyzed cyclization to furnish 7.31 Donor-acceptor copolymers incorporating these three acceptor comonomers (1-3) were synthesized by Pd2dba3/P(o-tol)3-catalyzed Stille polycondensation with donor comonomers bis(tri-methylstannyl)bithiophene or bis(trimethylstannyl)terthiophene (Figure 2). We found that using a slight excess of the dibrominated monomer (1.00:0.97 versus distannylated monomer) provided materials with higher molecular weight.32 Polymers were purified by Soxhlet extraction: highly soluble materials eluted in chloroform, while those with lower solubility were extracted with 1,1,2,2-tetrachloroethane (TCE). All bithiophene (2T) and terthiophene (3T) polymers were soluble when appended with 2-octyldodecyl (OD) side chains. Substitution of PII2T and PTII2T with 2hexyldecyl (HD) groups led to poorly soluble 11 and 16, respectively, while PTBII2T reached a solubility limit with 2-butyloctyl (BO) side chains (14). Likewise among HDsubstituted 3T-polymers, only PTBII3T-HD (19) was obtained in substantial yield (see SI). TBII-based copolymers were more soluble in all cases, likely due to a regiorandom structure compared to regiosymmetric II- and TIIpolymers. Although regiorandom polymers can display diminished thin-film nanostructural order, the use of a dissymmetric monomer may be an effective strategy for enhancing material processability.

Figure 1. Synthetic schemes of isoindigo (II, 1), thienobenzo-isoindigo (TBII, 2), and thieno-isoindigo (TII, 3) monomers. RESULTS AND DISCUSSION Syntheses of Monomers and Polymers. Isoindigo (II, 1) and its two related thienyl analogs, thieno-benzoisoindigo (TBII, 2) and thieno-isoindigo (TII, 3), were synthesized via acid-promoted condensation (Figure 1, see SI for detailed procedures). Based on published procedures, isoindigo (1) was obtained by condensing 6-bromoisatin (4) with 6-bromo-oxindole (5), followed by side-chain installation via alkylation.29 Similar coupling of 5 with thieno-dioxypyrrole (6), followed by alkylation and bromination, provided dissymmetric TBII (2). Symmetric TII

Figure 2. Synthetic scheme of II-, TBII-, and TII-polymers with bithiophene (2T, 10-16) or terthiophene (3T, 17-20) comonomers.


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Chemistry of Materials Figure 3. Normalized thin-film absorption spectra for ODsubstituted 2T- (10, 12, 15, solid lines) and 3T- (17, 18, 20, dashed lines) polymers spun from TCE.

Optical and Electrochemical Properties. In order to determine the effects of thienyl substitution on the optoelectronic properties of isoindigo-based polymers, UV-visNIR absorption spectra were obtained for each polymer in solution and thin film (Table 1, Figure 3). In films, PII2TOD (10) has a λmax at 705 nm and an absorption onset at 769 nm (Egfilm = 1.61 eV), PTBII2T-OD (12) has a λmax at 801 nm and onset at 905 nm (Egfilm = 1.37 eV), and

Table 1. Polymer molecular weight distributions and optoelectronic properties. Absorption Properties Polymer

Electrochemical Propera ties

Mn [kDa]

PDI

λmaxsol [nm]

λmaxfilm [nm]

Egsol [eV]

Egfilm [eV]

HOMO [eV]b

LUMO [eV]

EgCV [eV]c

59.0

1.4

708

705

1.64

1.61

-5.89

-3.98

1.91

Bithiophene Polymers PII2T-OD (10) PII2T-HD (11)

-

-

614

625

1.69

1.69

-5.85

-3.93

1.92

PTBII2T-OD (12)

55.9

1.5

781

801

1.41

1.37

-5.54

-3.84

1.70

PTBII2T-HD (13)

57.6

1.4

790

801

1.37

1.39

-5.43

-3.87

1.56

PTBII2T-BO (14)

-

-

749

769

1.33

1.37

-5.33

-3.86

1.47

PTII2T-OD (15)

-

-

900

889

1.15

1.12

-5.22

-3.78

1.44

PTII2T-HD (16)

-

-

897

903

1.10

1.09

-5.04

-3.82

1.22 1.83

Terthiophene Polymers

a

PII3T-OD (17)

77.0

1.3

628

634

1.68

1.64

-5.65

-3.82

PTBII3T-OD (18)

66.0

1.5

737

773

1.42

1.39

-5.30

-3.72

1.58

PTBII3T-HD (19)

-

-

743

781

1.41

1.39

-5.21

-3.72

1.49

PTII3T-OD (20)

-

-

829

826

1.25

1.23

-5.06

-3.77

1.29

Onsets, potentials vs. Fc/Fc+. b E = -5.13 eV – eEredox. c EgCV = e(Eox – Ered).

PTII2T-OD (15) has a λmax at 889 nm and onset at 1110 nm (Egfilm = 1.12 eV). Each successive thienyl substitution narrows the optical band gap (Eg) by ~ 0.2 eV, which leads to a redshift in λmax and absorption onset in all polymers (1020). This effect is likely due to the increased backbone coplanarity when phenyl groups are replaced by less sterically demanding thienyl moieties. Increased planarity extends effective conjugation length and delocalization of electron density, thereby narrowing the band gap (Eg). Interestingly, 3T- versus 2T-polymers display blueshifted absorption spectra. Although the extra thiophene per repeat unit is anticipated to enhance π-π interactions, spectroscopically it also appears to shorten effective conjugation length, possibly by increasing rotational freedom and disorder along the oligothiophene subunit.33 Evidence that thienyl versus phenyl substitution imparts higher coplanarity is also observed in the polymers’ solution absorption dependence on temperature (Figure S2-4). Strong π-π interactions, which are commonly exhibited by planar polymers, tend to promote redshifts in absorption due to solution-phase aggregation.34 When the temperature rises from 25 to 195°C, solutions of PTBII2TOD (12) and PTII2T-OD (15) show a blueshift in absorp-

tion that is consistent with high temperature dissolution of polymer aggregates. In contrast PII2T-OD (10), which exhibits a twisted backbone, does not appear to aggregate significantly at room temperature since there is no spectral change upon heating. This difference in absorption/aggregation behavior is one of several polymer properties that result from varying backbone coplanarity. The electrochemical properties of the polymers were characterized via cyclic voltammetry (CV) versus the Fc/Fc+ redox couple (Table 1). Among the OD-substituted 2T-polymers, the HOMO/ LUMO levels of PII2T-OD (10) are -5.89 and -3.98 eV, PTBII2T-OD (12) has levels of 5.54/-3.84 eV, and PTII2T-OD (15) has levels of -5.22/-3.78 eV. Although there is little difference between polymer LUMO levels, the HOMO levels rise significantly with each substitution of an electron-rich thiophene for a phenyl ring. The increase in the HOMO levels results in a decrease in the electrochemical band gaps (EgCV), providing 1.91 eV for 10, 1.70 eV for 12, and 1.44 eV for 15. In addition, minimization of side-chains (12 13 14) and replacement of bithiophene with terthiophene (12 vs. 18) cause the HOMO levels to rise. In both cases, band gap



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decreases are likely the result of increased charge delocalization via enhanced interchain π-π interactions.

Figure 4. Structures of (a) II-, (b) TBII-, and (c) TII-based trimers with mean average twist and torsion angles based on DFT calculations at the B3LYP/6-31G(d) level. Side-view of energy-minimized conformers depict the varying degrees of backbone coplanarity.

Figure 5. OFET output and transfer curve characteristics for films of OD-substituted polymers: a) PII2T (10), b) PTBII2T (12), c) PTII2T (15), d) PII3T (17), e) PTBII3T (18), f) PII3T (20). All transfer curves were obtained with a source-drain voltage (VDS) of -40 V.

Theoretical Calculations. Density functional theory (DFT) calculations of methyl-substituted trimers were used to simulate each polymer backbone and determine

the effects of thienyl substitution structural conformation. Computation was performed with Gaussian 09 using a hybrid B3LYP correlation functional and 6-31G(d) basis


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set. Due to the dissymmetry of TBII, we acquired data for all eight possible trimers in order to account for the regiorandom structure of TBII-polymers (see SI). Geometry-optimized structures reveal that increasing backbone coplanarity correlates with increasing thiophene substitution in the indigoid acceptor subunit (Figure 4). A single phenyl-thiophene substitution causes the intra-acceptor twist (θ1) to dramatically decrease from 12.82° with II2T to 0.11° with TBII2T. When the acceptor core contains two thienyl moieties, steric repulsion is further minimized and TII2T approaches complete coplanarity (0.03°). Additionally, thienyl substitution decreases the inter-monomer torsion (θ2) between the acceptor subunits and neighboring oligothiophenes. II2T displays average inter-monomer torsion angles of 20.73°, while the average torsion of TBII2T decreases to 12.38°, and TII2T is nearly coplanar with adjacent π-systems (0.45°). Notably, the systematic twist and torsion within each repeat unit of II2T cause its backbone to adopt a slight helical conformation. Since polymer backbone coplanarity often enhances π-π interactions, film crystallinity, and charge carrier transport, TBII- and TII-based materials appear promising for OFET and OPV applications. Field-Effect Transistor Fabrication and Characterization. Bottom gate/top contact devices were fabricated by spin-coating polymer solutions from TCE (3 mg/mL) onto octyltrichlorosilane (OTS)-treated SiO2 (300 nm)/n++-Si substrates, followed by thermal deposition of Au electrodes through a shadow mask (see SI). The highest performing devices were thermally annealed at 200°C for 30 minutes prior to top contact deposition, followed by vacuum annealing at 110°C (~ 1 mbar) for 5 days. All fabrication, annealing, and device testing was performed under inert conditions. Output curves of OFETs fabricated with 10-20 display saturation under p-channel operation with high Ion/Ioff ratios (≥ 105) and significant hole mobility (Figure 5). Films of PII2T-OD (10), PTBII2T-OD (12), and PTII2T-OD (15) provide hole mobilities (μh) of 0.46, 4.7 × 10-2 and 0.29 cm2 V-1 s-1, respectively (Table 2). Threshold voltages (Vth) are found to be proportional to the energy difference between electrode work function and polymer HOMO levels, where 10 has a Vth of -29 V, 12 has a Vth of -5 V, and 15 has a Vth of +10 V. These data infer that 10, due to a low HOMO level, requires a negative gate voltage to promote

hole accumulation at the dielectric interface, while the high HOMO level of 15 necessitates a positive gate voltage for transporting holes effectively. Despite the anticipated benefits of polymer planarity on OFET performance (i.e. tight π-π interactions, extended effective conjugation length),35 these hole mobilities do not correlate with backbone coplanarity. Instead, higher field-effect mobility is achieved with regiosymmetric polymers PII2T-OD (10) and PTII2T-OD (15). In this study it appears that polymer regiosymmetry, rather than backbone coplanarity, plays a significant role in facilitating charge transport, likely by increasing thin-film crystallinity and long-range order. Minimization of side chains (11, 13, 14, 16) and introduction of 3T as a comonomer (17-20) were also explored; however, none of these analogs showed improvement in OFET performance. Lower hole mobilities with 3T-polymers (17-20) are likely due to more grain boundary defects, as indicated by increased film roughness (vide infra), which can severely hinder charge transport. From analysis of the transfer curves, hysteresis is greater for II-polymers (10, 11, 17) than for more planar TBII(12 - 14, 18, 19) and TII- (15, 16, 20) polymers, regardless of side-chain substitution and processing conditions (Figure 5). Given that all devices were fabricated with the same dielectric layer within a controlled atmosphere (thereby minimizing contamination from impurities, water, or oxygen), we postulate that the increased hysteresis with IIpolymers is due to intrinsic polymer structure. The helical, twisting structures of II-polymers likely lead to increased structural disorder and charge trapping defects at the polymer-dielectric interface.36, 37 Photovoltaic Device Fabrication and Characterization. Bulk-heterojunction (BHJ) OPV devices were fabricated using polymers 10-15 and 17-20 as electron donors and [6,6]-phenyl-C71-butyric acid methyl ester (PC71BM) as the electron acceptor. Optimal performance was obtained with OPVs having the device architecture: ITO/PEDOT:PSS(40 nm)/polymer:PC71BM/LiF (1 nm)/Al(100 nm). The films were spun from 1,2dichlorobenzene, in blend ratios between 1:2 and 1:3 (polymer:fullerene, see SI). Devices fabricated with 10, 12, and 15 demonstrate average PCEs of 1.60%, 3.20%, and 0.40% and open-circuit voltages (VOC)

Table 2. Organic field effect transistor (OFET) and photovoltaic (OPV) device characteristics. OFET characteristics Polymer

OPV characteristics

μh[cm2/V·s]a

VOC [V]

JSC [mA cm2 ]

-29

0.46 (0.93)

0.65

-38

4.0 × 10-3 (5.0 × 103 )

0.36

-5

4.7 × 10-2 (8.8 × 102 )

Ion/Ioff

Vth [V]

PII2T-OD (10)

106

PII2T-HD (11)

103

PTBII2T-OD (12)

105

PTBII2T-HD (13)

105 - 106

FF

PCE [%]a

-4.82

0.51

1.60 (1.72)

-3.05

0.40 0.43 (0.44)

0.54

-10.40

0.57

3.20 (3.39)

Bithiophene Polymers

PTBII2T-BO (14) PTII2T-OD (15)

-9

5.1 × 10-2 (8.3 × 10-2)

0.52

-13.55

0.57

4.02 (4.27)

5

6

-14

1.7 × 10-2 (2.2 × 10-2)

0.45

-7.61

0.41

1.40 (1.50)

5

6

+10

0.29 (0.43)

0.19

-3.97

0.53 0.40 (0.42)

10 - 10 10 - 10


Chemistry of Materials PTII2T-HD (16)

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

106 – 107

+14

0.11 (0.18)

-

104 – 105

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-

-

Terthiophene Polymers -14

1.5 × 10-2 (3.8 × 10-2)

0.69

-11.45

0.44 3.48 (3.79)

PTBII3T-OD (18)

6

10

-8

6.2 × 10-2 (0.18)

0.49

-14.66

0.66 4.74 (4.79)

PTBII3T-HD (19)

106

-6

3.6 × 10-2 (4.7 × 102 )

0.42

-10.83

0.51

PTII3T-OD (20)

106

+11

7.5 × 10-2 (0.13)

0.18

-5.40

0.45 0.44 (0.47)

PII3T-OD (17)

a

2.32 (2.46)

Reported values are an average of at least 8 devices. Data in parentheses are maximum values.

of 0.65, 0.54 and 0.19, respectively (Table 2, Figure 6). Although backbone coplanarity does not trend with these efficiencies, it does appear to correlate with VOC.38 It has been shown that VOC is greatly influenced by polymer HOMO levels, which are highest with TII-polymers since they are the most planar and electron-rich polymers.39 Accordingly, the high HOMO levels of TII-polymers (15, 20) lead to the smallest VOC values and some of the poorest performing OPV devices. Side chain effects on PCE were also examined for PII2T and PTBII2T by fabricating devices with 2-hexyldecyl (HD, 11, 13) and 2-butyloctyl (BO, 14) groups. OPVs based on 13 achieve the highest PCEs (up to 4.27%) among 2Tpolymers, with short-circuit currents (JSC) of -13.55 mA cm-2. The limited processabilities of PII2T-HD (11) and PTBII2T-BO (14) led to poor films and devices. Overall, enhanced PCEs of PTBII2T-OD (12) and PTBII2T-HD (13) are due to greater fill factors (FF, 0.57 vs. 0.51) and JSC values that more than double those of PII2T-OD (10). Terthiophene (3T) polymers (17-20) exhibit even greater PCEs than the 2T-polymers (10-16); PII3T-OD (17) achieves an average PCE of 3.48% (compared to 1.60% for 10) and PTBII3T-OD (18) demonstrates an average PCE of 4.74% (compared to 3.20% for 12). PII3T-OD (17) again displays larger VOC values than PTBII3T polymers (18, 19) but suffers from lower JSC and FF. Since TBII-polymers demonstrate the highest OPV performance among the polymers studied, TBII appears to be a promising subunit for future materials development. The greater backbone coplanarity of TBII- versus II-polymers leads to improved optoelectronic properties without sacrificing too much VOC like TII-polymers. Additionally, it is likely that the increased planarity of TBII-polymers enhances OPV performance through improved film morphology. Thin-Film Morphology. In order to determine the effects of polymer planarity on nanoscale topography, polymer films were studied via atomic force microscopy (AFM). For comparison to devices, films were spun to model OFET (neat on OTS-treated SiO2) and OPV (blend with PC71BM on PEDOT:PSS-coated ITO) architectures (Figure 7). For OFETs to achieve high charge mobility it is favorable for thin films to have large interconnected domains in order to minimize crystallite grain boundaries.40 Among the OFET films of OD-

Figure 6. Characteristic J-V curves for BHJ OPV devices using blend films containing 2T- (10, 12, 13, 15, above) and 3T(17, 18, 20, below) polymers.

substituted polymers, PII2T-OD (10) and PTII3T-OD (20) display the largest visible domains. Additionally, it is observed that II- (10, 17) and TII- (15, 20) polymers generate films with higher roughness (RRMS) than their analogous TBII-polymers (12, 18). Although II- and TII-polymers vary greatly in backbone coplanarity, their regiosymmetric structure appears to favor well-ordered polymer packing and the formation of large crystalline domains, which are beneficial for field-effect hole mobility. In contrast, large domains in OPV films are undesirable since they can severely hinder exciton dissociation if their dimensions exceed average exciton diffusion lengths (~10 nm).41 Therefore ideal OPV film morphology is smooth with nanoscale features. Among the polymers studied, TBII-polymers 12, 13 and 18 form the smoothest OPV films (RRMS < 1.89 nm) and display some of the smallest visible domains, thereby suggesting that these polymers have the finest intermixing with PC71BM (see SI). In agreement with these morphological data, polymers 12, 13 and 18 are some of the highest performing OPV materials in this study. Thin-Film Nanostructural Order. Grazing incidence X-ray diffraction (GIXD) analysis can be used to determine the extent of cofacial π-stacking, crystallinity, and polymer-packing orientation relative to the substrate.42 It is often crucial to measure these solid-state parameters since they can have major implications on device perfor-


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mance. For example, a shorter cofacial π-π distance reduces the energetic barrier for charge hopping between polymer chains, thereby promoting charge carrier transport.43 In order to model OPV and OFET devices, samples were prepared by spin-coating neat films of polymers (10-20) on PEDOT:PSS-coated or OTS-treated SiO2 wafers. Polymer:PC71BM blends were also spun on PEDOT:PSS, but the resultant films were relatively amorphous compared to those consisting of neat polymer.

ing to π-stacking distances of 3.60 and 3.51 Å, respectively. Lamellar spacing was found to be inversely related to π-π spacing, where the most planar polymers displayed the largest lamellar packing distances. Long-range packing order is quantitatively measured by determining correlation length (LC), which is defined as the length over which a crystalline structure is preserved.44 Large LC corresponds to narrow diffraction peak breadth, which is the result of reduced variability in the polymer chain position and crystallite rotation. It should be noted that polymer properties other than π-stacking (i.e. regiosymmetry, molecular weight) likely contribute to π-π LC, which can explain why coplanarity-based analyses do not fully describe trends in LC values (Table 3). Nevertheless it is apparent that films consisting of 3Tversus 2T-polymers have larger π-π LC values, likely due to increased π-π interactions (12 vs. 18 and 15 vs. 20). Additionally, nearly every polymer film exhibits higher LC values when spun on a low energy OTS surface rather than polarizable PEDOT:PSS. Since OTS treatment often minimizes polymer-substrate interactions, we postulate these enhanced LC values result from π-π interactions dominating film formation, thereby allowing polymers to pack with longer-range order. Table 3. Polymer packing parameters based on GIXD. π-π spacing [Å] Polymer

Lamellar LC [nm]

PII2T-OD (10)

3.67

1.49

PTBII2T-OD (12)

3.64

1.28

21.32

1.28

PTII2T-OD (15)

3.59

2.01

22.63

7.97

PTBII3T-OD (18)

3.56

2.89

19.48

7.52

PTII3T-OD (20)

3.60

2.54

20.44

7.16

19.73

15.15

Amorphous

Polymer

Films on OTS (OFET model)

PII2T-OD (10)

3.63

3.08

20.47

8.13

PTBII2T-OD (12)

3.58

2.17

22.60

12.14

PTII2T-OD (15)

3.55

1.92

22.39

9.38

PII3T-OD (17)

Polymer packing distances are revealed in thin-film diffraction patterns through the identification of peaks that correspond to either cofacial π-stacking (π-π spacing, q ~ 1.7 Å-1) or side-chain packing (lamellar spacing, generally q ~ 0.3 Å-1). Neat films of PII2T-OD (10), PTBII2T-OD (12) and PTII2T-OD (15) spun on PEDOT:PSS show π-stacking distances of 3.67, 3.64 and 3.59 Å, respectively (Table 3). Films spun on OTS also display this packing trend since greater backbone coplanarity often promotes stronger π-π interactions and tighter cofacial polymer packing. Stacking distances are further minimized with PTBII2T-HD (13) and PTBII2T-BO (14) by decreasing side-chain size, lead-

Lamellar spacing [Å]

Films on PEDOT:PSS (OPV model)

PII3T-OD (17)

Figure 7. AFM images of OD-substituted polymer films spun ++ on OTS-treated SiO2/n Si (OFET, left) and polymer:PC71BM blends spun on PEDOT:PSS-coated ITO (OPV, right). The RRMS for each film is indicated in the lower left corner.

π-π LC [nm]

Amorphous

PTBII3T-OD (18)

3.54

4.92

20.04

4.33

PTII3T-OD (20)

3.55

5.21

20.32

12.09

If thin films are ordered, GIXD analysis can also reveal whether the polymer chains align edge-on or face-on relative to the substrate. Polymers that favor edge-on packing will show more intense π-π spacing peaks near the qxy plane (qz ≈ 0) due to in-plane π-stacking. Similarly, films that favor face-on packing will have out-of-plane πstacking and show more intense π-π spacing peaks along the qz plane (qxy ≈ 0). GIXD patterns of PII2T-OD (10) show a strong π-stacking peak near the qxy plane, which corresponds to a preference for edge-on packing (Figure 8). PTBII2T-OD (12) and PTII2T-OD (15) both display intense peaks near the qz plane, suggesting that these polymers favor face-on packing. Notably, these polymers



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show a trend where higher backbone coplanarity appears to promote face-on polymer alignment. This finding is significant since previous studies of polymer-packing orientation have focused on the effects of solubilizing sidechains, and not on π-conjugated backbone structure. Delongchamp et. al. recently observed that the orientation of diketopyrrolopyrrole (DPP)-based polymers could be modulated by varying side-chain density.45 Solubility appeared to control polymer-packing orientation, where polymers of higher solubility favored face-on packing and less soluble polymers favored edge-on packing. It was proposed that low polymer solubility led to aggregation in solution that, upon spin coating, formed a film composed of edge-on oriented polymers. Yet when data from our current study are considered, it is evident that solubility is not the sole factor that governs polymer-packing orientation. Within this work, polymers with high backbone coplanarity (12, 15), despite showing moderate levels of aggregation in solution (vide supra), favor face-on orientation. Contrarily PII2T-OD (10), which has a twisted, helical backbone, displays a preference for edge-on packing despite being highly soluble.

Figure 8. 2D GIXD patterns of neat OD-substituted polymers spun on PEDOT:PSS-coated (a-f) and OTS-treated substrates (g-l).

In order to integrate planarity effects into our understanding of polymer-packing orientation, we consider the forces that control thin-film growth.46 If a planar, π-

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conjugated polymer (12, 15) is highly soluble (i.e. interpolymer interactions are minimal), face-on packing is energetically favored due to strong Van der Waals interactions between the rigid π-conjugated backbone and substrate surface. Cofacial π-π interactions subsequently guide other polymer chains to adopt a similar alignment, thereby creating a film composed of face-on oriented polymers. In contrast, there is negligible energy benefit for a less planar polymer (10) to adopt such a face-on orientation since the twisting backbone precludes strong polymersubstrate interactions. In these cases interpolymer interactions (i.e. π-π and/or lamellar) dominate and promote aggregation that, upon deposition, form domains composed of edge-on oriented polymer chains. With this understanding, it is anticipated that relatively planar polymers (12, 15) will be induced to favor edge-on packing by tailoring their structures to have stronger interpolymer interactions. When OD-side chains are replaced by smaller HD-groups, the strong face-on packing preference of 12 erodes to a mixture of packing orientations with PTBII2T-HD (13), as revealed by the introduction of a minor peak in the qxy plane (see SI). Decreasing side-chain size to BO-groups leads to disappearance of the peak along the qz plane and appearance of a new peak along the qxy plane, indicating that PTBII2T-BO (14) strongly favors edge-on packing. This side-chain effect on packing orientation is observed with TII-polymers too, as PTII2T-HD (16) shows a strong preference for edge-on packing in contrast to the face-on orientation of PTII2TOD (15). Edge-on packing of polymers can also be promoted by enhancing π-π interactions via replacement of 2T with 3T. Films of PTBII3T-OD (18) show a new peak in the qxy plane (compared to 12) that corresponds to a mixture of packing orientations, while PTII3T-OD (20) demonstrates a clear preference for edge-on packing with an intense peak in the qxy plane. Significantly, these GIXD data support the observed trends in performance of polymer-based devices. The highest field-effect mobility in this study is obtained with II-polymer 10, likely benefitting from the edge-on packing of the twisted, non-coplanar polymer backbone. Conversely, the highest OPV efficiencies are obtained with TBII-polymers 13 and 18 that are more planar and display greater degrees of face-on polymer packing. CONCLUSIONS We have synthesized a series of π-conjugated polymers that display a range of backbone coplanarity by systematically replacing phenyl rings with thiophenes in isoindigo (II) conjugated subunits. In agreement with DFT calculations, polymers with more thienyl subsitution display greater backbone coplanarity, red-shifted absorption profiles and narrower optoelectronic band gaps. In OFETs PII2T-OD (10), despite having a low degree of backbone coplanarity, provides the highest hole mobility of 0.93 cm2 V-1 s-1. In OPVs, the highest power conversion efficiencies (up to 4.7%) are obtained with TBII-polymers (12, 13, 18). Upon AFM and GIXD analysis of these polymer thin films, we have discovered that highly soluble, planar polymers (12, 15) show a preference for face-on orienta-


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tion, while non-planar polymers (10) favor packing with an edge-on orientation. Planar polymers can be induced to adopt significant degrees of edge-on packing by reducing solubility, either through minimizing side-chain size (13, 14, 16) or enhancing π-π interactions (18, 20). Altogether, this work advances our understanding for how polymer-packing orientation is influenced by the interplay of backbone coplanarity and solubility. These key structure-property relationships promise to aid in the development of π-conjugated materials with rationally designed solid-state properties and enhanced electronic device performance.

ASSOCIATED CONTENT Supporting Information. Experimental details, syntheses of monomers and polymers, device fabrication, and complete characterization of materials and devices. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected], [email protected], or [email protected].

ACKNOWLEDGMENTS This work was supported in part by King Abdullah University of Science and Technology (KAUST) through the Center for Advanced Molecular Photovoltaics (CAMP) under Award No. KUS-C1-015-21, and the Fréchet “various gifts” fund for the support of research in new materials. Portions of this research were carried out at the Stanford Synchrotron Radiation Lightsource user facility, operated on behalf of the U.S. Department of Energy, Office of Basic Energy Sciences. M.S.C. thanks the Camille and Henry Dreyfus Postdoctoral Program in Environmental Chemistry for a fellowship; the assistance of Jessica C. Moreton is acknowledged with thanks.

ABBREVIATIONS OFET, organic field-effect transistor; OPV, organic photovoltaic; PCE, power conversion efficiency; TCE, 1,1,2,2tetrachloroethane; OD, 2-octyldodecyl; HD, 2-hexyldecyl; BO, 2-butyloctyl; Eg, band gap; CV, cyclic voltammetry; Fc, ferrocene; HOMO, highest occupied molecular orbital; LUMO, lowest unoccupied molecular orbital; DFT, density functional theory; OTS, octyltrichlorosilane; Vth, threshold voltage; BHJ, bulk heterojunction; PEDOT:PSS, poly(3,4ethylenedioxy-thiophene) poly(styrenesulfonate); ITO, indium tin oxide; VOC, open-circuit voltage; JSC, short-circuit current; FF, fill factor; AFM, atomic force microscopy; RRMS, root mean square roughness; GIXD, grazing incidence X-ray diffraction; LC, correlation length.

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