Ambipolar Charge Transport in Isoindigo-Based Donor–Acceptor

Article pubs.acs.org/cm

Ambipolar Charge Transport in Isoindigo-Based Donor−Acceptor Polymers Romain Stalder,† Sreenivasa Reddy Puniredd,‡,¶ Michael Ryan Hansen,§,∥ Unsal Koldemir,† Caroline Grand,⊥ Wojciech Zajaczkowski,‡ Klaus Müllen,‡ Wojciech Pisula,*,‡,# and John R. Reynolds*,†,⊥ †

The George and Josephine Butler Polymer Research Laboratory, Department of Chemistry, University of Florida, Gainesville, Florida 32611-7200, United States ‡ Max Planck Institute for Polymer Research, Ackermannweg 10, 55128 Mainz, Germany § Interdisciplinary Nanoscience Center (iNANO) and Department of Chemistry, Aarhus University, Gustav Wieds Vej 14, DK-8000 Aarhus C, Denmark ∥ Institute of Physical Chemistry, Westfälische Wilhelms-Universität Münster, Corrensstr. 28/30, D-48149 Münster, Germany ⊥ School of Chemistry and Biochemistry, School of Materials Science and Engineering, Center for Organic Photonics and Electronics, Georgia Institute of Technology, Atlanta, Georgia 30332-0400, United States # Department of Molecular Physics, Lodz University of Technology, Zeromskiego 116, 90-924 Lodz, Poland ¶ Institute of Materials Research and Engineering (IMRE), Agency for Science, Technology and Research (A*STAR), 2 Fusionopolis Way, Innovis, 138634 Singapore S Supporting Information *

ABSTRACT: A series of donor−acceptor isoindigo (iI)-based copolymers synthesized with increasing numbers of thiophene rings in the repeat unit (from zero to three thiophene rings, including silole and germole-bridged fused bithiophene units) is applied toward solutionprocessed OFET devices. Differential pulse voltammetry on thin films of the polymers recorded LUMO energy levels confined within a 0.1 eV range around 3.9 eV, while their bandgaps are estimated at 1.5 to 1.7 eV. The interchain π-stacking distance of each sample was evaluated from the 2D-WAXS diffraction patterns of annealed extruded filaments and the GIWAXS patterns of thin films, and were found to be all in the same range, between 3.65 and 3.75 Å for the thin films. Both p-type and n-type charge transport in thin film bottom gate, bottom contact transistor devices were recorded. In particular, the copolymer P(T-iI) containing one thiophene ring afforded wellbalanced ambipolar p-type and n-type mobilities of 0.04 cm2/(V s) and 0.1 cm2/(V s), respectively. Under our processing conditions, the charge transport properties evolved from exclusively n-type to solely p-type as the number of thiophene rings within the repeat unit is increased to three rings in the case of P(T3-iI). This was observed despite all polymers displaying similar LUMO energy levels, interchain π-stacking distances, and microscopic thin film morphology (all face-on arrangement on the dielectric surface). This prompted a molecular-scale morphological analysis of P(T-iI) and P(T3-iI) in particular, using solid-state NMR spectroscopy in order to further investigate the stark difference in n-type mobilities between these two polymers. Using the complete assignment of solution 2D-NMR spectra of a thiophene-iI-thiophene model compound as guideline, the analysis of proton−carbon correlations in the solid-state 2D 13C{1H} FSLG-HETCOR NMR spectra of P(T-iI) and P(T3-iI) revealed differences in the molecular environment surrounding each iI unit. The latter suggests a stronger correlation of neighboring iI units in P(T-iI), whereas a stronger intermixing of iI and thiophenes prevails in P(T3-iI). We conclude that, in this study, the choice of the donor unit length within the primary structure of the D−A polymer can be responsible for hindering its n-type character.

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INTRODUCTION

technology has evolved from almost exclusively p-type organic transistors to involving n-type transistors as well,3 although the latter type remains less-common among the solutionprocessable polymeric materials. This discrepancy was in part

Organic field-effect transistors (OFETs) have achieved high performance thanks to efforts to understand and control the requirements for efficient charge carrier transport. Both positive (p-type, holes) and negative (n-type, electrons) charge carriers can be generated in the OFET channel1,2 depending, in particular, on the energy of the frontier molecular orbitals (FMOs) of the conjugated material in the active layer. The © 2016 American Chemical Society

Received: October 10, 2015 Revised: January 24, 2016 Published: February 17, 2016 1286

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with high ionization potential (IP)) and bandgaps reduced to 1.5 eV, iI-based materials can display n-type and potentially ambipolar charge transport in OFETs. To this date, this has been documented using all-acceptor copolymers of iI conjugated with other common acceptors36 or with modified iI cores functionalized with halogen atoms at the 7,7′ positions.37 The latter approach significantly stabilized the EA of the corresponding D−A polymer and afforded ambipolar charge transport with p-type mobility exceeding 1 cm2/(V s) and n-type mobility approaching 0.5 cm2/(V s).38 The average electron mobility in the nonfluorinated equivalent was 2 orders of magnitude lower. Structural modification of the iI unit has thus been successfully employed to demonstrate ambipolar charge transport, but there is still limited knowledge of the ntype and ambipolar performance of D−A copolymers based on the original iI core. Since a great deal of iI’s attractiveness is its synthetic accessibility and scalability, further understanding of its n-type properties in OFETs is necessary. Herein, we synthesize a family of D−A polymers (Figure 1) based on iI, as well as its homopolymer, to be studied in

due to the lack of organic materials with electron affinities (EAs) high enough (deep LUMOs) to enable efficient electron injection and migration across the channel. Tailoring the energy of the frontier molecular orbitals using electron-deficient moieties to stabilize the LUMO eventually led to conjugated systems favoring n-type mobility over its p-type counterpart.4 This has been achieved most efficiently for solution-processed devices with conjugated polymers based on the donor− acceptor (D-A) approach, where electron-withdrawing functional groups such as imides5,6 or fluorocarbon substituents7−9 render the aromatic units in the backbone sufficiently electrondeficient to favor electron-injection and transport. In particular, conjugated polymers based on benzobisimidazobenzo-phenanthroline (BBL),10,11 diketopyrrolopyrrole (DPP),12,13 bithiophene imide (BTI),14,15 or the more extended perylene diimides (PDI),16−18 naphthalene diimides (NDI),18−20 bisindenofluorenes,21 and dithienocoronene-diimide22 have been reported with high n-type mobility, in some cases exceeding 0.1 cm2/(V s). In a material where both HOMO and LUMO levels are energetically favorable for the injection of positive and negative charges, respectively, into a transistor channel,23,24 ambipolar charge transport can take place depending on the applied bias. Both types of charge transport occurring in a single material is particularly attractive for the fabrication of simple organic (solution processed) complementary metal-oxide semiconductor (CMOS)-like devices.25 The CMOS-inverter capability allows lower power consumption and faster switching times: key features of technologically viable devices. OFET operation in ambient atmosphere is an essential requirement for technological viability, as the latter can suffer from the environmental instability of the generated charge carriers, particularly under n-type operation: the generated negative charges can react with atmospheric oxidants such as water or oxygen if the energy of the LUMO is below a threshold set at ca. 4 eV.26 Fortunately, since this is dependent on the energy of the material’s FMOs, it can be synthetically tailored using the D−A approach. Efficient charge transport in polymer OFETs is also dependent on the polymer chain organization within the transistor channel, which is a function of processing conditions and the chemical structure of the repeat unit. Polymers with carefully chosen repeat unit structures have allowed extensive intermolecular interactions between the polymer chains, with interchain π-stacking distances as low as 0.35 nm.27 This has been correlated with high charge carrier mobilities since increased ordering in between polymer chains creates suitable pathways for charge transport.28 The local arrangement of the stacked polymer chain lamellae on the surface of the channel substrate can range between face-on to edge-on, with various levels of disorder in between, which also influences the overall charge mobility and is a key parameter to monitor during device fabrication. Isoindigo (iI)-based compounds are one of the latest imidebased families of high performance materials for organic electronics applications. This class of acceptors is rapidly developing since their first report in the open literature,29 driven by their high performance in organic photovoltaics and OFETs.30,31 An extensive set of polymeric D−A conjugated materials have already been reported for transistor applications.32,33 In particular, the effect of backbone curvature32 and side-chain design33,34 has been studied and optimized, resulting in high p-type mobility reaching 2 cm2/(V s) in air-stable devices.33−35 With FMOs generally deep (high EA combined

Figure 1. Structures of the repeat units for each D−A polymer and that of the homopolymer of isoindigo used in this study.

transistor applications, aiming at both p-type and n-type charge transport. The electron-rich portion of the repeat unit is based on thienyl and bridged bithienyl units with increasing lengths. We report these polymers’ optical and electrochemical properties, and investigate their propensity for chain packing in the solid state using 2D and grazing incidence wide-angle Xray scattering (2D- and GIWAXS). Used in solution-processed OFETs, the measured charge carrier transport within the polymer series evolves from exclusively n-type for the homopolymer to exclusively p-type for P(T3-iI) under the device architecture studied. Importantly, we demonstrate ambipolar charge transport for the other D−A polymers, with well-balanced charge transport between p- and n-type up to 0.1 cm2/(V s) for P(T-iI). To further investigate the difference between the high n-type mobility of P(T-iI) and the lack thereof in P(T3-iI), we employ 2D solid-state NMR spectroscopy to probe the molecular scale morphology of both polymers in the bulk. Guided by the complete assignment of the 2D-NMR spectra of a T-iI-T model compound (described in the Supporting Information), the proton−carbon correlations in the solid state technique suggest a different molecular environment of iI units consistent with higher electron transport properties in P(T-iI) compared to P(T3-iI). 1287

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Chemistry of Materials Table 1. SEC Data, Bandgaps, and FMO Energies from Optical and Electrochemical Characterization polymer P(iI) P(T-iI) P(T3-iI) P(DTS-iI) P(DTG-iI)

Mn (kDa) 29 26 24 40 59

Đa 2.4 2.1 1.9 3.8 2.5

λmax sol. (nm) 690 694 645 719 731

λmax film (nm) 687 694 660 728 742

ΔEoptb (eV)

Eox/EHOMOc (V)/(eV)

Ered/ELUMOc (V)/(eV)

ΔEechem (eV)

1.70 1.61 1.58 1.54 1.50

d

−1.26/3.84 −1.24/3.86 −1.21/3.89 −1.15/3.95 −1.12/3.98

d 1.71 1.50 1.60 1.49

5.54 0.47/5.57 0.29/5.39 0.45/5.55 0.37/5.47

Polymer dispersity. bOptical bandgap calculated from the onset of absorption in a neutral polymer film. cEnergies calculated from a calibration of the Fc/Fc+ redox couple at 5.1 eV. dNo stable oxidation recorded for this polymer.

a

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RESULTS AND DISCUSSION Polymer Synthesis. The first polymer used in this study, the homopolymer of isoindigo, P(iI), is synthesized by the Suzuki cross-coupling of the dibrominated iI unit with its diborylated derivative as was previously reported.39 Each unit is N-alkylated with 2-hexyldecyl (HD) side chains for solubility purposes. The other four thiophene-containing D−A copolymers displayed in Figure 1 are synthesized by palladiumcatalyzed Stille cross-couplings of alkylated iI moieties with the corresponding distannylated thiophene-based comonomers. The preparation of P(T-iI) is based on the polycondensation of 6,6′-dibromoisoindigo with 2,5-bis(trimethyltin)thiophene, whereas the same iI monomer copolymerized with 5,5″bis(trimethyltin)-3,3″-dihexyl-2,2′:5′,2″-terthiophene affords the P(T3-iI) repeat unit. The ditin-terthiophene monomer is purified by preparative HPLC prior to polymerization to ensure optimal stoichiometry, as detailed in the Supporting Information (Figure S1). The 2-ethylhexyl (EH)-functionalized copolymer of iI and dithienosilole P(DTS-iI), for which we recently reported the synthesis,40 is prepared as an effort to planarize the ring system of the electron-rich moiety, thanks to the bridging silicon atom of the DTS unit. The silicon atom in P(DTS-iI) was substituted for a germanium atom in the bridged dithienogermole (DTG) unit,41 yielding the copolymer of iI and DTG following the same procedure. After completion of the polymerizations described above, each reaction mixture is treated with a palladium scavenger and subsequently precipitated in methanol. The solids are purified using a Soxhlet apparatus with a sequence of solvents ranging from hexanes to chloroform as described in the Supporting Information. For each polymer, the material from the chloroform fraction is also found to be soluble in THF, toluene, and chlorinated solvents in at least 8 mg/mL concentration for each solvent, and is used by casting from 2 mg/mL solutions in chloroform for devices used in the following study. Size exclusion chromatography is employed to estimate the molecular weight of each polymer, using THF as mobile phase and polystyrene standards. The molecular weights of all samples are found to be in the 24 to 59 kDa range with dispersities (Đ) between 1.9 and 3.8, as summarized in Table 1. For each sample, this method only compares a retention volume equivalent to known polystyrene samples. Since the hydrodynamic volume of fully conjugated polymers is different from that of polystyrene (given the rigidity of conjugated polymers backbones), the values reported here may be overestimated in light of recent size-exclusion quantification studies on poly(3-hexylthiophene).42 Optical and Redox Properties. In chloroform solution, the UV−vis spectra of all polymers (Figure S2) display most of the absorption in the low-energy region of the visible spectrum between 550 and 800 nm, which is consistent with experimental43 and theoretical44,45 results for such iI-co-

oligothienylene repeat units. The solid state UV−vis absorption of each polymer was recorded on films cast onto glass slides, as displayed in Figure 2. Again, all spectra display two absorption

Figure 2. Normalized solid-state UV−vis absorption spectra of the five polymers (P(iI), black stars; P(T-iI), green squares; P(T3-iI), blue diamonds; P(DTS-iI), yellow triangles; and P(DTG-iI), red circles) as thin films spray-coated onto glass slides from a ca. 2 mg/mL toluene solution. Solid-state onsets of absorption in the near-IR are used to calculate optical bandgaps.

bands with little deviation from the solution absorption, between 400 and 500 nm for the high-energy bands and between 600 and 750 nm for the low-energy ones. Compared to the homopolymer P(iI), the introduction of thienyl electronrich moieties into the backbone repeat units shifts the lowenergy onsets of absorption to longer wavelengths from 733 nm for P(iI) up to 826 nm for P(DTG-iI). Accordingly, the optical bandgaps for the series of polymers vary from 1.70 to 1.50 eV as summarized in Table 1. The greatest maximum absorption red-shift is found for P(DTG-iI), at 14 nm longer wavelength than for P(DTS-iI). This effect of substituting germanium for silicon in the bithienyl bridge is consistent with that observed for the D−A copolymer based on other acceptors such as thienopyrrolodione (TPD).41 The redox properties of each polymer were investigated using cyclic and differential pulse voltammetry (CV and DPV) on thin films of the materials dropcast onto Pt-button electrodes in 0.1 M TBAPF6 in acetonitrile under inert atmosphere. All potentials were calibrated against the ferrocene/ferrocenium (Fc/Fc+) redox couple. As previously reported,39 the oxidation of the homopolymer P(iI) is not stable to repeated voltammetric scans and quickly results in film degradation when the potential is scanned from 0 to 1.4 V. This is not surprising considering the lack of electron-rich aromatic units in the homopolymer’s backbone. Conversely, the D−A 1288

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Chemistry of Materials polymers display one-wave oxidations from 0 to 1 V, as shown in Figure S11, with no loss of current after 10 consecutive cycles. The redox behavior of the P(iI) film in the reductive region is markedly different from that in its oxidation region: when potentials are scanned negative of 0 V to −1.5 V under cyclic voltammetry, a distinct quasi-reversible reduction process centered at a half-wave potential of −1.30 V is recorded. As shown in the Supporting Information, quasi-reversible first reductions are also found within −1.5 V for the CVs all of the D−A polymers. A second reduction wave is observed centered around 1.65 V for all D−A polymers, albeit with less-defined peaks and reversibility, suggesting that the subsequent electrons injected into the polymers’ π-systems render it too energetic and lead to side reactions other than the targeted reversible redox process. Each polymer shows a distinct first reduction under DPV, with onsets of reduction confined between −1.12 V for P(DTG-iI) and −1.26 V for P(iI), as summarized in Table 1. Since the Fc/Fc+ redox standard is +0.38 V from that of the saturated calomel electrode (SCE) and the energy of SCE is found at 4.7 eV from vacuum, we calculated the energies of the HOMO and the LUMO levels from DPV oxidation and reduction onsets, respectively, for each polymer using the Fc/ Fc+ standard set at ∼5.1 eV vs vacuum.46,47 Also detailed in Table 1, the energies of the LUMO levels of all polymers are within a 0.2 eV range from 3.84 eV for P(iI) to 3.98 eV for P(DTG-iI). The energies of the HOMO levels are found electrochemically to be highest for P(T-iI) at 5.57 eV and lowest for P(T3-iI) at 5.39 eV, gradually decreasing as the thiophene content of the repeat unit increases from one to three rings. This is consistent with more stabilized positive charges on the oxidized backbone of the polymers containing larger amounts of electron-rich moieties. Overall, the electrochemical bandgaps calculated from the latter energies (except for P(iI) which only displayed stable reduction) follow the same energetic trend as that of the measured optical bandgaps, as they decrease from 1.71 eV for P(T-iI) down to 1.49 eV for P(DTG-iI). Such low bandgaps in conjunction with HOMO energies around 5.5 eV and LUMO energies around 3.9 eV offer good prospects when one envisions their potential for ambipolar behavior in field-effect transistors. Bulk Organization. The device performance of semiconducting polymers strongly depends on the packing motif of the conjugated backbones being driven by π-stacking interactions. X-ray scattering allows one to access detailed information about the supramolecular organization in thin film or in the bulk. The analysis of mechanically oriented fibers with an area detector for the X-ray scattering collection leads to further insight into polymer arrangement.48−51 The appearance of scattering intensities on specific planes of the area detector allows determination of the relative polymer organization within the more complex supramolecular assembly. For an analysis, the two-dimensional patterns are divided in equatorial and meridional planes on which the reflections can be located. The extruded fiber is placed vertically toward the detector as illustrated in Figure 3. In the specimen, the polymer chains are macroscopically aligned typically along the extrusion direction which is defined as the 00l axis of the lattice. Axes h00 and 0k0 are perpendicular to 00l and randomly distributed around this axis. This assignment of Miller’s indexes is in agreement with the GIWAXS analysis of thin films.52 Before measurement, the specimens were annealed for 1 h at 200 °C. This temperature is slightly higher than that applied for the thin films (150 °C vide infra) to ensure an even thermal

Figure 3. Schematic illustration of the setup for two-dimensional X-ray scattering experiments. The incident X-ray beam is diffracted as it passes through the extruded filament, and the diffraction patterns are collected on a back screen, yielding information on the chain organization in the fiber sample.

treatment over the full sample thickness of 0.7 mm. For all polymers, the patterns show a typical orientation of the backbones along the extrusion direction (Figures 4 and S13) with pronounced equatorial reflections. The structural parameters for the bulk are summarized in Table 2. From the equatorial small-angle 100 reflection, the interlayer distance between stacked polymers is calculated to be between 1.65 and 2.02 nm. The π-stacking spacing is determined from the equatorial wide-angle 010 scattering intensity. Meridional 001 reflections are attributed to backbone repeat units. In the examples in Figure 4, patterns of P(iI), P(T-iI), and P(T3-iI) are presented to understand the influence of thiophene incorporation in the backbone on the bulk organization. The three polymers reveal similar lattice constants, but slightly varying peak widths along the stacking direction and for interlayer organization. The increased background intensity in the patterns prevents the extraction of exact values for the coherence lengths, which are, however, determined for thin films (see below). Nevertheless, equatorial integrations of the 2DWAXS patterns enable a comparison of peak widths within this series of three polymers. The smallest peak width (highest coherence length) for the π-stacking is found for P(iI), while it is the highest width for P(T-iI) (Figure S14). Interestingly, the peak width for the layer structure follows the opposite order with highest coherence length for P(T-iI) and lowest for P(iI). This means that in bulk P(iI) backbones are well packed in layers but that these layer stacks are not well ordered with respect to one another on the long-range. As discussed below in the GIWAXS section, this trend is opposite in the thin film with P(iI) showing the highest coherence length also for the interlayer organization. Slight differences are also obvious for the meridional plane of the patterns related to the backbone repeat units. In contrast to P(iI) and P(T-iI), polymer P(T3-iI) does not reveal any 00l scattering intensities corresponding to backbone repeat units implying differences in the packing motif. These variations in packing have been studied in more detail by solid state NMR spectroscopy (see below). The shortest interlayer distance of 1.65 nm was found for P(DTS-iI) and P(DTG-iI). Typically, longer side chains and a 1289

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Figure 4. 2DWAXS patterns after annealing at 200 °C for 30 min for (a) P(iI), (b) P(T-iI), and (c) P(T3-iI). Reflections are assigned by Miller’s indexes.

Table 2. Summary of the Fiber 2D-WAXS Results and the Thin-Film GIWAXS and Transistor Data for the Polymer Series max. mobility (cm2/(Vs)) polymer P(iI) P(T-iI) P(T3-iI) P(DTS-iI) P(DTG-iI) a

interlayera (nm) 2.10 2.35 2.28 1.90 1.90

(2.02) (1.99) (2.03) (1.65) (1.65)

π-stackinga (nm) 0.375 0.370 0.365 0.365 0.370

interlayer co.b (nm)

π-stacking co.b (nm)

15.50 5.93 13.32 7.28 6.09

3.92 1.40 1.88 1.20 1.92

(0.375) (0.375) (0.370) (0.370) (0.390)

hole 4 5 1 2

× × × ×

10−2 10−2 10−3 10−3

electron 1 × 10−4 1 × 10−1 1 × 10−3 1 × 10−3

d-Spacing values in parentheses are related to the bulk, from the fiber 2D-WAXS analysis. bCoherence lengths.

Figure 5. GIWAXS patterns for films of (a) P(iI), (b) P(T-iI), and (c) P(T3-iI) after annealing at 150 °C for 1 h, performed directly on the thin film transistors.

0.39 nm for P(DTG-iI), which is the largest value of the series. Both polymers are also characterized by the lowest bulk order within the series, as evidenced by the broad shape of wide-angle equatorial reflections displayed in Figure S13. This low order is attributed to the four sterically demanding 2-ethylhexyl side chains attached to each monomer unit reducing the π-stacking interactions. Thin Film Organization. The organization in thin films was studied by grazing incidence WAXS (GIWAXS) on solution drop cast films, using the same procedure as for transistor preparation, to better understand the later discussed device performance (Figures 5 and S15). Thereby, an annealing step at 150 °C for 1 h was applied. The spacings found for the films are within the same range as that observed for the bulk and are summarized in Table 2. In all cases, the interlayer distance slightly increases, while the π-stacking remains almost unchanged except for P(DTG-iI). This polymer shows a decrease of the π-stacking from 0.39 nm for bulk to 0.37 nm for film. Within the copolymer series, P(T-iI) shows the lowest

pronounced backbone curvature result in increased interlayer distance.48−51,53 The polymers reported here are all based on axi-symmetric donor moieties with the highest degree of rigidity provided by two fused thiophenes at most; therefore, similar backbone curvature is likely across the copolymer series, and the influence of conjugated backbone curvature can be excluded in this case. For solubility purposes, P(iI), P(T-iI), and P(T3-iI) were functionalized with longer HD chains on the iI unit compared to that of P(DTS-iI) and P(DTG-iI) which bear shorter EH chains, likely accounting for the shorter polymer interlayer distances of 1.65 nm found for P(DTS-iI) and P(DTG-iI). From the wide-angle reflections in the equatorial plane, intralayer π-stacking distances in the range of 0.37 to 0.39 nm were found, as summarized in Table 2. Interestingly, the direct comparison between P(DTS-iI) and P(DTG-iI) indicates a strong influence of the semimetal atom in the fused bithienyl unit. Because of the larger size of the Ge atom compared to Si, the packing distance increases from 0.37 nm for P(DTS-iI) to 1290

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Chemistry of Materials coherence length for the π-stacking with a value similar to that determined for P(DTG-iI) and P(DTS-iI). Identically to the bulk, the highest coherence length values in the stacking direction, but this time also for the layer organization, are found for P(iI). Surprisingly, all polymers assemble in a face-on arrangement on the surface as indicated by the apparent meridional π-stacking 010 reflection. The two remaining lattice axes, 100 and 001, are oriented parallel to the surface. Such distinct uniformity of the polymer orientation is unusual since the drop-cast films are relatively thick and inhomogeneous. Typically, a face-on alignment is found mainly for spin-coated polymer layers which are thin (