Impact of the Crystallite Orientation Distribution on Exciton Transport in

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Impact of the Crystallite Orientation Distribution on Exciton Transport in Donor−Acceptor Conjugated Polymers Alexander L. Ayzner,†,‡,⊥ Jianguo Mei,†,# Anthony Appleton,†,¶ Dean DeLongchamp,§ Alexandre Nardes,∥ Stephanie Benight,† Nikos Kopidakis,∥ Michael F. Toney,*,‡ and Zhenan Bao*,† †

Department of Chemical Engineering, Stanford University, Stanford, California 94305, United States Stanford Synchrotron Radiation Lightsource, SLAC National Accelerator Laboratory, Menlo Park, California 94025, United States § National Institute of Standards and Technology, Gaithersburg, Maryland 20899, United States ∥ National Renewable Energy Laboratory, Golden, Colorado 80401, United States ‡

S Supporting Information *

ABSTRACT: Conjugated polymers are widely used materials in organic photovoltaic devices. Owing to their extended electronic wave functions, they often form semicrystalline thin films. In this work, we aim to understand whether distribution of crystallographic orientations affects exciton diffusion using a low-band-gap polymer backbone motif that is representative of the donor/acceptor copolymer class. Using the fact that the polymer side chain can tune the dominant crystallographic orientation in the thin film, we have measured the quenching of polymer photoluminescence, and thus the extent of exciton dissociation, as a function of crystal orientation with respect to a quenching substrate. We find that the crystallite orientation distribution has little effect on the average exciton diffusion length. We suggest several possibilities for the lack of correlation between crystallographic texture and exciton transport in semicrystalline conjugated polymer films. KEYWORDS: exciton diffusion, conjugated polymer, texture, crystallographic orientation, fluorescence quenching

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INTRODUCTION An organic solar cell functions by delivering an optically excited state (exciton) to an interface between a donor (D) and an acceptor (A) molecule.1−4 The D/A frontier molecular orbital energy levels are designed such that there is a thermodynamic driving force to convert an exciton into a pair of opposite charges localized across the heterointerface.5,6 Since excitons live for a finite amount of time, it is imperative that the excited state is able to reach the interface prior to decay back to the ground state. One of the major advances in the field of organic photovoltaics (OPV), which opened the doors to a flood in OPV research, was the formation of the D/A heterojunction spanning the bulk of the thin film. This architecture, called the bulk heterojunction (BHJ),5,7−12 tackled one of the limiting characteristics of singlet excited states in organic materials: the distance that excitons can diffuse during their (short) lifetime is fairly small, roughly 5−15 nm for a broad range of π-conjugated materials.13−27 Therefore, a D/A blend film where the domain size is comparable to the exciton diffusion length (EDL)the mean distance traveled by excitonsresults in a PV device with a large internal quantum efficiency, ensuring that most excitons © XXXX American Chemical Society

generate free charges that diffuse to the solar cell electrodes.28−32 For π-conjugated organic molecules, it has been established that the relative molecular orientation in a molecular crystal plays a large role in determining the rate of local charge transfer and long-range charge transport.33−37 Stemming from the orientational dependence of the electronic coupling, this is particularly true for small molecules. Moreover, there is also strong evidence that charge transport in conjugated polymers is anisotropic and thus depends on the relative chain orientation in the unit cell.38,39 Conjugated polymers are composed of a string of quasi-independent chromophore subunits, each spanning several monomers. Interaction between π-electron densities on adjacent molecules gives rise to an attractive interchain interaction, which in turn leads to formation of Special Issue: Advances towards Electronic Applications in Organic Materials Received: April 6, 2015 Accepted: August 7, 2015

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DOI: 10.1021/acsami.5b02968 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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Figure 1. (A) Chemical structure and GIXD pattern of PiI2T-ref. The XRD pattern shows that PiI2T-ref crystallites are oriented edge-on with respect to the substrate. (B) Chemical structure and GIXD pattern of PiI2T-Si. The GIXD image shows a coexistence of edge-on and face-on oriented crystallites.

substrates to the intensity for films on substrates capable of photoexcited electron transfer and thus quenching of emission. We show below that, although the ODF differs significantly, the thickness-dependent PL quenching is similar for the two polymers, suggesting that the EDL in these films is also quite similar. We conclude by considering the possible origins for the insensitivity of the EDL to the polymer crystallite orientation.

nanometer-sized crystallites (although with paracrystalline order or relatively strong packing disorder). These small crystals frequently exhibit preferred orientation, or crystallographic texture, in a thin film.40−42 What is less clear for polymers is whether there is an orientation dependence to the EDL: does the orientation of crystallographic planes in the nanometer-sized crystallites affect the average distance that an exciton can travel in its lifetime? An inherent complication is the fact that polymer crystallites are embedded in an amorphous matrix. Thus, exciton diffusion in the polymer phase takes place in a blend of amorphous chains and crystallites; the blend is orientationally heterogeneous at the nanometer scale but still exhibits an orientation distribution having one or more preferred orientations.39,43 Indeed, this has implications for the BHJ film, where the film morphology is composed of (at least) three phases: pure polymer crystallites, a largely amorphous intermixed phase, and a largely pure acceptor phase.44 We have recently used side chain engineering in an isoindigo−thiophene low-band-gap conjugated polymer to significantly improve the field-effect hole mobility of the polymer film.45 In the course of this work, we discovered that, depending on the side chain, the texture of the polymer film changed qualitatively. The polymer film displayed a dominant edge-on orientation with respect to the substrate (π-stacking inplane) in the case of a branched aliphatic side chain. Upon changing to a linear side chain with a terminal siloxy group, we found dual texture with both in-plane and out-of-plane πstacking. Since we are able to vary the crystallite orientation distribution (ODF) in such a significant manner, we have used these polymers to address the question: does the EDL depend on the mean crystallite orientation in donor/acceptor polymer films? To answer this, we perform steady-state photoluminescence (PL) quenching experiments as a function of film thickness, where we compare PL intensity for films on inert

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RESULTS AND DISCUSSION The chemical structures of the poly(isoindigo−bithiophene) (PiI2T) polymers used in this work are shown in Figure 1. In this work, both polymer films were annealed at 170 °C and slowly cooled on the hot plate to room temperature to increase the crystallite density. Figure 1A shows a zoomed-in view of the GIXD pattern of PiI2T-ref on a glass substrate. As we reported previously, the polymer shows pronounced Bragg peaks along the Qz axis (out-of-plane component of reciprocal space vector), which shows out-of-plane lamellar stacking of side chains. Thus, ordered polymer chains have their side chains approximately oriented perpendicular to the substrate, and the π-stacking is in-plane. Figure 1B shows the GIXD pattern for PiI2T-Si on glass. PiI2T-Si displays a preference for two dominant orientations: there is a population of crystallites oriented primarily edge-on with respect to the substrate, as with PiI2T-ref, indicated by the out-of-plane lamellar Bragg peaks. However, there is also a significant population of crystallites that are oriented face-on: their π-electron-rich backbones are oriented parallel to the substrate plane; that is, π-stacking takes place out-of-plane. In order to quantify the relative populations of crystallites adopting varying orientations, we construct a pole figure of the integrated (100) peak (along Q, the scattering vector magnitude) as a function of polar angle. The integrated intensity was then corrected with a weighting factor of sin(θ), where θ is the polar angle measured with respect to the Qz axis B


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tions, we have integrated the ODF (normalized to the curve integral) between ∼0 and 45°, which we take as a measure of the edge-on contribution, and the integral between 45 and ∼85° is taken as a measure of the face-on crystallite population. We find that, for PiI2T-ref, 78% of crystallites are oriented edge-on and 22% face-on; for PiI2T-Si, we find that 26% of the crystallites are oriented edge-on and 74% face-on. Thus, the crystallographic orientation averaged over the bulk of the film differs significantly between the two polymer films. Since diffraction is only sensitive to crystalline regions, we carried out near-edge X-ray absorption fine structure (NEXAFS) spectroscopy experiments at the C K-edge in order to gain a better understanding of the overall average polymer molecular orientation. The angular dependence of NEXAFS intensities can be used to estimate the average molecular tilt of the conjugated polymer plane within a few nanometers of the film/air interface. Since NEXAFS does not discriminate between crystalline and amorphous polymer, the results yield an average over the entire microstructure. Figure 3 shows NEXAFS spectra of PII2T-ref and PII2T-Si (A and B, respectively), collected in partial electron yield mode, as a function of X-ray incidence angle. As shown in the Supporting Information, the intensities in the low-energy π* regions as a function of angle correspond to an average tilt of the conjugated molecular plane of ∼28° with respect to the film normal for both polymers. However, we found significant variation in the tilt angle, with some samples yielding a tilt close to the magic angle, again for both polymers (see Supporting Information). These data suggest that, on average, near the film surface, polymer chains are oriented slightly edge-on. Thus, given the GIXD data, it is highly likely that the film surface has an average molecular orientation that is different from that of the bulk of the polymer film, and the majority of face-on oriented crystallites are likely located within the film bulk. This raises the possibility that the EDL may not be a constant as a function of film thickness. Having examined the polymer orientations, we continued to investigate the optical absorption and emission properties of the annealed polymer thin films, the spectra of which are displayed in Figure 4. Qualitatively, the shapes of the absorption spectra (solid lines) of both polymers are similar, though there are some relatively minor differences in the ratio

(i.e., the pole). This factor takes into account the fact that the number of crystallites contributing to the intensity of a given Bragg reflection for a 2D powder (isotropic crystallite orientation in-plane) decreases away from the pole toward the equator. This is a consequence of the fact that the in-plane radius swept out by the scattering vector increases toward the equator, and the number of crystallites oriented to intercept the Ewald sphere and thus fulfill the Bragg condition decreases.46,47 Thus, the raw measured intensity is not representative of the number density of crystallites at a given orientation with respect to the film normal and hence must be weighted by the sine factor to convert the polar angle dependence of the intensity to a quantity proportional to the orientation distribution function (ODF).48 When properly normalized, this quantity gives the fraction of crystallites having the scattering vector orientation in the angular range between θ and θ + dθ. Figure 2 shows the calculated ODF for PiI2T-ref and PiI2TSi thin films normalized to the peak intensity. Consistent with

Figure 2. Normalized crystallite orientation distribution functions for PiI2T-ref (red curve) and PiI2T-Si (black curve). The PiI2T-ref film has primarily edge-on oriented crystallites, whereas the majority of PiI2T-Si are oriented face-on with respect to the substrate.

the images, for PiI2T-ref, the dominant orientation is edge-on. Thus, the (100) reflection is peaked near zero polar angle. The ODF for PiI2T-Si shows that the primary chain orientation is face-on with respect to the substrate. There are clearly edge-on and isotropic contributions, but it is reasonably small compared to the face-on population. To estimate the relative contribu-

Figure 3. NEXAFS spectra as a function of incident angle of (A) PII2T-ref and (B) PII2T-Si films. The sharp peaks near 285 eV correspond to C 1s → π* transitions. The angular dependence of these peaks was used to calculate the tilt angle of the polymer backbone, averaged over all polymer chains in the excitation volume. We find that, for both polymers, the mean conjugated plane tilt with respect to the surface normal is ∼28°. Thus, in the vicinity of the film/air interface, both polymers are oriented mildly edge-on. C


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Figure 5. Experimental PL quenching ratios as defined in the text. Black squares correspond to PiI2T-Si, and red circles correspond to PiI2T-ref. The data make it clear that, given the similarity in the excited-state lifetime, the exciton diffusion length for both polymer films is similar. The cartoon shows the experimental geometry: the substrate (gray) is on the bottom, and the polymer film is excited through the air interface. The arrow shows the thickness that is varied in this experiment.

Figure 4. Normalized absorption (solid lines) and emission (dashed lines) spectra of PiI2T-ref (black) and PiI2T-Si (red). The spectral shift is consistent with the difference in the π-stacking distance.

of the vibronic overtones. However, the clearest difference is in the spectral position: the absorption peak for PiI2T-Si is clearly shifted to the red by ∼13 nm. This shift is consistent with the observation that the π-stacking distance for PiI2T-Si is smaller by ∼0.2 Å, as we reported previously; however, it is also possible that the small red shift is due to enhanced planarization of the chain backbone. PL spectra (dashed lines) are shifted with respect to each other by the same amount as absorption spectra, indicating similar rearrangements in the excited-state geometry, as largely expected for the same polymer backbone. We have also measured excited-state PL lifetimes using time-correlated single-photon counting. Interestingly, although we find differences in the positions of the steady-state spectra, the mean lifetimes of ∼330 ps are nearly identical for both polymers. Next, we address the question of whether the crystallographic differences in these polymer thin films translate into a difference in exciton quenching efficiency, which we do by measuring the thickness dependence of the PL intensity. This is done by coating a quenching titania layer, prepared via atomic layer deposition and subsequent thermal annealing. The energy level offset between the titania conduction band minimum and the polymer LUMO provides a driving force for electron transfer from the polymer excited state to the titania layer, which quenches polymer PL.18 A glass substrate is used to represent the case of no quenching. We compute the quenching ratio by integrating the PL spectrum for a given thickness on glass and on titania according to the following definition: ηq ≡ 1 − ∫ Iq(λ)dλ/∫ Inon‑q(λ)dλ. Here, Iq is the PL intensity on titania, Inon‑q is the intensity on glass, and λ is the wavelength. Given that the measured observable is the thickness-dependent quenching, this experiment gives a measure of the exciton diffusion length in the direction normal to the substrate. The quenching ratio as a function of film thickness for both polymers is shown in Figure 5. The thickness was determined using X-ray reflectivity, and to generate the abscissa error bars, X-ray reflectivity curves were collected at a number of lateral locations across each thin film. Error bars represent one standard deviation averaged across the film surface. It is clear that as the thicknesses approach 100 nm, film height variations become more pronounced for PiI2T-Si relative to PiI2T-ref. We believe that the reason for this can be traced back to the increased metastability of PiI2T-Si solutions: as the concentration is increased to produce thicker films, polymer aggregation is enhanced, leading to diminished in-plane uniformity.

For both samples, the quenching ratio decreases monotonically with thickness. This is because with increasing thickness, excitons that are created far from the quenching surface are unable to reach the quenching interface during the excited-state lifetime. The data in Figure 5 show that the quenching ratio decays over a similar length scale for both polymers. The large error bar for the highest thickness of PiI2T-Si is due to the large thickness nonuniformity as described above, and the fact that quenching persists for such a large nominal thickness is in fact likely due to quenching from the thinner regions of the film. Since the excited-state lifetimes are similar, the data imply that, in our polymer films, the EDLs for both polymers normal to the substrate plane must be similar. Therefore, in these films, the fact that the ratio of face-on to edge-on oriented crystallites differs substantially seems to have little effect on the vertical transport of excitons. We note that we attempted to model the thickness dependence of the PL quenching by solving the exciton diffusion equation with common perfectly quenching boundary conditions. The excitation profile was modeled with the transfer matrix formalism. As shown in the Supporting Information, we were unable to describe the experimental PL quenching ratio using this model satisfactorily. Given the comparison between NEXAFS and GIXD data, it is possible that a single thicknessindependent EDL is insufficient to characterize the diffusion of polymer excitons in a film with a microstructure that may depend on the vertical position within the film. However, GIXD as a function of thickness (shown in the Supporting Information) does indeed show that the dominant texture motif is preserved as a function of thickness for both polymers. Additionally, we find that the texture is independent of substrate type. Nevertheless, the relative comparison of the experimental data unambiguously shows lack of significant correlation between polymer crystallite orientation and the EDL in our thin films. Previous experimental and theoretical work has suggested a correlation between the excitonic coupling between chromophores and molecular orientation. Rand et al. found that for face-on oriented small-molecule Zn phthalocyanine films, the vertical EDL was smaller relative to the edge-on orientation.49 It was argued that the intermolecular exciton hopping length was larger in the edge-on orientation, which resulted in an D


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Figure 6. Possible scenarios which could result in a lack of correlation between the relative crystallite orientation and the EDL. (A) Exciton transport is dominated by intrachain motion as opposed to intercrystallite hops. (B) Crystallite volume density is significantly below the percolation threshold, resulting in inefficient intercrystallite excitonic communication. Purple strings represent single polymer chains bridging crystallites.

depends on the wavefunction extent.52,53 Although the transfer rate depends on relative orientation, relaxed emissive excitons residing in crystalline regions would be unable to electronically communicate due to their large mean separation, and thus the relative crystallite orientation would matter little. The possibility of varying the crystallite volume density to test this hypothesis is interesting but is beyond the scope of the present research. Our results have particular implications for engineering bilayer solar cells and cells with vertical electron donor/ acceptor concentration gradients, where the magnitude of the EDL may be of elevated significance relative to a BHJ. Assuming that the isoindigo−bithiophene system is representative of donor/acceptor low-band-gap polymers, it appears that modulating the polymer crystallite ODF is unlikely to result in an enhanced EDL in the current generation of donor/acceptor polymers. This finding is analogous to what Rand et al. found with bilayers based on small molecules, albeit owing to potentially different reasons.49 The polymer excitonic coupling was calculated to be significantly larger along the π-stacking direction relative to along the chain or through side chains38a characteristic which could lead to enhanced vertical transfer between percolated face-on oriented crystallites; we see no evidence that this is borne out in practice for our face-on oriented PiI2T-Si relative to edge-on PiI2T-ref. Frequently, polymer crystallites on weakly interacting (e.g., oxide) surfaces pack edge-on, which is thought to be unfavorable for charge transport in the vertical direction. However, this does not appear to be the case for exciton transport in our donor/ acceptor polymer films.

exciton hopping rate that was roughly a factor of 5 larger for the face-on orientation. The optical properties of conjugated polymer thin films differ substantially from small molecules. Since polymer films are composed of a distribution of conjugated segment lengths and thus a distribution of chromophore energies, it is widely thought that higher-energy excited states are funneled to extended, intrachain low-energy regions or to interchain states within crystallites. This mechanism is supported by the observation that the emission spectrum is found to be largely independent of excitation wavelength.50 There are two reasons that we believe could lead to the insensitivity of the EDL to the relative crystallite orientation, which we schematically illustrate in Figure 6. We speculate that the first possibility is exciton transport being largely driven by intrachain motion as opposed interchain hops. That is, if it is the polymer backbone that percolates through large regions of the film and in effect bridges different crystalline regions, then the rate of exciton motion along the polymer backbone may significantly exceed the rate of intercrystallite exciton hopping. In fact, there is evidence that some fraction of polymer backbones in polymer−fullerene BHJs can be oriented vertically, as a significant fraction of the transition dipole moment is oriented in that direction.51 This would provide a potential avenue for vertical exciton diffusion without the need for a face-on oriented crystallite population. Even in the absence of vertical backbone percolation, long regions of uninterrupted backbone conjugation could provide a largely lateral path between crystalline regions until a particularly facile vertical transport direction is sampled. The second possibility that we consider concerns crystallite percolation in the amorphous matrix. If the polymer film contained a large enough volume fraction of face-on oriented crystallites, it is quite possible that the exciton hopping rate (in the vertical direction) would reflect this fact when compared to the edge-on orientation. It is possible that the volume fraction of face-on oriented crystallites in our thin films is significantly below the percolation threshold, such that the intercrystallite exciton transfer rate normal to the substrate is sufficiently low. Generally, the energy transfer rate depends on the relative orientation between transition dipole moments, which in turn

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CONCLUSION To conclude, we have attempted to correlate the crystallite orientation distribution function of polymer films containing an identical backbone but varying side chains, which resulted in qualitatively different crystallographic texture. Although the population of face-on oriented crystallites was significant in the siloxy-functionalized polymer films, the thickness-dependent PL quenching was similar for both polymers. Thus, in these thin films, the crystalline orientation seems to exert little influence on the exciton diffusion length. Although a number of E


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scenarios could lead to this result, we believe it is more likely that this insensitivity arises due to an insufficient density of oriented crystallites, such that the intercrystallite exciton transfer rate is relatively small. This leaves open the possibility that by increasing the properly oriented crystallite volume fraction, an orientation-dependent EDL may be observed.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.5b02968. Experimental procedure, fits of angle-dependent NEXAFS data, exciton diffusion modeling, additional GIXD, and time-resolved PL curves (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. Present Addresses ⊥

(A.L.A.) Department of Chemistry and Biochemistry, UC Santa Cruz, Santa Cruz, CA 95064. # (J.M.) Department of Chemistry, Purdue University, West Lafayette, IN 47907. ¶ (A.A.) St. Petersburg College (Natural Science Dept.), St. Petersburg, FL 33733. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank the Bent group at Stanford University for help with ALD preparation of titania films. This work was partially supported by the Center for Advanced Molecular Photovoltaics, Award No. KUS-C1-015-21, made by King Abdullah University of Science and Technology. We also acknowledge support from the Global Climate and Energy Program at Stanford. GIXD measurements were carried out at the Stanford Synchrotron Radiation Lightsource, a national user facility operated by Stanford University on behalf of the U.S. Department of Energy, Office of Basic Energy Sciences. N.K. and A.M.N. acknowledge funding from the Energy Frontier Research Center “Molecularly Engineered Energy Materials (MEEMs)” funded by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Contract Number DE-SC0001342:001.

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