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Evidence of Molecular Structure Dependent Charge Transfer between Isoindigo-Based Polymers and Fullerene Tzung-Han Lai,† Iordania Constantinou,† Caroline M. Grand,‡ Erik D. Klump,† Sujin Baek,† Hsien-Yi Hsu,§ Sai-Wing Tsang,∥ Kirk S. Schanze,§ John R. Reynolds,‡ and Franky So*,†,⊥,# †

Department of Materials Science and Engineering and §Department of Chemistry, University of Florida, Gainesville, Florida 32611, United States ‡ School of Chemistry and Biochemistry, School of Materials Science and Engineering, Center for Organic Photonics and Electronics, Georgia Tech Polymer Network, Georgia Institute of Technology, Atlanta, Georgia 30332, United States ∥ Department of Physics and Materials Science, City University of Hong Kong, Hong Kong, China ⊥ Department of Materials Science and Engineering, North Carolina State University, Raleigh, North Carolina 27606, United States S Supporting Information *

ABSTRACT: The effects of the oligothiophene length of two thiophene-isoindigo copolymers on film morphology, charge transfer, and photovoltaic device performance are reported. Despite the similarities in their repeat unit structures, the two polymers show distinctly different film morphologies and photovoltaic performance upon blending with PC71BM. We found that there is a significant increase in the dielectric constant of the photoactive film upon blending fullerene with the polymer that exhibits a higher power conversion efficiency. Blend photoluminescence transients revealed a fast dissociation route in the better performing polymer followed by a slower decay. The fast decay in transient PL is attributed to a higher charge transfer efficiency when blending with the fullerene. We suggest that the charge transfer efficiency is determined not only by the microscopic morphology but also whether the polymer can accommodate the fullerene molecules in close proximity to the acceptor moiety to facilitate electronic coupling between the isoindigo acceptor and the fullerene molecule. We propose that the fast decay component seen in transient PL for the better performing polymer, along with the increase in dielectric constant, is a signature of enhanced electronic coupling between the polymer and the fullerene. The enhanced electronic coupling is thought to originate from a polymer chemical structure which allows the fullerene molecules to come to closer proximity for more efficient charge transfer.

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INTRODUCTION The combination of electron-rich and electron-deficient units to form donor−acceptor (D−A) polymers has been a widely used strategy for the synthesis of low bandgap conjugated polymers for organic photovoltaic (OPV) applications.1,2 D−A polymers are attractive due to the tunability of properties such as bandgap and carrier transport based on the variety of donor and acceptor units accessible for their design.3 Solar power conversion efficiencies for single junction OPVs, constructed using D−A low bandgap polymers, have been steadily increasing and are now exceeding 10% for a number of systems.4−7 Within this class of materials, isoindigo has been extensively used as an electron acceptor in D−A polymers due to its electron-deficient character, ability to yield materials with extensive light harvesting properties, facility for easy functionalization to control polymer solubility, and scalability for mass preparation as commercialization is considered.8 It has been widely demonstrated that the film morphology of polymer−fullerene blends can play an important role in device performance in bulk heterojunction (BHJ) OPVs. Morpho© 2016 American Chemical Society

logical parameters including face-on and edge-on interactions between the polymer, the accepting phase and the substrate, percent crystallinity, degree of polymer−fullerene intermixing, domain purity, and domain size, among others, have been shown to be the determining factors in the conversion of photogenerated excitons to electron−hole pairs and free carriers.9,10 In addition to processing parameters, such as thermal or solvent annealing and the use of processing additives, a subtle change in the polymer’s chemical structure has been shown to have profound effects on the film morphology and consequently the device performance.11−13 Even though multiple reports exist in the literature analyzing the effect of processing parameters on film morphology, a detailed relationship between the effect of the polymer’s chemical structure on film morphology, charge carrier generation, and device parameters is yet to be established.14 Received: February 26, 2016 Revised: March 23, 2016 Published: March 23, 2016 2433

DOI: 10.1021/acs.chemmater.6b00824 Chem. Mater. 2016, 28, 2433−2440

Article

Chemistry of Materials

Figure 1. Repeat unit structure of P(T1-iI) dimer and P(T3-iI) highlighting conjugation along the polymer backbone (in blue), torsion between the heterocycles (red arrows), and branched side chains (in green).

to the fullerene molecule. In addition, we used charge modulated electroabsorption spectra (CMEAS) and dielectric measurements to investigate the CT energies to probe possible coupling between the polymers and fullerene. Finally, based on molecular orbital calculations, a model based on molecular orbital interactions is presented to explain the difference in electronic coupling between the two polymers and the fullerene.

The exciton dissociation process in OPV is often a limiting factor determining the device performance, possibly attributable to the relatively low dielectric constants (εr ∼ 2−4) of most conjugated polymers when compared to inorganic semiconductors.15,16 As a result of the large exciton binding energies, the exciton dissociation process in OPVs relies on an energetic offset at the polymer/fullerene interface. An essential step in exciton dissociation is exciton diffusion from the polymer domains to the polymer/fullerene interface. In BHJ OPVs, exciton diffusion and dissociation efficiency highly depend on film morphology and domain size.17 At the polymer/fullerene interface, charge transfer (CT) excitons are formed with electrons localized in the fullerene molecules loosely bound to holes localized in the polymer.18 Therefore, the energetics of the CT manifold and the kinetics of CT excitons in OPVs determine the efficiency of exciton dissociation and the corresponding device performance.19−21 In this report, we describe the effect of the polymer’s chemical structure on film morphology, charge carrier generation, and device performance. Two isoindigo-based copolymers were used in this study where we have varied the number of thiophene donor units in the backbone as depicted in Figure 1. The small difference in the polymer repeat unit structure was found to cause significant differences in film morphology as well as device performance, despite the fact that the two polymers had nearly identical energy levels. Specifically, we observe a substantial difference in domain size in blends of the polymers with fullerenes and subsequently a large difference in the device short circuit current (JSC) and open-circuit voltage (VOC). Surprisingly, we found that the dielectric constant of the polymer−fullerene blend, which resulted in a higher Jsc, is higher than the dielectric constant of either the neat polymer or the fullerene. We attribute the higher dielectric constant to a strong electronic coupling of the polymer and the fullerene leading to a more efficient charge transfer. We investigated in detail the microscopic origin of the difference in charge transfer properties in the two polymer:fullerene blends with several spectroscopic techniques. We used sub-bandgap external quantum efficiency (EQE) along with transient photoluminescence (PL) measurements to study the charge transfer efficiency, which is a measure of the exciton dissociation efficiency resulting from exciton diffusion to the donor− acceptor interface leading to electron transfer from the polymer

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EXPERIMENTAL SECTION

Sample Preparation. P(T1-iI) and P(T3-iI) were blended with PC71BM in 1:1.5 ratio in 1,2-dichlorobenzene. PEDOT:PSS was spincoated onto indium tin oxide (ITO) substrate, precleaned with acetone and isopropyl alcohol, and annealed at 140 °C for 15 min. The polymer solutions were then spin-coated on top of the PEDOT:PSS layer. After the samples were dried in a nitrogen atmosphere, they were put in a thermal evaporator to deposit 1 nm LiF and 100 nm of Al. The device structure was ITO/PEDOT:PSS/P(T1-iI) or P(T3iI):PC71BM/LiF/Al, and the device area was 4.4 mm2. Sample Characterization. A HIOKI 3523-50 LRC meter was used to measure the device geometric capacitance. Film thicknesses were measured using atomic force microscopy (AFM). To eliminate the parasitic effect, the devices were calibrated at the short and open circuit condition. The capacitance of the device was measured at 100 kHz with a small 20 mV ac modulation. For transient photovoltage (TPV), the sample was kept under white light illumination. The optical perturbation was obtained using a LASER-EXPORT DTL319QT (527 nm 5 ns, 60 μJ/pulse) and a series of neutral density filters that were used to lower the power of the laser to ∼1 nJ. The photovoltage decay was obtained using a Tektronix TPS2024 oscilloscope. During the measurement, an external voltage equal to the open-circuit voltage of the devices was applied; this ensured there is not external current that can flow through the external circuit, and therefore TPV decay can be used to monitor the loss kinetics. For the measurement of CMEAS, the samples were probed by the incident monochromatic parallel beam into the sample through ITO with an incident angle of 45° and are reflected by the back Al electrode and captured by calibrated silicon and germanium photodetectors. The internal electric field of the sample was modulated by a DC bias superimposed with a small AC voltage at a modulation frequency of 1000 Hz. A current amplifier and a lock-in amplifier were connected to the detector to increase the signal-to-noise ratio. The sub-bandgap EQE was measured by connecting the device in series with a 120kΩ resistor and an SRS 830 lock-in amplifier. Monochromatic light from a monochromator with the light spot smaller than the device area was chopped and incident onto the device. The current output was measured using a lock-in amplifier. Transient PL was performed using 2434


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Chemistry of Materials Table 1. Electrochemical and Device Properties for P(T1-iI) and P(T3-iI) Taken over Tens of Devices P(T1-iI) P(T3-iI)

IP/EAff [eV]

Eg [eV]

VOC [V]

JSC [mA cm‑2]

FF [%]

PCE [%]

pristine εr

blend εr

−5.57/−4.02 −5.50/−3.97

1.55 1.53

0.91 ± 0.01 0.70 ± 0.00

5.1 ± 0.5 14.2 ± 0.1

56.3 ± 1.8 60.9 ± 0.3

2.6 ± 0.2 6.1 ± 0.1

2.37 ± 0.08 2.45 ± 0.10

4.22 ± 0.11 4.76 ± 0.18

Figure 2. a) J-V curves and b) EQE spectra for P(T1-iI):PC71BM and P(T3-iI):PC71BM devices.

reported for this system.25 The measured EQE spectra of the two polymer:PC71BM are shown in Figure 2b. The maximum EQE for P(T3-iI):PC71BM is 58%, while the maximum EQE for P(T1-iI):PC71BM is only 19%. In order to understand the origin of the difference in photovoltaic performance, we first investigated the film morphology for the two blends. Due to the excitonic nature and small exciton diffusion lengths in OPVs, a fine domain size on the order of 10−20 nm is desirable for optimum device performance.26 AFM was used to investigate the morphology for the two isoindigo-fullerene blend films. As shown in Figure 3, P(T1-iI):PC71BM and P(T3-iI):PC71BM films exhibited

a time-correlated single photon counting (TCSPC) spectrometer (Picoquant, Inc.) with a counting rate of 1.00 e+007(1/s). A pulsed laser with an average power of 1 mW, 70 ps duration, and 40 MHz frequency was used to excite the films. A 700 nm long pass filter was used to monitor the emission. The photoemission measurements were performed in situ in a VG ESCALAB 220i-XL ultrahigh vacuum (UHV) surface analysis system with a base vacuum of 10−10 Torr and an environmental chamber with a base pressure of 10−7 Torr. A He discharge lamp (photon energy 21.2 eV) with an instrumental energy resolution of 90 meV, as estimated from the Fermi edge of a cleaned Au sample, was used to measure the valence band structure of the sample. For efficient collection of secondary electrons, samples were negatively biased at −5.0 V with respect to ground. An Edinburgh Luminescence Spectrometer (Model: F900) equipped with a xenon lamp was used to measure the room temperature steady state photoluminescence (PL) spectra in the UV-NIR spectral region. The samples were excited with 2.02 eV excitation energy, and the PL spectra were measured using a red sensitive photomultiplier tube.

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RESULTS AND DISCUSSION The repeat unit chemical structures of the isoindigo polymers used in this study are shown in Figure 1, poly(thiophene-coisoindigo) P(T1-iI) (Figure 1a) and poly(terthiophene-coisoindigo) P(T3-iI) (Figure 1b).1,22−24 The polymer backbones are composed of the isoindigo acceptor unit along with different numbers of electron-donating thiophene units: one thiophene unit in P(T1-iI) and three thiophene units in P(T3iI). These two polymers have similar electrochemically measured ionization potentials (IP) and electron affinity (EAff) as shown in Table 1. The optical bandgaps of the pristine polymers extracted from the absorption spectra onsets are quite similar at 1.55 eV (800 nm) and 1.53 eV (810 nm), respectively, for P(T1-iI) and P(T3-iI) as shown in Figure S1. Figure 2a shows the current density−voltage (J-V) characteristics of ITO/PEDOT:PSS/polymer:PC71BM/LiF/Al devices under 100 mW/cm2 with AM 1.5 illumination. Despite the similarities in electronic and electrochemical properties, including the ionization potentials, bandgap energies, and dielectric constants of the pristine polymers, significant differences were observed in device performance in the resulting polymer:PC71BM devices. Optimized devices made with P(T1-iI):PC71BM active layers exhibited a VOC of 0.91 V and a JSC of 5.1 mA/cm2, while optimized P(T3-iI):PC71BM devices showed a smaller VOC of 0.70 V and a much higher JSC of 14.2 mA/cm2, consistent with what has been previously

Figure 3. AFM height images of (a) P(T1-iI):PC71BM and (b) P(T3iI):PC71BM blend films. The phase image of (c) P(T1-iI):PC71BM shows large domain size, while the phase image of (d) P(T3iI):PC71BM shows small domain size.

significantly different morphologies. In P(T1-iI):PC71BM films the root-mean-square (RMS) roughness was found to be 5.27 nm, indicating increased aggregation resulting in formation of relatively large domains (Figure 3a, 3c). Film morphology is known to be highly influenced by the solubility of the polymer and fullerene in the casting solvent, the degree of fullerene 2435


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Figure 4. (a) CMEAS spectra for P(T1-iI):PC71BM and P(T3-iI):PC71BM films. All measurements were done under reverse bias. (b) Sub-bandgap EQE plots for P(T1-iI):PC71BM and P(T3-iI):PC71BM.

indicating higher CT state recombination in the P(T1iI):PC71BM devices. Recently, we used CMEAS to study the CT state energetics in polymer:fullerene blends.34,35 In a typical CMEAS spectrum, electro-absorption (EA) signals for excitation energies above the polymer optical bandgap are due to the Stark effect, i.e. the coupling between the excitonic levels and higher forbidden states under electric field perturbation.36 Electroabsorption signals resulting from excitation energies below the bandgap were shown to signify the generation of carriers from the CT states. Figure 4 shows the CMEAS spectra for P(T1iI):PC71BM and P(T3-iI):PC71BM devices. For the P(T3iI):PC71BM devices, a large sub-bandgap signal was observed at energies below 1.5 eV. The sub-bandgap signal is attributed to the polarons generated by direct excitation of carriers in CT states at the heterojunction interface.34 Compared to the P(T1iI):PC71BM device, the stronger sub-bandgap signal in the P(T3-iI):PC71BM device suggests that the carrier generation at the CT states at the P(T3-iI)/PC71BM interface is more efficient than that at the P(T1-iI)/PC71BM interface. These results decouple the morphological factor and bring insight into the difference in electronic coupling efficiency between the polymers and fullerene at the heterojunction interface, which will be discussed more in detail below with consideration of the difference in molecular structure. Furthermore, regarding the CT state energetics, we obtained different sub-bandgap signal onsets in the two blends: 1.3 eV for P(T1-iI):PC71BM devices and 1.0 eV for P(T3-iI):PC71BM. This is consistent with the different electroluminescence energies as previously reported.25 In order to verify our findings on different CT energies, we used sub-bandgap EQE measurements to further probe the CT state energetics of the two isoindigo polymer:fullerene blends. The sub-bandgap EQE spectra for P(T1-iI):PC71BM and P(T3-iI):PC71BM devices are shown in Figure 4b. It is clearly shown that the sub-bandgap EQE for P(T3-iI):PC71BM device was almost an order of magnitude higher than that in the P(T1iI):PC71BM device. These results are in agreement with the CMEAS data presented above and further confirm that charge transfer and exciton dissociation are more efficient in P(T3iI):PC71BM devices. Moreover, it has been recently proposed by Vandewal et al. that the CT energy can be extracted by fitting the sub-bandgap EQE spectrum to the nonadiabatic electron transfer theory developed by Marcus.37 However, the fitting relies on a distinctive shoulder-like feature in the EQE spectrum, which is absent in both P(T1-iI):PC71BM and P(T3iI):PC71BM devices. Nevertheless, as shown in Figure 4b, the sub-bandgap EQE from the P(T3-iI):PC71BM blend extended

solubility in the polymer-rich domains, and the processing conditions for film formation. As such, it is hypothesized that the increased aggregation in the P(T1-iI):PC71BM films was due to the increased solubility of the P(T1-iI) polymer with higher density of branched side-chains along the polymer backbone as compared to P(T3-iI).27−29 This increase in polymer solubility allows for fullerene aggregation as the solvent evaporates and leads to the formation of large PC71BM-rich clusters.30 On the other hand, the RMS roughness (2.59 nm) and domain size for P(T3-iI):PC71BM films are significantly smaller. The smaller domain size facilitates polymer−fullerene intermixing for efficient exciton dissociation, possibly resulting in reduced geminate recombination and higher JSC for the P(T3-iI):PC71BM device. In order to investigate the contribution of different recombination processes to the difference in JSC for the two devices, we first measured the carrier lifetimes using the transient photovoltage (TPV) technique.31−33 As illustrated in Figure S2, the carrier lifetimes measured under open-circuit conditions for P(T1-iI):PC71BM and P(T3-iI):PC71BM are similar, 1.38 ± 0.5 and 1.25 ± 0.5 μs, respectively. The similar lifetimes between the two blends indicate a similar bimolecular recombination rate despite the morphology difference in these blends. Consequently, bimolecular recombination should not be the detrimental driving force for the lower JSC in P(T1iI):PC71BM devices. In fact, the large domain size observed in P(T1-iI):PC71BM is expected to reduce the probability of exciton dissociation at the heterojunction interface, as a result of higher geminate recombination and lower JSC. It should be noted that in organic bulk heterojunction photovoltaics, geminate recombination could originate either from excitons recombination in the bulk of the polymer or from charge transfer exciton recombination at the donor/acceptor heterojunction interface. However, it is experimentally challenging to separate the two processes. One approach is to compare the quantum yields of photoluminescence and electroluminescence of the blends.25 Figure S3 shows the normalized steady-state long wavelength photoluminescence spectra for both P(T1-iI) and P(T3-iI) along with the PL spectra for P(T1-iI):PC71BM and P(T3-iI):PC71BM blends relative to the corresponding pristine polymers. In both cases, more than 90% reduction in PL intensity is observed for the two blend films, indicating strong PL quenching. In the blend films the peak at 825 nm is attributed to emission due to the bulk, while the peak at 925 nm is attributed to emission due to the CT states. As you can in the figure, 925 nm emission is higher for P(T1-iI):PC71BM compared to P(T3-iI):PC71BM, 2436


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Chemistry of Materials around 0.2 eV to the lower energy than in P(T1-iI):PC71BM. This supports the lower CT energy obtained in P(T3iI):PC71BM by CMEAS as shown above. As shown in Figure S4, according to the photoemission results, assuming a vacuum level aligned at the polymer:PC71BM interface, we obtained effective bandgap Eeff, the energy difference between the HOMO of polymer and the LUMO of PC71BM, for P(T1iI):PC71BM and P(T3-iI):PC71BM to be 1.52 and 1.26 eV, respectively. The difference is in good agreement with the different CT energy determined by CMEAS and EQE. To determine the interface energy level alignment, we assumed vacuum level alignment with no considerable interface dipole.38,39 In order to gain deeper insight into the effectiveness of the CT states in the two BHJ blends, transient PL measurements were performed on these samples. Compared to the steady state PL, the transient PL gives more information on the decay process with corresponding quenching mechanism. Previously, it was shown by Bernardo et al. that increased delocalization of excitons in the CT states will lead to an additional fast decay component in the transient PL.40 Figure 5a shows the PL

CTeff = 1 −

τCT τex

(1)

Using eq 1, we obtained a charge transfer efficiency of about 85% in P(T3-iI):PC71BM blends which is in close agreement with the fast decay component amplitude (79%) obtained from the biexponential fitting mentioned above. Figure 5b depicts the above-discussed exciton decay via radiative (slow) and nonradiative (fast) pathways to the ground state. In P(T1iI):PC71BM, the absence of the fast decay component suggests a low charge transfer efficiency in this blend, similar to that of the pristine polymers. Despite the significant differences in the above observed CT energetics and charge transfer efficiency, it is still unclear if the differences are solely effects or molecular structure differences. In order to verify whether the difference in morphology alone can explain the difference in the CT state cutoff energies, we compared our findings to PTB7:PC71BM. It is well-known that the domain size in PTB7:PC71BM films is highly dependent on the amount of processing additive added to the solution during processing.44 Even though it is well-known that 1,8-diiodooctane (DIO) can significantly reduce the domain size in PTB7:PC71BM blends, the effect of the change in the domain size on the CT manifold has not previously been investigated.9,17,45 As shown in Figure S6, decreasing the domain size by adding DIO did not significantly change the energetic position of the CT manifold. Therefore, it suggests that solely altering the active layer morphology does not necessarily result in a large difference in the energetics of the CT manifold. The energetic position of the CT manifold derived from subbandgap EQE and CMEAS is a result of excitons generated from direct excitation into sub-bandgap energy levels. Since more tightly bound excitons are more difficult to dissociate and exciton binding energy is inversely proportional to dielectric constant, we sought to investigate any possible relationship between the dielectric constant and the CT manifold energy cutoff. The dielectric constants for our films were determined from the geometric capacitance measured as described in the Experimental Section, and the results are summarized in Table 1. While the dielectric constants for the pristine polymer films were nearly identical, the dielectric constant for P(T3iI):PC71BM was found to be approximately 10% higher than the dielectric constant for P(T1-iI):PC71BM. It is interesting to note that for both systems, especially for the P(T3-iI):PC71BM blend, the blend dielectric constant was higher than that of either of the two blend constituents. Figure S7 shows the dielectric constant in the blended films with different fullerene weight ratios. Specifically, we found that the dielectric constant of the blend is about 5 which is significantly higher than the dielectric constant of fullerene. The resulting higher dielectric constant in the P(T3-iI):PC71BM blend indicates enhanced electronic coupling between the polymer and fullerene. To further understand whether the higher dielectric constant is the result of a stronger electronic coupling in the P(T3iI):PC71BM blend, DFT calculations were performed. As shown in Figure S8, the electron density surface of the two polymers is mainly localized in the isoindigo acceptor moiety, in agreement with previous reports on the electronic structure of isoindigo compounds.46,47 Therefore, in order to dissociate the photogenerated excitons, where the electron will still be Coulombically attracted to a hole, the fullerene molecule has to be close to the isoindigo moiety. As illustrated in Figure 6a,

Figure 5. (a) Transient PL for P(T1-iI):PC71BM and P(T3iI):PC71BM blend films. P(T1-iI)):PC71BM shows a single exponential decay. P(T3-iI)):PC71BM shows a biexponential decay. (b) The Jablonski diagram showing the florescence decay (τex) and charge transfer lifetime (τCT) which represents the slow and fast decay observed in transient PL, respectively.

transients for the two isoindigo polymer:fullerene blends. For the P(T1-iI):PC71BM blend, the transient PL showed a single exponential decay with a lifetime of 987 ± 2 ps. On the other hand, for P(T3-iI):PC71BM blends a biexponential PL transient was observed where a fast decay component with a lifetime of 143 ± 1 ps was followed by a slow decay component with a lifetime of 975 ± 17 ps. The fast decay component accounted for 79% of the total PL signal amplitude, and the slow decay component accounted for 21% of the total PL signal amplitude. The slow decay component (τex) with a lifetime of about 1 ns observed in both BHJ blends is within the typical time scale for fluorescence decay in conjugated polymers.41,42 The slow decay lifetime values are almost the same as those obtained in pristine polymers where there is no charge transfer and PL decay is solely due to exciton decay as shown in Figure S5. The pristine PL lifetime for P(T3-iI) and P(T1-iI) were found to be close to 1 ns. The fast decay component (τCT) seen in the PL decay for P(T3-iI):PC71BM, which is not present in pristine polymer PL, is due to efficient charge transfer from a rapid decay of polymer excitons to the CT states at the DA interface. The efficiency of such a transfer process (CTeff) can be estimated using the following relationship previously developed to describe energy transfer in organic semiconductors43 2437


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dielectric constant is observed in the P(T3-il):fullerene blend. In fact, the dielectric constant of the blend is higher than that of the neat polymer or fullerene, indicating a strong donor− acceptor interaction in the blend. Second, the biexponential decay observed in the P(T3-il):fullerene blend indicates a more efficient charge transfer in the blend. These results suggest that a stronger electronic coupling in the P(T3-il):fullerene blend might result from the docking of the fullerene molecule to the isoindigo moiety. We propose that dielectric and transient PL measurements might be good vehicles to predict the photovoltaic properties of polymer fullerene blends. Further studies in other polymer systems are needed to confirm this point. Nevertheless, our results suggest that polymer design facilitating acceptor moiety-fullerene coupling is a key to efficient charge transfer for the future development of high efficiency organic photovoltaic devices.

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Figure 6. (a) Illustration of the molecular structural dependent interaction distance between fullerene and (i)P(T1-iI) and (ii) P(T3iI) polymers. (b) Graphic representation of the effects of the number of thiophenes on the energetics of (i) P(T1-iI):PC71BM and (ii) P(T3-iI):PC71BM systems.

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.6b00824. Normalized absorption spectra for P(T1-iI) and P(T3-iI) polymers, photovoltage transients for P(T1-iI) and P(T3-iI), steady state PL for P(T1-iI):PC71BM and P(T3-iI):PC71BM blends, photoemission spectroscopy of HOMO energies, transient PL for pristine P(T1-iI) and P(T3-iI), sub-bandgap EQE of PTB7:PC71BM with or without DIO, dielectric constants of P(T3-iI):PC71BM blend film with different weight percentage of PC71BM, and DFT calculations (PDF)

both polymers have branched side chains on the isoindigo moiety, which may prevent the fullerene acceptor from coming in close contact with the polymer. However, the longer thiophene donor in P(T3-iI) acts as a spacer between the bulky side-chains and may allow fullerene molecules to come in closer contact with an isoindigo acceptor moiety. The proximity of fullerenes to sterically hindered isoindigos is possibly driven by strong amide-fullerene interactions, resulting in better electronic coupling and more efficient charge transfer between the P(T3-iI) polymer and the fullerene compared to P(T-iI)based blends. This is in agreement with what has recently been reported by Graham et al. on a preferred intermolecular arrangement in high-efficiency polymer:fullerene systems, where the fullerene is generally docked to the electronaccepting moiety of the polymer.48 Subsequently, the stronger polymer−fullerene electronic coupling would be the cause of larger blend dielectric constant and more efficient charge transfer in P(T3-iI):PC71BM. More detailed investigation would be needed to pinpoint the underlying mechanism. Nevertheless, it has been previously shown that an increased dielectric constant can cause a CT state shift to lower energies.49,50 Our results are in agreement with the CT state manifold shift observed in this study. According to the above results, the effects of the number of thiophenes in the polymer backbone on the energetics of the two isoindigo blends are summarized in Figure 6b. The difference in electronic coupling between the two isoindigo polymers with fullerene, alters not only the CT energy but also the charge transfer efficiency at the heterojunction interface.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Present Address #

Department of Materials Science and Engineering, North Carolina State University, Raleigh, NC, USA. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors would like to acknowledge the support of Office of Naval Research (Award # N00014-14-1-0173). REFERENCES

(1) Cheng, Y.-J.; Yang, S.-H.; Hsu, C.-S. Synthesis of Conjugated Polymers for Organic Solar Cell Applications. Chem. Rev. 2009, 109, 5868−5923. (2) Lu, L.; Zheng, T.; Wu, Q.; Schneider, A. M.; Zhao, D.; Yu, L. Recent Advances in Bulk Heterojunction Polymer Solar Cells. Chem. Rev. 2015, 115, 12666−12731. (3) Beaujuge, P. M.; Amb, C. M.; Reynolds, J. R. Spectral Engineering in π-Conjugated Polymers with Intramolecular Donor− Acceptor Interactions. Acc. Chem. Res. 2010, 43, 1396−1407. (4) Chen, J.-D.; Cui, C.; Li, Y.-Q.; Zhou, L.; Ou, Q.-D.; Li, C.; Li, Y.; Tang, J.-X. Single-Junction Polymer Solar Cells Exceeding 10% Power Conversion Efficiency. Adv. Mater. 2015, 27, 1035−1041. (5) Constantinou, I.; Lai, T.-H.; Zhao, D.; Klump, E. D.; Deininger, J. J.; Lo, C. K.; Reynolds, J. R.; So, F. High Efficiency Air-Processed Dithienogermole-Based Polymer Solar Cells. ACS Appl. Mater. Interfaces 2015, 7, 4826−4832.

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CONCLUSIONS In conclusion, the effects of the thiophene donor length on film morphology and device performance were investigated for two thiophene-isoindigo copolymers. Despite the similar chemical structures with nearly identical energetics and photophysical properties, the device characteristics for the two polymer:fullerene blends were found to be significantly different. Here, we have made two key observations in the P(T3-il):fullerene device with a higher short-circuit current compared to the P(T1-il):fullerene device. First, a higher 2438


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