Alkyl Branching Position in Diketopyrrolopyrrole Polymers: Interplay

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Alkyl Branching Position in Diketopyrrolopyrrole Polymers: Interplay between Fibrillar Morphology and Crystallinity and Their Effect on Photogeneration and Recombination in Bulk-Heterojunction Solar Cells

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Rishi Shivhare,†,‡ Tim Erdmann,†,§,▲ Ulrich Hörmann,∥ Elisa Collado-Fregoso,∥ Stefan Zeiske,∥,△ Johannes Benduhn,⊥ Sascha Ullbrich,⊥ Rene ́ Hübner,# Mike Hambsch,†,‡ Anton Kiriy,§ Brigitte Voit,†,§ Dieter Neher,∥ Koen Vandewal,*,⊥,○ and Stefan C. B. Mannsfeld*,†,‡ †

Center for Advancing Electronics Dresden (cfaed), Technische Universität Dresden, Helmholtzstraße 18, Dresden 01069, Germany Faculty of Electrical and Computer Engineering, Technische Universität Dresden, Helmholtzstraße 18, Dresden 01069, Germany § Leibniz-Institut für Polymerforschung Dresden e.V. (IPF), Hohe Straße 6, Dresden 01069, Germany ∥ Department of Physics and Astronomy, University of Potsdam, Karl-Liebknecht-Straße 24−25, Potsdam-Golm 14476, Germany ⊥ Dresden Integrated Center for Applied Physics and Photonic Materials (IAPP) and Institute for Applied Physics, Technische Universität Dresden, Nöthnitzer Straße 61, Dresden 01187, Germany # Institute of Ion Beam Physics and Materials Research, Helmholtz-Zentrum Dresden-Rossendorf, Bautzner Landstraße 400, 01328 Dresden, Germany ‡

S Supporting Information *

ABSTRACT: Diketopyrrolopyrrole (DPP)-based donor−acceptor copolymers have gained a significant amount of research interest in the organic electronics community because of their high charge carrier mobilities in organic field-effect transistors (OFETs) and their ability to harvest near-infrared (NIR) photons in solar cells. In this study, we have synthesized four DPPbased donor−acceptor copolymers with variations in the donor unit and the branching point of the solubilizing alkyl chains (at the second or sixth carbon position). Grazing incidence wide-angle X-ray scattering (GIWAXS) results suggest that moving the branching point further away from the polymer backbone increases the tendency for aggregation and yields polymer phases with a higher degree of crystallinity (DoC). The polymers were blended with PC70BM and used as active layers in solar cells. A careful analysis of the energetics of the neat polymer and blend films reveals that the charge-transfer state energy (ECT) of the blend films lies exceptionally close to the singlet energy of the donor (ED*), indicating near zero electron transfer losses. The difference between the optical gap and open-circuit voltage (VOC) is therefore determined to be due to rather high nonradiative (≈ 418 ± 13 mV) and unavoidable radiative voltage losses (≈ 255 ± 8 mV). Even though the four materials have similar optical gaps, the short-circuit current density (JSC) covers a vast span from 7 to 18 mA cm−2 for the best performing system. Using photoluminescence (PL) quenching and transient charge extraction techniques, we quantify geminate and nongeminate losses and find that fewer excitons reach the donor−acceptor interface in polymers with further away branching points due to larger aggregate sizes. In these material systems, the photogeneration is therefore mainly limited by exciton harvesting efficiency.

1. INTRODUCTION In recent years, polymer solar cells have progressively become more efficient, with highest power conversion efficiencies (PCEs) reaching 13.1% for single-junction1 and around 15% © XXXX American Chemical Society

Received: June 28, 2018 Revised: September 10, 2018

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DOI: 10.1021/acs.chemmater.8b02739 Chem. Mater. XXXX, XXX, XXX−XXX

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Chemistry of Materials for multijunction devices.2 To some extent, this remarkable improvement can be attributed to the advancement in synthesis methods, creating an ever-growing repository of polymer semiconductors. Among newly developed materials, state-ofthe-art copolymers are based on the concept of a push−pull system with alternating electron rich and electron deficient units. This design concept is especially appealing because it allows fine-tuning of frontier energy levels and creates polymers with a desired optical gap. Within the class of push−pull copolymers, diketopyrrolopyrrole (DPP)-based low band gap materials have been studied extensively for photovoltaic applications.3,4 Using DPP polymers as an electron donor and a fullerene derivative as an electron acceptor, impressive PCEs of up to 9% have been achieved in single-junction devices.5,6 Moreover, their small optical gap (Eopt ≤ 1.5 eV) and their absorption spectra extending into the near-infrared (NIR) range makes them a suitable candidate to be used in tandem solar cells, which led to good PCEs in multijunction devices.7,8 Although DPP polymers have proven to be eminently suitable materials for solar cell applications their solution processing is often challenging, primarily due to their limited solubility in common organic solvents. The two main aspects which dictate the polymer solubility are the molecular weight and the regiochemistry of the solubilizing alkyl side chains. In the case of DPP polymers, branched side chains are required to render the polymers soluble. The overall chain length and the branching point of the alkyl chains are expected to have a significant impact on the polymer solubility and the film microstructure.8 Additionally, the branching point of side chains will also influence the crystallization of the polymer backbone.9−11 A branching point too close to the polymer backbone can lead to steric hindrance effects and deteriorate the planarity of the backbone,12 while branching further away from the backbone can cause irregularities in packing of the alkyl chain themselves. Thus, finding an optimum side chain configuration, which offers good solubility and promotes backbone crystallization, can be difficult and will depend on the nature of the conjugated core of the polymer. The bulk-heterojunction morphology of DPP:fullerene blends consists of semicrystalline polymer fibrils (donor phase), fullerene islands (acceptor phase) and amorphous mixed phases.13 Janssen et al. have empirically shown that the polymer solubility and the polymer fibril size are interdependent;14 with less soluble polymers forming narrower fibrils. The same research group investigated the influence of the side chain length on the morphology and the overall photovoltaic performance.15 Their results show that there is a direct correlation between the polymer fibril width (related to polymer solubility) and the maximum photovoltaic external quantum efficiency (EQEPV) values. Their study further suggests that photogeneration is limited by exciton diffusion, and high EQEPV values demand fibril widths lower than the exciton diffusion length. Meager et al. studied the influence of alkyl branching position in thieno [3,2-b]thiophene substituted DPP’s (DPPTT-T).16 In their study, solar cells showed a better performance upon moving the branching point away from the backbone mainly through improvement in photocurrent. They attributed the improved device performance to the enhanced crystallinity of the blended polymer films. Investigations on the alkyl branching position in other polymer systems have also revealed its critical influence on the polymer crystallinity and device performance.17,18 However, an overall understanding of the effect of branched alkyl chains on the device physics of solar cells is still lacking. In this work, we elucidate the impact of different alkyl chain configurations on

the blend morphology and how it impacts the overall energetics of the polymer:fullerene system and the photogeneration process. We find that the branching point of side chains affects the polymer fibril diameter which in turn influences the exciton harvesting efficiency at the donor−acceptor interface. However, field-independent free charge carrier generation as measured by time-delayed-collection-field (TDCF) measurements show that geminate recombination of CT excitons is not a limiting process. Besides exciton harvesting, the branching point of the side chains also influences the charge carrier mobility (μ) and carrier lifetime (τ) which impacts the sweep-out of the free charge carriers and determines the fill-factor of the photovoltaic devices.

2. RESULTS AND DISCUSSION 2.1. Materials and Device Characteristics. For this study, a series of DPP based donor−acceptor copolymers (Figure 1a) were synthesized using Stille coupling polycondensation, on the basis of previous reports.19−21 The DPP core unit was used as an acceptor moiety while thiophene (T) or thieno[3,2-b]thiophene (TT) units flanked by one additional thiophene unit per side were used as donor conjugation segments. Additionally, we varied the chain length and the branching point of the solubilizing alkyl chains. 2-octyldodecyl (2OD) or bulky 6-dodecyloctadecyl (6DO) side chains with branching at sixth carbon position were used. The main design idea was to move the branching point further away from the backbone [P(DPP2ODT2-X) vs (P(DPP6DOT2-X)] to facilitate stronger intermolecular π-orbital interactions between individual polymer chains and, thus to create copolymers with reduced π-stacking distances and enhanced degree of crystallinity. Number-average molecular weights (Mn) were derived from high-temperature gel permeation chromatography (HT-GPC) measurements performed in 1,2,4-trichlorobenzene at 150 °C. P(DPP6DOT2-T) and P(DPP2ODT2-T) have a comparable Mn of 19.4 kDa and 16.8 kDa, respectively. Thieno[3,2-b]thiophene-(TT)-substituted DPP copolymers exhibit significantly higher and very similar molecular weights. P(DPP6DOT2-TT) and P(DPP2ODT2-TT) showed a Mn of 92.0 kDa and 91.2 kDa, respectively. Additional thermal and optoelectronic material properties are summarized in Table S1; the thermal stability and the energetic position of frontier molecular orbitals showed no significant difference within the P(DPP-Y-T2-X) series. Figure 1b shows the absorption coefficients for the polymer: fullerene blend films. More pronounced vibronic features provide a proof for an increased tendency of aggregation for 6DO branched polymer backbones. This is related to the higher degree of interchain aggregation, facilitated by a farther branching point of the alkyl chains.22 Absorption and emission spectra were recorded to accurately calculate the optical gap of the polymer thin films (Figure S2).23 All four polymers have a similar optical gap (Eopt)of about 1.33 eV, obtained from the crossing point of the appropriately normalized reduced absorption and reduced emission spectra. The fact that the optical gap is not influenced by the side chain substitution suggests that alkyl groups have no significant influence on the electron density distribution across the conjugated core. We blended the polymers with the PC70BM-fullerene derivative (1:3 wt/wt) in a solvent mixture of chloroform and diphenylether (3 vol %) for use as active layer in organic solar cells. Devices were fabricated under ambient conditions in an inverted configuration, in the following device architecture: ITO/PEIE (5 nm)/Active Layer/MoOX (6.5 nm)/Ag (100 nm). Overall, devices that use the polymer with a -TT conjugation segment and 2OD side chains performed B

DOI: 10.1021/acs.chemmater.8b02739 Chem. Mater. XXXX, XXX, XXX−XXX

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Figure 1. (a) Chemical structure of the DPP polymers with different side chains and conjugation units. (b) Absorption coefficients of the polymer:PC70BM blend films. (c) Dark (dashed lines) and illuminated (solid line) J−V curves and the external generation efficiency (EGE) at 650 nm of the representative devices. (d) EQEPV of the optimized devices.

Table 1. Device Characteristicsa, Maximum EQEPV (polymer region), Photoluminescence Quenchingb of the Different DPP Polymers Blended with PC70BM (1:3) donor

Jsc (mA/cm−2)

VOC (mV)

fill factor (%)

PCE (%)

EQEPV,polymer (%)

PL-quenching (%)

P(DPP6DOT2-T) P(DPP2ODT2-T) P(DPP6DOT2-TT) P(DPP2ODT2-TT)

7.3 ± 0.6 11.3 ± 0.3 11.9 ± 0.2 17.2 ± 0.3

656 ± 3 680 ± 5 634 ± 3 648 ± 2

68.6 ± 1.5 64.8 ± 0.7 64.6 ± 0.3 60.0 ± 0.8

3.1 ± 0.3 5.0 ± 0.3 4.9 ± 0.1 6.5 ± 0.3

12.0 22.4 24.5 40.0

15 25 29 48

a

Values averaged over 18 devices. bExcitation wavelength (λ = 850 nm) to selectively excite polymer domains.

the best, achieving an average PCE ≈ 6.5%. P(DPP6DOT2-TT) and P(DPP2ODT2-T) reached a similar PCEs of around 5%, whereas P(DPP6DOT2-T) has the lowest PCE of 3.1% (Table 1). Open-circuit voltage (Voc) values for the four polymer systems were similar, all in the range of 630−680 mV. Interestingly, Jsc varied over a wide range for the four blends, despite the similar Eopt of the donor polymers. Devices based on P(DPP2ODT2-TT) had an average Jsc ≈ 17 mA cm−2; P(DPP6DOT2-TT) and P(DPP2ODT2-T) exhibit Jsc ≈ 11−12 mA cm−2, whereas P(DPP6DOT2-T) based solar cells had an average Jsc of only 7 mA cm−2 (Table 1). Figure 1d shows the EQEPV of the devices. In addition to the generation process, EQEPV values also take into account nongeminate recombination of charge carriers at short-circuit and reflect the quanta of charge carriers collected at the selective electrodes. In all four investigated blends, EQEPV values are higher in the spectral region between 300 and 650 nm, which is where mainly PC70BM absorbs. Relatively low EQEPV values on the polymer side imply that the DPP’s near-infrared absorption capability is not used to its full potential. This effect has previously been reported for solar cells using DPP-based polymers and was attributed to different charge-generation efficiencies for the photoexcited electron transfer from the donor to the acceptor and the hole transfer from the acceptor to the donor.24 To investigate the effect of geminate losses on the photocurrent, we carried out time-delayed collection field (TDCF) measurements (Figure 1c) at short times (≈6 ns) and low

excitation intensities. TDCF is an optoelectronic pump−probe technique described elsewhere.25,26 Briefly, our TDCF setup uses a 3.8 ns pulsed laser to photogenerate charges in a working device that is held at a certain prebias, Vpre. A high reverse bias (−Vcoll) is used as an electrical probe to collect the photogenerated charge carriers after very short delay times. In our experiments, the delay time between optical pump and the electrical probe was kept sufficiently small (≈6 ns) and the pulse intensity was very low (0.03 μJ cm−2, at 650 nm) to avoid any early nongeminate recombination. Therefore, the charges collected at reverse bias under these conditions can be used to calculate an external generation efficiency (EGE) at different prebiases and to assess the bias dependence of charge generation. As can be observed in Figure 1c, the EGE was found to be fairly independent of the applied bias, with all polymer:fullerene systems retaining more than 90% of the generation efficiency under open-circuit conditions (zero internal field) compared to the generation efficiency at −1.5 V. This suggests that geminate recombination of polaron pairs (or CT states) is not a limiting factor in photogeneration of charge carriers. In terms of the overall charge carriers generated, P(DPP2ODT2-TT) has the highest generation efficiency followed by P(DPP6DOT2-TT) and P(DPP2ODT2-T), with P(DPP6DOT2-T) showing the lowest generation efficiency, following the same trend that is observed in Jsc. Regarding the fill factor of the devices, the least efficient system: P(DPP6DOT2-T) exhibits the highest fill factors over a wide range of film thicknesses (Figure S8). Most C

DOI: 10.1021/acs.chemmater.8b02739 Chem. Mater. XXXX, XXX, XXX−XXX

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Figure 2. (a) Current density (J) transients measured after a delay time of 3 μs. (b) Mobility−lifetime products (μτ)eff as a function of charge carrier density. (c) Charge-carrier density plotted as a function of delay time (td). (d) Schematic depicting trapped electrons in shallow tail states. Trapped carriers need activation to become mobile and thus available for recombination.

values that have previously been reported for this material class.28 P(DPP6DOT2-T) has the highest μeff with an average value around 7 × 10−3 cm2V−1 s−1, while other polymers exhibit marginally lower values with P(DPP2ODT2-TT) having the lowest average mobility of around 2 × 10−3 cm2 V−1 s−1. DPPs with farther branching point (6DO) showed higher mobilities compared to the DPPs with closer branching position (2OD). This can be attributed to the closer π-stacking and a higher degree of crystallinity. Apart from effective mobility values, OTRACE also gives an estimate of the charge carrier lifetime (τeff). In Figure 2b, we plot mobility-lifetime products (μτ)eff as a function of carrier density. It is a useful figure of merit illustrating the competition between charge collection and recombination of carriers. These counteracting processes strongly influence the fill factor of the solar cell. Overall, the (μτ)eff product values obtained for these DPP polymers are among the highest reportedin polymer:fullerene blends with average values around 1 × 10−5 to 1 × 10−7 cm2 V−1 s−1. The (μτ)eff productfor P(DPP6DOT2-T) was an order of magnitude higher compared to P(DPP2ODT2TT). This explains higher average fill factors (FF) for P(DPP6DOT2-T) (FF ≈ 65−70%) as compared to P(DPP2ODT2-TT) (FF ≈ 55−60%) over a wide range of film thicknesses. The nongeminate recombination rate (R) can be expressed using the equation of the form:

likely, the reason for this lies in the material’s high degree of crystallinity (Figure 3c) and its high charge carrier mobility (see section on charge transport), which is in fact the highest among the four polymers. This points to an important fact, namely, that polymer-design-rules for organic solar cells should not exclusively focus on increasing the crystallinity and thus the charge carrier mobility of the blends, but should also consider the polymer’s ability to phase-separate at the appropriate length scale. 2.2. Charge Transport and Nongeminate Recombination Dynamics. We used OTRACE (open-circuit corrected charge carrier extraction) to investigate charge transport and recombination dynamics of the solar cells under real operation conditions. Usually, transient charge extraction measurements consist of three time periods: (i) illumination pulse: for photogeneration of charge carriers; (ii) delay-time (td): allowing for the recombination of a certain amount of charge carriers; and (iii) linear extraction pulse: to collect remaining photogenerated charge carriers. OTRACE27 is a state-of-the-art charge extraction technique where an adaptive-field control is used to inhibit charge injection or extraction during the delay-time (td) before the extraction pulse is applied. This limits the effect of capacitance and hence the RC time of the device to minimum (pure geometric capacitance) and provides better temporal resolution compared to photo-CELIV. Figure 2a shows representative current transients for devices measured after a specific delay time (td = 3 μs). From the peak position of the transient’s effective mobility of charge carriers can be extracted. For all the polymers, the effective charge carrier mobility (μeff) we obtained is on the order of 1 × 10−3 cm2 V−1s−1, which is in good agreement with

R = knα+ 1

(1)

where k is the recombination prefactor and n is the charge carrier density. By tracing the time evolution of charge carrier density (n) (Figure 2c), the recombination order (α+1) can be calculated. D

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Figure 3. (a) 2D-GIWAXS images for the blend films: P(DPP6DOT2-T), P(DPP2ODT2-T), P(DPP6DOT2-TT), P(DPP2ODT2-TT) (from left to right). (b) Radially integrated π-stacking peak positions. (c) (010) pole figures for P(DPP6DOT2-T):PC70BM and P(DPP2ODT2-T):PC70BM films supplying information on the relative degree of crystallinity (rdoc). For the P(DPP6DOT2-T) and P(DPP2ODT2-T), the rdoc value is 44%.

For delay times >10 μs, power-law decay (n ∼ t −1/ α) was observed. From a fit, we obtain recombination orders (α+1) of 3.56 for 6DO−DPP’s and 3.1 for 2OD-DPP’s. The difference in recombination cross-section can be related to the varying fibrillar morphology which changes the effective donor−acceptor interfacial area where mobile charge carriers can recombine. Moreover, recombination orders (α+1) > 2 (beyond the classical Langevin case) have previously been associated with disordered density of states (DOS), carrier concentration dependent mobility29,30 and recombination in the presence of traps.31−33 Recent model published by Hofacker et al.34 treats simultaneous thermalization and recombination of charge carriers. Thermal activation of localized carriers in tail states is what contributes to higher (>2) recombination orders (Figure 2d). Interestingly, decent fill factors were obtained for all the polymers (>60%), which points to the fact that trap-mediated recombination might not be a detrimental loss mechanism in systems with sufficiently high mobilities. 2.3. GIWAXS Analysis: π-Stacking and Degree of Crystallinity. We performed grazing incidence wide-angle X-ray scattering (GIWAXS) measurements to investigate the crystallinity of the neat polymer and blend films. For the neat films (Figure S3), all four polymers exhibit dominantly face-on orientation of crystallites with the π-stacking related scattering signal being most intense in the out-of-plane (qz) direction. For the polymers P(DPP6DOT2-TT) and P(DPP2ODT2-TT), there is also a weak π-stacking peak in the in-plane direction (qxy), indicating a small volume fraction of edge-on crystallites. Polymers with bulky 6DO side chains had a π-stacking distance ∼3.52 Å (q (010) = 1.79 Å−1) and lamellar spacing ∼28 Å, whereas polymers with 2OD side chains had a larger π-stacking distance ∼3.63 Å (q(010) = 1.73 Å−1) and a lamellar spacing ∼16 Å. These results indicate that for the polymers P(DPP6DOT2-T) and P(DPP6DOT2-TT) whose alkyl branching-point lies farther

from the backbone, the aromatic chains pack closer whereas the different DPP-backbone motifs did not apparently influence the π-stacking distance. Unlike the case of inorganic crystals, it is difficult to compute the exact crystallite size in semicrystalline polymer films from the scattering signals due to defects and paracrystalline disorder.35 Instead, a coherence length can be extracted that gives an approximate idea of the crystallite size. We used Schererr analysis, to compute the coherence length in the π-stacking direction. P(DPP6DOT2-T) exhibits the largest coherence length of around 41 Å, whereas P(DPP2ODT2-TT) had the shortest coherence length of around 23 Å. Polymers P(DPP2ODT2-T) and P(DPP6DOT2-TT) showed similar coherence length of about 30 Å. Interestingly, the π-stacking coherence length follows the same trend as the average polymer fibril diameter measured using transmission electron microscopy (TEM) images (Figure 4a). On this basis, one can speculate that the polymer chains are predominantly arranged such that the π-stacking direction roughly coincides with the short fiber axis. For the blend films, the π-stacking signal was considerably weaker than for the neat films. This might be due to the intercalation of fullerene within polymer domains leading to a reduction in crystalline order and registry between the polymer backbone planes. To calculate the orientation distribution of crystallites, we mapped the intensity of the π-stacking peak as a function of polar angle (χ). Here χ = 0° corresponds to strictly out-of-plane scattering signal resulting from face-on crystallites while χ = 90° corresponds to in-plane scattering related to edge-on crystallites.The detailed procedure of the pole figure analysis is described in earlier reports.36−38 For the thieno[3,2-b]thiophene-based DPPs, a reliable determination of (010) pole figures was difficult due to the weak scattering signal. However, pole figures of the π-stacking peak were reconstructed for terthiophene based DPPs [P(DPP6DOT2-T) and P(DPP2ODT2-T)] and reveal a bimodal distribution of E

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Figure 4. (a) TEM micrographs for spin-cast polymer:fullerene blend films. (b) PL-quenching: PL spectra of neat polymer and blend films. Excitation wavelength (λ) = 850 nm.

films at λ = 532 nm (Figure S6), the spectral region where mainly PCBM absorbs. As the appropriately normalized PL signal scales with average domain size, the PL measurements provide a relative measure for the PCBM domain sizes. Blend films with P(DPP2ODT2-TT) showed the lowest PL intensity followed by P(DPP6DOT2-TT) and P(DPP2ODT2-T) while P(DPP6DOT2-T) had the highest PL signal. This suggest that different fibrillar network also affects the size distribution of the fullerene domains, with fibril widths and fullerene domain sizes following the same trend. To study the effect of the fibrillar morphology and specifically the fiber-width and purity on photogeneration we performed PL-quenching measurements. The degree of PL quenching correlates with the fraction of excitons that dissociate at the donor−acceptor interface. To selectively excite polymer domains, we used an infrared pump LED with a peak emission at 850 nm. The PL signal was corrected for absorption over the spectral range of the LED in order to compare the relative PL count of the pure polymer and blend films. All the blend films showed poor luminescence quenching. The PL quenching efficiency of P(DPP6DOT2-T) was surprisingly low: only 15% of the neat polymer signal was quenched (Figure 4b). P(DPP2ODT2-T) and P(DPP6DOT2-TT) represent an intermediate case with quenching efficiencies of around 25% and 29% respectively. The best performing system, P(DPP2ODT2-TT), showed quenching efficiency of about 48% (Table 1). Although the PLquenching efficiency correlates well with the polymer fibril diameter as reported earlier,15 the overall PL quenching is rather low as compared to other high performing polymer:fullerene systems.40,41 There are two possible reasons for this poor PL quenching. First, DPP based polymers exhibit rather short exciton lifetimes (∼50 ps) which allows a large fraction of polymer excitons to relax back to the ground state before reaching the donor−acceptor interface.42 Second, in these low band gap polymers, the energy of the charge transfer state (ECT) lies close to the singlet energy of the donor. This may cause the CT states to repopulate the donor excited states, thus forming an indirect radiative pathway for CT states.43 In our case, the

crystallites with both face-on (χ = 0°) and edge-on crystallites (χ = 90°) contributing to the π-stacking peak. Because the scattering signal mainly arises from the crystalline part of the polymer film, the area under the pole figure curve gives an idea about the relative proportion of crystalline and amorphous regions in the film, i.e., the degree of crystallinity (DoC) of the film:35 DoC ∝

∫0

π /2

I(χ )sin(χ )dχ

(2)

Figure 3c showcases the higher crystallinity for P(DPP6DOT2T) relative to P(DPP2ODT2-T) corroborating the rationale that moving the alkyl branching point further away from the backbone (−6DO side chains) increases the tendency for aggregation and yields polymer phases with a higher DoC. 2.4. Fibrillar Morphology and Photoluminescence Quenching. TEM images (Figure 4a) depict the microstructure of the polymer:fullerene blend films. The bulk-heterojunction morphology consists of a dense network of semicrystalline polymer fibrils and interspersed fullerene domains. Earlier work by Janssen et al. suggests a direct correlation between polymer solubility and the fiber diameter of the polymer domains. The family of DPP-polymers investigated in this study also follow this empirical observation where polymers with bulky 6DO side chains (higher solubility) have wider fibers compared to polymers with smaller 2OD side chains. To analytically determine the fibril size, we analyzed TEM micrographs using the Image J software with the plug-in Diameter J39 to obtain statistical distributions of the fibril widths (Figure S4). Diameter J uses a series of image segmentation algorithms to yield binary images that are then processed using an Euclidean distance transformation algorithm, which yields a fiber diameter distribution. Mean fiber diameters for P(DPP6DOT2-T) and P(DPP6DOT2-TT) were around 25.5 and 16.2 nm, respectively, whereas for P(DPP2ODT2-T) and P(DPP2ODT2-TT), mean fiber diameters were 15.2 and 13.3 nm. Moreover, to qualitatively compare the average size of fullerene domains, we measured pholuminescence (PL) signal by excitation of blend F

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Chemistry of Materials field-independent PL yield (Figure S7) and field-independent charge generation measured via TDCF indicates an efficient dissociation of CT states into free carriers, which makes repopulation of singlets via CT states an unlikely process. Therefore, we conclude short exciton lifetimes to be the predominant cause for the observed poor PL quenching of all tested materials. Overall, even though DPP polymers absorb a large fraction of the electromagnetic spectrum, most of the excitons are not harvested. The situation is aggravated in systems with coarser morphology, i.e., bigger polymer fibrils, as shown by the data in Figure 4. 2.5. Charge-Transfer State Energy (ECT) and Voltage Losses. Even though the morphology and photocurrents are very different among these four donor polymers blended with PC70BM, despite their similar optical band gaps, the corresponding device Voc values are rather similar. The three physical processes that entail the overall potential loss of free charge carriers are (i) electron transfer losses (determined by the energetic offset between the lowest singlet energy (ED*) and ECT: ΔECT = ED* − ECT); (ii) radiative; and (iii) nonradiative voltage losses. In the following section, we determine ECT for the DPP:fullerene solar cells and quantify the voltage losses of the devices. (a) Determination of ECT: In recent years, it has been established that free-carrier generation and recombination in organic solar cells are mediated via the CT state.44,45 ECT dictates the electron transfer losses and also represents the theoretical upper limit for the open-circuit voltage VOC (at T = 0 K). Although the CT state is directly coupled to the ground state, the low oscillator strength demands sensitive measurement techniques for its determination. We use sensitive-EQEPV measurements in the sub-bandgap region combined with electroluminescence (EL) spectra of solar cells driven at low forward voltages to determine ECT, which is defined as the crossing point of sensitive-EQEPV and EL spectra. In our system, the singlet energy of the donor (ED*) lies relatively close to the charge-transfer energy (ECT). Thus, the broad emission and absorption bands associated with the charge-transfer transition are often masked by the emission and absorption tails of the neat materials. Figure 5 depicts the normalized sensitive EQEPV and EL spectra for optimized devices. The broad CT state transitions cannot be distinguished from the transition of the neat donor. Nevertheless, the CT energy can be identified from the crossing point of the reduced sensitive EQEPV and the reduced EL spectra. ECT for all the DPP:fullerene blends is nearly of the same magnitude ECT = 1.33 ± 0.05 eV. Intriguingly, the ECT values for the blend films are marginally higher (≈ 30−50 meV) than the optical gap of the neat polymers which were obtained similarly (see Figure S2). This can be attributed to less aggregation of the polymer chains in the blend films, shifting the effective gap of the blended films to higher energies. More importantly, these findings point to the fact that ECT lies very close to the donor singlet level (ED*), implying nearly zero electron transfer losses (ED* − ECT ≈ 0 eV). (b) Estimation of radiative and nonradiative losses: Apart from electron transfer losses, VOC is limited by the radiative and nonradiative recombination of free carriers. A previously developed formalism44,46 relates the VOC to ECT and radiative and nonradiative losses: VOC =

ECT − ΔVrad − ΔVnon − rad q

Figure 5. Normalized reduced sensitive EQEPV and EL spectra for the optimized devices. Dashed lines are the Gaussian fits (fit parameters: f, coupling strength; λ, reorganization energy; and ECT). ECT can be estimated from the crossing point of the sensitive EQEPV and EL spectra.

The second and third terms of the equation represent radiative (ΔVrad) and nonradiative (ΔVnonrad) voltage losses of free carriers, respectively: ΔVrad =

ECT − VOC,rad q

ΔVnon − rad = VOC,rad − VOC

(4) (5)

Voc,rad is the thermodynamic limit of VOC that would be reached in the absence of nonradiative recombination. VOC,rad can be calculated from the sensitive EQEPV spectrum by employing the reciprocity relation between absorption and emission.44,47 Radiative and nonradiative losses are nearly identical among all tested DPP systems: 255 mV (radiative) and ∼418 mV (nonradiative). These calculation results suggest that nonradiative recombination pathways dominate the overall voltage losses and are slightly higher than typical values of around 350 mV.48 According to the energy-gap law,49 high nonradiative decay rates in DPP polymers are related to their low optical band gap which might intrinsically limit the maximum achievable Voc in DPP solar cells.

3. CONCLUSIONS In conclusion, we have synthesized a series of DPP polymers with modifications of the alkyl chain length and the branching point of the side chains and studied their suitability for bulk heterojunction solar cells. The alkyl branching position

(3) G

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influences the aggregation tendency and polymers with the branching point farther away from the backbone exhibit a higher degree of crystallinity. The overall alkyl chain structure also influences the polymer’s solubility and affects the polymer fibril formation. TEM images reveal polymer fibrils with varying diameters. These structural differences have a strong impact on the photogeneration of charge carriers. By time-delayed collection field measurements, we were able to show field-independent generation, which points to the fact that the geminate recombination of charge-transfer excitons is not a major loss mechanism in the photogeneration process. The crystallinity of the polymer phase has a direct implication on the effective mobility of the blends. 6DO−DPPs show higher effective mobility-lifetime values over a broad range of charge carrier densities. Consequently, 6DO−DPPs also exhibit higher fill factors compared to 2OD-DPPs. Concerning voltage losses, the relative proximity of ED* and ECT points to the fact that there is nearly zero photon energy loss in the electron transfer process. Most of the photon energy loss in DPP polymers can be attributed to radiative and nonradiative decay pathways, especially nonradiative recombination losses are higher than for other donor: acceptor systems, which is related to the small optical gap of DPP polymers (according to the energy-gap law). OTRACE results indicate that nongeminate recombination in these DPP polymers is affected by traps and the overall recombination rate is influenced by the average size of the polymer fibrils. All these factors show that despite improving the crystallinity of the polymer phase and the charge carrier mobility, the solar cell performance was intrinsically limited by inefficient exciton harvesting. This work once again shows the complexity of bulk-heterojunction solar cells and that improving one aspect, in this case the charge transport properties, will not necessarily lead to an overall improvement in device performance, reminding the community that polymer design rules for photovoltaic applications need careful consideration and delineation from design rules for transistor applications.

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T.E. is currently at IBM Almaden Research Center, San Jose, CA, USA.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors acknowledge support by the German Excellence Initiative via the Cluster of Excellence EXC 1056 “Center for Advancing Electronics Dresden” (cfaed). For the GIWAXS measurements, the authors acknowledge KMC-2 diffraction beamline of the Photon source BESSY-II, Helmholtz Zentrum Berlin. Additionally, for TEM microscopy, the authors acknowledge Dr. Petr Formanek from the Leibniz-Institut für Polymerforschung Dresden. T.E. acknowledges support by the German Alexander von Humboldt foundation. J.B., S.U., and K.V. acknowledge support from the German Federal Ministry for Education and Research (BMBF) through the InnoProfile Projekt “Organische p−i−n Bauelemente 2.2” (03IPT602X). Furthermore, S.U. acknowledges support by the graduate academy of the TU Dresden, financed by the excellence initiative of the German federal and state governments. E.C.-F. and U.H. and D.N. acknowledge funding by the BMBF (UNVEIL, FKZ 13N13719) and the DFG (SFB 951 “HIOS”).

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.8b02739.

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REFERENCES

Instrumentation; materials and synthesis; UV−vis; optical gap determination; GIWAXS; fibril width distribution; TEM-EDX; PL data; JV data; OTRACE mobility (PDF)

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Elisa Collado-Fregoso: 0000-0002-2870-6983 Sascha Ullbrich: 0000-0003-1564-8355 Anton Kiriy: 0000-0002-4263-9377 Dieter Neher: 0000-0001-6618-8403 Stefan C. B. Mannsfeld: 0000-0003-0268-519X Present Addresses △

S.Z. is currently at Department of Physics, Swansea University, Singleton Park, Swansea, Wales SA2 8PP, UK ○ K.V. is currently at Instituut voor Materiaalonderzoek (IMO), Hasselt University, Wetenschapspark 1, BE-3590, Diepenbeek, Belgium H

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