Development of Highly Crystalline Donor–Acceptor-Type Random

Sep 26, 2017 - We developed donor–acceptor (D–A)-type random polymers based on 3,3′-difluoro-2,2′-bithiophene with various relative amounts of...
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Development of Highly Crystalline Donor−Acceptor-Type Random Polymers for High Performance Large-Area Organic Solar Cells Jae Hoon Yun,†,‡ Hyungju Ahn,§ Phillip Lee,† Min Jae Ko,∥ and Hae Jung Son*,†,‡ †

Photo-electronic Hybrids Research Center, Korea Institute of Science and Technology, Seoul 136-791, Republic of Korea Division of Energy and Environment, KIST School, University of Science and Technology (UST), Daejeon 34113, Republic of Korea § Pohang Accelerator Laboratory, Kyungbuk, Pohang 37673, Republic of Korea ∥ Department of Chemical Engineering, Hanyang University, Seoul 04763, Republic of Korea ‡

S Supporting Information *

ABSTRACT: We developed donor−acceptor (D−A)-type random polymers based on 3,3′-difluoro-2,2′-bithiophene with various relative amounts of 5,6-difluoro-4,7-bis(5bromo-(2-decyltetradecyl)thiophen-2-yl)-2,1,3-benzothiadiazole (2FBT) and 5,6-difluoro-4,7-bis(5-bromo-(2octyldodecyl)thiophen-2-yl)-2-(3,4-dichlorobenzyloxybutyl)2H-benzo[d][1,2,3]triazole (DCB-2FBTZ). Introducing small relative amounts of DCB-2FBTZ into the polymer was found to effectively enhance its solar cell performance, resulting in a power conversion efficiency of 9.02%, greater than the 7.29% that resulted from the PFBT-FTh copolymer. Moreover, when the active area of the BHJ film was increased to 1 cm2, the solar cell reproducibly showed a high performance, here with an efficiency of 8.01% even when the thickness of the active layer was 313 nm. Our studies revealed that including the DCB-2FBTZ group in the polymer simultaneously improved the solution processability and crystallinity of the polymer. These improvements resulted in the formation of highly homogeneous BHJ films throughout large areas with only minor amounts of defects resulting from overaggregation and hence with appropriate morphologies for effective charge generation and transport.

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INTRODUCTION Polymer solar cells (PSCs) garnered great attention over the past 10 years due to their advantages in the fabrication of efficient and lightweight devices with large flexible substrates using low-cost roll-to-roll printing techniques.1−5 So far, single bulk-heterojunction (BHJ) PSCs have achieved power conversion efficiencies (PCEs) of over 11% by using new conjugated polymers and fullerene derivatives as well as novel fabrication techniques.6−10 To realize this photovoltaic technology, research has focused mainly on enhancing device performance, extending device lifetime and developing techniques for scaling up small-area laboratory-scale devices to large-area industrial-scale modules, on which devices have serially connected or parallel-connected narrow-strip geometries (with strip width ≤1 cm).11−13 Although the impressive progress made in the past two decades has led to considerable improvements in both the efficiency and operational stability of organic solar cells (OSCs), the fabrication of large-area printed modules still suffers from significantly reduced PCEs, amounting to less than half the efficiency levels of small-sized laboratory cells.14−16 Recent research efforts have aimed to enhance the efficiency of organic photovoltaic modules. New module architectures have been developed for realizing high-efficiency printed © XXXX American Chemical Society

modules with minimal area loss (or aperture loss). There have been reports about the replacement of indium tin oxide (ITO) or SnO:F with higher conductivity materials to form new transparent conductive oxides (TCOs).17,18 Moreover, efforts are underway to develop a processing technology aimed at reproducibly preparing homogeneous large-area BHJ films with reduced variation in roughness and with thick active layers.14,19,20 The typical thickness values of the active layers of OSCs have been limited to ∼100 nm due to the relatively low charge mobilities of organic photoactive materials. The fill factors (FFs) of high efficiency BHJ-based solar cells generally decrease markedly with use for thicknesses >100 nm. This restriction on active layer thickness raises particular issues for the scaling up of solution-processed OSCs; for instance, thin active films on rough substrates are inclined to forming point defects, which scale with area. Relatively small variations in thin active layers result in large differences in device properties and thus in low performance reproducibility.21−24 Moreover, thin active films of ∼100 nm give limitations of utilizing high throughput deposition techniques for OSC manufacturing.21 Received: July 27, 2017 Revised: September 15, 2017

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DOI: 10.1021/acs.macromol.7b01613 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Some photoactive polymers with high crystallinity show high solar cell performance in an active layer with a thickness greater than 200 nm. For example, Yan and co-workers have developed highly crystalline polymers that exhibit high current density and fill factors even for an active layer film with a thickness of 300 nm.7,10,25,26 Also, a push−pull-type copolymer of benzodithiophene (BDT) and fluorinated benzotriazole (BTA) has exhibited high photovoltaic properties using a thick active layer.27−30 However, most highly crystalline polymers are poorly soluble even in chlorinated solvent, which makes room temperature film processing of active layers difficult and thus results in poor solar cell performance for large-area devices. Therefore, it is important to develop high performance polymers whose optimal BHJ morphologies can be reproducibly transferred to scaled-up areas with high thickness. One approach for improving the processability of polymeric active materials is to endow the polymer backbone with structural irregularity which interferes with intermolecular interactions and prevents molecular aggregation in solution.31−33 However, such decreased molecular interactions are at times not limited to the solution phase but also lead to poor solid-state order and inferior electrical properties.34,35 Another way is attaching aliphatic side chains. Although this approach has been successful, some highly aggregating polymers need long side chains to make them soluble, and the introduction of many long electrically insulating aliphatic side chains can reduce interchain charge transport and hinder charge transfer from the donor to the acceptor in organic photovoltaic devices.36−40 Therefore, it is important to find new solubilizing groups, i.e., other than long aliphatic side chains, that can improve the solubility of highly crystalline polymers. In the current work, we developed an o-dichlorobenzyloxybutyl group as a solubilizing group of conjugated polymers, which resulting in enhanced solubility of the polymer even without increasing the long aliphatic side chain on the polymer backbone. We synthesized D−A (D: electron donor; A: electron acceptor)-type polymers by copolymerizing a 3,3′difluoro-2,2′-bithiophene (2F2T) ditin compound with a 5,6difluoro-4,7-bis(5-bromo-(2-decyltetradecyl)thiophen-2-yl)2,1,3-benzothiadiazole (2FBT) compound and replacing a certain proportion of the 2FBT compound with 5,6-difluoro4,7-bis(5-bromo-(2-octyldodecyl)thiophen-2-yl)-2-(3,4dichlorobenzyloxybutyl)-2H-benzo[d][1,2,3]triazole (DCB2FBTZ). Random polymerization with a small relative amount of DCB-2FBTZ effectively prevented overaggregation of the highly crystalline D−A copolymer PFBT-FTh and increased its solution processability. Moreover, the blend film of the random polymers with PC71BM showed optimal morphology for charge generation and transport with high polymer crystallinity. As a result, the polymer PTAZDCB20, i.e., that containing 20 mol % DCB-2FBTZ units, was found to display a significantly better photovoltaic performance, with a value of 9.03%, than did the copolymer PFBT-FTh and achieved a high efficiency of 8.01% even for a BHJ film with an area of 1 cm2 and thickness of 313 nm.

Scheme 1. Synthetic Scheme for the Polymers

tolyl)3 as the catalyst. After carrying out the reaction, the polymers were collected by precipitation in methanol and purified by performing successive Soxhlet extractions with methanol, ethyl acetate, hexane, and dichloromethane to remove the byproducts and oligomers. Detailed descriptions of the syntheses and characterizations of the monomers and polymers are provided in the Supporting Information. The molecular weight (Mn) values of the polymers were estimated using gel permeation chromatography (GPC) at 80 °C to be 138, 149, and 115 kDa with polydispersity indices (PDIs) of 2.58, 2.51, and 2.96 for PFBT-FTh, PTAZDCB20, and PTAZDCB30, respectively. The thermal properties of the polymers were investigated by performing thermogravimetric analyses (TGA) of the polymers under a nitrogen flow (Figure S1). The 5% weight loss temperatures (Td) values of the PFBTFTh, PTAZDCB20, and PTAZDCB30 polymers were found to be 438, 430, and 417 °C, respectively. Differential scanning calorimetry (DSC) was carried out at a heating rate of 10 °C/ min under a nitrogen atmosphere. According to the DSC results (Figure S2), PFBT-FTh, PTAZDCB20, and PTAZDCB30 showed melting temperatures of 315.7, 316.1, and 292.1 °C and crystallization temperatures of 286.8, 286.8, and 258.6 °C, respectively. Notably, PTAZDCB20 showed sharper and more intense melting and crystallization peaks than did PFBT-FTh, indicating stronger interchain interactions of PTAZDCB20. The crystallization enthalpy ΔH c of PTAZDCB20 was determined to be 22.1 J/g, higher than 14.3 J/g determined for PFBT-FTh and 18.2 J/g for PTAZDCB30. Therefore, incorporation of the DCBZ group appeared to have effectively enhanced the crystallinity of the polymer. Optical and Electrochemical Properties. UV−vis absorption spectroscopy studies of the polymers were also carried out. Absorption spectra of the polymers in chlorobenzene and as thin films are shown in Figure 1a,b, and the characteristic parameters are summarized in Table 1. The polymers yielded two characteristic absorption peaks: one at ∼450 nm, which arose from the π−π* transition of the conjugated backbone, and the other spanning 550−750 nm,

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RESULTS AND DISCUSSION Synthesis and Characterization. Scheme 1 shows the synthesis of PFBT-FTh and random polymers with different feed molar ratios of 2FBT to DCB-2FBTZ (2FBT:DCB2FBTZ = 0.8:0.2 for PTAZDCB20; 0.7:0.3 for PTAZDCB30). The polymers were synthesized by carrying out Stille coupling reactions in a chlorobenzene solvent with Pd2(dba)3/P(oB

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Figure 1. UV−vis absorption spectra of the polymers (a) in chlorobenzene solution and (b) in the solid state. (c) Δλmax (Δλ = λmax/final − λmax/initial) plot with respect to the volume proportion of acetonitrile (VolAcCN); λmax/initial is the maximum absorption wavelength of the polymer solution without the addition of acetonitrile, and λmax/final is the maximum absorption wavelength of the polymer solution after each addition of acetonitrile. (d) Cyclic voltammogram of the polymers.

cosolvents of chlorobenzene and polar acetonitrile with various volume proportions of the acetonitrile (VolAcCN) in the range of 0−0.35 (Figure S4). Figure 1c shows a plot of degrees of wavelength shifts in λmax as a function of VolAcCN. Compared to the λmax of PFBT-FTh dissolved in pure chlorobenzene, the λmax value of PFBT-FTh dissolved in chlorobenzene with 0.05 proportion of acetonitrile was red-shifted, i.e., to the longer wavelengths, and this shift increased rapidly as more acetonitrile was included in the solvent; this shift was due to increased polymer aggregation and resulting decreased solubility. In contrast, PTAZDCB30 did not show any significant change in λmax until VolAcCN was 0.15, and PTAZDCB20 showed the red-shift when VolAcCN was greater than 0.1. Notably, when the same amount of acetonitrile was added into each polymer solution, the degree of red-shifting decreased as the relative amount of the DCB-2FBTZ unit in the polymer was increased; i.e., the red-shift for PFBT-FTh was greater than that for PTAZDCB20, which was greater than that for PTAZDCB30. This result indicated that the solubility of the DCB-2FBTZ-based polymers were greater than that of the PFBT-FTh copolymer despite the DCB-2FBTZ unit having the shorter 2-octyldodecyl group compared to the 2-decyltetradecyl group of PFBT-FTh. The electrochemical properties of the polymers were studied with cyclic voltammetry (Figure 1d). The highest occupied molecular orbital (HOMO) energy levels of the PFBT-FTh, PTAZDCB20, and PTAZDCB30 polymers were estimated from the onset oxidation potentials to be −5.49, −5.48, and −5.45 eV, respectively. That is the HOMO energy levels of the polymers increased slightly as the proportion of DCB-2FBTZ in the polymer was increased. The lowest unoccupied molecular orbital (LUMO) energy levels were estimated from

Table 1. Summary of Optical Properties of Polymers polymer

Mn (kDa)

PDI

λmax,film (nm)

PFBT-FTh PTAZDCB20 PTAZDCB30

138 149 115

2.58 2.51 2.96

684 618 617

a

Egopt (eV)

HOMO (eV)

LUMO (eV)

1.66 1.68 1.68

−5.49 −5.48 −5.45

−3.83 −3.80 −3.77

b

Maximum absorption wavelength in the polymer film. bEstimated from Egopt + HOMO.

a

which was composed of two vibronic peaks and which we attributed to intramolecular charge transfer between the electron-donating and electron-accepting units. In solution, introduction of DCB-2FBTZ made the vibronic peak at ∼600 nm more intense, resulting in broader absorption between 500 and 600 nm. Random copolymers including both DCB-2FBTZ and 2FBT as an accepting unit are expected to show the absorption properties coming from DCB-2FBTZ-co-terthiophene blocks as well as 2FBT-co-terthiophene blocks. It is reported that the copolymer based on benzotriazole and terthiophene showed the intramolecular charge transfer (ICT) absorption in the relatively short wavelength range 500−650 nm.26 Therefore, the random copolymers may show a broader absorption spectrum with enhanced absorption intensity in the lower wavelength range 500−650 nm as increasing the proportion of DCB-2FBTZ on the polymer backbone. The solid-state form yielded absorption spectra that showed features similar to those obtained from solution, but with red-shifted peaks. The optical band gap (Egopt) values estimated from the absorption onsets were similar for the polymers, with values of 1.66, 1.68, and 1.68 eV for PFBT-FTh, PTAZDCB20, and PTAZDCB30, respectively. We also recorded the UV−vis absorption spectra of the polymers dissolved at a concentration of 1 × 10−5 M in C

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Figure 2. (a) J−V curves and (b) EQE spectra of the solar cells (active area = 0.12 cm2). (c) J−V curves of the solar cells (active area = 1 cm2). (d) Histogram of the PCE deviations obtained from 20 solar cells (active area = 1 cm2).

Table 2. Photovoltaic Performances of the BHJ Solar Cells

a

polymer

VOC [V]

JSC (JSC,EQE) [mA cm−2]

FF [%]

PCEmax (PCEavg)a [%]

active area [cm2]

PFBT-FTh PTAZDCB20 PTAZDCB30 PFBT-FTh PTAZDCB20 PTAZDCB30

0.850 0.849 0.853 0.873 0.856 0.852

14.75 (14.12) 15.53 (15.31) 13.22 (13.24) 12.51 15.35 13.12

58.10 68.36 61.57 51.11 60.98 54.40

7.29 (7.20 ± 0.06) 9.02 (8.91 ± 0.09) 6.94 (6.86 ± 0.06) 5.58 8.01 6.09

0.12 0.12 0.12 1.00 1.00 1.00

Averaged PCEs from 10 solar cell devices.

the HOMO energy levels and Eg to be −3.83, −3.80, and −3.77 eV, respectively. Photovoltaic Properties. The photovoltaic properties of the polymers were studied with the device structure ITO/ polyethylenimine ethoxylated (PEIE)/polymer:[6,6]-phenyl C71 butyric acid methyl ester (PC71BM)/MoO3/Ag. The BHJ films were prepared using a solution of each polymer:PC71BM mixture in a o-dichlorobenzene (DCB) and a chlorobenzene (CB)/1-chloronaphthalene (CN) (97.5:2.5 vol %) cosolvent for PFBT-FTh and CB/CN (97.5:2.5 vol %) for PTAZDCB20 and PTAZDCB30. The optimal polymer:PC71BM ratios that provided the best solar cell performances were 1:2 for all polymers. The solar cell efficiency was measured under simulated AM 1.5G 1 sun illumination (100 mW/cm2). The current density−voltage (J−V) curves for the polymer solar cells are shown in Figure 2a, and the characteristic solar cell data are summarized in Table 2. Of the three polymers investigated, PTAZDCB20 achieved the highest performance with a PCE of 9.02%, short-circuit current (Jsc) of 15.53 mA/ cm2, fill factor (FF) of 68.36%, and open-circuit voltage (Voc) of 0.849 V, which is followed by PFBT-FTh and PTAZDCB30; PFBT-FTh showed a 7.29% PCE with a Jsc of 14.75 mA/cm2, FF of 58.10%, and Voc of 0.850 V, and PTAZDCB30 exhibited a 6.94% efficiency with a Jsc of 13.22 mA/cm2, FF of 61.57%, and Voc of 0.853 V. The 7.29% PCE of PFBT-FTh was comparable to that of previously reported.41 Notably, the

polymers incorporating DCB-2FBTZ achieved better FF values than did PFBT-FTh. Figure 2b shows the external quantum efficiency (EQE) spectra of the solar cell devices. PTAZDCB20 exhibited EQEs as high as 60−70% over the wavelength range 350−750 nm, with these EQEs ∼ 10% better than those of PFBT-FTh in the 350−650 nm range. These results are in good agreement with the observation of a higher Jsc value for PTAZDCB20 than for PFBT-FTh. The Jsc,EQE values calculated by integrating the EQE spectra were similar to the Jsc values from the J−V curves (Table 2), which supported the reliability of the J−V measurements. Transmission electron microscopy (TEM) images were acquired to study the morphologies of the polymer BHJ films and to explain the above photovoltaic results. Bicontinuous interpenetrating networks and well-developed polymer fibrils were observed throughout the images of the films (Figure 3a− c). The PTAZDCB30 fibrils were observed to be wider than were the fibrils of the other polymers. This increased width may have caused a decrease of the polymer/PC71BM interface area, resulting in less efficient exciton dissociation at the interface and, thereby, the reduced Jsc of PTAZDCB30 compared with that of PFBT-FTh. We also prepared polymer solar cells with increased photoactive areas of 1 cm2 but otherwise with the same device structures and fabrication conditions as used for the 0.12 cm2 area device. The thickness values of the large-area BHJ films D

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Figure 3. TEM and AFM images and the corresponding profiles of (a, d, g) PFBT-FTh, (b, e, h) PTAZDCB20, and (c, f, i) PTAZDCB30.

area. These performance drops were attributed primarily to the sheet resistance of the ITO film,42 and however, such significant Jsc reduction for PFBT-FTh was probably due to the suboptimal morphology of its large-area BHJ film, which displayed defects and heterogeneities.42 Figure 2d displays distributions and average PCEs of 20 solar cell devices each with an active area of 1 cm2. Here, of the three polymers tested, PTAZDCB20 yielded the narrowest distribution of PCE values, indicating its highest reproducibility in solar cell performance; PFBT-FTh showed the lowest such reproducibility. We would have expected the film inhomogeneity and extensive aggregation observed for PFBT-FTh to form charge recombination and trapping sites in its solar cell device. Such an effect may have caused the particularly large decrease in its efficiency, and the low reproducibility of its performance, upon increasing its active area to 1 cm2. Grazing-Incidence Wide- and Small-Angle X-ray Scattering Studies. We performed grazing incidence wideangle X-ray diffraction (GI-WAXD) measurements to study the molecular orientation and packing properties of the polymers in the polymer:PC71BM BHJ films. Figure 4a−f shows 2D GIWAXD patterns for the pristine polymer and blend films. Here, qxy and qz in the patterns are the in-plane and out-of-plane components of the scattering vector q, respectively; these components are normal to the plane of incidence and the sample surface plane, respectively. The in-plane and out-ofplane linecut profiles of the pristine polymers are shown in Figure 4g,h. All polymers yielded multiple (n00) and (010) diffraction peaks. In the out-of-plane intensity profile for PFBTFTh, the (100), (200), (300), and (400) peaks at qz = 0.247, 0.502, 0.757, and 1.00 Å−1 were due to a lamellar packing structure of the polymer, and the corresponding lamellar distance (d(100)= 2π/q(100)*) was estimated to be 25.5 Å. The

made from PFBT-FTh, PTAZDCB20, and PTAZDCB30 were 259, 313, and 281 nm for PFBT-FTh, respectively. Atomic force microscopy (AFM) measurements, as shown in Figure 3d−i, indicated the PTAZDCB20:PC71BM blend film to have a smooth surface (a root-mean-square (RMS) roughness of 2.83 nm) and continuous phase separation domains despite its considerable thickness of over 300 nm, whereas the PFBT-FTh BHJ film displayed a greater surface roughness (RMS = 21.5 nm) and lower film homogeneity with localized large spots having dimensions of several hundreds of nanometers. These spots may have been aggregates of active materials (polymer or PC71BM). Such inhomogeneous PFBT-FTh BHJ film is expected to show suboptimal morphology in a large area such as 1 cm2. The solar cell J−V curves and characteristic properties of the devices with large-area films are presented in Figure 2c and Table 2, respectively. For each of the polymers, the efficiencies derived from the films with the large active areas were less than those from the smaller films. Yet, of the three films with the large active areas, that of PTAZDCB20 still showed the highest performance, with a PCE of 8.01%, Jsc of 15.35 mA/cm 2 , FF of 60.98%, and V oc of 0.856 V. PTAZDCB30 achieved the next highest PCE, 6.09%, and an FF of 54.40% and Voc of 0.852 V, and PFBT-FTh yielded the poorest performance with a PCE of 5.58%. The lower performances resulting from the large active areas appeared to be mostly due to their lower FF values. Moreover, the extent to which the value of FF was reduced upon increasing the active area depended on the polymer donor: this decrease in FF was 10% for PTAZDCB20 but 12% for PFBT-FTh and PTAZDCB30. Notably, upon increasing the active area from 0.12 to 1 cm2, PFBT-FTh also showed a significant reduction in Jsc from 14.75 to 12.51 mA/cm2, whereas the Jsc values for PTAZDCB20 and PTAZDCB30 did not depend much on the E

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Figure 4. GI-WAXD images of (a, c, e) the pristine polymer films and (b, d, f) BHJ films. Linecut profiles in-plane and out-of-plane of (g, h) the pristine polymer films and (i, j) the BHJ films.

(010) peak at 1.76 Å−1 in the in-plane direction for PFBT-FTh was associated with π−π stacking, and the stacking distance was

calculated to be 3.56 Å. The other polymers, i.e., PTAZDCB20 and PTAZDCB30, produced (010) diffraction peaks at F

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the simple pole figures, the relative amounts of the face-on orientation in the PFBT-FTh, PTAZDCB20, and PTAZDCB30 blend films were estimated to be 58.5, 49.6, and 46.5%, respectively. As the proportion of DCB-2FBTZ in the polymer was increased, the polymers showed slightly less of the face-on orientation, but this orientation was still dominant for PTAZDCB20 and PTAZDCB30. Studies of Charge Generation, Charge Transport, and Charge Recombination Properties of the Solar Cells. The charge carrier mobilities of the pristine polymer and BHJ films were studied using a space charge limited current (SCLC) model. In the pristine film, PTAZDCB20 showed the highest hole mobility, 2.77 × 10−3 cm2 V−1 s−1, followed by PFBT-FTh (1.54 × 10−3 cm2 V−1 s−1) and PTAZDCB30 (3.04 × 10−4 cm2 V−1 s−1). The high hole mobility of PTAZDCB20 was due its high crystallinity, as shown by the GI-WAXD results, and due to its melting temperature and crystallization enthalpy being higher than those of the other polymers. The hole mobility of the PFBT-FTh BHJ blend film was calculated to be 2.56 × 10−4 cm2 V−1 s−1, considerably less than that of its pristine film, whereas PTAZDCB20 and PTAZDCB30 blend films showed hole mobilities of 1.03 × 10−3 and 2.29 × 10−4 cm2 V−1 s−1, respectively, each only a modest amount less than that of the corresponding pristine film. As a result, PTAZDCB20 achieved an almost 1 order of magnitude higher mobility than did the other polymers, which was related to its high crystallinity and preferential face-on orientation in the BHJ film. The electron mobility was determined to be highest for the PTAZDCB20 blend film, with a value of 6.34 × 10−4 cm2 V−1 s−1, and the PFBT-FTh blend film showed a slightly higher electron mobility (5.08 × 10−4 cm2 V−1 s−1) than did the PTAZDCB30 blend film (4.20 × 10−4 cm2 V−1 s−1). The high FF of the PTAZDCB20-based solar cell may have been due to high hole and electron mobility values of the BHJ film. To investigate carrier recombination in the active layers of the polymer solar cells, Jsc was measured for each device as a function of incident light intensity (Plight), and the data were fitted to the power law Jsc ∝ Plightα, where α is an exponential factor that approaches unity when the bimolecular recombination is weak, as shown in Figure 5a.47−49 PTAZDCB20 showed the highest value of the exponent α, 0.97, among the polymers, and the α values of PFBT-FTh and PTAZDCB30 were both 0.92. These results indicated the occurrence of less bimolecular recombination for the PTAZDCB20-based solar cell than for the cells of the other polymers under a short-circuit condition, which would be due to high charge carrier mobilities of the PTAZDCB20-based BHJ film. This result coincided with the high FF and Jsc values of the PTAZDCB20-based devices. In order to study the charge generation, dissociation, and extraction properties of the solar cells, we measured the photocurrent density (Jph) as a function of the effective voltage (Veff) for each device. Note that Jph = JL − JD, where JL and JD are the current densities under illumination of 100 mW cm−2 and in dark conditions, respectively, and that Veff = V0 − Va, where V0 is the voltage when Jph = 0 and Va is the applied bias.50−52 The extracted exciton dissociation probability (P(E,T)) under short-circuit conditions was calculated for each device by using the equation Jph = JsatP(E,T), where Jsat is the saturation photocurrent density. Under the short circuit conditions, the PTAZDCB20-based device achieved the highest P(E,T) value of 95%, which indicated that its charge carrier generation capability was higher than those of PFBT-FTh (P(E,T) = 90%) and PTAZDCB30 (P(E,T) = 94%).

positions similar to those of PFBT-FTh, with these results indicating π−π stacking distances of ∼3.6 Å. However, in the out-of-plane linecut profile, the lamellar distances for PFBTFTh, PTAZDCB20, and PTAZDCB30 were estimated from the (n00) peaks to be 25.5, 25.27, and 25.0 Å, respectively, i.e., to have gradually decreased as the relative amount of DCB2FBTZ was increased. This result was attributed to the shorter 2-octyldodecyl side chain in the DCB-2FBTZ polymer, which may have induced a closer lamellar structure, i.e., shorter such distance between the polymer chains, and thus denser polymer packing in the film. In the pristine PTAZDCB20 film, the polymer showed more intensive lamellar scattering peaks in the out-of-plane direction, and (010) peak is more intensive in the in-plane direction rather than in the out-of-plane direction. Compared with PTAZDCB20, PTAZDCB30 exhibited less pronounced lamellar peaks in the out-of-plane direction and but more intensive lamellar peaks in the in-plane direction. The (010) peak in the in-plane direction of PTAZDCB30 is decreased from the corresponding peak of PTAZDCB20, and however, the (010) peak in the out-of-plane direction of PTAZDCB30 is increased compared with PTAZDCB20. This result suggest that if we incorporate more DCB-2FBTZ units into the polymer from PTAZDCB20 to PTAZDCB30, the polymer crystallites in the film exhibit increased face-on character. Figure 4 displays 2D GI-WAXD patterns and the corresponding linecut profiles for the polymer:PC71BM blend film. All of these blend films yielded the same (n00) and (010) peaks in their diffraction patterns as did the pristine polymer film but also produced diffuse diffraction peaks at qxy or qz = 0.7 and 1.32 Å−1, which were due to the aggregation of PC71BM.43 The blending with PC71BM did not significantly affect the π−π stacking distances of the polymers, and all polymers showed similar stacking distance values of ∼3.6 Å. However, the lamellar spacing of the polymers in the blend films was observed to be 0.3−0.9 Å less than the corresponding values in the pristine films, with distances of 24.6, 24.98, and 24.4 Å for the blend films of PFBT-FTh, PTAZDCB20, and PTAZDCB30, respectively. The blending with PC71BM may have resulted in more compact packing of the polymer chain. From the full width at half-maximum (fwhm) values of the scattering peaks, coherence lengths of face-on crystallites (Lc(100) at qxy and Lc(010) at qz) were calculated using the simple Scherrer equation as shown in Table S1. The Lc(100) values for PTAZDCB20 and PTAZDCB30 along the in-plane direction (i.e., at qxy) were estimated to be 213.8 and 214.1 Å, respectively, higher than that of PFBT-FTh (Lc(100) = 162.4 Å). Such improved face-on crystalline properties of PTAZDCB20 and PTAZDCB30 were probably due to the compact molecular packing of the polymer. The Lc(010) values along the out-ofplane direction for PTAZDCB20, PTAZDCB30, and PFBTFTh were 45.7, 44.3, and 43.0 Å, respectively. These results showing PTAZDCB20 with the highest Lc(010), and PFBT-FTh with the lowest, suggested that the PTAZDCB20 blend film had the largest crystallites and the PFBT-FTh blend film the smallest. Larger crystallites are favorable for enhancing charge transport in polymer solar cells. Simple pole figure analyses were performed as functions of the azimuthal angle (or detector angle), as shown in Figure S3. Pole figure analyses provide information about the orientational distributions of the diffraction peaks as functions of all possible crystalline orientations in the blend film.44−46 In these pole figure analyses, we focused on the distributions of the (100) orientations as functions of the azimuthal angle (χ). From G

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without severe material aggregation, and therefore the solar cell device based on a 313 nm thick BHJ film with an active area of 1 cm2 achieved a high performance with a PCE of 8.01%. We expect our results to provide guidelines for developing donor polymers that can be reliably used for the preparation of large BHJ junctions with good photovoltaic properties.

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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/acs.macromol.7b01613. Experimental details for polymer synthesis and characterization, solar cell fabrication and measurements, and SCLC measurements (PDF)

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AUTHOR INFORMATION

Corresponding Author

*(H.J.S.) Tel +82-2-958-5320, e-mail [email protected]. ORCID

Hae Jung Son: 0000-0002-0912-3483 Notes

Figure 5. (a) Jsc−Plight and (b) P(E,T)−Veff curves of solar cells.

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the KIST institutional program, KIST Yong Fellow program, and New and Renewable Energy Program of the Korea Institute of Energy Technology Evaluation and Planning (KETEP) grant funded by the Ministry of Knowledge Economy (MKE) (20163030013620); this work was also supported from the Global Frontier R&D Program on Center for Multiscale Energy System (NRF2016M3A6A7945505, project title: Development of MEA with Durable High-Power Performance for PEMFC via Multiscale Technique).

As a result, we conclude that incorporating a small relative amount of DCB-2FBTZ into the polymer improved charge generation and transport in the polymer:PC71BM BHJ film by enhancing the polymer packing properties in the film. Thus, PTAZDCB20 yielded a PCE (9.02%) greater than that achieved by the corresponding PFBT-FTh homopolymer (7.29%). In particular, the simultaneously improved solution processability and crystallinity of PTAZDCB20 enabled it to form a highly reproducible and thick (>300 nm) BHJ film with an appropriate morphology over a large area. In many cases, highly crystalline donor polymers are often poorly soluble and thus cause the formation of polymer (or PCBM) aggregates, resulting in film inhomogeneities during BHJ film preparation over a large area. Consequently, when scaling up the active area from 0.12 to 1 cm2, PTAZDCB20 did exhibit a small decrease in PCE to a value of 8.01% when using a 313 nm thick BHJ film; but the Jsc and Voc values did not change much, and FF decreased a little, due to the sheet resistance of the ITO film. In contrast, upon increasing the active area, the PFBT-FTh solar cell device showed a large drop of Jsc from 14.75 to 12.51 mA/ cm2 and showed a relatively low PCE of 5.58%.

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REFERENCES

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CONCLUSION In the current research, we developed a series of random polymers by copolymerizing 2F2T with various relative amounts of 2FBT and DCB-2FBTZ. Introduction of a small relative amount of DCB-2FBTZ into the donor polymer was found to effectively enhance the packing of the polymer in the BHJ film and to simultaneously increase the solution processability of the polymer. Consequently, the PTAZDCB20-based solar cell exhibited a much greater efficiency (with a PCE value of 9.02%) than did the PFBTFTh device (PCE of 7.29%) due to improved charge generation and transport properties in the solar cell device. Moreover, the excellent solution processability of the highly crystalline PTAZDCB20 allowed the BHJ morphology optimized for a small active area to be well reproduced for a large active area H

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