Recent progress in the design of narrow bandgap conjugated polymers for high-efficiency organic solar cells

Progress in Polymer Science 37 (2012) 1292–1331

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Recent progress in the design of narrow bandgap conjugated polymers for high-efficiency organic solar cells Linyi Bian a , Enwei Zhu a , Jian Tang a , Weihua Tang a,∗ , Fujun Zhang b,∗ a Key Laboratory of Soft Chemistry and Functional Materials (Ministry of Education of China), Nanjing University of Science and Technology, Nanjing 210094, People’s Republic of China b Key Laboratory of Luminescence and Optical Information (Ministry of Education of China), Beijing Jiaotong University, Beijing 100044, People’s Republic of China

a r t i c l e

i n f o

Article history: Received 30 August 2011 Received in revised form 20 March 2012 Accepted 28 March 2012 Available online 3 April 2012 Keywords: Narrow bandgap polymer Polymer solar cells Bulk-heterojunction

a b s t r a c t The critical review on the recent development of novel narrow bandgap polymers for highefficiency polymer solar cells concentrates on (i) the structural design of narrow bandgap polymers, which occupy a central place in recent advances in high-efficiency polymer solar cells, (ii) the intrinsic physics and chemistry of special properties, such as absorption, bandgap and energy levels, and (iii) the correlation of polymer structure and device fabrication with their photovoltaic performances. The statistical summaries of their device parameters are also discussed. The description of these structure–property correlations may guide the rational design of polymer structures and the reasonable evaluation of their photovoltaic performance. © 2012 Elsevier Ltd. All rights reserved.

Contents 1. 2. 3. 4.

5.

6.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1293 Operation mechanism . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1294 Strategy in narrow bandgap polymer design . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1294 Fluorene-based narrow bandgap polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1294 4.1. Copolymers based on fluorene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1294 4.2. Copolymers based on silafluorene and germafluorene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1298 4.3. Copolymers based on ladder-type oligo-p-phenylenes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1298 Carbazole-based narrow bandgap polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1298 5.1. Copolymers based on carbazole . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1298 5.2. Copolymers based on Indolo[3,2-b]carbazole . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1302 Thiophene-based narrow bandgap polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1304 6.1. Copolymers based on thiophene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1304 6.2. Copolymers based on cyclopentadithiophene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1309 6.3. Copolymers based on dithieno-[3,2-b:2 ,3 -d]silole . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1310 6.4. Copolymers based on dithieno-[3,2-b:2 ,3 -d] thiophene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1311 6.5. Copolymers based on benzo[1,2-b:4,5-b ]dithiophene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1312 6.6. Copolymers based on dithieno[3,2-b:2 ,3 -d]pyrrole . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1317 6.7. Copolymers based on naphtho[2,1-b:3,4-b ]-dithiophene and dithieno[3,2-f:2 ,3 -h]quinoxaline . . . . . . . . . . . . . . . . . . . . . . . . 1320 6.8. Copolymers based on ladder-type coplanar thiophene-phenylene-thiophene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1321

∗ Corresponding author. Tel.: +86 25 84317311; fax: +86 25 84317311. E-mail addresses: [email protected] (W. Tang), [email protected] (F. Zhang). 0079-6700/$ – see front matter © 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.progpolymsci.2012.03.001

L. Bian et al. / Progress in Polymer Science 37 (2012) 1292–1331

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Data analysis of polymer and discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1322 7.1. HOMO–LUMO levels of polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1322 7.2. Correlation of HOMO levels of polymers with device Voc . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1323 7.3. Correlation of bandgap of polymers with device performance . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1324 7.4. Molecular weight issue . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1325 7.5. Device optimization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1326 Conclusions and perspective . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1327 Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1327 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1328

1. Introduction Harvesting energy directly from sunlight by photovoltaic (PV) technology is one of the most important ways to address the ever-growing world-wide energy needs and environmental concerns using renewable resources. Polymer solar cells (PSCs) have emerged as a promising alternative PV technology due to the potential of costeffective production of large area of flexible devices using solution-processing techniques, with low environmental impact and versatility in organic material design [1–4]. Within recent years, PSC technology has undergone dramatic progress with its key parameter, power conversion efficiencies (PCE, defined as the maximum power produced by a photovoltaic cell divided by the power of incident light) reaching new world record of 9.4% by Mitsubishi Chemical Co. in early April 2011 [5]. Meanwhile, three other companies – Konarka Technologies, Solarmer Energy Inc., and Heliatek – have reported cells with efficiencies greater than 8%. Projections indicate that the figure could soon top 10%, and possibly reach 15%. The PCE of solar cells is codetermined by the open circuit voltage (Voc ), the fill factor (FF) and the short circuit density (Jsc ). Great efforts have been made to develop ␲-functional materials with new molecular structures and optimize solar cell devices with innovative strategies, aiming to harvest solar irradiation in the visible light range. The most successful PSCs adopt a bulk-heterojunction (BHJ) architecture [6], in which a photoactive layer is cast from a mixture solution of polymer donor and soluble fullerene-based electron acceptor (e.g., PCBM: [6,6]-phenyl-C61-butyric acid methyl ester, PC71 BM: [6,6]-phenyl-C71-butyric acid methyl ester) and sandwiched between two electrodes with different work functions. A typical BHJ solar cell contains an indium tin oxide (ITO) coated glass substrate, covered by a layer of transparent conductive polymer, most often poly(3,4-ethylenedioxythiophene):polystyrene-sulfonate (PEDOT:PSS) [7]. The PEDOT:PSS [8] layer provides an improved interface between ITO and the active layer and, hence improves the device performance. A mixture of polymer and fullerene is spin-coated on the top of PEDOT:PSS layer, and a thin layer of metal cathode (e.g., aluminum or silver) is subsequently deposited on the active layer. For efficient charge collection, the work functions of anode and cathode should be matched to the highest occupied molecular orbits (HOMO) of donor and the lowest unoccupied molecular orbits (LUMO) of acceptor, respectively. When light shines on the device,

photons absorbed by the active layer lead to the formation of excitons (Coulomb-correlated electron–hole pairs). The excitons subsequently diffuse to the interface of donor–acceptor for which charge separation occurs and the excitons dissociate into electrons and holes. Finally, the free charge carriers move to their corresponding electrodes (electrons to metal cathode and holes to ITO) with the help of a built-in electric field. In this way, the active layer can be considered as a network of donor–acceptor heterojunctions that allows efficient charge separation and balanced bipolar transport throughout its whole volume [9]. Over the last decade, substantial progress has been made in understanding the structure–function relationships governing material performance of BHJ solar cells. On the material side, the most fundamental material properties of polymers, such as conjugation length, light absorption, carrier mobility, exciton dynamics, and processability may be tuned by manipulating the molecular structure of a ␲-conjugated polymer system. On the device side, the optimal intrinsic capabilities of the newly designed materials for PSC can be explored by optimization of the BHJ device architecture, device processing conditions and strategies in morphology control for the active layer [10–17]. Theoretically, the PCE of PSCs can be further improved to 10% and beyond for single-layer device [15] and 15% for tandem device [17] by implementing new materials [4,18–21], exploring new device architecture [17,22,23] and optimizing device processing approaches. Several fundamental challenging issues need to be addressed, including the bandgap (Eg ) of the polymer donor [18–21], interfaces [24], charge transfer [24,25], thin-film ordering in the active layer [11], device processing and stability [25] before translating the materials and device performance parameters obtained in laboratoryscale cells to large area solution-processed devices for practical applications. In this review, we try to comprehensively take into account all recent pursuits of high performance PSCs (PCE > 2%) with newly developed narrow bandgap polymers with Eg ranging from 1.1 to 2.1 eV from a viewpoint of material chemists. This review is organized as follows: following a brief introduction on the mechanism of PSCs in the section “Operation mechanism” an overview of the popular strategy for narrow bandgap polymers is presented in the section “Strategy in narrow bandgap polymer design”. The subsequent sections are devoted to update the state-ofthe-art novel polymers with PCE over 2%, which are mainly classified with the representative electron-donating units

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L. Bian et al. / Progress in Polymer Science 37 (2012) 1292–1331

Besides experimentally characterizing the performance of PSC with Eqs. (1) and (2), Voc is most often simply estimated as the difference between the donor HOMO level and the acceptor LUMO level as in Eq. (3), in the case of polymer:fullenrene based PSCs [10]: Voc = (1/e)(|E Donor HOMO| − |E PCBM LUMO|) − 0.3V

(3)

where e is the elementary charge, and a minimum energy difference of 0.3 eV between LUMO level of donor polymer and PCBM-like acceptor is estimated to facilitate exciton splitting and charge dissociation. Fig. 1. J–V curve for a typical solar cell. Geometrically, the fill factor (FF) can be visualized as the area ratio of the gray rectangle and the dashed rectangle (Pmax /Voc Jsc ).

along the backbone into “Fluorene-based narrow bandgap polymers”, “Carbazole-based narrow bandgap polymers” and, “Thiophene-based narrow bandgap polymers”. A further statistic analysis of characteristic data for all polymers and discussion is provided in Section 7. Through this analysis of structure–property relationship, we present the current activities in rational design of polymer structures and a reasonable evaluation of their photovoltaic performance. Finally, we conclude and offer our perspective on the challenges to achieving higher efficiencies in Section 8. 2. Operation mechanism The operation of PSCs mainly involves the five following steps: (i) light absorption by the active layer; (ii) the formation of exciton and subsequent diffusion to the interface of donor and acceptor; (iii) exciton dissociation into the electrons and holes; (iv) free charge carriers transport in their individual pathway or layer; (v) charge extraction by their corresponding electrodes [14]. The integrated target evaluating the performance of PSCs is the power conversion efficiency (PCE), which is strongly determined by the following key parameters: open circuit voltage (Voc ), short circuit current (Jsc ), and fill factor (FF). Here, Voc of a solar cell is defined as the voltage that compensates the current flow through the external circuit; Jsc is defined as the current during the external circuit without applied voltage; FF is a more sensitive parameter compared to Voc and Jsc , and depends on the mobility ()–lifetime () product of the bulk materials, thickness of the active layer and the morphology of the cathode-active layer interfaces [26]. The PCE () is calculated using the equation =

Voc Jsc FF Plight

(1)

where Plight is incident light power, and FF is calculated by the following formula FF =

Vmax Jmax Voc Jsc

(2)

The curvature of the J–V characteristics in the fourth quadrant can vary from convex (FF > 50%) to concave (FF < 12.5%) depending upon these parameters. The physical meaning of key parameters for solar cells is shown in Fig. 1.

3. Strategy in narrow bandgap polymer design Adopting the donor–acceptor (DA) approach [27], narrow bandgap polymers containing “push–pull” motifs have been designed by alternating electron-rich (donating) and electron-deficient (accepting) heterocylces along the same ␲-conjugated backbone, affording polymer chromophores with red-shifted absorption spectra towards wavelengths of 500–800 nm, where the solar photon flux is most intensive [11]. A large spectrum of narrow bandgap (1.1 eV < Eg < 2.1 eV) polymers have been developed and afford significantly high efficiencies (over 5% PCE) from their BHJ solar cells with fullerene acceptors [4,18,21]. Due to successful application in organic light-emitting diodes (OLEDs) and thin-film transistors (TFTs), fluorene, carbazole, thiophene and their analogue heterocycles as well as their fused ␲-conjugated moieties have become the first choice of electro-donating heterocycles. Importantly, several hybrid electro-donating heterocycles (Fig. 2) such as cyclopentadithiophene (CPD), naphtho[2,1-b:3,4-b ]dithiophene (NDT), benzo[1,2-b:4,5-b ]dithiophene (BDT) and thiophene-phenylene-thiophene (TPT) have been developed by incorporating different above-mentioned units. Meanwhile a variety of electro-accepting units such as 2,1,3-benzothiadiazole (BT), 4,7-dithien-2-yl2,1,3-benzothiadiazole (DTBT), diketopyrrolopyrrole (DPP), thieno[3,4-c]pyrrole-4,6-dione (TPD), dithieno[3,2f:2 ,3 -h]quinoxaline (DTQ) have been employed as the counterpart in designing DA polymers. Therefore, a good understanding of DA polymer structure and its impact on the active layer morphology, device processing, device performance of polymer:fullerene based BHJ cells have become very important. 4. Fluorene-based narrow bandgap polymers 4.1. Copolymers based on fluorene Fluorene (FL) and its derivatives been extensively investigated for their application in OLEDs due to its rigid planar molecular structure, excellent hole-transporting properties, good solubility, and exceptional chemical stability [28,29]. Along with their low-lying HOMO levels, polyfluorenes (PFs) are expected to achieve higher Voc and Jsc values in their PSC devices, which make FL unit a promising electron-donating moiety in DA polymer design. Meanwhile, the feasible dialkylation at 9-position and selective bromination at the 2,7-positions of FL allows versatile molecular manipulation to achieve good solubility and

L. Bian et al. / Progress in Polymer Science 37 (2012) 1292–1331

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Fig. 2. Fine tuning of absorption and bandgap of alternating DA polymers for high-performance BHJ solar cells, where ITO acts as anode, PEDOT:PSS layer improve the interface between the active layer (blend film of polymer donor and fullerene derivative) and ITO. A photoinduced charge generation occurs in the active layer, and the exciton dissociates into an electron and a hole at the donor–acceptor interface.

extended conjugation via typical Suzuki or Stille crosscoupling reactions. Fluorene-based homopolymers and copolymers containing electron-donating moieties are generally not suitable for solar conversion purpose, due to their weak absorption of visible light. One exception is the alternating copolymer of fluorene and electron-rich bithiophene [29]. Despite the fact that polymer:PCBM blend films only absorb light between 300 and 500 nm, a moderate PCE of 2.4% was achieved due to the high hole mobility of this polymer and well-defined nanoscale segregation morphology in the active layer [30]. This makes it a promising highenergy-absorbing polymer for use in tandem BHJ cells [31]. By incorporating electron-deficient BT unit into polymer backbone, PFs exhibit reduced LUMO levels and thus narrowed bandgaps to fine-tune the emission over the entire visible region [32]. Along this line, Andersson and coworkers [33] demonstrated the first breakthrough in the design of DA PFs, in which electro-accepting 4,7-dithien2-yl-2,1,3-benzothiadiazole (DTBT) was used to reduce the bandgap of P1 (P1–9 structure in Chart 1). Consequently, films of this polymer showed extended absorption, with the longest wavelength absorption maximum at ∼545 nm. By spin-coating from polymer blend with PCBM in CHCl3 solution, P1 showed an ∼50 nm bathochromic shift in absorption compared with poly[2-methoxy-5-(3 ,7 dimethyloctyloxy)-1,4-phenyl vinylene] (MDMO-PPV) [34], which helped in generating larger photocurrents due to a high photon flow in this energy range. A BHJ cell with the configuration ITO/PEDOT:PSS/P1:PCBM/LiF/Al presented a moderate PCE of 2.2% under the illumination of a simulated solar light AM 1.5G (100 mW/cm2 ), for which Voc was 1.04 V, Jsc 4.66 mA/cm2 and FF 0.46 (photovoltaic performance in Table 1). The device presented 0.17 V higher Voc and the same order Jsc as MDMO:PCBM devices [34]. Further optimization of the cell performance attempted via improved processing of the photoactive layer. Nevertheless, the relatively high Voc of P1 has attracted a great

deal of attention since poly(3-hexylthiophene) (P3HT) and poly[2,6-(4,4-bis-(2-ethylhexyl)-4H-cyclopenta[2,1b;3,4-b ]-dithiophene)-alt-4,7-(2,1,3-benzothiadiazole)] (PCPDTBT) devices can achieve high FF and Jsc while their highest PCE (5.5%) cannot be further optimized due to their low Voc (0.6–0.65 eV) [12,35]. By varying the side chains in the 9-position of fluorene for solubility considerations, a series of P1 analogue PFs, P2–P4, have been developed by Cao and coworkers [36] with subsequent comparison of the PV performances of the products. The device performance of PFs was found to be strongly dependent on the side chains, which affect the packing of polymer backbone and thin-film ordering in the active layer and consequently result in different cell performances. The octyl substituted P2 exhibited a slightly higher PCE (2.6%) than P1 (Table 1), while the 2-ethylhexylated P3 (BisEH-PFDTBT) obtained further improved PCE (2.7%) [37]. Hashimoto and coworkers [38] investigated the effect of side chains on device performance using bis-2-ethylhexylated P3 and bis-3,7-dimethyl-octyl substituted P4 (BisDMO-PFDTBT). The PV performance was found to be dramatically improved via increased ␲–␲ stacking between adjacent main chains, achieved by the reduction of “steric effect” of the non-conjugated side chains (bis-3,7-dimethyl-octyl in P4). Hence, the device ITO/PEDOT:PSS/P4:PCBM (1:4, w/w)/LiF/Al harvested a much higher PCE (4.2%). By blending higher molecular weight (Mn = 21 kDa) P3 and new electron acceptor PC71 BM, the PV performance of P3 was improved to 3.5% [39], due to the extended absorption of the active layer in the visible range with PC71 BM [40,41]. Similarly, P4 device can be further improved to 4.5% when blending with PC71 BM [39]. Both EQE (external quantum efficiency) and IQE (internal quantum efficiency) measurement for the bestperforming devices indicated efficient charge collection occurred (Fig. 3). The EQE for a BisDMO-PFDTBT:PC71 BM device reached, a maximum about 67%, while exceeding

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Table 1 The optical, electrochemical, hole mobility, and PSC characteristics of P1–18. Mn (kDa) (Mw /Mn )

abs max (nm)

Eg (eV)

h (cm2 V−1 s−1 )

HOMO/LUMO (eV)

P1 P2 P3

4.8 (2.9) 11 (1.7) 11.3 (1.4) 21 (1.5) 9.7 (3.6) 20 (1.3) 62 (2.3) 16.6 (2.4) 15.3 (2.2) 19.0 (2.2) 40.0 (2.4) 31.7 (1.4) 10.4 (1.7) 10.2 (1.6) 11.5 (2.7) 12.3 (2.8) 15 (1.3) 79 (4.2) 10 (2.4) 26.3 (2.4) 7.8 (2.7) 14.2 (1.8)

545 545 547

1.90 1.91 1.85 1.9 – 1.9 1.77 1.95 1.96 1.94 1.60 2.08 1.87 1.76 1.83 1.74 1.85 1.82 1.79 1.97 2.00 1.96

– – – a 6 × 10−6 – a 3 × 10−5 – –

– −5.70 −5.56/−3.47 −5.5/−3.6

P4 P5 P6 P7 P8 P9 P10 P11 P12 P13 P14 P15 P16 P17 P18

– 589 540 540 542 660 616 545 580 541 579 560 565 580 543 541 531

– – 4.5 × 10−5 a 5.3 × 10−4 a 1.2 × 10−3 a 1.8 × 10−4 a 2.1 × 10−4 – 1.0 × 10−3 1.1 × 10−4 1.8 × 10−4 2.3 × 10−4 6.1 × 10−4

−5.5/−3.6 −5.34/−3.22 −5.37/−3.42 −5.36/3.40 −5.34/−3.40 −5.00/−3.40 −5.71/−3.60 −5.30/−3.43 −5.26/−3.50 −5.32/−3.49 −5.35/−3.61 −5.70/−3.91 −5.39/−3.57 −5.58/−3.91 −5.47/−3.44 −5.45/−3.36 −5.45/−3.46

Polymer:PCBM 1:4 1:2 1:3 1:3b 1:4 1:3b 1:2 1:4b 1:4b 1:3 1:3 1:4 1:4b 1:4b 1:4b 1:4b 1:4 1:2 – 1:3.0 1:3.5b 1:4.0 1:4.0b

Jsc (mA/cm2 )

Voc (V)

FF

PCE (%)c

Ref.

4.66 5.18 5.12 8.4 7.7 9.1 5.86 9.72 4.05 6.0 8.88 8.67 9.62 9.61 6.69 6.22 2.8 9.5 6.9 6.1 7.57 7.8 10.3

1.04 0.95 1.02 0.95 1.0 0.97 0.86 0.99 0.91 1.00 0.59 0.65 0.99 0.99 0.85 0.90 0.97 0.90 0.79 1.00 1.00 1.06 1.04

0.46 0.35 0.51 0.44 0.54 0.51 0.52 0.57 0.47 0.63 0.42 0.36 0.5 0.46 0.37 0.45 0.55 0.51 0.51 0.40 0.40 0.44 0.42

2.2 2.6d 2.7 3.5 4.2 4.5 2.6 5.5 1.7 3.7 2.2 2.0 4.7 4.4 2.5 3.2 1.6e 5.4f 2.8 2.4 3.0 3.7 4.5

[33] [36] [37] [39] [38] [39] [42] [43] [43] [44] [45] [46] [47] [47] [48] [48] [49] [50] [51] [52] [52] [52] [52]

abs max : maximum absorption peak in film; Eg : optical bandgap; Jsc : short-circuit current density; Voc : open-circuit voltage; FF: fill factor; PCE: power conversion efficiency; –: no data available; h : hole mobility measured infield-effect transistor (FET) devices without substrate treatment. a h measured by SCLC. b Polymer:PC71 BM, in weight ratio, w/w. c PCE achieved at illumination of standard simulated solar light (AM 1.5G, 100 mW/cm2 ). d PCE achieved at AM 1.5 (78.2 mW/cm2 ). e PCE measured at AM 1.5 (90 mW/cm2 ). f PCE measured at AM 1.5 (80 mW/cm2 ).

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Polymer

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1297

Chart 1. Fluorene based narrow bandgap polymers.

50% for over half of the visible region; IQE was above 70% throughout the majority of the visible region with a maximum value of 87%, which validated the efficient photon conversion properties of this system. It should be noted that both devices exhibited higher IQE values over a much broader region as compared to the values reported for P3HT:PCBM devices [12]. Accordingly, a high IQE is regarded as one of the crucial factors contributing to the high performance PV cells with such a thin active layer. By incorporating two electro-rich 3-hexylthiophene moieties in the backbone of P3, DA polymer P5 [42] was developed with extended conjugation length. Broader absorption was observed with one maximum peaking at 400 nm from ␲–␲* transition of fluorene and another at

Fig. 3. IQE and EQE are plotted for BisEH-PFDTBT:PC71 BM (circle and diamond) and BisDMO-PFDTBT:PC71 BM (triangle and square) PV cells. The IQE values are calculated from the absorbance based on the reflectivity mode. Reproduced with permission from Ref. [38]. Copyright 2009 Wiley.

560 nm due to ␲–␲* transition of DBT. Nevertheless, P5 device exhibited similar PCE (2.6%) as P3. Using thiophene-quinoxaline-thiophene (TQT) as accepting unit, Kitazawa et al. [43] reported P6 and P7. The polymers showed identical HOMO energy levels and Eg values. P6 device performance was found to be dependent on the composition of co-solvent for film preparation, and a maximal Jsc was achieved with CHCl3 /chlorobenzene (CB) (2:3, v/v) co-solvent system. The optimized device exhibited 5.5% PCE by inserting 0.1 nm LiF layer between P6:PC71 BM film and Al cathode. However, P7 device showed only 1.7% PCE due to the larger phase separation morphology. Nanophase segregation morphology was observed for P6:PC71 BM films while a surface topography composed of domains with lateral dimensions of several hundred nanometers for P7:PC71 BM films. The PV performance of P7 was improved to 3.7% by blending higher molecular weight P7 with PCBM [44]. Inganäs and his coworker [45] further developed thieno[3,4-b]pyrazine containing P8. This polymer exhibited much lower Eg (1.6 eV) in comparison to P7. P8 device achieved an optimized 2.2% PCE with 140 nm-thick active layer. Another fluorene-containing copolymer P9 was reported by Cao and coworkers [46], in which 1,1dimethyl-3,4-diphenyl-2,5-bis(2 -thienyl)silole (TST) was incorporated into the polymer backbone as electron accepting moiety. The optimized device presented a 2.0% PCE when blended with PCBM. Besides linear DA polymer design, a new strategy for developing narrow bandgap of PV polymers was first reported by Jen and coworkers [47] in 2009, in which a fluorene-triarylamine copolymer (PFM) backbone with high-lying HOMO level was grafted with electronaccepting side chains to suppress its LUMO level. This strategy takes advantage of high hole transporting properties of main chains and internal charge transfer (ICT)

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behavior from side chains, which can not only fine-tune the Eg of the resultant polymers but promote an isotropic charge transport in the active layers. With styrylthiophene ␲-bridged malononitrile or diethylthiobarbituric acid as accepting side chains, polymers P10 and P11 were developed. Both polymers exhibited two obvious absorption peaks, with the first (∼385 nm) corresponding to the ␲–␲* transition of main chains and the other (∼600 nm) corresponding to the strong ICT of side chains. P10 and P11 maintained similar HOMO levels as that of the backbone, but much lowered LUMO levels (−3.43 and −3.50 eV, respectively), thus a depressed bandgap (<2 eV). P10 and P11 devices exhibited good PCE values of 4.7% and 4.4%, respectively, when blending with PC71 BM (Table 1). The high PV performance of P10 and P11 were attributed to their high Jsc (∼9.6 mA/cm2 ) for devices, mainly due to the isotropic charge transport ability of these polymers. Using the space-charge-limited-current (SCLC) method, the hole mobility of the active layer was determined as 5.27 × 10−4 cm2 V−1 s−1 for P10/PC71 BM blend and 1.16 × 10−3 cm2 V−1 s−1 for P11/PC71 BM blend, respectively, which are all higher than that of P3HT with similar device configuration [12]. 4.2. Copolymers based on silafluorene and germafluorene The introduction of a silicon or germanium bridge instead of a methylene bridge into the structure of biphenyl results in fluorene analogues: silafluorene or germafluorene. Shortly after Jen and coworkers had reported two-dimensional (2D) polymers P10–11, Cao and coworkers [48] developed silafluorene-based 2D copolymers P12 and P13 (Chart 2). P12 and P13 exhibited similar Eg as P10 and P11, respectively, due to identical polymer structures. However, P12 and P13 devices harvested only moderate PCEs of 2.5% and 3.5%, respectively. The lower PV performance of P12–13 was explained by the lower hole mobility of silofluorene-based P12–13 than fluorene-based P10 and P11. Leclerc and coworkers [49] reported a silafluorenebased linear copolymer P14, in which electron-accepting DTBT was employed. With an optical Eg of 1.85 eV, P14 films presented broader absorption than fluorene counterpart P1, for which two broad absorption peaks at 391 nm to 560 nm were observed. But P14:PCBM cells displayed only a low PCE of 1.6% under AM 1.5 (90 mW/cm2 ) illumination. Almost at the same time, the independent work from Cao and coworkers [50] reported much better PCE (5.4%) using 4-fold higher molecular weight P14 with a similar device structure. This significant improvement of P14 performance was explained by the higher hole mobility achieved with a much extended conjugation length, i.e., this mobility was nearly ten times higher than that of P1. Leclerc and coworkers further reported [51] germafluorene-based DA polymer P15. Spin-coated from o-dichlorobenzene (ODCB), a P15:PC71 BM device displayed a moderate PCE of 2.8%, a performance lower than that of fluorene-based P4 [38] and silafluorene-based P14 [50]. However, better performances might be expected by making devices with high molecular polymers (as for P14) and optimizing the thin-film ordering and ␲–␲

stacking of conjugated backbone by changing the side chains (as for P4). 4.3. Copolymers based on ladder-type oligo-p-phenylenes Ladder-type oligo-p-phenylene (indenofluorene) can be used to extend ␲-conjugation of the polymer backbone to (i) promote better hole transport, and (ii) absorb more solar radiation in the visible region than normal noplanar oligophenylenes. In addition, more solubilizing side chains can be introduced into this unique molecular backbone to improve the solution processability of resulted polymers. Katz and coworkers [52] reported a series of indenofluorene-containing copolymers P16–18, in which DTBT or 5,8-dithien-2-yl-2,3-quinoxaline (DTQ) were used as accepting units. Typical absorption spectra for DA copolymers were observed for P16–18, with two maximum absorptions with one peak (400 nm) corresponding to the ␲–␲* transition of indenofluorene and another (∼540 nm) from the ␲–␲* transition of the accepting unit. With one more phenylene moiety fused in the backbone, P18 presented the highest PCE (3.7%) among the three polymers with the same configuration, due to great improvement of Jsc with the extended conjugation length. The performance of P18 device was further improved to 4.5% when blending with PC71 BM in 1:4 weight ratio as the active layer. The length of the side chains is found to have impact on thin-film roughness as well as the phase separation in the spin-cast active layers. For P16–18, longer alkyl chains led to an increased roughness as well as enlarged domain sizes. The decylated P16 and P17 exhibited a film roughness between 3.5 and 4.0 nm, while hexylated P18 showed a lower roughness of 1.3–1.5 nm for polymer:PCBM films from atomic force microscopy (AFM) images (Fig. 4). Importantly, nanophase (>10 nm) separation morphology was found for both P16 and P17 blend films with PCBM, for which ∼100 nm sized clusters were observed in the films. P16 polymer chains were homogeneously distributed in the matrix with a larger domain size (∼20 nm) than P18 (10 nm). Importantly, interpenetrating networks were observed in P18:PCBM blend films, shown as the bright patterns in Fig. 4c. These interpenetrating networks enabled large interface areas for exciton dissociation as well as continuous percolated pathways for charge carrier transport to the corresponding electrodes [52]. This may be responsible for the improved performance of the devices. 5. Carbazole-based narrow bandgap polymers 5.1. Copolymers based on carbazole Structurally analogous to fluorene, carbazole has also been known as 9-azafluorene. The central fused pyrrole ring makes tricyclic carbazole fully aromatic and electronrich with its donating nitrogen. With solubilizing N-alkyl chain, carbazole derivatives have been widely used as electron-donor materials due to their excellent thermal and photochemical stabilities, as well as relatively high

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Chart 2. Fluorene and its analogue derived narrow bandgap polymers.

charge mobility. There are two approaches to functionalize carbazole, i.e., the first is on 3,6-positions to yield 3,6-carbazoles, and the second is on 2,7-positions towards 2,7-carbazoles. Structurally, 2,7-carbazoles are advantageous than 3,6-carbazoles for photovoltaic applications because of the extended conjugated backbone with carbazole linked at the 2,7-positions [53]. When two or more 2,7-carbazoles fused together to form indolo[3,2b]carbazole, it exhibited even stronger electron-donating

properties, higher hole mobility and better stability than carbazole [54,55]. Similar to fluorene, carbazole hompolymers or its copolymers with electron-donating moieties are not suitable as polymer donors for BHJ cells, due to their absorption mainly in UV region. Leclerc and coworkers [56–58] have made great contribution in the development of high-efficiency poly(2,7-carbazole) based solar cells. Based on donor–acceptor design strategy, a large spectrum of

Fig. 4. AFM topography images (top panel) and surface phase images (bottom panel) of (a) P16:PCBM, (b) P17:PCBM, and (c) P18:PCBM blend films spin-cast from a mixture of CB and ODCB. Adapted with permission from Ref. [52]. Copyright 2010 American Chemical Society.

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poly(2,7-carbazole)s have been developed by incorporating different strong electron-accepting units in polymer backbones. The polymers exhibited enhanced absorption and hence increased PV performance. The most successful poly(2,7-carbazole) is P19 (poly[N-9 -heptadecanyl-2,7-carbazole-alt-5,5-(4 ,7 di-2-thienyl-2 ,1 ,3 -benzothiadiazole)], PCDTBT) (Chart 3), in which heptadecanylated 2,7-carbazole alternated with DTBT in backbone [59]. Possessed an Eg of 1.88 eV and a HOMO level of −5.5 eV (Table 2), P19 device exhibited a high PCE of 3.6% under the illumination of 90 mW/cm2 simulated light, with a Jsc of 6.92 mA/cm2 and Voc of 0.89 V. The FET hole mobility of PCDTBT was 0.001 cm2 V−1 s−1 . The efficiency of P19 devices was further improved to 6.1% when blending with PC71 BM and inserting a hole-blocking layer of TiOx under aluminum cathode [60]. Research endeavor from Cao’s group [61] pushed the efficiency of P19:PC71 BM based devices further to 6.8% by inserting a PFN/Ca layer between the active layer and Al cathode. The insertion of PFN (structure in Chart 3) layer brought in simultaneous enhancement of Voc , Jsc and FF in PV devices. Cao and coworkers [62] claimed that PFN could establish better interfacial contacts by decreasing the series resistance due to the fact that poly(2,7-carbazole)s could supply the necessary N–N interactions at the interface between the polymer donor and PFN. These better interfacial contacts enhanced the electron collection at metal cathode and decrease the possibility of hole–electron recombination in active layer, resulting much-improved PV performance. By introducing a bulkier side chains onto 2,7-carbazole, P19, an analogous polymer to P20 was developed, exhibiting a moderate PCE of 3.1% for the device [37], with all photovoltaic parameters similar to those of P19. Leclerc and coworkers [57] further reported another poly(2,7-carbazole) P21, with 4,7-dithien-2-yl2,1,3-benzothiadiazoxaline (DTBX) as accepting unit. Due to its symmetrical backbone, P21 showed good structural organization, resulting in good hole mobility and thus resultant improvement of Jsc and FF to achieve 2.4% PCE for BHJ cells. As discussed above, polymer donors with broad absorption spectra and high hole carrier motilities are preferable for high-efficiency BHJ cells, while high hole mobility requires a good thin-film ordering in the active layers [11,42,50,57,60]. Therefore, most polymer donors were incorporated with lateral flexible solubilizing chains to reduce the steric hindrance that prevents the stacking of adjacent conjugated backbones. Under this design concept, Inganäs and coworkers [63] designed planar copolymer P22 with BT moiety substituted with two lateral octyloxy chains. P22 showed good solubility at elevated temperature. Spin-coated from ODCB mixture solutions with addition of 2.5% low-vapor-pressure processing additive, 1,8-diiodooctane (DIO), P22:PC71 BM (1:2.5, w/w) device showed an excellent PCE of 5.4%. Cao and coworkers [62] reported another poly(2,7carbazole) P23, in which alkylated benzotriazole was used as electron-accepting unit. Possessed with a wider Eg (2.18 eV), P23 showed slightly broader absorption than P3HT. P23 device exhibited an initial PCE of 1.5%, while

Fig. 5. UV–vis spectra of polymers P24–P29 in thin films. Adapted with permission from Ref. [64]. Copyright 2010 American Chemical Society.

much improved PCE (2.8%) was obtained by sandwiching a thin layer of PFN (Chart 3) between Al cathode and active layer as for P19. Inspired by the successful design in 2D polyfluorene (P10 and P11) [47], Jen and coworkers [64] reported 2,7carbazole-based 2D conjugated copolymers P24–29, in which electron-accepting side chain was “grafted” onto a hole transporting carbazole-triarylamine backbone. Similar to P10–11, P24–29 shared identical HOMO levels with the polymer backbone. With their LUMO levels suppressed by accepting side chain including malononitrile and diethylthiobarbituric acid, P24–29 exhibited low Eg (1.74–1.88 eV). All polymers showed two obvious absorption peaks (Fig. 5), with the first absorption peak (at ∼380 nm), corresponding to the ␲–␲* transition of their conjugated main chains and their second broad absorption band (500–700 nm) attributed to the strong ICT interaction between their conjugated backbone and the pendant acceptor groups [47]. For polymers sharing the same backbones (P24 and P25, P26 and P27, P28 and P29), diethylthiobarbituric acid-functionalized polymers (P25, P27 and P29) exhibited obviously red-shifted ICT absorbance peak than malononitrile-functionalized counterparts, because of more-intense ICT interactions occurred in the former polymers. P28–29 exhibited the broadest absorption spectra with the strongest absorption intensity due to extended conjugation in the main chain. The PV characteristics of P24–29 devices indicated negligible impact of the side chains on carbazole on cell’s Voc , however the bulkier heptadecanyl chains led to a significant decrease in Jsc and FF. Among all six copolymers, P24:PC71 BM devices exhibited the highest PCE (4.2%). The bulkier alkyl chains of P26–P29 possibly were thought to result in imbalanced electron transporting and consequently decreased FF of polymer. A comparison with fluorene-based (P10, P11) and silafluorene-based (P12, P13) 2D copolymers, carbazolebased P24–29 exhibited almost one magnitude lower hole mobility. For P24–29, their hole-transporting abilities were strongly dependent upon the side chains on carbazole. P24 and P25 with short side-chains

Table 2 The optical, electrochemical, hole mobility, and PSC characteristics of P19–33. h (cm2 V−1 s−1 )

HOMO/LUMO (eV)

Polymer:PCBM

Jsc (mA/cm2 )

Voc (V)

FF

PCE (%)c

Ref.

576

1.87

–

−5.50/−3.60

16.5 (1.5) 26 (2.0) 16.6 (3.1) 107 (2.8) 2.7 (2.7) 2.7 (2.3) 7.4 (1.7) 5.6 (1.2) 13.6 (2.0) 8.3 (1.4) 50.7 (2.1)

551 570 579 523 548 586 545 575 546 598 557

1.72 1.87 1.95 2.18 1.83 1.74 1.88 1.77 1.83 1.74 1.70

– 1 × 10−4 1 × 10−4 – 2.4 × 10−4 1.1 × 10−3 5.3 × 10−5 1.5 × 10−4 8.8 × 10−5 1.6 × 10−4 –

−5.48/−3.46 −5.47/−3.65 −5.21/−3.35 −5.54/−3.36 −5.23/−3.40 −5.25/−3.51 −5.29/−3.41 −5.30/−3.53 −5.20/−3.37 −5.24/−3.50 −5.15/−3.45

8.3 (1.7) 14.4 (1.8) 19.5 (2.1)

551 431 –

1.85 1.51 2.02

– – –

−5.43/−3.62 −5.32/−3.59 −5.17/3.15

1:4 1:4b 1:4 1:3 1:4 1:2.5b 1:2 1:4b 1:4b 1:4b 1:4b 1:4b 1:4b 1:1b 1:1b 1:2b 1:2b 1:4b 1:4b 1:4b 1:3 1:3 1:2

6.92 10.6 12.7 7.66 3.7 9.8 4.68 8.94 8.18 7.51 6.23 7.18 6.53 5.83 6.82 8.13 8.32 9.10 11.4 9.84 4.83 3.94 9.17

0.89 0.88 0.90 0.80 0.96 0.81 0.9 0.91 0.92 0.91 0.89 0.83 0.84 0.70 0.85 0.70 0.85 0.78 0.81 0.80 0.90 0.70 0.69

0.63 0.66 0.59 0.50 0.60 0.69 0.65 0.51 0.47 0.44 0.38 0.39 0.40 0.41 0.50 0.46 0.61 0.56 0.66 0.57 0.48 0.46 0.57

3.6e 6.1 6.8 3.1 2.4 5.4 2.8 4.2 3.5 2.9 2.1 2.3 2.2 1.7g 2.9h 2.6g 4.3h 4.0g 6.1h 4.5i 2.1 1.3 3.6

[59] [60] [61] [37] [57] [58] [62] [64] [64] [64] [64] [64] [64] [66] [66] [66] [66] [66] [66] [66] [69] [37] [73]

Mn (kDa) (Mw /Mn )

P19

37 (2.0)

P20 P21 P22 P23 P24 P25 P26 P27 P28 P29 P30

P31 P32 P33

abs max (nm)

Definitions for all parameters are the same as those in Table 1. b Polymer:PC71 BM, in weight ratio, w/w. c PCE achieved at illumination of standard simulated solar light (AM 1.5G, 100 mW/cm2 ). e PCE measured at AM 1.5 (90 mW/cm2 ). g Al cathode. h PFN/Al bilayer cathode. i Ca/Al bilayer cathode.

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Eg (eV)

Polymer

1301

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Chart 3. Carbazole and indolo[3,2-b]carbazole based narrow bandgap polymers.

demonstrated much higher hole mobility (2.4 × 10−4 and 1.1 × 10−3 cm2 V−1 s−1 , respectively) than P26–P29 with bulkier side chains. As observed for P4, the side chains on the N atom of carbazole are crucial for improving the molecular organization of polymers in thin films and consequently affect their carrier transporting abilities in BHJ devices [65]. Cao and coworkers [66] reported a new 2,7-carbazole based DA polymer P30, in which 4,5-ethylene-bridged 2,7-carbazole was used as the donating moiety while modified DTQ was used as the accepting unit. Prepared with extremely high Mn (50.7 kDa) but moderate PDI (2.1), P30 films presented broad absorption with two maximum peaks at 401 and 557 nm. P30:PC71 BM device performance was found to be strongly dependent on the content of PC71 BM and the cathode composition in device processing (Table 2). Due to contribution of PC71 BM in the absorption [40,41], the device PV performance increased from 1.7% to 4.0% when the weight ration of P30:PC71 BM increased from 1:1 to 1:4 for the blend film. A bilayer cathode is also beneficial for devices, i.e., the device exhibited remarkably

higher PCE 4.5% with a Ca/Al cathode [66]. The performance of best-performing device was further improved to 6.1% when 5 nm PFN interlayer was inserted between the active layer and Al cathode. The insertion of PFN layer increased the interface contact between the active layer and metal cathode, reflecting in the improvement of Jsc and FF for the devices (Fig. 6).

5.2. Copolymers based on Indolo[3,2-b]carbazole As discussed for ladder-type oligo-p-phenylene (indenofluorene), rigid coplanar fused aromatic rings can efficiently enhance not only the intermolecular ␲–␲ stacking of main chains to improve charge mobilities [67,68], but the absorption with extended conjugation. Indolo[3,2b]carbazole (IC) can be formed by connecting two or more 2,7-carbazole at 2- and 7-positions with nitrogen bridges. This large-size fused aromatic conjugated system is helpful in making closer ␲–␲ packing of conjugated backbones.

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Fig. 6. (a) J–V curves and (b) external quantum efficiency (EQE) curves of solar cells with P30:PC71 BM (1:4) as the active layer when using different cathodes. Reproduced with permission from Ref. [66]. Copyright 2011 Wiley.

Hashimoto and coworkers [69] reported the first IC-based DA polymers P31 and P32. With DTBT as the accepting unit to tune the energy level alignment, P31 exhibited a narrow bandgap of 1.90 eV. P31:PCBM (1:3, w/w) device demonstrated a moderate PCE of 2.1%. The high Voc of P31 device (0.90 V) is close to those of BT based alternating polymers with donating units like fluorene (1.0 V) [38], dibenzosilole (0.9–0.97 V) [55] or carbazole (0.88 V) [59]. And this Voc is much higher than those of D–A derivatives with donating units such as cyclopenta[2,1-b:3,4-b ]dithiophene (0.6 V) [69,70], dithienosilole (0.44 V) [71] or dithieno[3,2-b:2 ,3 d]pyrrole (0.52 V) [72]. The device optimization revealed that performance of BHJ cell can be improved by forming interpenetrated network of donors and acceptors inside the active layer via adjusting the composition of the active layer and applying post-annealing treatment or solvent evaporation. Using a stronger electron-deficient moiety than DTBT, Hashimoto and coworkers [38] further developed a series

of IC-based copolymers. The effect of DTBT and 5,7-dithien2-yl-thieno[3,4-b]pyrazine (DTTP) on polymer properties and device performance was compared (Fig. 7). Compared with the DTBT-based polymers [PF-DTBT (P3), PC-DTBT (P20), and PIC-DTBT], the 2,3-dimethyl-DTTP-based polymers (TP1, TP3, and TP5) and the 2,3-diphenyl-DTTP-based polymers (TP2, TP4, and P32) had broader absorption ranges. TP1, TP3, and TP5 exhibited absorption maxima in the longer wavelength range at 577–590 nm with absorption coefficients that were ∼20% lower, and TP2, TP4, and P32 had the maxima at 635–651 nm with absorption coefficients that were ∼40% lower, which were red-shifted by about 50 and 100 nm in comparison with the DTBTbased counterpart polymers. The shape and the position of the absorption peaks and the absorption coefficients seemed to be determined mainly by the identity of the accepting segment, and the effects of the donor segment were small. Spin-coated from CHCl3 solution, the PV cell performance of all DTBT-based and DTTP-based polymers are

Fig. 7. UV–vis absorption spectra of polymer films on a quartz plate. Adapted with permission from Ref. [37]. Copyright 2010 American Chemical Society.

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Fig. 8. J–V curves of nine polymer solar cells based various polymers under illumination of AM 1.5 at 100 mW/cm2 . Adapted with permission from Ref. [38]. Copyright 2010 American Chemical Society.

compared in Fig. 8. The best device from DTTP-based P32 exhibited only a PCE of 1.3%, with nearly 0.1 eV lower Voc than all DTBT-based polymers. Compared with DTBTbased polymers (such as P4, P16–20, P31), the lower PCEs of all devices with DTTP-based polymers could be understood from two aspects: (i) Voc , which decreased due to an increase in the HOMO level of DTTP-based polymers, and (ii) Jsc , which decreased probably because of the poor charge separation or transport of DTTP-based polymers. By alternating hexylthiophene substituted DTBT and alkylated IC units, Lu et al. [73] developed polymer P33, for which a film showed strong broad absorption spectra with two maximum absorption peaks at 420 and 638 nm. Fabricated by slow solvent evaporation process, P33 device achieved 3.6% PCE with LiF/Al bilayer cathode. The high Jsc observed from this polymer confirms the finding that high crystallinity of the photoactive material is important to achieve high device performance. As discussed above, both polyfluorenes and polycarbazoles exhibit relatively low-lying HOMO level and thus high Voc in BHJ cells. Compared with DTTP and DTQ accepting units, DTBT-based copolymers often have deeper HOMO levels to give higher Voc values in devices. The employment of PFN layer to improve interfacial contacts [62,66] and slow solvent evaporation processing in making better interpenetrating networks in the active layer [73] proved to be effective strategies to achieve good Jsc and FF in devices. 6. Thiophene-based narrow bandgap polymers 6.1. Copolymers based on thiophene Polythiophenes (PTs) are considered to be the most important conjugated polymers for a broad spectrum of optoelectronic applications such as light-emitting diodes, field-effect transistors, and PSCs due to their excellent optical and electrical properties [68,74,75]. As the previous workhorse donor polymer for BHJ cell processing and optimization study, regioregular poly(3-hexylthiophene)

(P3HT) has demonstrated good self-assembly ability and crystallinity [76,77], which endow P3HT with a hole mobility high as 10−3 cm2 V−1 s−1 [78], and a field-effect transistor carrier mobility as high as 0.1–0.2 cm2 V−1 s−1 [79,80]. The material combination of P3HT and PCBM has given the highest reported PCE values of 4–5% in several research groups after the optimization of device fabrication process [81]. However, P3HT presents a large bandgap (2.0–2.1 eV) and thus lacks a broad absorption profile to collect a large fraction of the solar spectrum (ca ≤30%) [82,83]. Also, the high-lying LUMO of P3HT is 0.5 V higher than necessary for electron transfer to fullerene [10,83]. Moreover, the relatively high-lying HOMO limits the maximum Voc (∼0.65 V) and causes oxidation instability of the resulted cells at ambient conditions. Janssen and coworkers [84] reported the first successful terthiophene-containing polymer P34 (P34–54 structure in Chart 4) for photovoltaic devices, in which diketopyrrolopyrrole (DPP) was used as the acceptor. Due to its strong electron-deficient nature, a DPP core has been exploited in developing extremely narrow Eg conjugated polymers/oligomers for high-efficiency PV cells. Furthermore, the planarity of DPP skeleton and its ability to construct hydrogen bonds facilitate the resulting conjugated polymers in ␲–␲ stacking along conjugation backbone. Possessed with an extremely low Eg of 1.4 eV (Table 3), P34 films cast from ODCB solution showed extended absorption to 900 nm, 200 nm red-shifted relative to its films from CHCl3 solution. Fast evaporation from CHCl3 solution was found to be a failure since P34 did not crystallize, which resulted in poor morphology when making P34:PCBM films. A certain degree of crystallization could be achieved upon heating the film at 130 ◦ C. Spin-coated from CHCl3 :ODCB (4:1, v/v) mixture solution, P34:PCBM device reached an optimized PCE of 4%, higher than that of device processed from a single solvent, i.e., 1.1% from CHCl3 and 3.2% from ODCB. The performance improvement is attributed to better improvement of nanophase segregation morphology in thin films using mixture solvent (Fig. 9). P34:PCBM films processed from CHCl3 :ODCB mixture solvent (Fig. 9b) displayed smaller features (<100 nm), whereas films from CHCl3 presented domains with lateral dimensions of several hundred nanometers. The unfavorable morphology with larger dimension of phase-separated domains reduces the generation of charges and thereby the effectiveness of these layers. Inspired by Janssen’s work, Fréchet and coworkers [85] developed P34 analogous DPP-containing copolymers P35 and P36, in which furan rings replaced thiophenes. Furans and thiophenes used in organic dyes produced dyesensitized solar cells with similar optical and electronic properties as [86]. P35:PCBM device exhibited a PCE of 3.4% when casting the active layer from CB solution. Surprisingly, the P35:PC71 BM (1:3) device performed poorly, with PCE only 0.86%. The addition of high-boiling-point 1chloronaphthalene (CN) was found to be effective in device optimization of P35. An addition of 1% CN (by volume) helped to improve P35 device performance from 3.4% to 3.8% with device structure ITO/PEDOT:PSS/P35:PCBM (1:3, w/w)/LiF/Al. The P35:PC71 BM (1:3) device exhibited over

Table 3 The optical, electrochemical, hole mobility, and PSC characteristics of P33–45. Mn (kDa) (Mw /Mn )

abs max (nm)

Eg (eV)

h (cm2 V−1 s−1 )

HOMO/LUMO (eV)

Polymer:PCBM

Jsc (mA/cm2 )

Voc (V)

FF

PCE (%)c

Ref.

P34 P35 P36 P37 P38 P39 P40 P41

20 (3.4) 66 (2.1) 29 (2.0) 54 (3.2) 14 (5.4) 16 (4.9) 9.7 (1.4) 35 (1.3)

650 789 767 – 795 784 572 775

1.4 1.41 1.35 1.36 1.38 1.28 1.82 1.46

– – – 0.04 1.42 0.037 a 1.0 × 10−4 2.0 × 10−2

– −5.40/−3.80 −5.50/−3.80 −5.17/−3.61 −5.06/−3.68 −5.04/−3.76 −5.56/−3.10 −5.30/−3.57

P42 P43 P44 P45 P46 P47 P48 P49 P50 P51 P52 P53 P54

17.7 (1.8) 64 (4.7) 20 (1.4) 11.9 (2.1) 18 (1.2) 8 (1.2) 28 (1.6) 48 (1.7) 35 (1.63) 5.6 (1.7) 8.7 (1.6) 80 (5.0) 111 (2.3)

560 688 590 577 652 644 – 679 – 558 605 546 585

1.84 1.67 1.60 1.61 1.45 1.51 1.73 1.69 1.59 1.85 1.81 1.93 1.74

a

5.2 × 10−4 8.9 × 10−3 – – 3.0 × 10−3 a 3.0 × 10−6 1.0 × 10−4 – 0.20 3.1 × 10−4 3.56 × 10−3 2.7 × 10−4 3 × 10−4

−5.07/−3.55 −5.26/−3.59 −5.37/−3.51 −5.35/−3.52 −5.05/−3.27 −5.02/−3.19 −5.57/−3.88 −5.60/−3.50 −5.18/– −5.18/−3.09 −5.06/−2.81 −5.05/−3.12 −5.66/−3.99

1:2b 1:3b 1:3b 1:2b 1:2b 1:2b 1:1.5 1:1b 1:2–3b 1:3.6b 1:2 1:2b 1:2b 1:2b 1:1b 1:1b 1:2b – 1:1b 1:1b 1:1b 1:3b 1:1b

11.5 11.2 9.1 11.8 15.0 8.9 8.0 11.0 16.2 15.7 8.0 8.1 12.5 11.1 12.7 9.8 12.2 12.6 12.3 7.9 11.9 8.5 9.0

0.61 0.74 0.73 0.65 0.58 0.57 0.95 0.70 0.62 0.61 0.70 0.76 0.75 0.74 0.68 0.60 0.88 0.85 0.72 0.68 0.77 0.63 0.61

0.58 0.60 0.58 0.60 0.61 0.59 0.62 0.47 0.55 0.53 0.54 0.50 0.59 0.48 0.55 0.50 0.68 0.68 0.70 0.54 0.61 0.60 0.54

4.0 5.0 4.1 4.7 5.4 3.0 4.7 3.2 5.5 5.1 3.0 3.1 5.5 3.9 5.1 3.0 7.3 7.3 6.3 2.9 5.6 3.2 3.0

[84] [85] [85] [89] [90] [90] [92] [68,95] [35] [96] [97] [98] [99] [99,100] [102] [103] [105] [107] [108] [109] [110] [118] [119]

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Polymer

Definitions for all parameters are the same as those in Table 1. a h measured by SCLC. b Polymer:PC71 BM, in weight ratio, w/w. c PCE achieved at illumination of standard simulated solar light (AM 1.5G, 100 mW/cm2 ).

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Chart 4. Thiophene-based narrow bandgap polymers.

4-fold improvement in PCE (average 4.7%, highest 5.0%) with the addition of 9% CN in CB solution (Fig. 10). With the same device architecture, P36:PC71 BM (1:3, w/w) presented a PCE of 4.1% when spin-coated from CB solution with 5% CN inside. The dramatic difference in device performance with and without CN additive is most likely due to the optimization of blend morphology (Fig. 11). As clearly observed from AFM images, P35:PC71 BM (1:3) blend films processed without additives exhibited coarse phase separation between polymer and PC71 BM, in which large micrometer-sized domains were formed (Fig. 11a). In contrast, the addition of CN led to much finer phase separation between the two materials and the formation of fiber-like interpenetrating morphologies at the length scale of ∼20 nm (Fig. 11b), which is close to the ideal domain size, assuming an exciton diffusion length of 5–10 nm [87,88].

Janssen and coworkers [89] reported DPP-thiophene alternating polymer P37 with two different Mn (54 kDa and 10 kDa). An Eg of 1.3 eV was calculated from the onset of optical absorption of P37. Though different in Mn , two polymers of P37 exhibited identical and nearly balanced hole and electron mobilities of 0.04 and 0.01 cm2 V−1 s−1 , respectively, by FET measurement. The device of P37 was optimized using DIO as a processing additive. For high Mn polymer, P37:PCBM device exhibited a PCE of 3.8% when adding 25 mg/mL DIO in CB solution; while P37:PC71 BM exhibited an enhanced PCE of 4.7% by making blend film from CHCl3 solution with 100 mg/mL DIO. In contrast to mobility, the performance of P37 device was affected by the molecular weight of the polymer (Fig. 12). Devices with low Mn P37 did not exceed PCE of 1.3% with PCBM or 2.7% with PC71 BM, respectively, under the same device processing conditions. The differences were caused by a

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Fig. 9. Surface topography of P34:PCBM film processed from CH3 Cl (a) and from CH3 Cl:ODCB (b). z-range is 42 and 35 nm, respectively, for (a) and (b). Reproduced with permission from Ref. [84]. Copyright 2010 Wiley.

Fig. 10. (a) J–V curves of optimized P35:PC71 BM devices spin-coated out of chlorobenzene (with no additive and with 9 vol% CN). (b) External quantum efficiency spectra of optimized devices. Reproduced with permission from Ref. [85]. Copyright 2010 American Chemical Society.

Fig. 11. AFM phase images of P35:PC71 BM (1:3) blend films spin-coated from (a) CB only and (b) CB + 9% CN. Inset: height images of the same films, data scale is 0–60 nm. Reproduced with permission from Ref. [85]. Copyright 2010 American Chemical Society.

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Fig. 12. (a and b) J–V curves and (c and d) EQEs for P37:PCBM bulk-heterojunction solar cells with (a and c) high and (b and d) low Mn . Reproduced with permission from Ref. [89]. Copyright 2009 American Chemical Society.

reduction in photocurrent (Fig. 12b, d vs. a, c) and exemplified the importance of molecular weight as a crucial parameter in the formation of efficient bulk heterojunctions [89]. By replacing thiophene or bithiophene in P37 with larger fused thieno[3,2-b]thiophene to increased the intermolecular association, Bronstein et al. [90] developed DPP-based copolymers P38 and P39. With a similar HOMO value ∼−5.05 eV, P38 and P39 exhibited low Eg of 1.38 and 1.28 eV, respectively. Both polymers presented ambipolar charge-transport characteristics in FET studies, i.e., P38 showed a maximum saturation hole and electron mobilities of 1.95 and 0.03 cm2 V−1 s−1 , respectively, while P39 showed only modest hole mobilities but higher electron mobilities. The maximum hole mobility observed in P38 is the highest value for polymer-based FET reported to date. The superior charge mobilities are attributed to the existence of coplanar thieno[3,2-b]thiophene, which endows P38–39 with more delocalized HOMO

distribution along the backbone and may be expected to enhance intermolecular charge-carrier hopping. Spincoated from CHCl3 /ODCB (4:1) solution, P38 devices displayed similar Voc and FF values but remarkably higher Jsc than P39, resulting in a much higher PCE (5.4%) than P39 (3.0% PCE). The difference in PV performance was attributed to the LUMO levels of two polymers and the morphology of the active layers. The slightly lower LUMO level of P39 may be too close in energy to that of PC71 BM, which resulted in less efficient charge separation [90]. A more favorable morphology may exist in the solid-state packing effects, in which reduced distance between two DPP repeat units of P39 could in turn reduce the amount of fullerene intercalation. This behavior may be easily explained with the discovery of McGehee and coworkers [91], that the intercalation of fullerene derivatives between the side chains of conjugated polymers was controlled by adjusting the fullerene size intercalated and compared with the properties

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Fig. 13. Molecular structures of pBTTT, PC71 BM, and bisPC71 BM, schematics showing possible structures for intercalated pBTTT (a), and a space-filling ChemDraw model of pBTTT, PC71BM, and bisPC71BM to show their relative sizes (b). The second side group on bisPC71BM can attach to the fullerene at a number of different locations. Reproduced with permission from Ref. [91]. Copyright 2009 American Chemical Society.

of non-intercalated poly(2,5-bis(3-hexadecylthiophen2-yl)thieno[3,2-b]thiophene) (pBTTT):fullerene blends (Fig. 13). Wei and coworkers [92] reported bithiophenecontaining P40 using thieno[3,4-c]pyrrole-4,6-dione (TPD) as the accepting unit. The electro-deficient TPD adopts a symmetric, rigidly fused, coplanar structure to promote intramolecular interactions and to reduce Eg by lowering the HOMO level when incorporating into s DA polymer backbone [93]. With an Eg of 1.82 eV, P40 exhibited a maximum absorption at 572 nm and a vibronic shoulder at 628 nm. The X-ray diffraction pattern revealed highly ordered arrangement of P40 in the solid films, with strong ␲–␲ stacking between adjacent backbones. P40:PCBM device presented the best PCE of 4.7% with controlled thin-film thickness within 90–100 nm. A moderately homogeneous surface without significant phase segregation was observed for the P40:PCBM film. 6.2. Copolymers based on cyclopentadithiophene Recent years have witnessed the emergence of several important fused coplanar heterocycles in replacement of thiophene/bithiophene rings. For example, when biothiophene rings are tied together in a coplanar fashion by methylene or silane, to obtain cyclopentadithiophene (CPDT) or dithieno-[3,2-b:2 ,3 -d]silole. respectively. These two fused thiophene moieties have become popular donating blocks in design DA polymers for photovoltaic applications. A facile synthetic approach for solution-processable CPDT was reported by Turner and coworkers [94] early in 2003. It was not until 2007 that Brabec and coworkers [70,95] reported the first successful CPDT-containing polymer, P41 (PCPDTBT) for PSC applications. Possessed with low Eg (1.46 eV), P41 exhibited a wide absorption, ranging from 300 to 850 nm [68]. Without any processing additive for film casting, P41 exhibited a 3.2% PCE for P41:PC71 BM devices and 2.7% PCE for P41:PCBM ones [68]. Thermal

annealing of blend films was ineffective in morphological optimization of P41 [35]. Importantly, PCPDTBT-based PSCs have presented a wide scope for device optimization via controlling the thin-film morphology by the use of processing additives such as alkanedithiols [35] and 1,8-dihalo-octanes [96]. By controlling the thin-film morphology of PCPDTBT:PC71 BM (1:2–3) blends using 0.8–1% alkanedithiols for film spin-coating, the device achieved significant enhancement in PCE from 2.8% to 5.5%. With 2.5% 1,8-diiodo-octane as a processing additive, PCPDTBT:PC71 BM (1:3.6) cells achieved significant improvement in PCE from 3.4% to 5.1% [96]. As depicted in Fig. 14 [96], processing additives were used to fine-tune the nanosphase segregation morphology in PCPDTBT:PC71 BM blend films. Wei and coworkers [97] reported another CPDTcontaining DA polymer P42 with bithiazole as the accepting unit. P42 exhibited high crystallinity in thin films, which enabled the formation of highly selfassembled ␲–␲ stacking in the solid state. By optimizing P42:PCBM weight ratio as 1:2 to balance the charge transporting in the active layer, P42 device contributed a PCE of 3.0%. Using thieno[3,4-c]pyrrole-4,6-dione (TPD) as the accepting unit, Jenekhe and coworkers [98] prepared TPD-containing copolymer P43. Its device showed a moderate PCE of 3.1%, with Jsc of 8.12 mA/cm2 , Voc of 0.76 V, and FF 50% by blending with PC71 BM in 1:2 weight ratio. Ding and coworkers [99,100] reported two solutionprocessable CPDT copolymers P44 and P45, in which electron-deficient s-tetrazine (Tz) was used as accepting unit. Prepared via Stille coupling reaction, P45 showed almost 50% lower Mn than P44, due to the steric hindrance from the bulkier side chains. The side chain showed no impact on the energy level alignment and Eg value for P44 and P45. Both polymers were stable up to 220 ◦ C before they decomposed into dinitrile compounds by breaking Tz linkage at 250–300 ◦ C. P44:PC71 BM (1:2) device showed an impressive PCE of 5.5%, with Jsc of 12.5 mA/cm2 , Voc of 0.75 V, and FF of 59% when processed with 3.0% DIO in ODCB

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Fig. 14. Schematic depiction of the role of the processing additive in the self-assembly of bulk heterojunction blend materials (a) and structures of PCPDTBT, PC71 BM, and additives (b). Reproduced with permission from Ref. [96]. Copyright 2008 American Chemical Society.

solution to control the blend morphology [98]. Under similar conditions, P45 showed a lower PCE of 3.9% and lower Jsc and FF, which resulted from its lower Mn and hindered ␲–␲ stacking due to the longer side chains [38]. 6.3. Copolymers based on dithieno-[3,2-b:2 ,3 -d]silole Dithieno-[3,2-b:2 ,3 -d]silole-containing polymers have recently attracted considerable attention due to their potential applications in the field of optoelectronic devices. Many silole-containing polymers have been proven to achieve high Voc by reducing the LUMO levels of polymers [101]. It is feasible to realize a low-lying HOMO energy level and high Voc simultaneously under the conditions that a sufficient energy level offset (0.3–0.5 V) between the LUMO levels of silole-containing donor and PCBM is ensured. Furthermore, silole-containing polymers have also proved to be effective in improving hole mobility and rendering higher efficiencies compared to CPDT in PSC applications [50,102]. Yang and coworkers [102,103] reported the first dithieno-[3,2-b:2 ,3 -d]silole-containing DA polymers P46 and P47 for successful photovoltaic applications. Possessed with a low Eg of 1.45 eV, P46 exhibited a broad absorption ranging from 350 to 800 nm [102]. Though similar to P41 in Eg , P46 presented 3-fold higher FET hole mobility (3 × 10−3 cm2 V−1 s−1 ) than P41. P46:PC71 BM (1:1) devices achieved an average PCE of 4.7% out of 100 devices, with the best PCE reaching 5.1%, featuring much improved Jsc and FF than P41. By franking thiophenes on BT sides, a DTBT-containing polymer P47 was prepared with a lower Mn of 8 kDa [103]. P47 devices achieved lower PCE (3.0%) than P46. The research group also reported P47 analogues poly[(4,4 -bis(2-dodecyl)dithieno-[3,2-b:2 ,3 polymer d]silole)-2,6-diyl-alt-(2,1,3-benzothiadiazole)-4,7-diyl]

(PSiDTBT12) by introducing dodecyl chains onto dithieno[3,2-b:2 ,3 -d]silole [103]. Though sharing similar Eg and energy levels with P47, PSiDTBT12 device presented better performance (3.4% PCE) than P47. This indicates device performance can be further improved by the modification of side chains on dithieno-[3,2-b:2 ,3 -d]silole. Recent studies have shown that “push–pull” structures can not only lower the bandgap of DA copolymers, but also enhance the charge carrier mobility with reduced interchain ␲–␲ stacking distance [104]. In this regard, a large library of DA polymers have been developed for successful photovoltaic applications by alternating the above mentioned electron-rich units such as 2,7-carbazole (P19–29), IC (P31–33), CPD (P41–45), dithieno-[3,2-b:2 ,3 -d]silole (P46–50), and BDT (P53–72), with electron-deficient units like BT, DPP, ester- or ketone-substituted thieno[3,4b]thiophene, and TPD. Tao and coworkers [105] recently reported a dithieno[3,2-b:2 ,3 -d]silole-containing polymer P48 with N-octylthieno[3,4-c]pyrrole-4,6-dione as accepting unit. In comparison to CPDT counterpart P43 [98], P48 exhibited better hole-transport properties and lower HOMO energy level. Hence, P48 achieved a low Eg (1.73 eV) and a low-lying HOMO level (−5.57 eV) simultaneously. By processing from dichlorobenzene (DCB) solution, the P48:PC71 BM (1:2, w/w) device exhibited a Voc of 0.90 V, Jsc of 10.95 mA/cm2 and FF of 63%, affording an overall PCE of 6.2%. The device configuration was ITO/PEDOT:PSS (30 nm)/P48:PC71 BM (1:2, w/w) (90 nm)/BCP (5 nm)/Al, in which BCP stands for bathocuproine and functions as a hole/exciton blocking layer [106]. By changing the device fabrication solvent from DCB to CB with 3% DIO, the best device demonstrated a landmark PCE of 7.3% under the same device processing conditions.

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Fig. 15. Height (top panel) and phase (bottom panel) AFM images of P48/PC71BM (1:2) films spin-cast from (a) DCB, (b) DCB + 3% DIO, (c) CB and, (d) CB + 3% DIO. Reproduced with permission from Ref. [105]. Copyright 2011 American Chemical Society.

Similarly to device optimization of P44 [99], the processing additive DIO was found to play an important role in improving the nanoscale morphology of P48:PC71 BM thin films (Fig. 15). Large isolated PC71 BM domains (∼0.4 ␮m diameter) were observed in P48:PC71 BM films when processing devices without DIO (Fig. 15a and c). These large isolated domains are unfavorable for efficient exciton dissociation and charge transport [106], reflected by the significant drop of Jsc from 12.2 to 2.6 mA/cm2 and resultant device PCE less than 1.0%. With the addition of 3% DIO in DCB solution (Fig. 15b), the active layer showed a much more uniform and finer domain structure with an average dimension of 20–40 nm, ideal for forming effective polymer:PC71 BM interpenetrating networks. The device performance was thus greatly improved to over 7%. This finding further highlights the importance of morphology control of the active layer in pursuit of high-performance solar cells. By bridging two thiophenes rings with a germole ring, Reynolds and coworkers [107] reported dithienogermolecontaining polymer P49 (P Ge). Compared with P48 (P Si), the longer C Ge bond further removed the bulky side-chains from the planar heterocycle and allowed stronger ␲-stacking interactions to occur. BHJ cell (ITO/ZnO/P49:PC71 BM/MoO3 /Ag) fabricated using 5% DIO as processing additive achieved a best PCE of 7.3%, with both Jsc and FF significantly increased. TEM images (Fig. 16) showed the addition of 5% DIO led to phase separation on the order of tens of nanometers, with no large aggregates of PC71 BM or polymer observed. This small-scale phase separation was on the order of the exciton diffusion length. Chen and coworkers [108] reported a quaterthiophenecontaining P50 with BT as accepting units. P50:PC71 BM device exhibited an impressive PCE of 6.3%, due to efficient charge separation and transport in the active layer. Long fibers (∼20−25 nm in diameter) were observed in the morphology of active layer after annealing. Li and coworkers [109,110] reported two dithieno[3,2-b:2 ,3 -d]silole based DA copolymers using dithiazole

(P51) or thiazolothiazole (TTz) (P52) as accepting fragments. Compared with bithiazole, TTz has a more rigid and coplanar fused ring structure to ensure a highly extended ␲-electron system and strong ␲–␲ stacking. Possessed with higher Mn , P52 showed 10-times higher hole mobility than P51, i.e., 3.56 × 10−3 cm2 V−1 s−1 for P52 and 3.1 × 10−4 cm2 V−1 s−1 for P51. P51:PC71 BM (1:2, w/w) device achieved a moderate PCE of 2.9%. However, P52:PC71 BM (1:1, w/w) demonstrated excellent performance with PCE of 5.6% via thermal-annealing of blend films at 100 ◦ C for 15 min. This performance is even better than its carbazole analogous PCDTTz reported by Moon and coworkers [111], which delivered a PCE of 4.9%, Voc of 0.86 V, Jsc of 9.15 mA/cm2 , and FF of 62% [111]. 6.4. Copolymers based on dithieno-[3,2-b:2 ,3 -d] thiophene Dithieno[3,2-b:2 ,3 -d]thiophene (DTT) can be prepared by bridging bithiophene rings with sulfur atom. This fused conjugated DTT moiety has been widely utilized in the development of a variety of conductive materials for electroluminescence, two-photon absorption and excited fluorescence, nonlinear opticals, photochromism, TFT, and dye-sensitizied solar cells [112–117]. Ong and coworkers [118] reported another DTTcontaining copolymer P53. With an Eg of 1.93 eV, the amorphous P53 presented a maximum absorption at 546 nm and shoulder peak at 577 nm. The best-performing P53:PC71 BM (1:3, w/w) device exhibited a PCE of 3.2%, the best device performance reported to date for DTTcontaining polymer donors. By alternating DTQ with thiophene flanked DTT, a DTTcontaining polymer P54 was synthesized with extremely high Mn (111 kDa) and relatively narrow PDI (2.3) [119]. The introduction of flanking thiophene on DTT core is beneficial in extending conjugation to improve hole transport property and enhance absorption. Long linear alkyl chains were introduced in the DTT ring for solubility

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Fig. 16. Cross-sectional TEM images of P48:PC71 BM and P49:PC71 BM-based cells without any additives (top) and with 5% DIO (bottom). Reproduced with permission from Ref. [107]. Copyright 2011 American Chemical Society.

considerations. The P54:PC71 BM (1:1, w/w) demonstrated a Voc of 0.61, Jsc of 9.03 m/A cm2 , and a FF of 0.54, giving a moderate PCE of 3.0%. 6.5. Copolymers based on benzo[1,2-b:4,5-b ]dithiophene The benzo[1,2-b:4,5-b ]dithiophene (BDT) has emerged as an attractive building block for making DA photovoltaic polymers due to its large planar conjugated structure, ideal for efficient ␲–␲ stacking. Alternating BDT and bithiophene copolymers exhibited high FET mobilities between 0.15 and 0.25 cm2 V−1 s−1 and enhanced stability [120], leading to great potential for FET and PSC applications [121]. Yu and coworkers [122,123,125–127] have made great contributions in pushing BDT-based DA polymers for excellent photovoltaic solar cells. A large library of highefficiency polymers has been developed in recent 3 years. Using ester substituted thieno[3,4-b]thiophene as the accepting unit, a series of BDT copolymers P55–P59 and PTB6 (X = H, R1 = n-butyloctyl, R2 = n-octyloxy) were developed, with the effect of side chains on backbones investigated in details (P55–79 structure in Chart 5) [122,123]. Thieno[3,4-b]thiophene contributed to the stability of the quinoidal structure of the backbone, and hence narrowed the energy gap of the resulting polymers [124]. Though different in side chain lengths, all polymers exhibited similar Eg (∼1.60 eV) and wide absorption, which covered the whole visible light region (Fig. 17). With the large planar conjugated structure, all polymers demonstrated high hole mobility (>4.0 × 10−4 cm2 V−1 s−1 ) with SCLC model. All polymers exhibited both good solubility and oxidative stability. P55 exhibited an absorption maximum at 690 nm with the onset at 784 nm for thin films, which coincided with the corresponding maximum photon flux region of the solar spectrum. The device P55:PCBM (1:1, w/w) achieved an impressive PCE of 4.8% (Table 4). Further improved to 5.3% PCE by blending with PC71 BM with the same device architecture. The absorption spectra of the two devices almost covered the whole visible light range. High EQEs were obtained for both devices: the maximum EQE reached over 60% at 650 nm and 50% in the 550–750 nm range (for P55:PCBM); while a higher EQE almost all over 60% in the 400–750 nm range for the later. Preferred thin-film morphology, featuring nanophase separation with very fine domains and interpenetrating networks was observed in the blend films, in which small

nanofibers (∼5 nm width) distributed in the entire films were found in TEM images. All results have shown that the effective interpenetrating networks formed in the blend film favor the exciton dissociation in the interfacial area and transport in the corresponding domains. The effect of chain length was further investigated with P55–P59 and their analogous polymer PTB6 (X = H, R1 = n-butyloctyl, R2 = n-octyloxy) [123]. Though different in side chain lengths, all polymers exhibited quite similar Eg (∼1.60 eV) and wide absorption, which covered the visible light region (Fig. 17). With the large planar conjugated structure, all polymers demonstrated high hole mobility (>4.0 × 10−4 cm2 V−1 s−1 ) with SCLC model. With electron-withdrawing fluorine in polymer backbone, P58 exhibited significantly lower lying HOMO level compared with P59. By spin-coating a blend film from DCB solution, the devices P56–59:PCBM (1:1, w/w) exhibited a high PCE of 5.1%, 5.5%, 3.1% and 3.0%, respectively. The better device performance of P56 and P57 than P55 can be explained with their finer segregation morphology (Fig. 18a–c), which may be due to the increased miscibility of the polymer with PCBM after shortening the dodecyl side chain into 2-ethylhexyl side chain. By adding 3% (vol) DIO in DCB solution to control thin-film ordering [98], P57–P59 devices harvested dramatic improvement in PCE to 5.9%, 6.1% and 4.1%, respectively. The significant improvement in performance from 3.1% to 6.1% for P58 could be clearly explained with the dramatic enhancement in thin-film ordering in the active layer (Fig. 18g vs. h). Using branched 2-ethylhexyl side chain, Yu and coworkers [125] further developed P60. By spin-coating P60:PC71 BM film with weight ratios of 1:1, 1:1.5 and 1:2, P60 devices harvested a corresponding PCE of 5.7%, 6.2% and 5.6%, with the same device structure as for P55–59. With addition of 3% DIO in DCB to optimize the blend morphology, P60:PC71 BM (1:1.5, w/w) device demonstrated a landmark PCE as high as 7.4%, with a Jsc of 14.5 mA/cm2 , Voc of 0.74, and FF of 0.69. Uniform morphology of blend film was formed to facilitate charge transport. P60:PC71 BM cell was further optimized by Cao and coworkers to a new world-record PCE of 8.4% by insertion of a PFN layer to improve the interfacial contacts [61]. The effect of fluorination of a BDT block was investigated by Yu and coworkers [126] with polymers P58 and P60–P62. As the HOMO–LUMO energy level alignment shown in Fig. 19, the fluorination of thienothiophene and

Table 4 The optical, electrochemical, hole mobility, and PSC characteristics of P53–72. Mn (kDa) (Mw /Mn )

abs max (nm)

Eg (eV)

h (cm2 V−1 s−1 )

HOMO/LUMO (eV)

Polymer:PCBM

Jsc (mA/cm2 )

Voc (V)

FF

PCE (%)c

Ref.

P55 P56 P57 P58 P59 P60

18.3 (1.3) 16.8 (1.4) 15.9 (1.5) 14.6 (1.3) 16.1 (1.4) 46.4 (2.1)

687 683 682 682 677 675

1.62 1.59 1.60 1.63 1.62 1.61

4.5 × 10−4 4.0 × 10−4 7.1 × 10−4 7.7 × 10−4 4.0 × 10−4 5.8 × 10−4

−4.9/−3.2 −4.94/−3.22 −5.04/−3.29 −5.12/−3.31 −5.01/−3.24 −5.15/−3.31

P61 P62 P63 P64 P65 P66 P67 P68 P69 P70 P71 P72 P73 P74 P75 P76 P77 P78 P79

11.2 (2.4) 30 (2.6) 30 (1.8) 13 (2.6) 5.6 (–) 21.9 (4.1) 27.4 (1.8) 41.2 (1.7) 33.8 (2.6) 19.2 (2.1) 22.7 (2.6) 47.2 (3.8) 42 (2.5) 39 (3) 35 (2.7) 47.6 (2.6) 42.2 (2.4) 40.5 (3.2) 9 (–)

611 613 675 610 595 644 596 – – 634 661 658 608 616 627 – – – 538

1.75 1.73 1.61 1.80 1.70 1.70 1.75 1.70 1.70 1.67 1.63 1.68 1.75 1.70 1.73 1.98 2.00 1.58 2.00

1.8 × 10−4 7.0 × 10−5 2 × 10−4 – a 1.6 × 10−5 a 3.8 × 10−5 – – – – – – – – – 3.34 × 10−6 6.67 × 10−5 1 × 10−5 1.67 × 10−5

−5.41/−3.60 −5.48/−3.59 −5.12/−3.55 −5.56/−3.75 −5.26/−2.96 −5.33/−3.17 −5.31/−3.44 −5.40/−3.13 −5.54/−3.33 −5.32/−3.58 −5.41/−3.72 −5.34/−3.45 −5.48/−3.73 −5.57/−3.87 −5.4/−3.67 −5.29/−2.87 −5.36/−3.05 −5.19/−3.26 −5.30/−3.20

1:1.2b 1:1 1:1 1:1 1:1 1:1.5b 1:1.5b 1:1.5b 1:1.5b 1:1.5b 1:2a 1:1 1:1 1:2b 1:1 1:1 1:2b 1:2b 1:1.5b 1:2 1:1.5 1:1.5 1:2 1:2 1:1b 1:3b

15 12.8 13.9 13 10.7 14.5 15.75 11 9.1 14.7 9.8 9.7 7.8 10.7 10.03 12.91 12.53 12.05 13.6 8.1 9.7 11.5 11.1 11.8 11.7 10.4

0.56 0.60 0.72 0.74 0.66 0.74 0.76 0.68 0.75 0.70 0.85 0.81 0.83 0.92 0.87 0.91 0.82 0.86 0.89 0.87 0.81 0.85 0.70 0.79 0.80 0.85

0.63 0.66 0.58 0.61 0.58 0.69 0.70 0.43 0.39 0.64 0.66 0.55 0.60 0.58 0.57 0.61 0.55 0.60 0.51 0.56 0.67 0.68 0.55 0.73 0.61 0.59

5.3 5.1 5.9 5.9 4.1 7.4 8.4 3.2 2.7 6.6 5.5 4.3 3.9 5.7 5.0 7.2 5.6 6.2 6.1 4.0 5.7 6.8 4.4 7.1 6.0 5.2

[122] [123] [123] [123] [123] [125] [61] [126] [126] [127] [93] [128] [128] [129] [130] [130] [131] [131] [132] [133] [133] [133] [134] [134] [135] [136]

L. Bian et al. / Progress in Polymer Science 37 (2012) 1292–1331

Polymer

Definitions for all parameters are the same as those in Table 1. P78 measured under the illumination of AM 1.5G at 95 mW/cm2 . a h measured by SCLC. b Polymer:PC71 BM, in weight ratio, w/w. c PCE achieved at illumination of standard simulated solar light (AM 1.5G, 100 mW/cm2 ).

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Chart 5. Benzo[1,2-b:4,5-b ]dithiophene based narrow band gap polymers.

benzodithiophene in backbone lowered both the HOMO and LUMO energy levels of polymers and resulted in slightly widened Eg (0.1–0.2 eV). Similar behavior was observed for P62, which exhibited 0.2 eV larger Eg than P61 (1.73 eV). Spin-coated from DCB solutions, P61–62:PC71 BM devices delivered a much lower PCE of 2.7% (P61) and 3.3%

(P61) than P55–60 with their optimized weight ratios of 1:1.5 and 1:1. The lower device performance was attributed to the poor compatibility between PC71 BM molecules and perfluorinated polymer backbone. The poor compatability was explained by the enhanced self-organization properties of the polymer chains and the fluorophobicity effect.

Fig. 17. (a) HOMO and LUMO energy levels of the polymers. Energy levels of PC61BM are listed for comparison. (b) UV–vis absorption spectra of the polymer films. Reproduced with permission from Ref. [123]. Copyright 2009 American Chemical Society.

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Fig. 18. TEM images of polymer/PCBM blend films P55 (a), P56 (b), P57 (c), P58 (d), P59 (e), and PTB6 (f) and polymer/PCBM blend films prepared from mixed solvents dichlorobenzene/diiodooctance (97/3, v/v) P57 (g), P58 (h), and P59 (i). Reproduced with permission from Ref. [123]. Copyright 2011 American Chemical Society.

Fig. 19. HOMO and LUMO energy levels of the polymers P58 and P60–62. Adapted with permission from Ref. [126]. Copyright 2011 American Chemical Society.

Furthermore, it was found that perfluorination of the polymer backbone resulted in poor photochemical stability against singlet oxygen attack [126]. It is commonly known that alkoxy groups have much stronger electron donating effects than alkyl chains. Conjugated polymers with alkoxy groups as a substituent usually exhibit higher HOMO levels than their alkyl-substituted counterparts. These observations suggest that it is possible to further reduce the HOMO levels of BDT-based copolymers P55–P60 by the removal of the oxygen atom on the ester group in thieno[3,4-b]thiophene, Yang and coworkers [127] reported a new BDT copolymer P63, in which ketone-substituted thieno[3,4-b]thiophene was used as accepting unit. With further reduced HOMO level to −5.12 eV, P63 exhibited an Eg of 1.61 eV. The best PSC based on P63:PC71 BM showed a high PCE of 6.6%. By changing the electron-accepting unit, TPD-based copolymer P64 was reported by Leclerc and coworkers [93]

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and alklyated DTBT based P65 and P66 were reported by You and coworkers [128]. The device P64:PC71 BM showed a high PCE of 5.5%. P65 and P66 exhibited one magnitude lower hole mobility than P55–60, which in turn resulted in lower Jsc and FF in P65 and P66 devices, giving lower PCE values (4.3% for P65 and 3.9% for P66). However, P66 presented over double mobility due to high Mn when compared with P65. By introducing thiopene moiety onto BDT, Yang and coworkers [129] reported polymer P67 with almost 4-fold Mn (27.4 kDa) as P65. This higher Mn endowed P67:PC71 BM (1:2, w/w) device with an optimized PCE of 5.7% without annealing treatment, due to much improved Jsc in comparison to P65. A recent effective strategy in DA polymer design proposed fluorination of accepting units to tune the energy level alignment of the polymers [126]. Several BDTcontaining polymers P68–71 were reported along this line. You and coworkers [130] reported P68–P69, in which fluorinated BT and benzo[d][1,2,3]triazoles (TAZ) were employed as new electron-accepting units. The lone pair on the nitrogen atom in TAZ is more basic than the lone pairs on sulfur in BT, hence it is more easily donated into the triazole ring. This endowed P69 with more electron richness, leading to a higher LUMO energy level. Therefore, wider Eg are observed for TAZ-based P69 than the BT-based P68. P69 cells showed a landmark PCE of 7.2%, while a P68:PCBM device showed a PCE of only 5.0%. The formation of a near optimal morphology in the active layer without annealing or processing additives offers P69 great potential in developing high performance solar cells. Similarly, by alternating mono-alkoxylated or fluorinated BT with BDT, Dai and coworkers [131] recently reported copolymers P70 and P71. The fluorination of BT in P71 led to even lower HOMO and LUMO, indicating its effectiveness in controlling both energy levels for the resultant polymers. P70–71:PC71 BM devices afforded excellent PCE values (5.6% for P70 and 6.2% for P71), probably due to deeper HOMO and LUMO levels. The surface morphology of P70–71:PC71 BM blend film showed low roughness (Ra) averaging 1.86 nm and 2.30 nm, respectively, from AFM images. Using 2,1,3-benzoxadiazole as the electron-accepting unit, Coffin and coworkers [132] reported another BDTbased crystalline copolymer P72. The device performance was found to be strongly dependent upon film-casting solvent. Devices with the active layer cast from pristine CB exhibited an average PCE of only 1.4%. The device harvested significant enhancement in performance (6.1% PCE) when making the devices with 2% CN as processing additive [85]. However, devices processed with DIO as the additive led to poorly reproducible device performance. Fréchet and coworkers [133] reported P64 analogue polymers P73–75 by changing the alky substituent in TPD. All polymers were developed with more than three-fold higher Mn than P64. The alkyl chain exerted negligible effect on the Eg and energy level alignments of P73–75. Spin-coated from CB solution, P73–75:PCBM devices achieved a PCE of 2.8%, 3.9%, 6.4%, respectively at the optimized weight ratios. Clearly, the device performance increased with the decrease in the side chain length for P73–75. By adding 2% DIO as processing additive to control

Fig. 20. 2D GIXS patterns of blends of P73 (a), P74 (b), and P75 (c) with PCBM in the optimized condition spin-coated from chlorobenzene and P73 (d), P74 (e), and P75 (f) prepared from mixed-solvent chlorobenzene/1,8-diiodooctane. Adapted with permission from Ref. [133]. Copyright 2010 American Chemical Society.

the nanoscale morphology in the blend films, P73 and P74 devices were significantly improved to 4.0% and 5.7%, respectively. However, for P75, the addition of DIO only led to slight PCE improvement, attributed to the original high level of ordered film morphology even without DIO. The use of DIO is effective in promoting the packing of polymer backbones by preventing excessive crystallization of PCBM. The promotion of thin-film ordering in the polymer domains was proved by the increased intensity of the ␲-stacking peaks from GIXS images (Fig. 20). You and coworkers [134] reported another two BDT-containing polymers P76 (PBnDT-HTAZ) and P77 (PBnDT-FTAZ), with benzotriazole or its fluorinated counterpart as accepting unit. Like P61, fluorination of the accepting units in polymer backbone of P76 lowered the HOMO by 0.07 eV in P77. By spin-coating from 1,2,4trichlorobenzene (TCB) solution, P77:PCBM (1:2, w/w) device showed an extremely high efficiency of 7.1% with an active layer thickness of 250 nm. Under the same conditions, P76 showed a PCE of only 4.4%. The performance improvement can be explained with the fluorinated benzotriazole backbone in P77, which resulted in a 0.09 V increment in Voc and a great increase in FF from 55 to 73%. It is noted that P77:PCBM cells achieved over 6% efficiency even at an unprecedented thickness of 1 ␮m for the active layer (see Fig. 21). Cao and coworkers [135] reported another BDTbased polymer, P78, using naphtho[1,2-c:5,6-c]bis[1,2,5]thiadiazole (NT) as the accepting unit. Compared with BT in P67, NT has an enlarged planar aromatic structure with two fused 1,2,5-thiadiazole rings, facilitating the interchain packing of the resultant polymer and further enhancing the carrier mobility. In addition, the electron-withdrawing capability of NT is slightly stronger than that of BT, which lowered the Eg of P67 by 0.17 eV for P78, resulting in more efficient solar energy harvesting while still maintaining enough driving force for the charge separation between the polymer and PCBM. By spin-coating from ODCB solution and further annealing P78:PC71 BM (1:1, w/w) blend film

L. Bian et al. / Progress in Polymer Science 37 (2012) 1292–1331

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Fig. 21. (a) J–V curves for the highest-performing cells for each polymer. The fluorine atoms cause increases in every performance category. P77 overall performs 76% better than P76. (b) Incident photon to current efficiency and solid film absorption of each blend of polymer:PCBM. (c) Dependence of the FF and Jsc on the thickness of the active layer. (d) SEM of 1 ␮m active layer that showed 6% power conversion efficiency (scale bar: 1 ␮m). Adapted with permission from Ref. [134]. Copyright 2011 American Chemical Society.

(40 nm) at 130 ◦ C for 7 min, its PSC showed an excellent PCE of 6.0%. Hou et al. [136] reported a new BDT copolymer P79 using thiophene flanked TTZ as the accepting unit. Similarly to P52 [110], TTz-based conjugated polymers exhibited excellent hole mobility and excellent stability [137]. However, TTz-containing P79 presented almost one magnitude lower hole mobility than thienothiophene-based P55–P60. DCB appeared to be more favorable than CB in device processing of P79. When adding 3% DIO as processing additive, smoother films with less phase separation were observed, and thus P79:PC71 BM (1:3, w/w) device was optimized to a high PCE of 5.2%. The optimal weight ratio of P79:PC71 BM was found to be 1:2 or 1:3 (Fig. 22). The champion devices exhibited a Voc of 0.85 V, Jsc ∼10.3 mA/cm2 and FF ∼0.59.

in Chart 6), in which BT was employed as accepting fragment. With extremely low Eg ∼ 1.42 eV (Table 5), the phenyl substituted DTP based P81 exhibited stronger absorption in the 600–900 nm range. P81:PCBM (1:3, w/w) device showed a PCE of 2.8%, slightly high than that for P80 at their optimized blend film thicknesses (100 nm for P80, 90 nm for P81). The relatively low efficiency device of P80–81 was attributed to the extremely low Voc (0.4–0.54 V) of PC

6.6. Copolymers based on dithieno[3,2-b:2 ,3 -d]pyrrole With the successful exploration of tricyclic cyclopentadithiophene and dithienosilole rings as excellent donating blocks for DA polymer design, dithieno[3,2-b:2 ,3 d]pyrrole (DTP) has emerged as another appealing fused bithiophene member, due to its good planarity and stronger electron-donating ability with nitrogen atom. The homopolymer of DTP was first prepared by Rasmussen and coworkers [138] via oxidation polymerization to give a low Eg (1.7 eV) polymer with high red-emitting quantum efficiency. Wang and coworkers [139] reported two nearinfrared absorbing copolymers P80–81 (P80–100 structure

Fig. 22. J–V curves of the PSCs based on P79/PC71 BM with different donor/acceptor ratios (1:1, 1:2, 1:3, and 1:4) under illumination of AM 1.5G, 100 mW/cm2 . Adapted with permission from Ref. [136]. Copyright 2011 American Chemical Society.

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Chart 6. Fused thiophene based and ladder type narrow bandgap polymers.

devices, caused by the high-lying HOMO levels (−4.80 eV) of P80–81. By flanking thiophene on BT in P81, Hashimoto and coworkers [72] reported P82 with extremely low Mn (1.76 kDa). P82 exhibited an Eg of 1.46 eV, the lowest value observed among all DTBT-containing DA copolymers with donating segments, such as fluorene [36], silafluorene [49], carbazole [37,59], and CPDT [140]. P82 device showed a PCE of only 2.2% and an extremely low Voc (0.52 V), which is much lower than that of the corresponding devices from

DTBT-based polyfluorene (1 V), polysilafluorene (0.9 V) or polycarbazole (0.9 V). Hashimoto and coworkers [141] further reported a DTP-containing polymer P83 using 3,6-dithien-2-yl2,5-dialkylpyrrolo[3,4-c]pyrrole-1,4-dione (DTDPP) as the acceptor unit [84], With Eg extremely lowered down to 1.13 eV, P83 exhibited broad absorption spanning from 500 to 1100 nm. P83 also showed a high hole mobility of 0.05 cm2 V−1 s−1 , which is 1–2 orders of magnitude higher than that of other DA polymers. Spin-coated blend films

Table 5 The optical, electrochemical, hole mobility, and PSC characteristics of P73–93. Mn (kDa) (Mw /Mn )

abs max (nm)

Eg (eV)

h (cm2 V−1 s−1 )

HOMO/LUMO (eV)

Polymer:PCBM

Jsc (mA/cm2 )

Voc (V)

FF

PCE (%)c

Ref.

P80 P81 P82 P83 P84 P85 P86 P87 P88 P89 P90 P91 P92 P93 P94 P95 P96 P97 P98 P99 P100

15.6 (1.4) 17.9 (1.6) 1.76 (3.4) 18.9 (2) 5.2 (1.9) 6.8 (2.1) 8 (2.1) 9.6 (2.3) 28.7 (3.6) 25.7 (3.9) 12.9 (2.3) 22.7 (2.1) 17 (16) 22.2 (1.7) 68.2 (2.4) 82.4 (2.8) 38.1 (1.6) 14.1 (2.2) 21 (2.8) 16 (2.2) 18.1 (2.6)

764 764 697 790 547 548 580 650 675 681 647 928 520 590 581 644 645 599 651 720 664

1.42 1.43 1.46 1.13 – – 1.53 1.56 1.51 1.61 1.7 1.34 1.76 1.7 1.81 1.67 1.64 1.66 1.61 1.48 1.69

– – – 0.05 – – – – – a 1.7 × 10−5 a 7.2 × 10−6 – 7.0 × 10−4 3.4 × 10−3 4.6 × 10−3 2.4 × 10−2 – – 5.6 × 10−2 1.5 × 10−2 a 1.7 × 10−4

−4.86/−3.07 −4.81/−3.08 −5.00/−3.43 −4.90/−3.63 −5.30/−3.21 −5.34/−3.26 −5.36/−3.42 −5.50/−3.44 −5.47/−3.44 −5.34/−3.29 −5.46/−3.28 −5.21/−3.63 −5.46/−3.56 −5.43/−3.66 −5.33/−3.52 −5.28/−3.61 −5.31/– −5.31/– −5.23/−3.53 −5.26/−3.69 −5.04/−3.28

1:3 1:3 1:1 1:2b 1:1 1:1 1:1 1:1 1:1 1:0.8 1:1.2 1:2b 1:3b 1:3b 1:3b 1:3b 1:3b 1:3b 1:3b 1:3b 1:1

11.1 11.9 9.47 14.87 10.93 10.67 14.2 13.49 12.78 14.20 11.38 10.1 8.7 10.1 10.9 11.2 10.3 9.2 11.6 10.1 15

0.43 0.54 0.52 0.38 0.59 0.69 0.71 0.75 0.85 0.67 0.83 0.68 0.84 0.8 0.87 0.87 0.74 0.74 0.83 0.74 0.66

0.43 0.44 0.44 0.48 0.46 0.46 0.62 0.55 0.58 0.54 0.46 0.63 0.53 0.53 0.6 0.64 0.6 0.52 0.63 0.43 0.58

2.1 2.8 2.2 2.7 3 3.4 6.2 5.6 6.3 5.1 4.3 4.3 3.9 4.3 5.7 6.2 4.6 3.5 6.1 3.2 5.6

[139] [139] [72] [141] [143] [143] [144] [144] [144] [146] [146] [147] [150] [150] [152] [152] [153] [153] [154] [154] [155]

L. Bian et al. / Progress in Polymer Science 37 (2012) 1292–1331

Polymer

Definitions for all parameters are the same as those in Table 1. a h measured by SCLC. b Polymer:PC71 BM, in weight ratio, w/w. c PCE achieved at illumination of standard simulated solar light (AM 1.5G, 100 mW/cm2 ).

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Fig. 23. (a) Weak donor–strong acceptor concept and energy levels. (b) Calculated energy-conversion efficiency of P3HT and “ideal” polymer, assuming FF and IPCE at 65%. Adapted with permission from Ref. [142]. Copyright 2011 American Chemical Society (a); adapted with permission from Ref. [10]. Copyright 2006 Wiley (b).

from CHCl3 and ODCB (4:1, v/v) mixture solutions, P83 device showed a PCE of 2.3% when blending with PCBM and 2.7% with PC71 BM. Similarly, such low performance is mainly attributed to the lower Voc values, resulting from its high-lying HOMO level (−4.90 eV). 6.7. Copolymers based on naphtho[2,1-b:3,4-b ]-dithiophene and dithieno[3,2-f:2 ,3 -h]quinoxaline Recently, an exciting design strategy for DA polymers was proposed by You and coworkers [142], in which “weak donor” and “strong acceptor” units alternated to achieve ideal polymers (Fig. 23a) with both low HOMO energy level and narrow Eg , hence both high Voc and Jsc could be achieved for PV devices. The “weak donor” should help maintain a low HOMO energy level, while a “strong acceptor” should reduce the bandgap via ICT behavior. Assuming a FF value of 0.65, an external quantum efficiency of 65%, and an optimal morphology, one can estimate the overall PCE from the optical bandgap and the LUMO/HOMO of donor polymers in a polymer:PCBM BHJ solar cell [142] (Fig. 23b). Under the guidance of this design concept, a library of high-efficiency photovoltaic polymers P84–90 have been developed, in which DTBT and thiadiazolo[3,4c]pyridine were employed as strong acceptors while naphtho[2,1-b:3,4-b ]-dithiophene (NDT), dithieno[3,2f:2 ,3 -h]quinoxaline (QDT) and BDT were employed as weak donors [143–146]. For NDT-containing P84–85, two 4-(2-ethylhexyl)thiophene flanked BT was used to reduce the steric hindrance of the polymer backbone and hence to achieve near identical Eg and energy levels [143]. It was believed that Eg and energy level alignment of conjugated polymers were primarily determined by the molecular structure of backbone, while the solubilizing alkyl chains should have a negligible impact on these properties, and subsequently on Jsc and Voc for devices. However, Yu and coworkers [123] showed the variations of length and shape of alkyl chains in BDT had significant influence on device performance: long and branched

side-chains weaken the intermolecular interaction, benefiting for the improvement of Voc , while short and straight side-chains promote the intermolecular interactions, rendering a large Jsc value. The balance between Voc and Jsc can be adjusted through adopting branched side-chains with suitable length. P84–85:PCBM (1:1, w/w) devices showed best PCEs of 3.0% and 3.4%, respectively. By flanking thiadiazolo[3,4-c]pyridine (PyT) with two alkylated thienyls, You and coworkers [144] developed a new soluble and stronger acceptor (DTPyT) to construct P86–88 with weak donors including NDT, QDT or BDT. P86–88 exhibited identical LUMO levels and agreed with the general discovery that the LUMO of DA polymers was primarily located in the acceptor unit [57,142,145]. The incorporation of the stronger acceptor DTPyT in P86–88 lowered the LUMO level by approximately 0.2 eV compared with their DTBT counterparts [128,145]. Hence, P86–88 exhibited an Eg of 1.53, 1.56 and 1.51 eV, respectively, which is 0.09–0.19 eV lower than that of corresponding BT counterpart polymers [128,145]. P86–88:PCBM (1:1, w/w) devices achieved remarkable PCE values (ca 6.2% for P86, 5.6% for P87, and 6.3% for P88). P86 device showed the highest current of 14.16 mA/cm2 , which is one of the highest Jsc values obtained for BHJ devices consisting of polymer donor and PCBM acceptor [122,123]. The low-lying HOMO (ca −5.4 eV) and high Jsc values suggested the potential of P86 for further device optimization, which includes using PC71 BM as the acceptor, the addition of processing additives and thermal annealing of blend films (see Fig. 24). By alternating DTBT acceptor with NDT or QDT weak donors, P89 and P90 were developed with an Eg of 1.61 eV and 1.70 eV, respectively [146]. The introduction of electron-withdrawing nitrogen atoms in QDT was demonstrated to lower the HOMO level of the resulting polymer while exerted negligible impact on its LUMO level, and thus widened the Eg when compared with NDT. P89–90:PCBM devices achieved a PCE of 5.1% and 4.3%, respectively. The better device performance of P89:PCBM blends was attributed to finer phase-separation, indicating stronger interchain ␲–␲ interactions in P89 which correlated to

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Fig. 24. (a) J–V curves of polymer/PCBM based solar cell devices under AM 1.5G illumination (100 mW/cm2 ). (b) EQE curves of polymer/PCBM-based solar cell devices. Adapted with permission from Ref. [144]. Copyright 2010 Wiley.

high Jsc . However, no clear phase separation was observed for P90:PCBM films. The employment of even weaker donors and stronger acceptors via innovative structural modification is thus recommended for future research in order to concurrently achieve both higher Voc and higher Jsc values for PV cells. Yang and coworkers [147] reported another BDTcontaining polymer, P91, by adopting DPP as the accepting unit. The efficiency of a P91:PC71 BM (1:2, w/w) device improved from 4.3% to 4.5% after annealing the blend film at 110 ◦ C for 30 min. P91 was shown to be a promising candidate for highly efficient PSCs owing to its relatively low Eg and suitable HOMO level. 6.8. Copolymers based on ladder-type coplanar thiophene-phenylene-thiophene Ladder-type coplanar thiophene-phenylene-thiophene (TPT) structure tends to lead to strong molecular ␲–␲ interactions [148] and thus its polymer exhibit a remarkable hole mobility, as high as of 10−3 cm2 V−1 s−1 [149]. Ting and coworkers [150] reported TPT-containing P92 and P93, in which accepting DTBT was employed. With an Eg ∼1.7 eV, P92–93 exhibited a wide absorption band of 350–700 nm and 350–730 nm, respectively. Spin-coated from DCB solution, P92–93:PCBM (1:3, w/w) device showed moderate PCE values of 2.0% and 2.5%, respectively, much lower than that (3.9% PCE) of a P3HT:PCBM (1:1, w/w) device under the

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same conditions. Using PC71 BM as the acceptor, the performance of devices was improved to (3.9% for P92, 4.3% for P93), better than that of a P3HT device (3.4%) under the same conditions [7]. Indacenodithiophene (IDT) is also an attractive fused ring structure, which can enhance the coplanarity of polymer backbone with reduced energetic disorder. A high and stable FET hole mobility (∼1 cm2 V−1 s−1 ) was found to be another feature of IDT-containing polymers [151]. Jen and coworkers [152] reported IDT-containing DA copolymers P94–95 with accepting quinoxaline. With two phenyl rings connected by a single bond between the ortho positions in P95, the planarity of quinoxaline was significantly increased to facilitate both intermolecular packing and charge transport. Moreover, the extended ␲-conjugation of the fused qunioxaline functioned as a stronger electronacceptor, which lowered the Eg of P95 by 0.14 eV compared with P94. P95 device showed better performance (6.2% PCE) than P94 (5.7%), in which the blend films were made from ODCB solutions and thermally annealed at 110 ◦ C for 10 min. The high performance of P95 devices was attributed to its high mobility in smooth morphology of blend films. A novel carbazole-based coplanar ␲-conjugated system, carbazole-dicyclopentathiophene (CDCT), was developed by covalently connecting the 3-positon of two outer thiophenes with the 3,6-position of central carbazole cores. Hsu and coworkers [153] reported four CDCT-containing copolymers using different accepting units including BT (P96), DTBT (P97), DPP and quinoxaline. A comparison of the differences of absorption maximum and Eg of the resulted polymers revealed the order of acceptor strength is DPP > BT > QX. The best-performing P96 device exhibited an impressive PCE of 4.6%, while a PCE of only 3.5% was achieved for P97. The better device performance of P96 was attributed to its high hole mobility (ca.1.2 × 10−3 cm2 V−1 s−1 ), leading to high Jsc and FF values. Jen and coworkers [154] further reported two IDTcontaining copolymers P98 and P99 using dithienobenzoquinoxaline (M1) and dithienobenzopyridopyrazine (M2) as electron-deficient acceptors. The large fused rings present two advantages over other quinoxalines: (i) the fused dithiophene ring decreases the steric hindrance by forming a planar structure to improve the intermolecular stacking of polymers; (ii) the fused qunioxaline with extended ␲-conjugation acts as a strong electron-deficient acceptor. Since M2 exhibited even stronger electronaccepting ability than M1, P99 achieved 0.16 eV deeper LUMO level and thus 0.13 eV lower Eg than P98. Spincoated from ODCB solution, P98 device showed a high PCE of 6.1%, almost double of that for P99. The lower PCE of P99 may be attributed to the non-irradiative recombination of excitions at the interface, leading to lower Voc and Jsc values since the pyridine may play a role as an electron trap compared with benzene in P98 [57]. Tetrathienoanthracene has also emerged as a promising donating unit in developing narrow bandgap polymers for high-efficiency PV cells. With extended ␲-conjugation to enhance ␲–␲ stacking between polymer backbones, tetrathienoanthracene-based polymers can promote more

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Chart 7. HOMO–LUMO levels based on fluorene polymers.

Chart 8. HOMO–LUMO levels based on the cabarzole polymers.

efficient hole transport in the active layer to improve the performance of their PSCs. Yu and coworkers [155] reported the first example of tetrathienoanthracene-based polymer P100, in which ester substituted thieno[3,4b]thiophene accepting unit was used. Spin-coated from 2% DIO contained CHCl3 solution, P100 device showed a high PCE, up to 5.6%, with a Jsc of 15.0 mA/cm2 , Voc of 0.66 eV, and FF of 58%. Using other co-solvent systems like CB/DIO and ODCB/DIO, the as-fabricated devices showed only moderate Jsc and FF values, resulting in PCE ∼3%. TEM images showed that the blend films cast by CHCl3 /DIO solutions possessed relatively small domain size, which increased the donor–acceptor interfacial areas and enhanced the PV performance. 7. Data analysis of polymer and discussion 7.1. HOMO–LUMO levels of polymers Ideal polymers for use in bulk BHJ solar cells should have a suitable alignment of HOMO and LUMO levels. For airstability concerns, polymers should have a HOMO energy level below the air oxidation threshold (−5.27 eV). For DA alternating polymers, both HOMO and LUMO energy levels are mainly determined by the polymer backbones (including DA structure, conjugation length and configuration, etc.): the HOMO energy level is generally fixed by the donating moiety, whereas the LUMO energy level is directly related to the nature of electron-accepting unit. As with the HOMO–LUMO level alignment for PFs shown in Chart 7, P1–P17 exhibited a HOMO ∼−5.5 eV

and a LUMO ∼−3.6 eV, with energy level alignment varied with backbone structure. A comparison study of P3 and P5 indicates DTBT could maintain a deep HOMO level while increasing the LUMO more effectively than DTBT. Comparisons of P3 with P6–P7 and P16 with P17 indicates that DTBT is more effective than TQT in narrowing polymer Eg , which lowers the LUMO level deeper than it does for the HOMO. For 2-D conjugated polymers P10–P11 and P12–P13, each pair of polymers present similar HOMO levels due to the same donating polymer backbone but different LUMO levels due to different accepting side chains. As the HOMO–LUMO level alignment for PCs shown in Chart 8, P19–P33 presented similar energy levels to their PF analogues. HOMO energy levels of polymers are mainly determined by polymer backbone, whereas LUMO energy levels seem to be strongly dependent upon the nature of electron-accepting units. Stronger electron-deficient units generally result in lower LUMO levels. A close look at the energy level alignment for 2-D copolymers P24–P26 and P27–P29 indicates that they shared similar HOMO levels, but different LUMO levels due to different pendant accepting side chains. Compared with indolo[3,2-b]carbazole containing P31–P33, P31 exhibits most low-lying HOMO energy level and hence the highest Voc in PV cell. The HOMO–LUMO level alignment of PTs is shown in Chart 9. DPP-containing P37–P39 showed significant low Eg ∼1.35 eV, while P35–36 showed higher HOMO levels due to the presence of a furan donating unit instead of thiophene. TPD-containing P40, P43, P48 and P49 presented much low-lying HOMO energy levels than BT-containing P41 and P46, indicating that TPD moiety is

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Chart 9. HOMO–LUMO levels based on the thiophene polymers.

less electron-withdrawing than BT. The much lifted HOMO level of P42 can be explained with the insertion of strong accepting bithiazole moiety in backbone. A comparison of dithieno-[3,2-b:2 ,3 -d]silole-containing P46–P48 and P51–52, P48 shows the “ideal” HOMO and LUMO alignment with TPD accepting unit and great potentials in PSC applications (7.3% PCE) [105]. Although P54 has similar Eg as P48, its LUMO level is too low-lying to facilitate exciton splitting and charge dissociation. As the HOMO–LUMO level alignment for PBDT shown in Chart 10, most PBDT exhibited relatively lower Eg than PF and PC analogues. Comparing P56 with P57, it is noted that octyloxy side chain lowered backbone’s HOMO level from −4.94 to −5.04 eV. Comparisons of P58 with P59, P61 with P62, and P76 with P77 indicate that the introduction of electron-withdrawing fluorine in polymer backbone can effectively lowered their HOMO levels. P64 presented similar energy level alignment as P73–P75, due to the same TPD-containing backbone. A comparison of P67 with P78 shows that stronger accepting BT fused-rings (in P78) narrowed Eg more effectively than BT, since the raising of HOMO and lowering of LUMO level occurred simultaneously in P67.

polythiophenes P80–P100 present a Fused HOMO–LUMO level alignment as shown in Chart 11. P80–81 exhibited similar HOMO levels (∼4.86 eV) as P3HT but a lower LUMO level [156], leading to narrower Eg . P86–88 maintain low HOMO levels around the ideal HOMO level (−5.4 eV) due to the presence of weak donors: NDT, QDT and BnDT. It is worth noting that substitution naphthalene (P90) with more electron-deficient quinoxaline as the donating moiety (P91) weakened the electron-donating ability of the donor, leading to the observed lower HOMO level of −5.46 eV (P90) compared with −5.34 eV (P91). Comparing P89 with 90, P92 with 93, P98 with P99, it is found that thiophene–benzene fused donating units maintain very low-lying HOMO, hence ideal Eg values were achieved by successfully lowing-down the LUMO levels with strong accepting units. 7.2. Correlation of HOMO levels of polymers with device Voc For devices, deep low-lying HOMO levels of polymer donors would help to achieve higher Voc values [126,145], since Voc is closely related to the energy difference between

Chart 10. HOMO–LUMO levels based on the thiophene polymers.

Chart 11. HOMO–LUMO levels based on the thiophene polymers.

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the HOMO level of the polymer donor and the LUMO level of the acceptor [10] (Eq. (3)): A minimum energy difference of 0.3 eV between LUMO level of polymer donor and PCBM-like acceptor is estimated to be required to facilitate exciton splitting and charge dissociation. Taking PCBM as an example, the ideal lowest possible LUMO levels of polymer donors would be near 3.9 eV due to a LUMO level of 4.2 eV possessed by PCBM. Generally, the lower the HOMO level of the donor, the higher the Voc of its device. However, a deeply low HOMO level of polymer donor would lead to an increased Eg and less efficient light absorption. An optimal Eg of 1.5 eV is proposed to be a compromise between these two contradictory factors. This positions the HOMO level of the “ideal” polymers around −5.4 eV. A statistical plot of the cell Voc against the HOMO lever of donor polymer is presented in Fig. 25, shown with a fitting line according to Eq. (3). For simplification, polymers P1–100 are classified as polyfluorenes (PF), polycarbazoles (PC), polythiophenes (PT), BDT-containing polymers (PBDT) and fused thiophene-containing polymers (PTFT). It is obvious that the Voc of most PSCs do not agree well with Eq. (3). Significant scattering from linear correlation of Voc with energetics of the DA pair indicates that factors (e.g., BHJ film morphology) other than the energy difference between donor HOMO and acceptor LUMO levels may play important role in these DA polymers. A number of recent reports have suggested that besides the energetics of the DA pair the generated photovoltage is also controlled by the saturation dark current [157–160], which would be intimately connected to bulk morphology, degree of microstructural order, and intermolecular interactions in active layer. In typical small-molecule:C60 bilayer BHJ solar cells, Thompson and coworkers [157] reported the inverse dependence of device Voc on the saturation dark-current density. A lower device Voc was observed for donor molecules showing aggregation and polycrystallization in the thin-films, in which higher Jsc occurred. In contrast, the less ordered, more amorphous donor materials with structures that hindered the

intermolecular ␲–␲ interactions produced lower Jsc and therefore higher Voc values. Such correlation of Voc and morphology-dependant Jsc was further reinforced in polymer:C60 bilayer solar cells by Frisbie and coworkers [158], in which regioregular P3HT, regiorandom P3HT, and poly(3-hexyl-2,5-thienylene vinylene) (P3HTV) were used as acceptors and the active layer morphology was enhanced via post-position thermal annealing. As the donor and acceptor layers intermixed, coincident improvement in Voc and reduction of Jsc were observed until C60 fully penetrated the polymer phase and reached the ITO/PEDOT:PSS anode, thus causing a reduction in Voc . Under these conditions, introduction of a pentacene electron-blocking layer at the interface between ITO/PEDOT:PSS anode and photoactive heterojunction leads to a drastic increase in Voc (150%) and shunt resistance (×103 ), with reduced Jsc (10−3 ), and improved device performance (Fig. 26) [11,158]. Vandewal et al. [159] further underlined the importance of the weak ground-state interactions between polymer donor and fullerene since Voc was determined by the formation of these states. They validated the use of the detailed balance approach [161] in estimating Voc values of polymer:fullerene BHJ solar cells, in which PCBM and PC71 BM were used as acceptors and MDMO-PPV, P3HT, PCPDTBT and poly[2,7-(9-di-octylfluorene)-alt5,5-(4 ,7 -di-2-thienyl-2 ,1 ,3 -benzothiadiazole)] (APFO3) and poly[1,4-(2,5-dioctyloxybenzene)-alt-5,5-(5 ,8 -di-2thienyl-2 ,3 -di-(3 -butoxyphenyl)quinoxaline)] (LBPP5) were used as the polymer donors. The results showed that electroluminescence and photovoltaic external quantum efficiency spectra in the low-energy, charge-transfer region were connected, and that the corresponding spectra were related to Voc . Given the logarithmic dependence of Voc on Jsc , it is concluded that Voc depends linearly on the spectral position of the charge-transfer band [11,159]. A more recent study by Ohkita and coworkers [160] further revealed that the Voc is mainly dependent on the saturation current, i.e., Voc increased with the decrease in the saturation current density J0 for P3HT:fullerene BHJ device. The activation energy and the pre-exponential factor J00 for Jsc were evaluated from the temperature dependence of Jsc . They proposed that a large electronic coupling in the charge separation between fullerene/fullerene and a small electronic coupling in the charge recombination between polymer/fullerne were the key to improve Voc effectively without loss of the charge generation efficiency, in addition to large energetics of the DA pair [160,161]. 7.3. Correlation of bandgap of polymers with device performance

Fig. 25. Plot of Voc versus (1/e)(|EDonor HOMO| − |EPCBM LUMO|) − 0.3V for P1–100.

A critical goal of DA polymer design is to achieve extended absorption to match the solar terrestrial radiation. The correlation of the cell PCE with the Eg values of polymer P1–100 is shown in a statistical plot (Fig. 27). A reference vertical line at Eg of 1.97 eV (Eg of P3HT) [18] and a horizontal line at PCE of 3% (P3HT cell) are inserted to divide the plot into four zones. As clearly shown, most polymers with an Eg below 2 eV have achieved high-efficiency cells

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Fig. 26. Pentacene/P3HTV/C60 (40 nm) bilayer devices annealed at 170 ◦ C. The pentacene layer, evaporated at the anode between PEDOT:PSS and P3HTV, was varied from 0 to 70 A˚ thick. P3HTV was spin-coated from a 15 mg/mL solution in 1,2-dichlorobenzene for all samples, which gave a film thickness of 23 nm for the device with no pentacene. (a) Schematic of device structure. (b) Semilog plot in the dark. Symbols are experimental data; the dashed lines represent fits to the forward current using Eq. (1). (c) Energy level diagram. (d) Under 100 mW/cm2 light intensity. Adapted with permission from Ref. [158]. Copyright 2009 American Chemical Society.

(PCE > 3%). And PT, PDBT and PTFT present higher possibilities than PF and PC in achieving high-performance cells. It should be noted that low Eg value is not the only prerequisite for polymers in realizing high PCE devices, since a large number of polymers with an Eg < 1.97 eV exhibited PCE values far below 3%. This correlation study shows that factors besides polymer Eg also contribute for cell

Fig. 27. Plot of cell performance PCE versus bandgap for P1–100.

performance, which may include polymer’s Mn , type of acceptor and its content, interchain ␲–␲ stacking, and morphology of BHJ films [11]. 7.4. Molecular weight issue In the design of molecular structures of polymer backbones, the molecular weight (and its distribution) of the DA polymers is another important issue for their PSC applications [11,89]. As known, the performance of BHJ devices is strongly dependent upon the Voc and Jsc values. While higher Jsc values in BHJ devices require relatively higher carrier mobilities and exciton diffusion lengths for polymer donors [11]. For the case of the well-documented P3HT, somewhat conflicting experimental results have been reported on the effect of polymer molecular weight on its hole mobilites in FET and photovoltaic performance in PV cells [162–164]. Brabec and coworkers [162] found the performance of P3HT:PCBM based BHJ devices increased dramatically with the increment of the molecular weight of P3HT (Fig. 28). The improvement of device performance with polymer Mn was also observed for before-mentioned DA polymers P3 (Mn = 21 kDa) [39], P4 (Mn = 20 kDa) [39], P7 (Mn = 19 kDa) [44], P14 (Mn = 79 kDa) [50] and P37 (Mn = 54 kDa) [89]. The enhancement in performance is attributed to more light absorption and efficient hole transport along the backbones.

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Fig. 28. (a) Photocurrent spectra (external quantum efficiency) of bulk heterojunction solar cells with different P3HT fractions blended into PCBM. (b) Typical J/V curves of bulk heterojunction solar cells comprising P3HT fractions with different molar mass distributions and under light in a linear) plot. Adapted with permission from Ref. [162]. Copyright 2005 American Chemical Society.

In contrast, Bradley and coworkers [164] reported that the hole mobilities in thin films of higher regioregular P3HT remained constant for polymers with Mn ranging from 13 kDa to 18 kDa and decreased by 1 magnitude when Mn increased to 34–124 kDa. A subsequent drop in performance of P3HT:PCBM BHJ devices was observed due to the dramatic changes in blend film morphology with higher Mn polymers. It is thus postulated that higher Mn fractions possessed a higher degree of entanglement, hindering intrachain carrier transport by increasing density of traps, while also impeding interchain hopping via diminished ␲–␲ overlap between the backbones. Hence, more careful consideration of both the positive and negative effects of molecular weight on device performance is a must before selecting suitable processing strategies for device optimization. 7.5. Device optimization As discussed by Fréchet and coworkers [3], the ideal high-performance BHJ solar cell should contain an active layer with a bicontinuous composite of donor and acceptor with a maximum interfacial area for exciton dissociation and a mean domain size commensurate with the exciton diffusion length (5–10 nm). The two components should phase-segregate on a suitable length scale to allow maximum ordering within each phase and thus effective charge transport in continuous pathways to the electrodes to minimize the recombination of free charges. Control of the morphology of the active layer, which is also crucial for BHJ solar cells [165,166], depends on the interplay between a number of intrinsic and extrinsic variables. The intrinsic properties are those that are inherent to the polymer and the fullerene, as well as the fundamental interaction parameters between the two components [3]. These include the crystallinity of the two materials, as well as their relative miscibility. How to take most advantage of these intrinsic variables of both donor and acceptor materials is the major concern for device fabrication in order to optimize the efficiencies of BHJ devices. The extrinsic factors include all the external

influences associated with device fabrication, such as solvent choice [52,84], weight ratio of the polymer:fullerene blend [35,60,61,66,96,136], the use of high-boiling-point processing additives [85,96,99,107,111,123,133,167], solvent evaporation rate, the use of cathode interfacial layer [62,66,106], control of the film thickness [92,134], as well as thermal and/or solvent annealing [166]. The application of these device optimization strategies has been well documented for P3HT and MDMO-PPV devices [1,3,12,25,35,41]. It has become increasingly apparent that a balance among the competing effects of solution processing, miscibility with the fullerene component, and solid-state packing need to be established. Generally the polymer/PCBM blends are dissolved in chloroform, toluene, DCB, CB, or a mixture these, with the addition of lowvapor-pressure small-molecule additives, such as 1,8octanedithiol, 1,8-diodioctane, or CN. Processing additives produce remarkable effects on device performance via morphology control, in which particularly ordered and homogeneous interpenetrating networks with nanoscale phase-separated domains are promoted to match the modest exciton diffusion length (5–20 nm) [85,96,105]. Because of the large difference in vapor pressures between solvents and co-solvents or additives, it is believed that the final morphology is essentially dominated by the evaporation of the lower-vapor-pressure co-solvents or additives. This process allows the polymer chains to slowly crystallize and continue arranging after the higher-vapor-pressure solvents evaporation. Care should be taken in the selection of processing additives. For example, DIO has been shown to promote polymer aggregation and thus increase the domain sizes within the films [96], while CN has been shown to prevent polymer aggregation and decrease phase separation [132,165]. The results are usually highly dependent on both the system and the processing method selected. In the context of BHJ optimization, the determination of the ideal polymer:PCBM blend ratios has been a matter of trial and error, with 1:1 and 1:(3–4) ratios used most frequently [3]. PCBM loading is believed to control the

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formation of adequate electron and hole percolation paths, leading to efficient charge collection to electrodes, in addition to efficient dissociation of the excited states. The recent report by McGehee and coworkers [91] exploring BHJs of PBTTT with PCBMs of various sizes suggests that fullerenes can intercalate between the side chains of ␲-conjugated polymers. Many empirical observations have shown that the presence of PCBM can improve the molecular packing of polymer donors in blends, and ultimately device operation. However, larger PCBM loadings may result in decreased optical absorption of the resulting active layer at long wavelengths for which the solar photon flux is the most intense. This alone may ultimately constitute the limiting factor with respect to device performance. 8. Conclusions and perspective With the extensive research and successful DA polymers that we have outlined in this review, the guidelines to pursuing high-performance BHJ solar cells have become evident. The requirements of specific intrinsic properties necessary for an ideal donor material include (1) sufficient solubility to guarantee solution-processability as well as miscibility with n-type materials, (2) a low optical bandgap (1.2–1.9 eV) for broad and strong absorption to capture more solar energy, (3) a HOMO energy level (ca −4.8 to −5.4 eV) to ensure efficient charge separation while maximizing Voc , and (4) high hole mobility to allow adequate charge transport, which in turn allows a thicker blend films beneficial for increased light harvesting as well as reduced charge recombination and series resistance. In combination with a fixed electron acceptor, specific characteristics between a DA pair are required, which include (1) appropriate energy level matching between the HOMO of polymer donor and LUMO of acceptor to ensure a large Voc and a downhill energy offset for exciton dissociation and (2) formation of interpenetrating networks with the optimum morphology to create distinctly bicontinuous pathways for the transport of the free charge carriers. Structural analysis of the current successful low bandgap polymers reveals that “push–pull” motif is essential for DA polymer design. Research efforts are required to develop newly designed donating segments composed of multicyclic aromatic rings with enforced planarity for the combination with suitable accepting units for rationally fine-tuning of the energy level alignment of the resulting copolymers [142,166]. A recent trend in DA polymer design is directed toward hybridizing different electronrich aromatic units into mutually fused structures, with the anticipation of combining the individual intrinsic advantages on one hand and introducing strong electronwithdrawing atoms or moieties to existing accepting units to promote both efficient ICT transition and electron transfer along conjugated backbone on the other. The choice of side chains with appropriate lengths and structures is also critical in terms of polymer design and device engineering. Branched solubilizing groups introduced into the polymers to ensure solubility sometimes can attenuate intermolecular interactions in the solid state [37]. In particular, the branched and bulky solubilizing side chains appended to the ␲-conjugated backbones may substantially affect ␲–␲

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stacking distances and density of ␲–␲ stacked backbones. In addition, the density and spatial configuration of side chains along polymer backbones are another concern when selecting suitable fullerene acceptors. As one benefit of DA polymer design, enhanced absorption abilities via bandgap engineering is only a general requirement for high-efficiency PSCs, since a large number of newly produced low Eg polymers with suitable energy level alignments still exhibited inferior device performance compared to that of P3HT. Morphology control in the polymer:fullerene blend films is the key issue for BHJ solar cells. Without optimal control of the morphology on the nanoscale level, it becomes difficult to translate the microscopic intrinsic properties of the conjugated polymer into macroscopic device performance. It should be emphasized that the power conversion efficiency of solar cell is more a device parameter than an intrinsic material parameter. This is because too many factors can affect the performance, i.e., the blending ratio with the acceptor, casting solvent, processing additive, thermal annealing of BHJ film and insertion of holetransporting layer, device structure [23], and morphology control of BHJ film all play important role in device optimization. High-efficiency photovoltaic devices are the combination of research pursuit of not only the material development, but also require judicious and careful device optimization. There are still lots of challenges in translating highperformance laboratory-scale photovoltaic devices into large-area devices via solution-processable roll-to-roll or printing techniques for practical applications [168]. It is, therefore, essential that all researchers and engineers involved in the development and fabrication of organic solar cells effectively collaborate with each other to elucidate the fundamental structure–property-processing relationships for high-performance devices, and push them for engineering scale-up in the end. With the fast development of polymer solar cells, a great number of research papers and reviews have recently been reported, with the topics covering materials design [2,3,7,14,17–21,56,85,166,169], morphology control [1,11,25,165], device physics [3,10,16,81,84,170], device architectures [3,11,17,24,60,65,66,170], device stabilities [25,166], large-area solution processing [2,4], and economical aspects [15]. This review summarizing most of the high-performance photovoltaic DA polymers with novel donating and accepting building blocks should complement the breadth of existing literature in this exciting research field.

Acknowledgements The authors acknowledge the financial supports from the National Natural Science Foundation of China (Grant nos. 21074055 and 10804006), Doctoral Fund of Ministry of Education of China (no. 20103219120008), Foundation of Key Laboratory of Luminescence and Optical Information (Grant no. 2010LOI04), Beijing Natural Science Foundation (Grant no. 2122050) and NUST Funding (Grant no. 2010ZDJH04).

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