New conjugated alternating benzodithiophene-containing copolymers with different acceptor units: synthesis and photovoltaic application

Abstract

Two new alternating low band gap D–A copolymers with different acceptor structures of 4,8-bis-(5-bromothiophene-2-yl)-benzo[1,2,5]thiadiazole (P1) and 4,8-dithiophene-2-yl-benzo[1,2-c;4,5-c′]-bis-[1,2,5]thiadiazole (P2) and a common BDT donor segment have been synthesized under Stille reaction conditions and characterized. The polymers showed good solubility, broad absorption bands and optical band gaps of 1.62 eV and 1.16 eV for P1 and P2, respectively. Bulk heterojunction (BHJ) polymer solar cells based on P1 and P2 as electron donors and fullerene derivatives (PC₆₀BM and PC₇₀BM) as acceptor were fabricated and their photovoltaic response was investigated. The overall power conversion efficieny (PCE) achieved for BHJ solar cells based on P1:PC₆₀BM, P2:PC₆₀BM, P1:PC₇₀BM and P2:PC₇₀BM blends cast from THF solvent is about 2.17%, 0.80%, 3.45% and 1.19%, respectively. The higher PCE for the device based on P1 has been attributed to the high value of hole mobility for P1 as compared to P2 and a larger driving force i.e. LUMO–LUMO offset, for photo-induced charge transfer for P1:PCBM BHJ active layer. The PCE has been further increased up to 5.30% and 1.58% for P1:PC₇₀BM and P2:PC₇₀BM blends cast from DIO/THF solvent, which is attributed to the improved crystallinity and a more balanced charge transport in the device.

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1. Introduction

Polymer solar cells (PSCs) have attracted considerable attention as renewable energy sources due to their unique advantages such as low cost, low weight and the possibility of flexible large-area device fabrication.¹ So far the most successful PSCs are based on bulk heterojunction (BHJ) active layers which consist of a phase separated blend of organic electron donor (polymers or small molecules) and acceptor (mostly fullerene derivatives), sandwiched between two electrodes with different work functions.²

To date, poly(3-hexylthiophene) (P3HT) and [6,6]-phenyl-C₆₁-butyric acid methyl ester (PCBM) are the most representative donor and acceptor materials, respectively. The power conversion efficiency (PCE) of the PSC based on P3HT/PCBM has reached more than 4%.³ However, P3HT has a relatively large band gap (∼2.0 eV) and absorbs photons with wavelengths shorter than 650 nm, in addition, the relatively high location of the HOMO (−4.80 to −4.90 eV) leads to a lower open circuit voltage (∼0.6 V).⁴ Thus, the design and synthesis of new donor conjugated polymers with broad absorption bands and low HOMO energy level is one of the most important issues.⁵ In order to achieve high efficiency polymer solar cells, it is necessary to prepare conjugated polymers with the following desired features: good optical and mechanical properties, long term stability and processability and efficient charge transfer. For a BHJ active layer consisting of fullerene as electron acceptor and polymer as electron donor, the ideal polymer should have a band gap between 1.2 and 1.9 eV with a broad absorption spectrum to absorb as much sunlight as possible and to generate the high short circuit current, and the LUMO level should be in the range of 3.5–3.8 eV to create enough driving force to promote an efficient exciton dissociation at the D–A interface.

As is well known, the PCE of solar cells is proportional to the open circuit voltage (V_oc), short-circuit current (J_sc) and the fill factor (FF) of the PSC. Thus, the key issues that must be taken into account in the molecular design of conjugated polymers are: broad absorption spectra to improve the collection of sunlight with a high J_sc, suitable energy levels i.e. highest occupied molecular orbital (HOMO) and lowest unoccupied molecular orbital (LUMO) matching with the fullerene acceptor, low lying HOMO with maximum V_oc, high hole mobility, and appropriate compatibility with fullerene acceptor to form nanoscale bi-continuous interpenetrating networks.

Among the many possible directions, one feasible approach towards broadening the absorption spectrum and tuning the energy levels is to design alternating donor–acceptor copolymers, in which the electron rich (donor) monomers are copolymerized with electron deficient (acceptor) monomers. The incorporation of alternating electron donating units (push) and electron withdrawing units (pull) within the polymer backbone results in intra-molecular charge transfer (ICT) which lowers the band gap. Moreover, the HOMO and LUMO energy levels of D–A copolymers can also be independently controlled by their donor and acceptor components, respectively, which provides a means for narrowing the band gap of the copolymer with appropriate energy level alignments with fullerene derivative acceptors.⁶ To date, the efficiency of PSCs based on donor–acceptor (D–A) copolymers has reached ∼5–8%.⁷ Recently, PCEs more than 10% in single junction and tandem solar cells have been achieved by development of new materials and device architectures.⁸ Benzo[1,2-b:3,4-b′]dithiophene (BDT) is one of the most successful donor units in the construction of the D–A copolymers for high performance photovoltaic materials.⁹ The advantages of BDT-coupled planar structure are easy to purify BDT monomers, and a high hole mobility (for example, the hole mobility of the copolymer of BDT and thiophene reaches 0.25 cm² V⁻¹ s⁻¹).¹⁰

For these reasons, we opted to synthesize copolymers P1 and P2 derivatives of benzo[1,2-b:4,3-b′]dithiophene (BDT) as the donor and two new acceptor structures for the copolymerization with 4,8-bis-(5-bromothiophene-2-yl)-benzo[1,2,5]thiadiazole and 4,8-dithiophene-2-yl-benzo[1,2-c;4,5-c′]-bis-[1,2,5]thiadiazole, respectively. Acceptor structures are easily synthesized and have a relatively simple and flat structure, which may be useful for the delocalization of the electrons when they are included in the D–A copolymers. In addition, their relatively strong electron acceptor ability will lead to a decrease in the HOMO and LUMO levels of D–A copolymers, which will lead to a high open-circuit voltage (V_oc) in the PSC. Due to the presence of alkyl substituents the copolymers exhibit good solubility in common organic solvents. We have used these copolymers as electron donors along with fullerene derivatives, i.e. PC₆₀BM and PC₇₀BM, as electron acceptors for the fabrication of BHJ polymer solar cells. The PCE of the devices based on P1:PC₆₀BM, P2:PC₆₀BM, P1:PC₇₀BM and P2:PC₇₀BM is about 2.17%, 0.80%, 3.45% and 1.19%, respectively. The higher value of PCEs for solar cells based on P1 as electron donors in the BHJ active layer was attributed to the relatively higher value of hole mobility and larger driving force for photoinduced charge transfer at the D–A interface present in the BHJ active layer. The PCE has been further improved up to 5.30% and 1.58%, for P1:PC₇₀BM and P2:PC₇₀BM blends processed from DIO/THF solvent.

2. Experimental details

2.1 Instruments and characterization methods

¹H and ¹³C NMR spectra of the starting compounds and polymers were recorded on the spectrometer “Bruker Avance-400” with a working frequency of 400.13 and 100.62 MHz, respectively. IR spectra were recorded by a FT-IR spectrometer “Perkin-Elmer 1720-X”, TGA and DSC analysis was performed on “Perkin-Elmer TGA-7” and “Perkin-Elmer DSC-7” devices with heating rate of 20 deg min⁻¹.

The absorption spectra in the range 190–1100 nm were recorded on the spectrophotometer “Varian Cary 50.” The source of the exciting light was a xenon lamp L8253, which is part of the block radiator with optical fiber output radiation “Hamamatsu LC-4.” Cyclic voltammetry measurements were performed on a potentiostat-galvanostat μAUTOLAB Type III equipped with standard three-electrode scheme in an acetonitrile solution of 0.1 mol l⁻¹ tributylammonium hexafluorophosphate (n-Bu₄NPF₆) at a potential scan rate of 50 mV s⁻¹. Films of investigated polymers were deposited on a glass surface coated with ITO and then dried and used as working electrode. Ag/Ag⁺ and platinum were used as reference and counter electrodes, respectively.

2.2 Synthesis of copolymers P1 and P2

4,7-Di-(thiophene-2-yl)-benzo[1,2,5]thiadiazole (2). Synthesized by the method of ref. 11. Yield 2.78 g (71%). Mp = 122–124 °C. NMR ¹H (CDCl₃, 400 MHz, δ, ppm): 8.11 (d, J = 2.7 Hz, 2H); 7.86 (s, 2H); 7.45 (d, J = 4.1 Hz, 2H); 7.20 (t, J = 3.8 Hz, 2H); NMR ¹³C (CDCl₃, 100 MHz, δ, ppm): 152.52, 139.23, 127.90, 127.39, 126.70, 125.87, 125.67. Found, %: C 55.88; H 2.61; N 9.40. For C₁₄H₈N₂S₃ calculated, %: C 55.97; H 2.68; N 9.32.

4,7-Bis-(5-bromothiophene-2-yl)-benzo[1,2,5]thiadiazole (3). Synthesized by the method reported earlier.¹² Yield 2.2 g (66%). Mp = 252–253 °C. NMR ¹H (CDCl₃, 400 MHz, δ, ppm): 7.81 (d, J = 3.9 Hz, 2H); 7.80 (s, 2H); 7.15 (d, J = 4.2 Hz, 2H); NMR ¹³C (CDCl₃, 100 MHz, δ, ppm): 152.5, 140.9, 131.1, 127.7, 125.8, 125.4, 115.1. Found, %: C 36.88; H 1.29; Br 34.79; N 6.08. For C₁₄H₆Br₂N₂S₃ calculated, %: C 36.70; H 1.32; Br 34.88; N 6.11.

4,7-Dibromo-5,6-dinitrobenzo[1,2,5]thiadiazole (4). Synthesized by the method reported earlier.¹³ Yield 11.7 g (36%). Mp = 184–186 °C, lit. Mp = 198 °C.¹³ NMR ¹³C (CDCl₃, 100 MHz, δ, ppm): 151.24, 144.81, 110.18. Found, %: C 18.62; Br 41.75; N 14.38. For C₆Br₂N₄O₄S calculated, %: C 18.77; Br 41.62; N 14.59.

5,6-Dinitro-4,7-dithiophene-2-yl-benzo[1,2,5]thiadiazole (5). Synthesized by the method reported earlier.¹³ Yield 5.4 g (62%). Mp = 260–261 °C, lit. Mp = 259–260 °C.¹⁸ NMR ¹H (CDCl₃, 400 MHz, δ, ppm): 7.73 (d, J = 6.3 Hz, 2H); 7.51 (d, J = 4.9 Hz, 2H); 7.23 (t, J = 5.0 Hz, 2H); NMR ¹³C (CDCl₃, 100 MHz, δ, ppm): 179.43, 136.05, 132.48, 131.30, 130.84, 127.87, 124.11. Found, %: C 43.11; H 1.44; N 14.24. For C₁₄H₆N₄O₄S₃ calculated, %: C 43.07; H 1.55; N 14.35.

4,7-Dithiophene-2-yl-benzo[1,2,5]thiadiazole-5,6-diamine (6). Synthesized by the method reported earlier.¹³ Yield 2.2 g (71%). Mp = 219–220 °C, lit. Mp = 239–240 °C.¹³ NMR ¹H (CDCl₃, 400 MHz, δ, ppm): 7.55 (d, J = 6.0 Hz, 2H); 7.36 (d, J = 3.4 Hz, 2H); 7.24 (t, J = 3.6 Hz, 2H); 4.38 (s, 4H). Found, %: C 50.80; H 3.00; N 17.06. For C₁₄H₁₀N₄S₃ calculated, %: C 50.89; H 3.05; N 16.95.

4,8-Dithiophene-2-yl-benzo[1,2-c;4,5-c′]bis[1,2,5]thiadiazole (7). Synthesized by the method reported earlier.¹⁴ Yield 1.9 g (79%). Mp = 332–334 °C, lit. Mp = 334–336 °C.¹⁴ NMR ¹H (DMSO-d6, 400 MHz, δ, ppm): 7.72 (d, J = 4.2 Hz, 2H); 7.51 (d, J = 2.4 Hz, 2H); 7.23 (t, J = 3.6 Hz, 2H). Found, %: C 46.70; H 1.61; N 15.49. For C₁₄H₆N₄S₄ calculated, %: C 46.91; H 1.69; N 15.63.

4,8-Bis-(5-bromothiophene-2-yl)-benzo[1,2-c;4,5-c′]-bis[1,2,5]thiadiazole (8). Synthesized by the method reported earlier.¹⁵ Yield 1.72 g (74%). Mp > 300 °C. NMR ¹H (DMSO-d6, 400 MHz, δ, ppm): 7.86 (d, J = 4.2 Hz, 2H); 7.18 (d, J = 4.2 Hz, 2H). Found, %: C 32.6. For C₁₄H₄Br₂N₄S₄ calculated, %: C 32.57; H 0.78; Br 9.14; S 7.23; N 10.85.

2,6-Bis-trimetylstannyl-4,8-didodecyloxybenzo[1,2-b;4,5-b′]dithiophene (9). Synthesized by the method reported earlier.¹⁶ Yield 7.1 g (90%). Mp = 48–50 °C. NMR ¹H (CDCl₃, 400 MHz, δ, ppm): 7.50 (t, J = 14.1 Hz, 2H); 4.29 (t, J = 6.5 Hz, 4H); 1.88 (m, 4H); 1.60–1.53 (m, 4H); 1.40–1.24 (m, 32H); 0.88 (t, J = 6.8 Hz, 6H); 0.44 (t, J = 27.7 Hz, 18H). Found, %: C 53.98; H 8.03; S 7.24. For C₄₀H₇₀O₂S₂Sn₂ calculated, %: C 54.31; S 7.25; H 0.83.

Synthesis of polymer P1. In a 25 ml three-necked flask equipped with a reflux condenser and magnetic stirrer, 0.7311 g (0.8265 mmol) 2,6-bis-trimetylstannyl-4,8-didodecyloxybenzo[1,2-b;4,5-b′]dithiophene (9), 0.3785 g (0.8265 mmol) 4,7-bis-(5-bromothiophene-2-yl)-benzo[1,2,5]thiadiazole (5) and 0.065 g (0.056 mmol) Pd(Ph₃P)₄ were placed in a stream of argon and 15 ml dry toluene and 2 ml dry DMF were added. The reaction mixture was stirred at 115 °C for 48 h under argon, after which 0.02 g 2-bromothiophene and 0.02 g 2-(tributylstannyl)thiophene was added and stirring was continued for 5 h. The mixture was then cooled to room temperature and the product was poured into 200 ml methanol and filtered. The polymer was dissolved in chloroform and reprecipitated in methanol, then purified by extraction with methanol, hexane, and acetone in a Soxhlet apparatus and dried in a vacuum. Yield 62%. NMR ¹H (CDCl₃, 400 MHz, δ, ppm): 8.00–6.84 (m, 8H); 4.16–3.61 (m, 4H); 2.02–1.88 (m, 2H); 1.74–0.90 (m, 72H); 0.83–0.53 (m, 6H).

Synthesis of copolymer P2. Synthesized similarly to copolymer P1, using 2,6-bis-trimetylstannyl-4,8-didodecyloxybenzo[1,2-b;4,5-b′]dithiophene (9) and 4,8-bis-(5-bromothiophene-2-yl)-benzo[1,2-c;4,5-c′]-bis-[1,2,5]thiadiazole 8. Yield 60%. NMR ¹H (CDCl₃, 400 MHz, δ, ppm): 7.85–6.33 (m, 6H); 2.37–0.35 (m, 50H).

2.3 Methodology of the theoretical calculations

In addition to the experimental measurements, we have studied the structures theoretically within the framework of density functional theory (DFT) and time-dependent density functional theory (TD-DFT). The P1 and P2 monomers exhibit structural flexibility which led us to design several stereoisomers as initial geometries for our calculations. The initial stereoisomers also included cis- and trans-configurations of the alkane units.

The calculations have been performed employing the gradient corrected functional PBE¹⁷ and the hybrid exchange–correlation functional PBE0^18a (without adjustable parameters) of Perdew, Burke and Ernzerhof, as well as the hybrid functional of Becke, Lee, Yang and Parr, B3LYP.^18b,c Jacquemin et al. have reported an extensive comparative study¹⁹ on the performance of hybrid functionals with and without long-range corrections for the calculation of the visible absorption spectra of a wide range of organic dyes. They find that among the functionals under study, while also taking into account solvent effects, the PBE0 functional provides un-calibrated excitation energies with the smallest average absolute deviation of 0.14 eV. The long-range corrected functionals may outperform PBE0 on average only when the excitation energies are calibrated with (per case) linear corrections. Compared to our experimental results we find that the widely used B3LYP functional performs very well, and can also serve as a reference to facilitate comparison with the literature.

For each structure we have used several stereoisomers as initial geometries for the geometry optimizations. This first round of calculations was performed employing the PBE functional. The resolution of the identity method²⁰ was used for the treatment of the two-electron integrals. The TZVP basis set²¹ which is of triple-ζ quality was used throughout. The resulting structures were further optimized using both the hybrid PBE0 and B3LYP functionals, using the same basis set and also taking into account solvent effects. The polarization continuum model (PCM)²² was employed to account for effects from solute–solvent interactions. The importance of accounting for solvation effects when calculating the absorption spectrum has been extensively shown for a wide range of systems, including, but not limited to organic dyes, dye-sensitized TiO₂ nanoparticles and Ru(II)–polypyridyl complexes.^23–27 Tight convergence criteria were in place for the SCF energy (up to 10⁻⁷ Eh) and the one-electron density (rms of the density matrix up to 10⁻⁸) as well as for the norm of the Cartesian gradient (residual forces both average and maximum smaller than 1.5 × 10⁻⁵ a.u.) and residual displacements (both average and maximum smaller than 6 × 10⁻⁵ a.u.).

The TD-DFT excited state calculations have been performed, using the same functionals, on the corresponding ground state structures. In each case we have calculated the optical gap in the presence of solvents. For solvents we have used tetrahydrofuran (THF) and ortho-dichlorobenzene (o-DCB or 1,2-DCB) using the self-consistent reaction field solvent parameters as defined in the Gaussian 09 electronic structure package.²⁸ The first round of geometry optimization was performed using the Turbomole package.²⁹ All of the follow up calculations were performed using the Gaussian package.²⁸

2.4 Device fabrication and characterization

Indium tin oxide (ITO) coated glass substrates were ultrasonically cleaned sequentially in detergent, deionized water and isopropyl alcohol and dried. A layer of poly(3,4-ethylenedioxythiophene):poly(styrene sulphonate) (PEDOT:PSS) (thickness about 65 nm) was deposited on the ITO coated glass substrates by spin coating from an aqueous solution and then annealed on a hot plate at 90 °C for 20 min in air and subsequently cooled to room temperature. The blends of P1 or P2:PC₆₀BM or PC₇₀BM of different weight ratios were prepared by dissolving the copolymers and fullerene derivatives in THF solution and stirring for about 2 h. For the solvent additive, i.e. DIO/THF, a small amount of DIO (3% by volume) is added to the THF solution. The photoactive layer of P1 or P2:PC₆₀BM or PC₇₀BM (in different ratios with and without solvent additive DIO) was deposited onto the PEDOT:PSS layer by spin coating and then dried in ambient atmosphere. Finally, an aluminum (Al) metal top electrode was deposited in vacuum on the active layer at a pressure of ca. 5 × 10⁻⁵ Pa. The active area of the device was ca. 0.14 cm². The current density–voltage (J–V) characteristics were measured on a computer-controlled Keithley 236 source meter unit. A xenon lamp coupled with AM 1.5 solar spectrum filter was used as the light source, and the optical power at the sample was 100 mW cm⁻².

3. Results and discussion

3.1 Synthesis and characterization of the copolymers

Benzo[1,2,5]thiadiazole is one of the most attractive electron-withdrawing fragments. Its functionalization by different substituents makes it possible to obtain a wide range of heteroaromatic compounds that are promising as “building blocks” for new electroactive polymers. In this regard, on the basis of key compound – 4,7-dibromobenzo[1,2,5]thiadiazole 1 – we have developed a number of thiophene-containing derivatives presented in Scheme 1. Herewith, compound 6 has been previously described in the literature.

Scheme 1

By the interaction of 4,7-dibromobenzo[1,2,5]thiadiazole 1 with 2-tributylstannylthiophene in Stille coupling reaction using Pd(PPh₃)₂Cl₂ as catalyst, followed by bromination of product 2 by N-bromosuccinimide, we obtained 4,7-bis-(5-bromothiophene-2-yl)-benzo[1,2,5]thiadiazole 3 in 66% yield,¹² the purity, composition and structure of which are confirmed by melting point, NMR spectroscopy and elemental analysis. Nitration of 4,7-dibromobenzo[1,2,5]thiadiazole 1 by fuming nitric acid in 100% sulfuric acid yielded dinitroderivative 4, which is then converted to the 5,6-dinitro-4,7-dithiophene-2-yl-benzo[1,2,5]thiadiazole 5 by Stille cross-coupling in THF.¹³ Reduction of 5 by iron powder in acetic acid resulted in 4,7-dithiophene-2-yl-benzo[1,2,5]thiadiazole-5,6-diamine 6.¹³ By condensation of compound 6 with N-thionylaniline in the presence of (CH₃)₃SiCl, 4,8-dithiophene-2-yl-benzo[1,2-c;4,5-c′]-bis-[1,2,5]thiadiazole 7 was obtained,¹⁴ which proved to have a quite high melting point (Mp = 332–334 °C) and is a barely soluble product.

Further bromination in the α-positions of the thiophene fragments was performed using N-bromosuccinimide in a mixture of chloroform–acetic acid (1 : 1) at room temperature. The desired product 8 was purified by recrystallization from DMF.¹⁵

On the basis of the synthesized heteroaromatic thiophene-containing monomers, a series of new low band gap copolymers with the quinoid π-conjugation nature and a strict alternation of donor–acceptor units has been developed, obtained by the Stille cross-coupling reaction. Polycondensation was carried out under argon in a mixture of toluene and DMF at 115 °C for 48 h, using tetrakis(triphenylphosphine) palladium(0) as a catalyst. The role of electron-donor component in the structures of polymers P1 and P2 is served by 2,6-bis-trimetylstannyl-4,8-didodecyloxybenzo[1,2-b;4,5-b′]dithiophene 9, and the electron acceptor fragments are monomers 3 and 8 (Scheme 2).

Scheme 2

The resulting polymers were purified from the residual metal catalyst, organotin and low molecular weight impurities by re-precipitation twice from solution in methanol and subsequent extraction with methanol, hexane and acetone. The yield of the desired high molecular mass products was 60–76%.

The composition and structure of the low band gap polymers P1 and P2 were confirmed by ¹³C NMR spectroscopy as shown in Fig. 1a and b for P1 and P2, respectively.

Fig. 1 ¹³C NMR spectra of copolymer (a) P1 and (b) P2.

For P1¹³C NMR (100 MHz, CdCl₃, δ) 14.10 (82, 100); 22.72 (81, 99); 26.48 (73, 91); 28.24 (74, 92), 29.36 (79, 97); 29.54 (78, 96); 29.58 (77, 95); 30.09 (75, 93); 30.22 (76, 94); 30.50 (72, 90); 31.98 (80, 98); 69.47 (71, 89) (Aliph); 118.00 (63); 118.38 (69); 118.85 (29, 57); 121.52 (60); 121.92 (66); 122.01 (30, 56); 123.66 (51, 52); 131.26 (68); 131.66 (62); 132.98 (44, 50); 134.71 (26, 55); 137.06 (61, 67); 137.48 (64); 139.21 (58); 144.92 (28, 53); 151.08 (45, 49) (Ar).

For P2¹³C NMR (100 MHz, CdCl₃, δ) 14.10 (85, 103); 22.72 (84, 102); 26.48 (76, 94); 28.24 (77, 95), 29.36 (82, 100); 29.54 (81, 99); 29.58 (80, 98); 30.09 (78, 96); 30.22 (79, 97); 30.50 (75, 93); 31.98 (83, 101); 69.47 (74, 92) (Aliph); 111.65 (44, 50); 113.75 (29, 60); 118.00 (66); 118.38 (72); 121.52 (63); 121.92 (69); 122.28 (30, 59); 131.26 (71); 131.66 (65); 134.98 (26, 58); 137.06 (64, 70); 137.48 (67); 140.71 (61); 141.32 (28, 56); 150.79 (45, 49); 153.19 (51, 55) (Ar).

Average molecular weights and polydispersity of the polymers were determined by GPC (eluent – THF, standard – polystyrene) and varied in the range of (1.22–1.43) × 10⁴ and 1.48–1.72, respectively. Both the copolymers P1 and P2 are soluble in common organic solvents such as DMF, N-methylpyrrolidone, chloroform, THF, and o-dichlorobenzene. The glass transition temperature of polymers determined by TMA is in the range of 344–348 °C (Table 1).

Table 1 Molecular mass and thermal properties of copolymers P1 and P2

Polymer	–Y–	M n (×10^−4)	M w (×10^−4)	M w/Mn	T g (°C)	T 10% (°C)
a In the numerator – the air, in the denominator – the atmosphere of argon.
P1		1.22	1.81	1.48	196	341, 344
P2		1.43	2.46	1.72	333	338, 348

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As can be seen from Table 1, the presence of long alkyl substituents in the structure of macromolecules generally leads to lower glass transition temperature of the polymers. However, a larger number of annelation of heterocyclic fragments into monomeric units, which reduces the flexibility of individual structural fragments, tends to increase T_g.

T _10% values were determined by TGA both in air and argon in the range 295–341 °C and 316–348 °C, respectively. Fig. 2 shows the typical TGA curves of polymers P1 and P2 in the air. Both the synthesized copolymers exhibit relatively high thermal stability.

Fig. 2 TGA curves of copolymers P1 and P2 in the air.

3.2 Optical and electrochemical properties

One of the main properties of the copolymers resulting in their perspectives for application as photoactive materials in solar cells is their absorption coefficient in the visible and near IR ranges.³⁰ The electronic absorption properties of the synthesized copolymers were measured in both THF solution and as thin films spin coated onto quartz substrates and are shown in Fig. 3a (P1) and b (P2). The results of the spectral data are presented in Table 2.

Fig. 3 The normalized absorption spectra of copolymers (a) P1 and (b) P2 in solution and thin film form cast from THF solution.

Table 2 Optical properties of copolymers P1 and P2

Polymer	λ ^absmax solution (nm) (molar extinction coefficient)	λ ^absmax film (nm) (molar extinction coefficient)	λ ^absonset film (nm)	E ^optg (eV)
P1	398 (3.20 × 10^4 cm^−1); 539 (2.34 × 10^4 cm^−1)	419; 588	786	1.58
P2	479 (3.14 × 10^4 cm^−1); 866 (1.56 × 10^4 cm^−1)	481; 896	1042	1.18

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The UV-visible spectrum of the copolymers is similar to that of the D–A copolymers³¹ with two characteristics absorption bands observed near 410–481 and 588–1050 nm (Table 2, Fig. 3). We assign the high energy absorption bands to π–π* transitions from the BT unit and the longer wavelength absorption bands to the intramolecular charge transfer (ICT) between the BT donor and different acceptor units in the copolymers. An interesting observation is the impressive bathochromic shift (415 nm) of the absorption band of polymer P2 with respect to polymer P1 (Fig. 3a and b), which differs from the first fragment by the absence of a second thiadiazole ring in co-monomer 3. This band of polymer P2 is almost completely located in the near-infrared area and determines the ability of this material to absorb photons of the solar spectrum with the lowest energy.³² These results also indicate that the conjugated side chains can easily tune the absorption behavior of BT based D–A copolymers, which is beneficial for the design and preparation of efficient photovoltaic materials with enhanced absorption and suitable energy levels. The peaks of the long-wavelength absorption bands of copolymers P1 and P2 (which are based on benzo[1,2,5]thiadiazole derivatives) are bathochromically shifted in the order λ^abs_max (P1) < λ^abs_max (P2). This observed shift is consistent with the general ideas on the nature of electron-withdrawing condensed heterocyclic systems as the number of “pyridine type” nitrogen atoms increases. As shown in Fig. 3a and b, the absorption peaks in the longer wavelength region are red shifted relative to their counterparts in the solution. The absorption bands become broader than that in solution presumably due to the better planarity of the polymer backbone, the formation of π-stacked structures as well as a strong electronic interaction between the energy values of the optical band gaps of P1 and P2 that were calculated from the absorption spectra of the polymers in the thin films, specifically from the onset absorption edge in the longer wavelength region. The values have been complied in Table 2. The optical band gap of polymer P2 (1.18 eV) is lower than P1 (1.58 eV) which indicates an increase in the ability of the monomer to lower the HOMO–LUMO energy gap of polymers. The small band gap observed in P2 can be attributed to the high electron affinity and large quinoid contribution of the BDT unit.³³

The electrochemical properties of polymers P1 and P2 were examined by cyclic voltammetry. It was observed that both copolymers exhibit reversible or partially reversible redox properties due to the high electro-activity. As shown in Fig. 4, both copolymers show a predominant oxidation peak (p doping) due to the electron donating benzothiophene and aryl-vinylene segment and a minor reduction peak (n doping) related to the electron withdrawing benzothiadiazole unit. The onset oxidation potential (E^oxd_onset) and onset reduction potential (E^red_onset) observed in the cyclic voltammogram for P1 and P2 are summarized in Table 3. The values of HOMO and LUMO energy levels for copolymers were estimated from the following expressions

HOMO = −q(E^ox_onset + 4.71) eV

LOMO = −q(E^red_onset + 4.71) eV

Fig. 4 Cyclic voltammograms of P1 andP2 in 0.1 M solution of lithium perchlorate in acetonitrile at a scan rate 50 mV s⁻¹.

Table 3 Electrochemical properties of copolymers P1 and P2

Polymer	E ^redonset (V)	E ^oxonset (V)	E ^HOMO (eV)	E ^LUMO (eV)	E ^ecg (eV)
P1	−1.14	0.60	−5.24	−3.48	1.76
P2	−0.92	0.34	−5.10	−3.82	1.28

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From the onset potentials of the p-doping process, the HOMO energy levels were estimated to be −5.24 eV and −5.10 eV for P1 and P2, respectively. On the other hand, the LUMO energy levels were estimated from the onset reduction potentials of n doping, to be −3.48 eV and −3.82 eV for P1 and P2, respectively. The values found for the energy band gap (E^ec_g) of the polymers using cyclic voltammetry according to the equation E^ec_g = (HOMO − LUMO) are 1.76 eV and 1.28 eV for P1 and P2, respectively. The values and trends are consistent with those estimated by the optical method described above. The higher electrochemical band gap is a common phenomenon for the conjugated polymers because of the energy barrier of charge transfer at the electrodes during electrochemical measurements.

3.3 Theoretical calculations

The stereoisomers for each of the monomers (P1 and P2), optimized at the PBE/TZVP level of theory, are found to be nearly isoenergetic. This is to be expected for rotamers and cis/trans isomers. The maximum energy difference for each case does not exceed 65 meV (1.5 kcal mol⁻¹), with the trans isomers energetically lower. This energy difference can be considered as the largest vertical rotational barrier in the gas phase. Corresponding adiabatic values would of course be larger and would account for possible steric effects.

The PBE functional belongs to the generalized gradient approximation (GGA) family of functionals which are known to underestimate the highest occupied molecular orbital (HOMO) − lowest unoccupied molecular orbital (LUMO) gap (HL gap). The calculated HOMO and LUMO energy levels, HL gap, optical band gap, oscillator strength and dipole moment computed at different levels of theory in the gas phase as well as THF and o-DCB solvent phases are shown in Tables 4 and 5. From Tables 4 and 5 we can see that the PBE functional systematically underestimates the HL gap, compared to the experimental values, on average by about 0.25 eV.

Table 4 Calculated properties of the P1 and P2 structures. Specifically HOMO and LUMO energies (eV), HOMO–LUMO gap (eV) and dipole moment (D)

		Gas phase
		HOMO (eV)	LUMO (eV)	HL gap (eV)	|Dipole moment| (D)
P1	PBE	−4.73	−3.52	1.21	0.50
	PBE0	−5.50	−2.85	2.65	0.26
	B3LYP	−5.23	−2.93	2.30	0.21
P2	PBE	−4.65	−4.13	0.52	1.72
	PBE0	−5.30	−3.73	1.58	0.80
	B3LYP	−5.04	−3.75	1.30	1.09

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Table 5 Calculated properties of the P1 and P2 structures. Specifically HOMO–LUMO gap (eV), optical gap (eV) with corresponding oscillator strength, and dipole moment (D). Solvent effects are taken into account for tetrahydrofuran (THF) and ortho-dichlorobenzene (o-DCB)

		HOMO (eV)	LUMO (eV)	HL gap (eV)	Optical gap (eV)	f, osc. strength	|Dipole moment| (D)
Solvent THF
P1	PBE	−4.84	−3.55	1.28	1.43	0.23	0.75
	PBE0	−5.63	−2.89	2.75	2.22	0.80	0.51
	B3LYP	−5.44	−3.02	2.42	2.06	0.68	0.57
P2	PBE	−4.72	−4.15	0.57	0.87	0.28	1.76
	PBE0	−5.44	−3.76	1.68	1.29	0.57	0.78
	B3LYP	−5.25	−3.86	1.38	1.16	0.52	0.94
Solvent o-DCB
P1	PBE	−4.85	−3.56	1.29	1.43	0.26	0.77
	PBE0	−5.64	−2.89	2.75	2.21	0.83	0.54
	B3LYP	−5.45	−3.02	2.43	2.05	0.72	0.60
P2	PBE	−4.73	−4.15	0.57	0.86	0.31	1.74
	PBE0	−5.45	−3.76	1.69	1.28	0.60	0.79
	B3LYP	−5.26	−3.87	1.39	1.15	0.55	0.94

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The energetically lowest isomers of each case were further optimized, initially in the gas phase, employing the hybrid functionals PBE0 and B3LYP. No significant structural differences arise from these optimizations which suggests that the PBE functional produces high quality structures. For the structure P2 the calculated values for the HL gaps are in excellent agreement with the experimental values, as is shown in Tables 2–5. A discrepancy is noted for the P1 structure for which the calculated HL gaps are overestimated by both hybrid functionals with regards to the experiment. More so, the overestimation is not the same for the two functionals but differs by about 0.16 eV.

The HOMO and LUMO of the P1 and P2 monomers are shown in Fig. 5, plotted using an isovalue of 0.02. In both cases the HOMO is delocalized over the dithiophene, thiophene and thiadiazole units of the structures as is shown in Fig. 5a and b. Contribution of alkane groups to the HOMO is negligible. In both cases the LUMO is highly localized on the thiadiazole units extending slightly over the linking thiophene units. The localization of the LUMO on the thiadiazole is more pronounced in the case of P2, compared to P1.

Fig. 5 HOMO and LUMO isosurfaces of the P1 and P2 monomers.

The UV/visible absorption spectra, calculated at the DFT/B3LYP level of theory and using THF as a solvent, are shown in Fig. 6 for both the P1 and P2 monomers. The spectra have been produced by convoluting Gaussian functions with HWHM = 0.1 eV centered at the excitation energies. The P2 spectrum exhibits three clear bands with high intensities, centered at 1.2 (1033 nm), 2.8 (443 nm) and 3.6 eV (344 nm) respectively. All remaining peaks have significantly lower intensities. The most prominent difference between the two spectra is the existence of a gap between the first and second significant peaks in the spectrum of P2 which is about 1.6 eV (775 nm). The spectrum of the P1 monomer exhibits two well defined main bands centered at 2.0 (620 nm) and 3.0 eV (413 nm). A third band is also present which is much broader and has a fine structure of two lower intensity peaks centered at 3.8 eV (326 nm) and 4.2 (295 nm). Secondary peaks of low intensities are also found between the bands. However, the intensities are larger that the corresponding secondary peaks of the P2 spectrum.

Fig. 6 Theoretical UV/vis absorption spectra of the (a) P1 and (b) P2 monomers estimated from DFT.

3.4 Electrical properties of the copolymers

We have investigated the current–voltage characteristics of the copolymers P1 and P2 to investigate their semiconductor nature, in the dark employing the ITO/PEDOT:PSS/P1 or P2/Al device structure and illustrated in Fig. 7. The J–V characteristics of the devices in the dark showed a rectification effect when positive potential was applied to the ITO/PEDOT:PSS electrode with respect to the Al electrode. Since the HOMO level of both copolymers (in the range of −5.02 to −5.24 eV) lies very close to the work function of PEDOT:PSS (−5.1 eV), it forms a nearly ohmic contact for hole injection from the ITO/PEDOT:PSS into the HOMO level of copolymer. However the LUMO level of the copolymers (in the range of −3.56 to −3.80 eV) is very far from the work function of the Al electrode (−4.2 eV) and they form a Schottky barrier for electron injection from Al into the LUMO level. Therefore the rectification effect observed in the devices is due to the formation of the Schottky barrier at the Al–copolymer interface which indicates that the copolymer behaves as a p-type organic semiconductor, and thus can be used as electron donors for the BHJ organic solar cells. We have also estimated the hole mobility of the pristine P1 and P2 from the current–voltage characteristics in the dark using the SCLC method and found them to be 6.62 × 10⁻⁵ cm² V⁻¹ s⁻¹ and 2.14 × 10⁻⁶ cm² V⁻¹ s⁻¹ for P1 and P2, respectively. We assume that the polymer solar cell based on P1 as electron donor in the BHJ active layer may achieve higher power conversion efficiency compared to its P2 counterpart.

Fig. 7 J–V characteristics of ITO/PEDOT:PSS/P1 or P2/Al devices in the dark at room temperature.

3.5 Photovoltaic properties

The energy band diagrams of the copolymers along with fullerene derivatives are shown in Fig. 8. As can be seen from this figure, the LUMO offset for the P1:PCBM BHJ is higher than that for the P2:PCBM BHJ active layer, therefore we assume that the photo-induced charge transfer in the P1 based device is more efficient that for the P2 based device.

Fig. 8 Energy band diagrams of the P1, P2 and PCBMs.

The bulk heterojunction PSCs were fabricated with a device structure ITO/PEDOT:PSS/P1 or P2:PC₆₀BM (1 : 1 by wt)/Al to investigate the photovoltaic properties of the copolymers. Fig. 9 shows the typical current–voltage (J–V) characteristics of the devices under illumination of AM 1.5, 100 mW cm⁻². The corresponding photovoltaic parameters such as short circuit current (J_sc), open circuit voltage (V_oc), fill factor (FF) and overall power conversion efficiency (PCE) are summarized in Table 6. The BHJ photovoltaic device based on P1:PC₆₀BM showed a V_oc of 0.86 V, a J_sc of 5.26 mA cm⁻² resulting in a PCE of 2.17% and FF of 0.48. Under the same conditions, the P2:PC₆₀BM device exhibits a lower PCE of 0.80% with a V_oc of 0.61 V, a J_sc of 3.28 mA cm⁻², and FF of 0.40. The higher value of V_oc for the device based on P1 is due to its low lying HOMO energy level, since V_oc is directly related to the difference between HOMO energy level of donor and LUMO energy level of the acceptor.

Fig. 9 J–V characteristics of BHJ photovoltaic devices based on P1:PC₆₀BM (1 : 1 wt ratio) and P2:PC₆₀BM (1 : 1 wt ratio) blends cast from THF solvent, under illumination (100 mW cm⁻²).

Table 6 The results of photovoltaic measurements of solar cells based on low band gap polymers P1 and P2 as electron donor and PC₆₀BM and PC₇₀BM as electron acceptors (optimized blend ratio is 1 : 1 w/w)

Blend	J sc (mA cm^−2)	V oc (V)	FF	PCE (%)
a Cast from THF solvent. b Cast from DIO/THF solvent.
P1:PC60BM^a	5.26	0.86	0.48	2.17
P1:PC70BM^a	7.94	0.84	0.52	3.45
P2:PC60BM^a	3.28	0.61	0.40	0.80
P2:PC70BM^a	4.48	0.62	0.43	1.19
P1:PC70BM^b	10.52	0.84	0.60	5.30
P2:PC70BM^b	6.05	0.58	0.45	1.58

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It was anticipated that copolymer P2 should show a high J_sc value, since it absorbed strongly in both visible and NIR regions, as confirmed by the absorption spectra. However, results gave a lower J_sc value than P1. One of the reasons for the low J_sc value was attributed to the strong electron accepting ability of BDT. The high electron withdrawing capability of the BDT would result in a low lying LUMO level, and thus excitons generated after the photo-excitation might not have enough driving force to allow the generated exciton to dissociate.³⁴ Another possibility can be ascribed to the small LUMO–LUMO offset between P2 and PCBMs which would not provide enough driving force for the electron injection process to occur.

The solar cell based on polymer P2:PC₆₀BM showed a low value of efficiency (0.80%). Despite the low band gap, the short-circuit current of the device based on this polymer was significantly lower than that for polymer P1. J_sc is defined as the product of the density of photoinduced charge carriers on their mobility in organic semiconductors. The mobility of charge carriers are related to the parameters of the device, and not to the photoactive material and in a more significant extent depends on film nano-morphology. The morphology, in turn, is affected by the conditions of solar cell manufacturing: solvent evaporation time, the temperature of the solution and application method. Probably, the smaller solubility of macromolecules P2 in THF (effective solvent for the formation of bulk heterojunction) was the reason which led to lower values of the characteristics of photovoltaic devices based on this polymer.

The solar cells’ morphology of the active layer plays an important role in improving the photovoltaic characteristics. We have measured the images of the surface morphology of the active layer P1:PC₆₀BM and P2:PC₆₀BM obtained by AFM. The rms roughness of the P1:PC₆₀BM and P2:PC₆₀BM films are 0.76 nm and 2.71 nm, respectively. The P1:PC₆₀BM blend film has well-distributed small-sized domains with small roughness of 0.76 nm, suggesting that the polymer is suitably compatible with the PC₆₀BM. Film P2:PC₆₀BM shows relatively large domains and phase separation. The rough morphology and large size of the phase separation of the P2:PC₆₀BM film does not promote effective separation of the excitons, leading to a lower short-circuit current. Nano-sized aggregated domains and the improved phase separation of the P1:PC₆₀BM film contributes to a better separation of the charge and increase in PCE.

The PCE of the BHJ solar cells which are based on a blend of polymer P1 or P2 (donor) and PC₆₀BM (acceptor) is quite low. To further improve the PCE, we have used PC₇₀BM as the electron acceptor instead of PC₆₀BM. PC₇₀BM has a similar electronic structure and energy levels, but has considerably stronger absorption in the visible region.³⁵ Therefore, the copolymer based BHJ polymer solar cells can be expected to exhibit improved performance when PC₇₀BM is used in place of PC₆₀BM as the acceptor due to the compensation of the poor absorption of the copolymer:PC₆₀BM blend in the 400–550 nm range. We have fabricated the BHJ devices with the P1:PC₇₀BM and P2:PC₇₀BM blends and also investigated their photovoltaic properties. The J–V characteristics of these devices are shown in Fig. 10a and photovoltaic parameters are listed in Table 6. As shown in Fig. 10a and Table 6, both devices exhibit improvement in the PCE compared to the corresponding PC₆₀BM based devices. The P1 and P2 devices show V_oc values of 0.84 V and 0.62 V, J_sc values of 7.94 mA cm⁻² and 4.48 mA cm⁻², and FFs of 0.52 and 0.43, which results in PCEs of 3.45% and 1.19%, respectively. The enhancement in the PCE of the devices based on the PC₇₀BM acceptor is mainly due to the increased value of J_sc. This increase is attributed to the strong absorption of PC₇₀BM in the range of 400–550 nm, which is missing in the case of PC₆₀BM. The increased absorption of PC₇₀BM enhances the light harvesting efficiency when PC₇₀BM is used instead of PC₆₀BM, which results in further improved PCE. The V_oc values in these device change slightly, since both PC₆₀BM and PC₇₀BM have almost the same LUMO energy level.

Fig. 10 (a) J–V characteristics under illumination intensity of 100 mW cm⁻² and (b) IPCE spectra of BHJ photovoltaic devices based on P1:PC₇₀BM (1 : 2 wt ratio) and P2:PC₇₀BM (1 : 2 wt ratio) blends cast from THF solvent.

The IPCE values for the P1:PC₇₀BM based solar are significantly higher than that for the P2:PC₇₀BM counter part (Fig. 10b) indicating a more efficient generation and transport charge carrier in the P1:PC₇₀BM BHJ active layer. Similar results were also observed for the BHJ solar cells with PC₆₀BM as electron acceptor. The significant increase in the J_sc in the device based on P1 as electron donor in the BHJ active layer either with PC₆₀BM or PC₇₀BM as electron acceptor, is due to the high value of IPCE. The IPCE values for the P2:PC₇₀BM device are relatively very small as compared to P1:PC₇₀BM, although both P1 and P2 have a broad absorption spectral response, even P2 has a broader and low band gap. The difference in the values of IPCE can be ascribed to the higher value of LUMO level of P1, leading to a large LUMO–LUMO offset towards the PCBM and an increased driving force for exciton dissociation.

3.6 Effect of solvent additive on the photovoltaic response

The values of the PCE for these copolymers as donor and fullerene derivatives are still low. To improve the PCE of BHJ polymer solar cells, it is necessary to properly form continuous pathways in the active layer to transfer the electrons and holes towards the collection electrodes without leading to recombination. In addition, the morphology of the BHJ active layer plays an important role in determining the performance of the BHJ polymer solar cells, because charge carrier generation, recombination and transport are influenced by the spontaneously formed micro/nanostructure of the active layer during solution processing.³⁶ The morphology of the active layer can be significantly improved through different processing methods to prepare the active layer i.e. thermal annealing^36b or solvent annealing.^36c For the P3HT based BHJ solar cells, the thermal annealing of the BHJ active layer has led to a remarkable enhancement in the PCE compared to that of its pristine film because P3HT consists of an ordered crystalline phase.^36b,37 On the other hand, D–π–A type copolymer does not usually exhibit higher crystallinity than P3HT even after thermal annealing,³⁸ therefore the thermal annealing method does not always produce the optimal intermolecular self organization in amorphous morphologies.³⁹ As an alternative, Heeger et al. introduced a solvent mixture method in which the BHJ active layer was prepared through a mixture of two different solvents having different boiling points.⁴⁰ The addition of small amounts of high boiling point additive into the host solvent for the preparation of BHJ active layer also improves the morphology of the active layer.^37,41 In recent years, remarkable progress has been achieved for the BHJ polymer solar cells through the combination of D–A copolymers as electron donor and solvent additive⁴² and recently a PCE of 8.62% has been reported for BHJ polymer solar cells using this processing method.⁴³ We have also prepared the P1:PC₇₀BM and P2:PC₇₀BM BHJ active layers by adding (1%, 3% and 5% by volume) 1,8-diiodooctane (DIO) into the host solvent. The optimized concentration was found to be 3% by volume. The current–voltage characteristics of P1:PC₇₀BM and P2:PC₇₀BM solar cells processed from the DIO/THF solvent are shown in Fig. 11a and the photovoltaic parameters are summarized in Table 6. The J_sc was increased from 4.48 to 6.05 mA cm⁻² and from 7.94 to 10.52 mA cm⁻² for the P2:PC₇₀BM and P1:PC70BM based solar cells, respectively. Similarly the FF has been increased from 0.43 to 0.45 and from 0.52 to 0.60 for the P2:PC₇₀BM and P1:PC₇₀BM based solar cells, respectively. However, the V_oc was slightly lowered for the devices processed with DIO/THF solvent. The increase in J_sc and FF leads to enhancement of the PCE from 1.19 to 1.58 and from 3.45 to 5.30% for the P2:PC₇₀BM and P1:PC₇₀BM based solar cells, respectively. To investigate the reason for the increase in J_sc with the addition of DIO in the host THF solvent for the processing of the BHJ film, we have measured the series resistance (R_s) of the device from the slope of the J–V characteristics under illumination near the open circuit voltage. We found that the value of R_s has been reduced from 6.4 Ω cm² (THF processed) to 4.7 Ω cm² (DIO/THF processed). The decrease in the R_s for the device processed from DIO/THF solvent as compared to that of THF indicates that the mobility of charge carriers increases. The higher IPCE value (as shown in Fig. 11b) of the device processed from DIO/THF than that processed from THF indicates an increased charge carrier generation and transportation and also supports the enhancement in the value of J_sc.

Fig. 11 (a) J–V characteristics under illumination intensity of 100 mW cm⁻² and (b) IPCE spectra of BHJ photovoltaic devices based on P1:PC₇₀BM (1 : 2 wt ratio) and P2:PC₇₀BM (1 : 2 wt ratio) blends cast from DIO/THF solvent.

To acquire information about the increase in the PCE of the devices by adding the DIO as solvent additive, the effect of the additive on the absorption and morphology of the active layers was investigated. The UV-visible absorption spectra of the P1:PC₇₀BM films with and without DIO are shown in Fig. 12. As shown in Fig. 12, the absorption of the blend film with DIO was increased, which could enhance the efficiency of charge carrier generation due to a better photon harvesting. It was already reported earlier that the additives were able to influence the polymer absorption profiles.^37a,44 When 3% DIO was added, an increase in the absorbance over the wavelength region 500–750 nm was observed, because of the π–π transition among the copolymer chains. It contributes to the increase in the PCE compared to the device without processing additive. We assume that the crystallinity of P2:PC₇₀BM blend cast from DIO/THF may be improved as compared to the blend cast from THF without DIO.

Fig. 12 Normalized absorption spectra of P1:PC₇₀BM thin films processed from THF and DIO/THF solvents.

In order to examine the ordering of the copolymer chain by addition of DIO additive, XRD patterns were recorded and are shown in Fig. 13. The increase in the intensity of (100) diffraction peak was observed around 2θ = 7.58° in the P1:PC₇₀BM blend film processed from DIO/THF compared to the blend processed from THF. The increase in the intensity means increase in the crystallinity in copolymer. Therefore addition of DIO as additive has an impact on the phase separation of copolymer and PC₇₀BM, because PCBM selectively dissolved in the DIO. In addition, the DIO has a higher boiling point than the host THF solvent, PC₇₀BM tends to remain in the solution mixture (during drying) longer than copolymer, thereby enabling control of the phase separation and morphology of the BHJ active layer. We have estimated the crystalline size (d) of copolymer (100) using the Scherrer's equation (d = kλ/Bcos θ), where B is the full width at half maximum (FWHM) of the (100) peak, λ is the wavelength of incident X-ray (0.154 nm), θ is the angle of refraction and k is Scherrer's constant (0.9). The crystalline size of the copolymer P1 was estimated at 11.2 nm and 13.4 nm for the film processed from THF and DIO/THF solvents, respectively. Increase of crystalline size in P1 leads to an increase in the hole mobility created in the BHJ active layer, as a result the J_sc and PCE of the device improves significantly.

Fig. 13 XRD patterns of P1:PC₇₀BM active layers processed from THF and DIO/THF.

Besides the absorption, charge carrier mobility and balance between the hole and electron mobilities are also crucial factors for achieving high PCE with BHJ organic solar cells. Both hole and electron mobility were calculated using the space charge limited current model⁴⁵ using hole only and electron only devices, respectively. The J–V characteristics in the dark under forward bias were fitted with the model (Fig. 14a and b for hole and electron only devices using P1:PC₇₀BM, respectively), described by the Mott–Gurney Law. We found about over an order of magnitude increase in the hole mobility (μ_h) for the device processed with DIO/THF as compared to that processed from THF, with a mobility increased from 1.7 × 10⁻⁵ cm² V⁻¹ s⁻¹ to 3.6 × 10⁻⁴ cm² V⁻¹ s⁻¹. However, the electron mobility (μ_e) increases from 4.8 × 10⁻⁴ to 5.3 × 10⁻⁴ cm² V⁻¹ s⁻¹. Moreover, the device processed from DIO/THF shows relatively well balanced mobility (μ_e/μ_h = 1.47) compared to the device processed from THF (μ_e/μ_h = 28.23). The more balanced mobility contributes to higher J_sc and FF as observed in the device processed from DIO/THF solvent because the accumulated space charge limited current (SCLC) charges and hence recombination processes are reduced by the increase in the hole mobility and enhanced charge collection efficiency.⁴⁶

Fig. 14 Current–voltage (J–V) characteristics of (a) hole only and (b) electron only devices, in the dark, V_app is applied voltage and V_bi is built in potential.

The nanoscale morphology plays an important role in the PCE of the BHJ organic solar cells. Proper morphology is essential not only for exciton dissociation but also for the charge transport to respective electrodes for their efficient collection.^35b,47 An ideal domain size of 10–20 nm of copolymer and PCBM with an interpenetrating bi-continuous network is needed for achieving high PCE of BHJ polymer solar cells. However, both larger (>10–20 nm) and smaller (<10–20 nm) domain sizes of BHJ active layers may limit both charge transfer and separation. The morphology of the BHJ active layer processed from THF and DIO/THF was investigated through atomic force microscopy (AFM) images as shown in Fig. 15. From the height images we found that the root mean square (rms) of P1:PC₇₀BM cast from THF (0.95 nm) was lower than that processed from DIO/THF (1.6 nm). The surface roughness indicates the self-organization ordering of the polymer. Therefore, the increase in surface roughness and highly ordered crystalline domain may effectively provide separate pathways for electrons and holes, and increase the J_sc. With the addition of DIO, the nanoscale domain can provide a pathway through which electrons and holes are separated after forming an interpenetrating network.

Fig. 15 AFM phase images (image area 3 μm × 3 μm) of P1:PC₇₀BM blended thin films cast from (a) THF and (b) DIO/THF solvents.

4. Conclusion

We have synthesized two novel D–A copolymers with different acceptor structures of 4,8-bis-(5-bromothiophene-2-yl)-benzo[1,2,5]thiadiazole (P1) and 4,8-dithiophene-2-yl-benzo[1,2-c;4,5-c′]-bis-[1,2,5]thiadiazole (P2) with a common BDT donor segment and investigated their optical and electrochemical properties. Both donor–acceptor copolymers exhibit good solubility in common organic solvents and a broad absorption from 350 nm to the near infrared region. The optical bands of P1 and P2 estimated from the onset absorption edge are 1.58 eV and 1.16 eV, respectively. These copolymers were used as electron donors along with the fullerene derivatives (PC₆₀BM and PC₇₀BM) as electron acceptors for the fabrication of BHJ polymer solar cells. Preliminary results show that the solar cells based on P1: PC₆₀BM, P2:PC₆₀BM, P1:PC₇₀BM and P2:PC₇₀BM exhibit PCEs of 2.17%, 0.80%, 3.45%, 1.19%, respectively. The higher value of PCE for P1 as electron donor has been attributed to its relatively high mobility and high open circuit voltage due to the deep lying HOMO level and high J_sc due to the large driving force for exciton dissociation at the D–A interface in the BHJ active layer. The PCE has been further improved up to 5.30% and 1.58% for P1:PC₇₀BM and P2:PC₇₀BM blends cast from DIO/THF solvent. This increase in the PCE is attributed to the improved crystallinity of copolymer phase, more balanced charge transport in the device and larger D–A interface in the active layer for exciton dissociation, when the blend was processed from the DIO/THF solvent. These results show that the new copolymers are useful materials for organic photovoltaic applications.

Acknowledgements

This work was supported by the Presidium of the Russian Academy of Sciences under Program P-8: Polyfunctional Materials for Molecular Electronics; by the Ministry of Education and Science of the Russian Federation under State Contract agreement no. 8848; and by the Division of Chemistry and Materials Sciences, Russian Academy of Sciences, under programs of fundamental research OKh-3 (Creation and Study of Macromolecules and Macromolecular Structures of New Generations) and OKh-2 (Creation of New Metal, Ceramic, Glass, Polymer, and Composite Materials).

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