Effect of alkyl-chain branching position on nanoscale morphology and performance of all-polymer solar cells

Article abstract of DOI:10.1039/c4ra13351j

Manipulating the alkyl-chain branching position afforded two naphthalene diimide-based polymers (C1 and C3). The optoelectronic properties, including the absorption spectra and electronic energy levels, conformations and photovoltaic properties of the two

Full text of DOI:10.1039/c4ra13351j

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Effect of alkyl-chain branching position on
nanoscale morphology and performance of all-
Cite this: RSC Adv., 2015, 5, 10072
polymer solar cells†
a
a
Fangbin Liu,ab Hui Li,ab Chunling Gu and Hongbing Fu
*
*
Manipulating the alkyl-chain branching position afforded two naphthalene diimide-based polymers (C1 and
C3). The optoelectronic properties, including the absorption spectra and electronic energy levels,
conformations and photovoltaic properties of the two polymers, were fully studied and compared. The
effect of alkyl-chain branching position on their optoelectronic properties and crystallinity was
investigated in detail. Moving the branching position away from the backbone resulted in superior
crystallinity of C3. All-polymer solar cells (all-PSCs) based on C1 or C3 as the acceptor and polymer
PBDTTT-C-T as the donor were fabricated and optimized. The power conversion efficiencies (PCEs) of
the optimized all-PSCs based on C1 and C3 reached 1.68% and 1.92%, respectively. The improved
photovoltaic performance may be due to the clear nanoscale morphology of C3-based blend film. The
blend film exhibits favorable phase separation and well-ordered structure of polymer C3. This
Received 28th October 2014
Accepted 18th December 2014
substantially increased the electron mobility (up to 1.0 ꢀ 10ꢁ4 cm2 Vꢁ1
s
ꢁ1), which contributes to
balancing the charge transport in C3-based blend film. This is the first report on the realization of
favorable phase separation and high electron mobility through improvement of acceptor crystallinity in
all-PSCs.
DOI: 10.1039/c4ra13351j
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acceptors to replace fullerene derivatives is to improve the
phase separation and charge transport ability.14 Therefore,
Introduction
acceptor polymers with high crystallinity or aggregation
ability are highly desired, similar to fullerene derivatives,14–16
to realize high phase separation and electron mobility in all-
PSCs.
All-polymer solar cells (all-PSCs) have received much atten-
tion in recent years, as they provide a facile and low-cost way
to utilize solar energy.1–4 p-Conjugated polymers as acceptor
materials in all-PSCs have potential advantages. In compar-
ison with fullerene derivatives, p-conjugated polymers can be
easily chemically synthesized and modied. They have broad
absorption spectra and high absorption coefficients, and
possess ne-tuned energy levels to match the donor poly-
mer.4–7 However, all-PSCs are still being advanced slowly with
low power conversation efficiency (PCE) with, for example,
performance based on backbones of naphthalene diimide
(5.7%),8 perylene diimide (4.4%),9 uorene (2.0%),7,10 and
carbazol (3.68%).11 The relatively low efficiency has been
considered to be due to the unfavorable morphology of bulk
heterojunction (BHJ) blend lms such as weak phase sepa-
ration12 and reduced ordering of polymer chains.13 A signif-
icant challenge for p-conjugated polymers as alternative
From the viewpoint of molecular design, polymer crystal-
linity is signicantly related to the planarization of polymer
backbones which can be achieved by modifying the molec-
ular structure of p-functional backbones and the pattern of
alkyl side chains.17–23 The latter has been widely accepted to
impact on the morphology (crystallization, p-stacking and
packing orientation), charge transport and photovoltaic
performance of devices.24–27 Takimiya's group found that
different linear and branched alkyl side chains altered the
polymer orientation, thus affecting the PCEs of devices.28 It
should be mentioned that previous investigations were
limited to the study of donor polymer alkyl chains, including
alkyl chain length and branching point;29–35 however, the
effect of alkyl-chain branching position in acceptor polymers
on the performance has not been reported for all-PSCs. Detail
studies on the critical role of alkyl chains in photovoltaic
performance of all-PSCs are thus highly desired. Further-
more, the manipulation of alkyl-chain branching position
should be an important strategy to design polymers with
good crystallinity and achieve high performance as acceptors
in all-PSCs.
aBeijing National Laboratory for Molecular Sciences (BNLMS), Institute of Chemistry,
Chinese Academy of Sciences, 100190, Beijing, P. R. China. E-mail: hongbing.fu@
iccas.ac.cn; chunlgu@iccas.ac.cn
bUniversity Chinese Academy of Sciences, Beijing 100049, P. R. China
† Electronic supplementary information (ESI) available: Details of for the
measurement of cyclic voltammogram and charge mobilities, TGA images and
AFM images. See DOI: 10.1039/c4ra13351j
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Due to the highly electron-decient character of naphthalene
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Experimental section
diimide (NDI),36,37 NDI-based derivatives possess deep low
unoccupied molecular orbital (LUMO) energy levels, which
make them good electron transporting semiconductors and
suitable acceptor materials for PSCs. As a result, NDI-based
copolymers have attracted much attention in the elds of
PSCs and organic eld effect transistors (OFETs).4,37–43 In
previous studies, a typical NDI-bithiophene copolymer (PNDI-
TT) with good aggregation ability was initially proved to
exhibit low PCEs of 0.2–1.4% in all-PSCs.42–46 Very recently, Ito's
group reported high PCE of above 4% based on the same
polymer as acceptor via signicantly enhancing charge-carrier
generation and collection efficiencies.8,47 At the same time,
Liu et al. reported that the introduction of vinyl groups into the
copolymer backbone of PNDI-TT could improve the charge
mobility of neat lm in OFETs.38 It is proved that vinyl groups
can extend the polymer coplanarity as well as the intermolecular
p–p stacking, which would enhance the polymer crystallinity
with negligible inuence on the low-lying LUMO energy level.
Herein, we focus on the manipulation of the branching position
of alkyl chains to improve polymer crystallinity, which would be
benecial to increase the PCE via favorable phase separation in
all-PSCs.
Materials
The polymers PBDTTT-C-T were purchased from Solarmer
Materials Inc. and used as received. The number-average
molecular weight (Mn) and polydispersity index (PDI; given by
Mw/Mn, where Mw is the weight-average molecular weight)
provided on the Certicate of Analysis, were 54.9 kg molꢁ1 and
1.9, respectively. 2-(Tributylstannyl)thiophene, NBS, and trans-
1,2-bis(tributylstannyl)ethene were purchased from Sigma-
Aldrich. Compounds 2-octyldodecan-1-amine (2a) and 4-
octyltetradecan-1-amine (2b) were purchased from LYN (Bei-
jing) Technology Co. Ltd. Compound 2,6-dibromo-1,4,5,8-
naphthalene tetracarboxylic dianhydride (1) was synthesized
according to the literature.36,37 Toluene was dried over Na/
benzophenoneketyl and freshly distilled before use. The other
reagents and solvents used in this work were commercially
purchased and used without further purication. All air- and
water-sensitive reactions were performed under nitrogen
atmosphere. The detailed synthetic processes of the other
compounds are as follows.
Synthetic procedures
In this study, we successfully designed and synthesized
polymers C1 and C3, where Cm refers to the number of linear
N,N0-Bis(4-octyltetradecyl)-2,6-dibromo-1,4,5,8-naphthalene
carbon atoms between the alkyl-chain branching position diimide (3b). A mixture of 1 (2.12 g, 5.0 mmol), 4-octylte-
and the alkylated nitrogen (Fig. 1). A systematic study of their tradecylamineꢂ(4.07 g, 12.5 mmol), and acetic acid (100 mL) was
optical, electronic, and conformation properties, and the stirred at 110 C for 12 h. Aer cooling to room temperature,
characteristics of all-PSCs based on these polymers are most of the solvent was removed under vacuum, and the residue
described. For comparision, a detailed structural analysis of was puried by column chromatography on silica gel with a
these polymers was performed using grazing incidence X-ray mixture of ethyl acetate/petroleum ether (1 : 60, v/v) as eluent,
diffraction (GIXD) and transmission electron microscopy affording a slightly yellow solid as the product (1.30 g, 1.25
1
(TEM) to elucidate the effect of the branching position on mmol, yield 25%). H NMR (CDCl3, 400 MHz) d: 8.98 (s, 2H),
polymer crystallinity and morphology of blend lms. It was 4.16 (d, J ¼ 7.6 Hz, 4H), 1.70 (m, 4H), 1.32–1.41 (m, 6H), 1.15–
found that the morphologies of all-polymer BHJ lms 1.35 (m, 64H), 0.85–0.89 (m, 12H); 13C NMR (CDCl3, 400 MHz) d:
strongly depend on the alkyl-chain branching position. 160.60, 138.97, 128.27, 127.60, 125.28, 124.01, 77.45, 77.03,
Better device performance was obtained from polymer C3 76.61, 41.49, 37.22, 33.54, 31.93, 30.96, 30.10, 29.71, 29.66,
with superior crystallinity where a PCE of 1.92% was obtained 29.37, 26.67, 25.15, 22.69, 14.13; MS-TOF [M]+: 1040.5 (m/z).
under AM 1.5G illumination with 3% 1,8-diiodooctane (DIO)
N,N0-Bis(2-octyldodecyl)-2,6-dibromo-1,4,5,8-naphthalene
as processing additive. The improved performance arises diimide (3a). A mixture of 1 (1.92 g, 4.5 mmol), 2-octyldodecan-
from the clear nanoscale morphology with favorable phase 1-amine (3.35 ꢂg, 11.25 mmol), and acetic acid (90 mL) was
separation and well-ordered domains. Our results provide an stirred at 110 C for 12 h. A slightly yellow solid (1.37 g, 1.40
example for rational design criteria of polymer acceptors to mmol, yield 31%) was obtained using the same processing
realize nanoscale phase separation morphology and high method as for compound 3a. 1H NMR (CDCl3, 400 MHz) d: 8.95
electron mobility.
(s, 2H), 4.12 (d, J ¼ 7.5 Hz, 4H), 1.97 (m, 2H), 1.20–1.40 (m, 64H),
0.84–0.89 (m, 12H); MS-TOF [M]+: 984.5 (m/z).
N,N0-Bis(4-octyltetradecyl)-2,6-bis([2,20]-bithiophenyl-5-yl)-
1,4,5,8-naphthalene diimide (4b). Compound 3b (1.04 g, 1.0
mmol), 5-tributylstannyl-2,20-bithiophene (0.94 g, 2.5 mmol),
Pd(PPh3)2Cl2 (98 mg, 0.1 mmol) and 20 mL of anhydrous
toluene were added to a ask under argon. The mixture was
reuxed for 16 h and then cooled down to room temperature.
Aer quenching with CaF2$12H2O, the solvent was removed
under vacuum and the product was puried by column ash
chromatography using an ethyl acetate/petroleum ether solvent
mixture (10 : 1). The pure product was obtained as a red solid
Fig. 1 Chemical structures of polymers C1, C3 and donor polymer
PBDTTT-C-T.
1
(943 mg, 0.9 mmol, yield 90%). H NMR (CDCl3, 400 MHz) d:
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8.77 (s, 2H), 7.78 (d, 4H), 7.31 (d, 2H), 7.21 (d, 2H), 4.10 (d, J ¼ (100 mL), and ltered off to afford a dark blue solid (77.5 mg,
7.6 Hz, 4H), 1.70 (m, 4H), 1.32–1.41 (m, 6H), 1.15–1.35 (m, 64H), yield 61%).
0.85–0.89 (m, 12H); 13C NMR (CDCl3, 400 MHz) d: 162.13,
Poly[N,N0-bis(2-octyldodecyl)-2,6-bis(thiophenyl-5-yl)-1,4,5,8-
161.96, 140.73, 140.28, 136.62, 128.85, 127.94, 127.46, 127.39, naphthalene diimide-2,20-diyl]-alt-trans-ethylene (C1). Compound
125.43, 123.47, 77.34, 77.03, 76.71, 41.48, 37.31, 33.58, 31.92, 5a (229.8 mg, 0.20 mmol), trans-1,2-bis(tributylstannyl)ethane (13.2
31.01, 30.11, 29.69, 29.65, 29.35, 26.68, 25.34, 22.68, 14.10; MS- mg, 0.20 mol), Pd2(dba)3 (18.2 mg, 0.02 mol), P(o-tol)3 (24.2 mg, 0.08
TOF [M]+: 1046.8 (m/z).
mol) and 5 mL of toluene were added to a Schlenk tube. The same
N,N0-Bis(2-octyldodecyl)-2,6-bis([2,20]-bithiophenyl-5-yl)- processing method as for polymer C3 was used during polymeri-
1,4,5,8-naphthalene diimide (4a). Compound 3a (1.37 g, 1.4 zation and purication. Aer Soxhlet extraction, the product was
mmol), 5-tributylstannyl-2,20-bithiophene (1.32 g, 3.5 mmol), nally collected with chlorobenzene. Dark blue solid, 147.4 mg,
Pd(PPh3)2Cl2 (137.2 mg, 0.14 mmol) and 20 mL of anhydrous yield 58%.
toluene were added to a ask under argon. A red solid (1.24 g,
1.25 mmol, yield 89%) was obtained with the same processing
method as for compound 4a. 1H NMR (CDCl3, 400 MHz) d: 8.74
(s, 2H), 7.54 (d, J ¼ 5.1 Hz, 2H), 7.28 (d, J ¼ 3.6 Hz, 2H), 7.18 (m,
2H), 4.05 (d, J ¼ 7.2 Hz, 4H), 1.93 (m, 2H), 1.23 (m, 64H), 0.86
(m, 12H); MS-TOF [M]+: 990.8 (m/z).
Measurements and characterization
1H NMR spectra and 13C NMR spectra were measured with a
Bruker DMX-400 spectrometer in CDCl3 (1H, d: 7.26 ppm; 13C, d:
77 ppm) with tetramethylsilane (TMS) as the internal standard.
Gel permeation chromatography (GPC) was performed with a
Polymer Laboratories PL 220 at 150 ꢂC using 1,2,4-tricholor-
obenzene as eluent. The obtained molecular weight was cali-
brated with polystyrene standards. Thermogravimetric analysis
(TGA) was done using a TA Instruments NETZSCH STA 409 PC/
PG analyzer operated at a heating rate of 10 ꢂC minꢁ1 and under
a nitrogen atmosphere. UV-visible (UV-vis) absorption spectra
were obtained with a Shimidazu UV-3600 UV-vis-NIR spectro-
photometer. The structural properties of polymers were deter-
mined by GIXD measurements on thin lms produced by
dropping 1,2-dichlorobenzene (DCB) solutions of polymer onto
Si substrate. The thickness of the functional layers was
measured using a KLA-T encore P-6 prolometer (Tencor). The
morphology and surface roughness of the layers were deter-
mined by atomic-force microscopy (AFM; Digital Instruments
Nano Scope III). TEM images were obtained with a JEM-2100
transmission electron microscope operated at an acceleration
voltage of 200 kV. The sample used in AFM and TEM
measurement was from the naked active layer in devices. Cyclic
voltammetry measurements were performed using a Zahner
IM6e electrochemical workstation with a three-electrode system
in a solution of 0.1 M Bu4NPF6 in anhydrous acetonitrile at a
N,N0-Bis(4-octyltetradecyl)-2,6-bis(5-bromothiophenyl-5-yl)-
1,4,5,8-naphthalene diimide (5b). Compound 4b (524 mg, 0.5
mmol) was dissolved in 25 mL of chloroform. NBS (214 mg, 1.2
mmol) was added to the reaction mixtures. Aer stirring in the
dark for two days, the product was extracted with dichloro-
methane and washed with brine, and the organic layer was
evaporated. The red solid was puried by column chromatog-
raphy using an ethyl acetate/petroleum ether (1 : 50) solvent
mixture to give the nal product as a red solid (493 mg, 0.41
1
mmol, yield 82%). H NMR (CDCl3, 400 MHz) d: 8.72 (s, 2H),
7.15 (d, J ¼ 3.9 Hz, 2H), 7.08 (d, J ¼ 3.9 Hz, 2H), 4.10 (d, J ¼ 5.6
Hz, 4H), 1.55 (m, 4H), 1.24 (m, 64H), 0.85–0.87 (m, 12H); 13C
NMR (CDCl3, 400 MHz) d: 161.89, 142.05, 139.11, 136.44,
130.18, 128.85, 127.50, 125.63, 123.21, 115.50, 77.47, 77.05,
76.62, 41.57, 37.27, 33.55, 31.94, 30.99, 30.12, 29.71, 29.67,
29.37, 26.67, 25.32, 22.70, 14.14; MS-TOF [M]+: 1205.4 (m/z).
N,N0-Bis(2-octyldodecyl)-2,6-bis(5-bromothiophenyl-5-yl)-
1,4,5,8-naphthalene diimide (5a). Compound 4a (1.09 g, 1.1
mmol) was dissolved in 25 mL of chloroform. NBS (235.0 mg,
1.32 mmol) was added to the reaction mixture. A red solid (1.08
g, 0.935 mmol, yield 85%) was obtained with the same pro-
cessing method as for compound 5a. 1H NMR (CDCl3, 400 MHz)
d: 8.77 (s, 2H), 7.44 (d, J ¼ 3.9 Hz, 2H), 7.08 (d, J ¼ 3.9 Hz, 2H),
4.07 (d, J ¼ 7.2 Hz, 4H), 1.94 (m, 2H), 1.24 (m, 64H), 0.85 (m,
12H); MS-TOF [M]+: 1148.4 (m/z).
scan rate of 50 mV sꢁ1
.
PSC device fabrication and measurement
Preparation of poly[N,N0-bis(4-octyltetradecyl)-2,6-bis([2,20]- The PSC devices were fabricated with a structure of ITO/
bithiophenyl-5-yl)-1,4,5,8-naphthalene diimide-50,5%-diyl]-alt- PEDOT:PSS/active layer/Ca/Al, where the active layer was
trans-ethylene (C3). Compound 5b (120.5 mg, 0.1 mmol), trans- composed of C1 or C3 as acceptor and polymer PBDTTT-C-T as
1,2-bis(tributylstannyl)ethane (6.6 mg, 0.1 mol), Pd2(dba)3 (9.2 donor. Patterned ITO-glass substrates were used as the anode in
mg, 0.01 mol), P(o-tol)3 (12.2 mg, 0.04 mol) and 3 mL of toluene PSCs. ITO substrates were rstly pre-cleaned sequentially by
were added to a Schlenk tube. The tube was charged with sonicating in a detergent bath, deionized water, acetone, and
nitrogen through a freeze–pump–thaw cycle for three times. The isopropanol each for 20 min, and then dried in an oven for 30
mixture was stirred for 48 h at 110 ꢂC. The mixture was cooled to min, and lastly cleaned in a UV/ozone cleaner for 60 min. A thin
room temperature and precipitated into methanol (100 mL). layer of poly(3,4-ethylenedioxythiophene) (PEDOT):poly(styrene
The precipitate was ltered through a nylon lter, puried via sulfonate) (PSS) was spin-cast onto the ITO surface. Aer being
ꢂ
Soxhlet extraction for 12 h with methanol, acetone, hexane, baked at 150 C for 20 min in ambient air, the substrates were
chloroform, and chlorobenzene, and nally collected with transferred into a nitrogen-lled glove box (<0.1 ppm O2 and
dichlorobenzene. The dichlorobenzene solution was then H2O). C1 or C3 was blended with polymer PBDTTT-C-T and
concentrated by evaporation, precipitated into methanol dissolved in DCB. The active layers were then spin-cast from the
blend solutions (20 mg mLꢁ1, polymer weight concentration)
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with various rotational speeds on the ITO/PEDOT:PSS onset temperatures (Td) of C1 and C3 with 5% weight loss are
substrates. A small amount of DIO was used as an additive to over 430 ꢂC.
improve the photovoltaic performance of the devices if neces-
sary. Beside this, no further treatment to the blend lm was Optoelectronic and electrochemical properties
performed before the deposition of the metal electrode. Finally,
The optical properties of C1 and C3 were measured in toluene
solution and as thin lms (Fig. 2), and the detailed absorption
data are collected in Table 1. Those polymers show similar
broad absorption spectral features across the entire UV-vis
a Ca/Al metal top electrode was deposited under vacuum onto
the active layer at a pressure of ca. 5 ꢀ 10ꢁ5 Pa. The active area
of the device was ca. 6.0 mm2. The current density–voltage (J–V)
characteristics were measured with a computer-controlled
region in toluene solution. As for polymers comprised of
Keithley 2400 Source Measure Unit. A xenon lamp coupled
donor and acceptor units, the absorption spectrum is charac-
with an AM 1.5 solar spectrum lter was used as the light
terized by two spectral features, a high-energy peak and a broad
low-energy band.41,51 The high-energy peak located at approxi-
mately 400 nm can be attributed to the p–p* transition in the
polymeric backbone. The low-energy band with major absorp-
source, and the optical power was 100 mW cmꢁ2. The EQE
measurements of the encapsulated devices were performed in
air (PV Measurements Inc., model QEX7). The hole and electron
mobilities were calculated by tting the dark J–V curves for the
tion peaks at 678 nm for C1 and 684 nm for C3 is ascribed to the
hole-only and electron-only devices to the space-charge-limited-
charge-transfer transition between electron-donating ethylene
current (SCLC) model at low voltages.
units and electron-decient NDI units. In the solid state, the
main absorption bands become broader, and red shis occur of
the maximum absorption peaks at 716 nm and 726 nm by 38
nm and 42 nm, relative to those in solution, for C1 and C3
respectively. These large red shis from solution to thin lm are
indicative of a planarization of the conjugated backbone and
Results and discussion
Synthesis and characterization
The synthetic routes of the monomers and corresponding strong interchain p–p stacking in the solid state. Notably, the
polymers are outlined in Scheme 1. Polymers C1 and C3 were absorption spectrum of C3 lm exhibits a sharp shoulder peak
synthesized by Stille coupling reaction between the dibromide at 805 nm while that of C1 lm does not, which may originate
monomers (3a and 3b) and the trans-1,2-bis(tributylstannyl) from more efficient stacking in polymer C3. The absorption
ethene using Pd2(dba)3 as the catalyst. The obtained polymers edges (lonset) of the two polymer lms are 890 nm for C1 and 895
were precipitated in methanol and then puried by Soxhlet nm for C3, from which the optical band gaps (Eogpt) of the
extraction to remove oligomers and other impurities. The polymers can be calculated according to Eogpt ¼ 1240/lonset
.
number-average molecular weights (Mn) and polydispersity Eogpt values of C1 and C3 are 1.40 and 1.38 eV, respectively.
indices (PDIs) were 23.0 kDa, 3.5 and 74.5 kDa, 2.7 for C1 and
Cyclic voltammetry measurements were carried out to eval-
ꢂ
C3 respectively, evaluated by high-temperature GPC at 150 C uate the electrochemical properties and estimate the HOMO/
using 1,2,4-tricholorobenzene as eluent. The two polymers had LUMO energy levels of the two polymers (Fig. S1†). Our results
molecular weights of more than 20 kg molꢁ1, which is consid- are similar to the data of previous work by Liu et al.38 As we
ered to be quite high for low bandgap polymers.48–50 In our work, expected, the changes of alkyl-chain branching position have a
the effect of molecular weight of the two polymers is a negligible negligible effect on polymer energy levels.28–31,35 As listed in
factor according to the similar performance of BHJ solar cells in Table 1 and shown in Fig. 2b, the HOMO/LUMO energy values
unoptimized devices. The TGA measurements showed that the of C1 and C3 are ꢁ5.61/ꢁ3.93 eV and ꢁ5.61/ꢁ3.94 eV. The
Scheme 1 Synthesis routes of the monomers and conjugated polymers C1 and C3.
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Fig. 2 (a) Normalized UV-vis absorption spectra of C1 and C3 in toluene solution and as thin films. (b) HOMO and LUMO energy levels of the
polymers and PC70BM.
energy levels of the polymers are well matched with those of T : polymer was 1 : 1 (w/w). Fig. 4a displays the current density–
polymer PBDTTT-C-T52 and the LUMO energies are signicantly voltage (J–V) curves of the devices under illumination of AM
higher-lying than that of PC70BM (Fig. 2b); thus, a large Voc 1.5G simulated white light with an intensity of 100 mW cmꢁ2
.
could be expected to occur in the corresponding devices.
The Voc, short circuit current density (Jsc), ll factor (FF) and
PCE of the devices are shown in Table 2 and Fig. 4. The other
device data are summarized in the ESI.†
GIXD characterization
Devices fabricated from a blend of C1 with PBDTTT-C-T gave
a Jsc of 3.64 mA cmꢁ2, a Voc of 0.81 V and a FF of 0.49, corre-
sponding to a PCE of 1.21%. When the linear chain increased to
three carbon atoms in C3, devices showed a slight improvement
in the photovoltaic performance, with a PCE of 1.27%, which
indirectly demonstrated that molecular weight has little effect
on photovoltaic performance. In order to further improve the
PCE, DIO was used as a processing additive.54–56 When 1% DIO
was added, the photovoltaic performance based on C1 was
further improved to 1.68%, with a Jsc of 3.64 mA cmꢁ2; the
highest PCE of 1.92% was achieved by adding 3% DIO to
To study the effect of alkyl-chain branching position of the
polymers on the packing orientation and crystallinity of lms,
GIXD was performed. As seen in Fig. 3 and S3,† an edge-on
orientation was found for both thin lms, with evident (100)
diffraction peaks in the out-of-plane direction. The d100 spac-
ings corresponding to out-of-plane space distance between
˚
˚
polymer backbones are 23.2 A for C1 and 26.0 A for C3. This
kind of spacing related to the side-chain length is much shorter
than that of fully extended alkyl chains, indicating that side
chains of polymer in adjacent layers were closely interdigitated
in lamellar structures.23,53 Noticeably, an obvious (200) diffrac-
tion at 6.80ꢂ and (001) diffraction halos of alkyl chains38 at 5.68ꢂ
for polymer C3 demonstrate that polymer C3 possesses superior
crystallinity ability compared with polymer C1. Therefore,
moving the branching position away from the conjugated
backbone may reduce the steric effect of bulk octyldodecyl
chain and promote the ordering of lamellar structure of the
polymer.
PBDTTT-C-T/C3 devices, with a highest Jsc of 4.70 mA cmꢁ2
.
However, further increasing the ratio of DIO resulted in a
gradual decrease in photovoltaic performance for both C1 and
C3 (Table S1†). Charge mobilities of the optimal blend lms
based on polymers C1 and C3 were measured using the charge-
only SCLC method as shown in Fig. S4.† Hole and electron
mobilities of PBDTTT-C-T : polymer BHJ ꢁd1evices were estimated
to be, respectively, 1.3 ꢀ 10ꢁ4 cm2 Vꢁ1
s
and 6.6 ꢀ 10ꢁ5 cm2
Vꢁ1 sꢁ1 for C1, and 1.8 ꢀ 10ꢁ4 cm2 Vꢁ1 sꢁ1 and 1.0 ꢀ 10ꢁ4 cm2
Vꢁ1 sꢁ1 for C3. It is noteworthy that electron mobility in C3-
based active layers is much higher than that in layers based on
C1, and equal to the hole mobility value. The balanced hole and
electron motilities in the C3-based blend lm can largely
explain the high performance of the BHJ solar cells.2 On account
Photovoltaic properties
All-PSCs were fabricated by employing C1 or C3 as acceptor
material and PBDTTT-C-T as donor material (Fig. 1), with a
conventional device structure of ITO/PEDOT:PSS/PBDTTT-C-
T : polymer/Ca/Al. The best weight ratio of PBDTTT-C-
Table 1 Basic properties of polymers (molecular weights, thermostability, optical and electrochemical properties)
Cn
Mn (kDa)
Mwa (kDa)
PDIa
lmax (nm)
lmaxd (nm)
lonsetd (nm)
Eogpte (eV)
EHOMO (eV)
ELUMO (eV)
a
b
Td (ꢂC)
c
C1
C3
23.0
74.5
79.9
197.8
3.5
2.7
433
445
395, 678
396, 684
403, 716
405, 726, 805
890
895
1.40
1.38
ꢁ5.61
ꢁ5.61
ꢁ3.93
ꢁ3.94
a
Determined by GPC in trichlorobenzene at 150 ꢂC using polystyrene standards. b Decomposition temperature, determined by TGA in nitrogen at a
heating rate of 10 ꢂC minꢁ1, based on 5% weight loss. c Measured in dilute toluene solution. d Spin-cast from 4 mg mLꢁ1 dichlorobenzene solution.
e
Optical band gap calculated from the onset of lm absorption.
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of the very close Voc and FF, improvement of the efficiency is
mainly from the enhanced Jsc and high electron mobility, which
is probably linked to the optimized morphology of the blend
lm (vide infra).
Moreover, improvement of Jsc is supported by the external
quantum efficiency (EQE) responses (Fig. 4b), in which the
addition of DIO yields an increase of the EQE response. As for
the PBDTTT-C-T/C1 device, 1% DIO addition greatly improved
the maximum monochromic EQE from 13.4% to 18.0 % at 720
nm. Aer further manipulating the alkyl-chain branching
position, the maximum value reached 19.6% at 640 nm under
the optimal condition, consistent with the highest Jsc.
Fig. 3 2D GIXS patterns of C1 (a) and C3 (b) thin films.
Morphology characterization
TEM and AFM were used to study the phase separation in blend
lms (processed with DCB or DCB : DIO solvent). TEM images
of the blend lms are shown in Fig. 5, and the corresponding
AFM images are provided in Fig. S3.† In TEM images, the dark
regions could be attributed to the acceptor (C1 or C3) domains,
while the bright regions could be attributed to the donor
(PBDTTT-C-T) domains (depicted in Fig. S4†). When no DIO is
added, blend lms based on C1 and C3 show both dark
aggregation domains and small phase separation, which indi-
cates that the poor performances for polymers C1 and C3 are
related to the non-optimized morphologies. The PBDTTT-C-T/
C3 blend lm exhibits more and bigger aggregation domains
(ca. 15 nm) than the PBDTTT-C-T/C1 blend lm. The
Fig. 4 The current density (J)–voltage (V) curves (a) and the typical
EQE curves (b) of the devices under different conditions.
Fig. 5 TEM images of blend films: PBDTTT-C-T:C1 with (a) 0% DIO, (b)
1% DIO and (c) 2% DIO; PBDTTT-C-T:C3 with (d) 0% DIO, (e) 3% DIO
and (f) 4% DIO. The scale bars are 200 nm.
Table 2 Photovoltaic performances of all-PSC devices based on the two polymers with DCB as solvent and DIO as processing additive
Cna
C1
Solvent (DCB : DIO)b
Jsc (mA cmꢁ2
)
Voc (V)
FFc (%)
PCEc (%) (bestd)
c
c
100 : 0
99 : 1
100 : 0
97 : 3
2.88 ꢃ 0.13
3.56 ꢃ 0.04
3.17 ꢃ 0.09
4.75 ꢃ 0.11
0.81 ꢃ 0.01
0.81 ꢃ 0.00
0.78 ꢃ 0.01
0.79 ꢃ 0.01
47.63 ꢃ 1.65
55.04 ꢃ 1.91
50.65 ꢃ 0.60
49.29 ꢃ 1.33
1.11 ꢃ 0.08 (1.21)
1.59 ꢃ 0.06 (1.68)
1.24 ꢃ 0.03 (1.27)
1.85 ꢃ 0.05 (1.92)
C3
a
Cn : donor ¼ 1 : 1 (w/w). b DCB : DIO (v/v). c Only the optimized recipes were considered for the estimation of the average PCE; data have been
averaged for different devices (see ESI). The performance of the best device is given in parentheses. The active area of the device was 0.06 cm2.
d
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aggregation domains may originate from the better crystallinity, blend lm shows much clearer diffraction arcs. This indicates
while polymer PBDTTT-C-T is non-crystalline.52,57 The small that polymer C3 in the BHJ blend is more highly ordered than
improvement of PCE for polymers C1 and C3 indicates that polymer C1, which well matches with the crystallinity of the two
polymer crystallinity cannot be the sole factor inuencing the polymers. The thin lms processed from DCB and DCB : DIO
performance of devices. Other factors, like the phase domain mixed solvent yield similar results, which suggests that DIO is
sizes and ordering of polymer aggregation in the BHJ blends, benecial to the aggregation of the two polymers. When the lm
may also be important because the exciton separation and is cast from DCB : DIO (97 : 3, v/v), several arcs appeared in
charge mobility should be efficient across the entire active layer. Fig. 6d, demonstrating that polymer C3 distributes with highly
It is well known that addition of the small-molecule additive ordered structure in the BHJ blend lm, which enhances the
DIO into the solvent would optimize the blend morphology in electron transport ability. The clear diffraction arcs in Fig. 6d
BHJ polymer solar cells.56,58–60 Aer 1% DIO is added, PBDTTT- also provide additional evidence for the clear nanoscale phase
C-T/C1 blend lm exhibits a continuous interpenetrating separation and aggregation of acceptor C3. Consequently,
network with nanoscale phase segregation. Simultaneously, moving the branching position further from the main chain
PBDTTT-C-T/C3 blend lm exhibits clear nanoscale phase could promote the performance of devices by realizing a
separation and many dark spots (ca. 20 nm in size), which nanoscale morphology and highly ordered structure of the
relates to the solubility and crystallization of polymer C3. In all- polymer.
polymer systems, the polymer crystallinity would be an impor-
tant driving force to form the acceptor domains in this work.
The slow rate of evaporation of DIO provides sufficient time for Conclusions
polymer C3 to aggregate gradually and form a nanoscale phase
In conclusion, two polymers based on NDI with different alkyl-
chain branching positions, C1 and C3, were synthesized and
characterized. It was proved out that the location of the
branching point had a signicant effect on optical absorption,
conjugated backbone orientation, crystallinity, device
morphology and charge mobility. Moving the branching posi-
tion away from the backbone resulted in superior crystallinity of
C3. From C1 to C3, the morphology in an optimal device showed
separation.56,61 The nanoscale phase separation provides large
donor/acceptor interfaces and efficient channels for charge
transport, thus leading to high Jsc and charge mobility, espe-
cially electron mobility. Moreover, the observation in TEM
images can also be conrmed by AFM measurements (see
Fig. S3†). These results demonstrate that the alkyl-chain
branching position signicantly inuences the nanoscale
morphology of blend lms by improving polymer crystallinity. A
improved nanoscale morphology and high ordering structure of
polymer with superior crystallinity could easily aggregate in
the polymer. Hence, all-PSCs based on C3 showed a higher PCE
of 1.92%, whereas it was 1.68% for C1. Therefore, manipulation
of alkyl-chain branching position provides new insights into the
design of optimum polymers as acceptors for high performance
blends, thus forming a nanoscale morphology.
GIXS was used to provide the structural ordering of the two
polymers in blend lms (processed with DCB or DCB:DIO mixed
solvent) as shown in Fig. 6. The PBDTTT-C-T:C1 blend lm
shows ambiguous diffraction arcs; however, the PBDTTT-C-T:C3
in all-PSCs.
Acknowledgements
The GIXD data were obtained at 1W1A, Beijing Synchrotron
Radiation Facility. The authors gratefully acknowledge the
assistance of scientists of the Diffuse X-ray Scattering Station
during the experiments. This work was supported by the
National Natural Science Foundation of China (no. 20925309,
21190034, 21221002, 21203212) and the National Research
Fund for Fundamental Key Project 973 (2011CB808402,
2013CB933500).
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