2,1,3-Benzothiadiazole-5,6-dicarboxylicimide-Based Polymer

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2,1,3-Benzothiadiazole-5,6-Dicarboxylicimide-Based Polymer Semiconductors for Organic Thin-Film Transistors and Polymer Solar Cells Jianwei Yu, Joshua Loroña Ornelas, Yumin Tang, Mohammad Afsar Uddin, Han Guo, Simiao Yu, Yulun Wang, Han Young Woo, Shiming Zhang, Guichuan Xing, Xugang Guo, and Wei Huang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b11863 • Publication Date (Web): 13 Nov 2017 Downloaded from http://pubs.acs.org on November 14, 2017

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ACS Applied Materials & Interfaces

2,1,3-Benzothiadiazole-5,6-Dicarboxylicimide-Based Polymer Semiconductors for Organic Thin-Film Transistors and Polymer Solar Cells Jianwei Yu,†,‡ + Joshua Loroña Ornelas,‡,§, + Yumin Tang,‡ Mohammad Afsar Uddin,# Han Guo,‡ Simiao Yu,† Yulun Wang,‡ Han Young Woo,# Shiming Zhang,*,† Guichuan Xing,*,& Xugang Guo*,‡ and Wei Huang†,∥ †

Key Laboratory of Flexible Electronics (KLOFE) & Institute of Advanced Materials

(IAM), Jiangsu National Synergetic Innovation Center for Advanced Materials (SICAM), Nanjing Tech University (NanjingTech), 30 South Puzhu Road, Nanjing 211816, China ‡

Department of Materials Science and Engineering and The Shenzhen Key

Laboratory for Printed Organic Electronics, South University of Science and Technology of China, No. 1088, Xueyuan Road, Shenzhen, Guangdong 518055, China §

Department of Chemistry

and

Pharmacy,

Friedrich-Alexander-University

Erlangen-Nürnberg (FAU), Egerlandstr. 3, 91058 Erlangen, Germany #

Research Institute for Natural Sciences, Department of Chemistry, Korea

University, Seoul 136-713, South Korea & Institute of Applied Physics and Materials Engineering, University of Macau, Macao SAR 999078, China ∥

Shaanxi Institute of Flexible Electronics (SIFE), Northwestern Polytechnical

University (NPU), 127 West Youyi Road, Xi'an 710072, China +

The two authors contributed equally to this work

Corresponding Author * [email protected] (S.Z.)

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* [email protected] (G.X.) * [email protected] (X.G.) KEYWORDS: conjugated polymer, organic thin-film transistors, organic solar cells, ambipolar, energy loss

ABSTRACT:

A

series

of

polymer

semiconductors

incorporating

2,1,3-benzothiadiazole-5,6-dicarboxylicimide (BTZI) as strong electron withdrawing unit and alkoxy-functionalized head-to-head linkage containing bithiophene or bithiazole as highly electron rich co-unit are designed and synthesized. Due to the strong intramolecular charge transfer characteristics, all three polymers BTZI-TRTOR (P1), BTZI-BTOR (P2), and BTZI-BTzOR (P3) exhibit narrow bandgaps of 1.13, 1.05, and 0.92 eV, respectively, resulting in a very broad absorption covering from 350

nm

up

to

1400

nm.

The

highly

electron-deficient

2,1,3-benzothiadiazole-5,6-dicarboxylicimide and alkoxy-functionalized bithiophene (or thiazole) lead to polymers with low-lying LUMOs (−3.96 to −4.28 eV) and high-lying HOMOs (−5.01 to −5.20 eV). Hence P1 and P3 show substantial and balanced ambipolar transport with electron mobilities/hole mobilities of up to 0.86/0.51 and 0.95/0.50 cm2 V-1 s-1, respectively, and polymer P2 containing the strongest donor unit exhibit unipolar p-type performance with an average hole mobility of 0.40 cm2 V-1s-1 in top-gate/bottom-contact thin-film transistors with gold as the source and drain electrodes. When incorporated into bulk heterojunction polymer solar cells, the narrow bandgap (1.13 eV) polymer P1 shows an encouraging power conversion efficiency of 4.15% with a relative large open-circuit voltage of 2 ACS Paragon Plus Environment

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0.69 V, which corresponds to a remarkably small energy loss of 0.44 eV. The power conversion efficiency of P1 is among the highest reported to date with such small energy loss in polymer:fullerene solar cells.

Introduction Over the past two decades, π-conjugated polymers have emerged as promising semiconductors for various opto-electronic devices, particularly organic thin-film transistors (OTFTs)1-5 and polymer solar cells (PSCs).6-9 Solution-processable polymers show several distinctive advantages over the conventional inorganic-based semiconductors, such as mechanical flexibility, light-weight, cost-effectiveness, facile large area device production in a high throughput fashion.9-14 In order to achieve improved device performance, the semiconducting materials should possess appropriate bandgaps with well-tailored frontier molecular orbitals (FMOs), e.g. highest occupied molecular orbitals (HOMOs) and lowest unoccupied molecular orbitals (LUMOs).1,

15-18

A tremendous amount of conjugated polymers with

donor-acceptor (D-A) architectures have been synthesized for achieving high charge carrier mobilities in OTFTs.19-22 To date, the high-performance polymers typically are p-type, showing large hole mobilities, and the development of ambipolar polymers with substantial electron mobility falls short compared to p-type polymers.21, 23-30 In comparison to unipolar semiconductors, ambipolar semiconductors enable the fabrication of complementary-like inverters without using advanced patterning techniques,31-32 and the fabricated circuits using this complementary approach show lower power dissipation, wider noise margins, and higher robustness, versus the 3 ACS Paragon Plus Environment


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circuits fabricated using the unipolar approach.33-34 In addition, the recombination of electron and hole within the transistor channel can lead to light emission, which is an intriguing subject in organic light-emitting transistors (OLETs).35 Therefore it is important to develop ambipolar semiconductors with large and balanced OTFT mobilities from both scientific perspective and practical application. In OTFT devices with gold (Au) as the source and drain electrodes, polymer semiconductors with HOMO and LUMO levels close to −5.0 and −4.0 eV, respectively, are beneficial to hole and electron injection from the Au electrodes.31 Donor-acceptor (D-A) type copolymers, specifically strong acceptor-strong donor polymeric semiconductors, have been regarded as excellent semiconducting materials for enhancing ambipolar OTFTs performance due to the well-hybridized FMOs and the strong intermolecular D-A interaction.36-37 The molecular orbital hybridization in the strong donor-acceptor copolymers can yield low-lying LUMOs and high-lying HOMOs, which are essential for facile electron and hole injection, regarded as an effective approach to achieve high-performance ambipolar semiconducting polymers.21, 30, 33, 38-39 In the PSCs field, strong intramolecular charge transfer between the donor and the acceptor co-units will lead to the resulting polymers with narrow bandgaps, which can broad solar absorption spectra and achieve large short-circuit current (Jsc) for PSC application.40 Scheme 1. Chemical Structures of Phthalimide (PhI), Benzothiadiazole (BTZ), and 2,1,3-Benzothiadiazole-5,6-dicarboxylicimide (BTZI).


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In this work, the 2,1,3-benzothiadiazole-5,6-dicarboxylicimide (BTZI, Scheme 1) is chosen as the strong electron-withdrawing unit which was achieved by introducing a cyclic imide into the 5 and 6 positions of the benzothiadiazole (BTZ) moiety. The attachment of the imide group leads to BTZI with a stronger electron-withdrawing capability versus the well-known acceptor BTZ.15, 41-45 Furthermore, the solubility of BTZI-based polymers can be greatly increased compared to that of polymers based on BTZ units due to the incorporation of an alkyl solubilizing group on the imide moiety in the BTZI unit. In comparison to phthalimide, the incorporation of thiadiazole affords BTZI with improved physicochemical properties. Although some BTZI-based polymers have been reported, showing good performance in PSCs with PCEs exceeding 8%,46 very few works focus on OTFTs with the highest mobility of ~4× 10-3 cm2 V-1 s-1 reported by Woo and co-workers. Due to the backbone planarity enabled by the intramolecular non-covalent (thienyl)S···(alkoxy)O interaction47-49 and the strong electron donating ability afforded by the alkoxyl side chains on the 3,3’-position,

three

alkoxy-functionalized

head-to-head


linkage

containing


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bithiophene or bithiazole donor co-units are synthesized and copolymerized with phthalimide, the resulting polymers show good charge carrier mobilities in OTFTs and encouraging PCEs in PSCs.48, 50-52 By replacing the phthalimide with the stronger electron acceptor BTZI, the resulting polymer semiconductors reported here should have appropriately-lying FMOs for more facile charge carrier injection and improved mobility in OTFTs. Therefore, BTZI is copolymerized with three strong electron donor

co-units,

3-alkyl-3’-alkoxy-2,2’-bithiophene

3,3’-dialkoxy-2,2’-bithiophene(BTOR),50

and

(TRTOR),51

4,4’-dialkoxy-5,5’-bithiazole

(BTzOR),52 to generate a series polymer semiconductors (Scheme 2), BTZI-BTRTOR (P1), BTZI-BTOR (P2), and BTZI-BTzOR (P3), for applications in OTFTs and PSCs. Due to the highly electron deficiency of BTZI and the electron rich characteristic of BTOR, P2 has a LUMO of −3.96 eV and a high-lying HOMO of −5.01 eV with a narrow optical bandgap of 1.05 eV. The absorption of P2 covers a broad range of the solar spectrum of up to 1200 nm. To achieve lower-lying FMOs, we replaced one alkoxy substituent with a less electron-donating alkyl one for P1 and substituted thiophene with a more electron-deficient thiazole for P3. The LUMO levels of P1 and P3 are decreased to −4.07 and −4.28 eV, respectively, with low-lying HOMO levels of −5.20 eV. On the basis of the absorption onsets, the optical bandgap of P1 is 1.13 eV, and polymer P3 shows the lowest bandgap of 0.92 V in the series. The relationships between molecular structures and opto-electrcial properties were investigated in detail. Among the polymers, P1 and P3 show balanced ambipolar transport with substantial electron/hole mobilities of up to 0.86/0.51 and 0.95/0.50 cm2 V-1 s-1 in OTFTs with Au


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as the source and drain electrodes, which are the highest values among polymers based on BTZI. In spite of its narrow bandgap (1.13 eV), P1 shows an encouraging PCE up to 4.1%, with a large open-circuit voltage of 0.69 V and the lowest photon energy loss of 0.44 eV in polymer:fullerene solar cells. Scheme 2. Synthetic Route to the BTZI monomer and the BTZI-Based Polymer Semiconductors.

NC

NC

CN

R N

O

CN

O

R N

O

O

Br H2N

NH2

N

1

N

N

N

Br N

N

S

S

S

2

3

4

OR1

R N

O

R2 O

SnMe3

S

S

S

n

S

Me3Sn

OR1

R2 N

N S

P1 R N

OR1 SnMe3

S

4

+

Me3Sn

O

R1O O S

S R1O

S

n

OR1 N

N S

P2 R N

OR1 N Me3Sn

SnMe3

S S

O

R1O

N

S

S

N

OR1

O

N R1O N

n

N S

1

2

P1: R= 2-ethylhexyl; R = n-heptyl; R =n-octyl P2: R= 2-ethylhexyl; R1= n-octyl P3: R= 2-ethylhexyl; R1= n-octyl

P3

Reagents and Condition: (i) SOCl2, Et3N; (ii) 2-ethyl-1-hexylamine, ZnBr2, o-DCB; (iii) Fe, Acetic acid; then Br2, SOCl2, Et3N; (iv) Pd2(dba)3, P(o-tolyl)3, Toluene. Results and discussion Materials Synthesis 7 ACS Paragon Plus Environment


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We simplified the synthetic steps of the BTZI monomer43, 45 and the detailed synthetic route is shown in Scheme 2. The three BTZI-based polymer semiconductors P1, P2, and P3, were prepared via Stille coupling-based polymerization between the brominated BTZI monomer and a series of alkoxy-functionalized donor co-units, BTRTOR, BTOR, and BTzOR, using microwaves as the heating source. After polymerization, the polymer chains were endcapped with mono-functionalized 2-(tributylstannyl)thiophene first and then 2-bromothiophene. Product polymers were collected by precipitation into methanol and were then subjected to purification via Soxhlet extractions using methanol, acetone, hexane, dichloromethane, and chloroform as the solvent sequence. After drying, the final chloroform fractions were used for the device fabrication and characterization. Due to the solubilizing chains on the BTZI and the side chains on the head-to-head bithiophene units, all polymers exhibit good solubility in common organic solvents for device fabrication. Polymer molecular weights were measured by gel permeation chromatography (GPC) at 150°C using 1,2,4-trichlorobenzene as the eluent versus polystyrene standards. The number average molecular weights (Mns) of P1, P2, and P3 are 63.6, 22.6, and 83.9 kDa with a polydispersity index (PDI: Mw/Mn) of 1.6, 2.0, and 2.1, respectively (Table 1). Table 1. Molecular Weights, Thermogravimetric Data, Optical Absorption, and Electrochemical Properties of Polymers P1, P2, and P3.

P1

Mn [kD a]a) 63.6

P2

22.6

2.0

365

890

939

−5.01

−3.49

−3.96

1.05

P3

83.9

2.1

340

726

954

−5.20

−3.68

−4.28

0.92

Polymer

a)

PDI

Td [°C]b)

1.6

391

λmax (soln)c) [nm] 663

λmax (film)d) [nm] 851

EHOMO [eV]e) −5.20

ELUMO [eV]g)

Egopt [eV]h)

−3.64

−4.07

1.13

ELUMOf) [eV]

GPC versus polystyrene standards, trichlorobenzene as the eluent at 150°C;


b)

decomposition

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temperature defined as the temperature at which weight has 5 % decrease; spectra (1 × 10-5 M in chloroform);

d)

c)

solution absorption

absorption spectra of polymer film on a quartz glass;

e)

EHOMO = −e(Eoxonset + 4.80) eV, and Eoxonset determined electrochemically using Fc/Fc+ internal standard; f) EHOMO = −e(Eredonset + 4.80) eV, and Eredonset determined electrochemically using Fc/Fc+ internal standard; g) ELUMO = EHOMO + Egopt; h) optical bandgap estimated from the absorption onset of polymer films: Egopt = 1240/λonset eV.

Polymer Thermal Properties Thermogravimetric analysis (TGA) was carried out to investigate the polymer thermal properties with a 5% mass loss defined as the thermal decomposition threshold. As shown in Figure S1, the TGA curves reveal that the onset decomposition temperatures of polymers P1, P2, and P3 under N2 are 391, 365, and 340°C, respectively, and hence all BTZI-based polymer semiconductors show sufficient thermal stability for device fabrication and optimization over a broad range of annealing temperatures. The data indicates that monoalkoxy-functionalized P1 has the highest thermal stability in the polymer series. Differential thermal calorimetry (DSC) reveals no distinctive peaks associated with crystallization/melting transitions of these BTZI-based polymers, which is similar to the phthalimide-based polymer semiconductors. The lack of distinctive transition peaks indicates that the crystallinity of these BTZI-based polymers are not high, which is likely due to a certain degree of torsion angles (24-27°) between the BTZI unit with the neighboring thiophene moiety.



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Figure 1. (a) Optical absorption spectra of polymer semiconductors P1-P3; (b) cyclic voltammograms of P1-P3 films in 0.1 M (n-Bu)4N·PF6 acetonitrile solution, Fc/Fc+ redox couple was used as an external standard having an oxidation potential of +0.42 V vs SCE. Polymer Optical Properties The UV−vis absorption spectra of BTZI-based polymers P1, P2, and P3 in diluted chloroform (CF) solutions and as thin-films (spin-coated from 5 mg/mL CF solutions) are shown in Figures 1a, and the relevant data are collected in Table 1. As illustrated in Figure 1a, all BTZI-based copolymers show two characteristic absorption peaks with the higher energy one in the wavelength range of 350-500 nm, arising from a π-π* transition. While the broader one in the wavelength range of 500-1400 nm can be attributed to the strong intramolecular charge transfer between the electron-deficient BTZI and the electron-rich alkoxy-based bithiophene (or bithiazole) units. The absorption maxima of P1, P2, and P3 in chloroform solution are located at 663, 890, and 726 nm, respectively. From solution to film state, P1 and P3 show substantial bathochromic shifts (~200 nm) due to increased backbone planarity and enhanced aggregation in film state. However, P2 appears to have a greatly smaller red-shift


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from solution to film than the other two copolymers. The results indicate that, compared with P1 and P3, P2 is likely to have a higher degree of aggregation in solution. In film state, the respective absorption maxima of polymers P1, P2, and P3 are 851, 939, and 954 nm and the optical bandgaps (Egopts) estimated from the absorption edges are 1.13, 1.05, and 0.92 eV, respectively. It is evident that replacing dialkoxy BTOR with monoalkoxy TRTOR leads to a wider Egopt due to the weaker TRTOR electron-donating ability. When copolymerized with highly electron-deficient acceptor co-unit, 2,1,3-benzothiadiazole-5,6-dicarboxylicimide, the stronger donor unit BTOR (versus TRTOR) leads to polymer P2 (versus P1) with a higher-lying HOMO and hence smaller bandgap through orbital hybridization.53 Although BTzOR is less electron-rich than BTOR, P3 shows a smaller bandgap than P2, which is likely attributed to the higher P3 backbone coplanarity due to the elimination of C-H moiety on the thiazole. The higher degree of backbone planarity results in reduced polymer bandgap, as shown in phthalmide-based polymer analogues,52 and leads to red-shifted absorption profile as well as decreased solubility of polymer P3. In addition, the thiazole has different aromatic resonance energy than thiophene, which can result in varied population of quinoidal form and reduced bandgap.53 Electrochemical Properties The electrochemical properties of the polymers P1, P2, and P3 were determined by cyclic voltammetry versus ferrocene/ferrocenium (Fc/Fc+) as the internal standard (Figure 1b). The FMOs were calculated using the equation: EHOMO= −e(Eoxonset + 4.80) eV, where Eoxonset is the onset oxidation potential versus Fc/Fc+. ELOMO= −e(Eredonset +



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4.80) eV, where Eredonset is the onset oxidation potential versus Fc/Fc+. Among the three copolymers, P2 shows the highest-lying HOMO (−5.01 eV), reflecting the strongest electron-donating nature of the BTOR unit. In comparison to BTOR-based polymer P2, P1 and P3 have lower HOMO levels (−5.20 eV) due to the removal of one oxygen (P1) or the replacement of thiophene with the less electron-rich thiazole unit (P3). Low-lying HOMOs can lead to increased hole injection barrier from the source/drain electrode, therefore the hole mobilities of P1 and P3 should, in principle, be suppressed. Through incorporating the strong electron-withdrawing acceptor (BTZI) into the blocking building, the LUMOs are −3.64, −3.49, and −3.68 eV for P1, P2 and P3, respectively. Such low-lying LUMOs of P1 and P3 should facilitate electron injection from Au source electrode with a work function of −5.10 eV, and promote n-channel performance in OTFTs. Among the polymers, the strongest donor co-unit BTOR leads to the highest-lying LUMO (−3.49 eV) for polymer P2, which results in hole-dominating transport characteristics in OTFTs (vide infra). According to the FMOs of the three polymers, P1 and P3 are likely to obtain balanced ambipolar transport than P2 in OTFTs.


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Figure 2. Theoretically calculated top views, side views, and FMO levels of the trimers of the repeating units of polymers P1，P2, and P3. The calculations were performed at the DFT//B3LYP/6-31G* level and the torsion angles are indicated in the figure. The plots shown are the local minima and the side chains are truncated to simplify the calculation. Theoretical Computation Density functional theory (DFT) was used to study the molecular geometry, electronic structures, and opto-electrical properties of BTZI-based polymers using a hybrid B3LYP correlation function and 6-31G* basis set. For simplified computation, the branched alkyl side chains are replaced with methyl groups. Figure 2 shows the optimized geometries and the FMO energy levels of the dimers of the repeating units with the topologies shown in Table S1. In P1 and P2, the computation results reveal that the dihedral angles between thiophene and BTZI are > 24o, indicating a certain degree of steric hindrance. The introduction of thiazole units into the polymer backbones slightly reduces the steric hindrance between thiophene and BTZI to 24.19o in P3 (versus the angle of 27.18o in P2) due to the removal of the steric demanding C-H moiety on the thiazole. In the donor units, the dihedral angles between two thiophenes or thiazoles are very minimal (0.09-2.52o) in all polymers, which are attributed to the conformation locking enabled by the intramolecular non-covalent S···O interaction.49 Particularly, by using two alkoxy chains the dihedral angles are close to 0° in P2 and P3. As seen in the side views of the polymer geometries, it was found that the whole backbone planarity of P3 is greatly higher than the other two polymers P1 and P2, which accounts for its smallest bandgap in the 13 ACS Paragon Plus Environment


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polymer series. On the basis of the DFT calculation, the substitution of bithiophene with bithiazole leads to ~0.2 eV lower-lying LUMO/HOMO (−2.94/−5.17 eV) of the trimer of the P3 repeating unit versus that (−2.79/−4.91 eV) of the trimer of the P2 repeating unit, which is attributed to the electron-deficient thiazole core.50, 52 By replacing one alkoxy chain with a less electron-donating alkyl chain,51, 54 the LUMO/HOMO of the trimer of the P1 repeating unit was lowered to −2.95/−5.14 eV versus the those (−2.79/−4.91 eV) of the trimer of the P2 repeating unit. The high-lying HOMO levels will contribute to the hole mobility in OTFTs for all polymers. Furthermore, compared with that of P2, the LUMOs of both P1 and P3 are decreased by ~0.2 eV, which should be beneficial to the electron injection and hence promote n-channel performance for polymers P1 and P3. Organic Thin-Film Transistors Performance In order to investigate the charge transport properties of these BTZI-based polymers, top-gate/bottom-contact (TG/BC) OTFT devices were fabricated and characterized. For this purpose, glass substrates with Au source/drain (S/D) electrodes patterned by photolithography were used. The OTFTs have a channel width of 5 mm and varied channel lengths of 10, 20, 50, or 100 µm. The semiconducting layer was deposited under nitrogen atmosphere by spin-coating from a 5 mg/mL solution either in chlorobenzene for P1, 1,2-dichlorobenzene for P2, or in chloroform for P3. After spin-coating, the semiconducting films are subjected to thermal annealing at 200°C for 10 min. Then low-k CYTOP film was deposited through spin-coating as the


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dielectric layer, and concluded by thermal annealing at 100°C for 10 min in nitrogen. The Al gate electrode was subsequently deposited by thermal evaporation using shadow masks to complete the device fabrication. The fabricated OTFTs were characterized in nitrogen atmosphere using semiconductor analyzer (Keithley 4200-SCS). Representative transfer and output characteristics of the OTFT devices based on the three polymers are shown in Figure 3, and OTFT performance parameters are summarized in Table 2 and Table S2. To extract the field-effect mobility in the saturation regime (µsat), standard equation Isd = µsat Ci (W/2L)(Vg − Vth)2 was used, wherein Isd is the source-drain current, Ci is the dielectric capacitance per unit area, W is the channel width, L is the channel length, Vg is the gate voltage, Vth is the threshold voltage. Ci = 4.1 nF cm-2 was used as the capacitance per unit area of the CYTOP layer. For top-gate devices with this series of polymers, we observed the existence of high gate-leakage current (Ig), evidenced by the high Isd at Vd = 0 V in the output plots. Nonetheless, the devices are still operating well below the break-down field of the CYTOP dielectric layer, shown clearly by the Ig plots in the transfer curves. Meanwhile, the gate leakage current part in the total Isd was carefully subtracted off first before the mobility extraction process, to get the actual current between the source and drain to ensure that reliable results were consistently obtained.



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Figure 3. Transfer (a, b, c, g, and h) and output (d, e, f, i, and j) characteristics of P1 (a, b, d, and e); P2 (c and f); P3 (g, h, I, and j) top-gate/bottom-contact OTFTs (L = 20 µm, W = 5 mm) annealed at the optimal temperature of 200°C. P-type and n-type regimes are indicated in the figures. The dashed lines represent the leakage current between gate and drain electrodes. Table 2. Top-Gate/Bottom-Contact (TG/BC) OTFT Performance Parameters of the BTZI-based Polymer Semiconductors. Polymer P1 P2 P3 a)

Ta (°C) 200 200 200

n-channel µe (cm V-1 s-1) a 0.86(0.56) NA 0.95(0.50) 2

Ion/Ioff 2 x 103 NA 2 x 105

p-channel VT (V) 64 NA 43

µh (cm V-1 s-1)a 0.51(0.32) 0.70(0.40) 0.50(0.34) 2

Ion/Ioff 3 x 104 300 2 x 105

VT (V) −40 −39 −37

Data represent the best mobilities with average mobilities from > 5 devices shown

in parentheses.


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Strong D-A type copolymers with relatively low-lying LUMOs and high-lying HOMOs should promote both hole and electron injection from S/D electrodes, leading to ambipolar transport. As seen in Figure 3, P1 and P3 TG/BC OTFTs with Au S/D electrodes exhibit distinctive ambipolar transport characteristics. Under the optimized fabrication condition, P1 OTFTs show an average electron mobility (µe) of 0.56 cm2 V-1 s-1 and an average hole mobility (µh) of 0.32 cm2 V-1 s-1. The highest µe/µh of 0.86/0.51 cm2 V-1 s-1 was obtained for P1-based OTFTs. After device optimization, P3 demonstrates comparable OTFT performance with an average µe/µh of 0.50/0.34 cm2 V-1 s-1 and the highest µe/µh of 0.95/0.51 cm2 V-1 s-1. For the n-channel operation, both P1 and P3 show large threshold voltages (|VTs| > 40 V), indicating the presence of large electron injection barriers, i.e. Schottky barrier in the OTFT devices. The Schottky barrier can be approximated by the difference between the polymer LUMOs and the work function (~ −5.10 eV) of Au electrode. The lower-lying LUMO of P3 leads to smaller Schottky barrier than P1, as reflected by the relatively lower VT (43 V) when comparing to VT (64 V) of P1. Among the polymers, P2 contains the strongest electron-donating BTOR co-unit, leading to the highest-lying LUMO (−3.96 eV) and HOMO (−5.01 eV) in the series (Table 1). The Schottky barrier for hole injection is almost negligible while it is ~1.0 eV or even higher for electron injection when estimated from the standard work function of Au. As a result, P2-based device exhibits unipolar p-channel performance with an average µh of 0.40 cm2 V-1 s-1 and a peak µh of 0.70 cm2 V-1 s-1. The performance variation can be correlated to the evolution of polymer FMOs. P2 contains the strongest electron-donating BTOR



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co-unit, leading to the highest-lying LUMO (−3.96 eV) and HOMO (−5.01 eV) in the series (Table 1). The high-lying HOMO also results in small Ion/Ioff ratio (3 x 102) for P2 OTFTs. Bottom-gate/top-contact (BGTC) OTFTs with low work function Al S/D electrodes were also fabricated to further study the effect of Schottky barrier on charge carrier injection. Heavily p-doped silicon with 200 nm thermally grown oxide was used as substrate, which was treated with self-assemble layer (SAM) prior to device fabrication. Trimethoxy(n-octadecyl)silane (OTMS) SAM was used for P2 and P3-based devices, while hexamethyldisilazane (HMDS) was used for P1-based device since P1 could not form film properly on OTMS-treated SiO2 surface. The semiconductor layers were spin-coated and thermally annealed under the same conditions as the top-gate devices. Finally, 40 nm Al was evaporated through shadow mask on top to form source/drain electrodes with channel lengths of 50, 75, 100 µm and a channel width of 1 mm. The representative I-V curves are plotted in Figure S2-4 and the device performance is summarized in Table S3. In bottom-gate OTFTs with Al S/D electrodes, we observed ambipolar transport for all polymers, including P2-based devices, which show an average µe/µh of 0.095/0.18 cm2 V-1 s-1 and the highest µe/µh of 0.10/0.20 cm2 V-1 s-1. The appearance of n-type transport of P2-based OTFTs is attributed to the reduced Schottky barrier for electron injection when low work function Al is used as the S/D electrodes. Polymer Solar Cell Performance The wide absorption of these BTZI-based polymers indicates that they are


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promising materials for solar cell application, which should be beneficial to short-circuit currents (Jscs). It should be pointed out that narrow bandgap polymers can suffer small open-circuit voltages, leading to mediocre power conversion efficiencies. Bulk heterojunction (BHJ) PSCs are then fabricated using P1, P2, and P3 as the electron-donating materials blended with (6, 6)-phenyl-C71-butyric acid methyl ester (PC71BM) as the electron-accepting material with a device structure of glass/ITO/PEDOT:PSS/polymer:PC71BM/Ca/Al. The active layers are spin-coated from a warm o-dichlorobenzene solution (total concentration: 25 mg/mL). In order to achieve optimized photovoltaic performance, the polymer:PC71BM blend ratios were systemically varied from 1:1.5 to 1:2.5 and to 1:3.5. Table 3. Device Performance Parameters of PSCs under Optimized Conditions and the Charge Transport Mobilities of Hole and Electron-Only Devices Based on Polymer:PC71BM (1:2.5) Blend Films. Polymer

Voc (V)

Jsc (mA cm−2)

FF (%)

PCE (%)

µh (SCLC) (cm2 V-1 s−1)

µe (SCLC) (cm2 V-1 s−1)

P1

0.69

11.11

54.15

4.15

7.2 x 10-6

4.0 x 10-4

P2

0.50

7.69

55.17

2.10

6.4 x 10-5

2.4 x 10-4

P3

0.52

0.50

44.48

0.12

1.12 x 10-5

7.5 x 10-5

It was found that the blend ratio of 1:2.5 leads to optimal performance (Table 3). The current

density−voltage

(J−V)

curves

of

the

PSCs

having

the

optimal

polymer:PC71BM ratio under the illumination of air mass (AM) 1.5G, 100 mW cm−2 and the external quantum efficiencies (EQEs) of the optimized solar cells are shown in Figure 4. During the fabrication of the BTZI-based PSCs devices, it was found that processing additive and thermal annealing show minimal impacts on the PSC performance.43 Polymer P3 shows nearly no PSC performance with a PCE of 0.12%, 19 ACS Paragon Plus Environment


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which is likely due to its very low-lying LUMO, resulting in insufficient charge transfer from polymer to PC71BM. Polymer P2 shows improved solar cell performance with a PCE of 2.10% with a Voc of 0.50 V. The smallest Voc is in good agreement with its highest-lying HOMO (-5.01 eV) in the series. In comparison to the PSCs based on P2 and P3, P1 shows the highest PEC of 4.15% with a Jsc of 11.11 mA cm−2, a FF of 54%, and a Voc of 0.69 V. The EQE curve of P1-based PSCs shows broad photovoltaic response in the wavelength range from 300-1000 nm with maximum EQE approaching 50%.

Figure 4. (a) J−V characteristics of the optimized inverted P1-P3-based polymer solar cells under simulated AM 1.5 G illumination (100 mW cm−2); (b) corresponding EQE plots of the optimized solar cells. The charge carrier mobilities of polymer:PC71BM blend films were also investigated using the space charge limited current (SCLC) model, which is widely applied to measure the hole and electron transporting ability of active layers between electrodes in the PSC field. All hole-only or electron-only devices show good mobilities, specially the hole mobility/electron mobility (µh,SCLC/µe,SCLC) is 7.2 x 10-6/4.0 x 10-4, 6.4 x 10-5/2.4 x 10-4, 1.12 x 10-5/7.5 x 10-5 cm2 V-1 s−1 for P1, P2, and 20 ACS Paragon Plus Environment

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P3-based PSCs, respectively. The mobility measurements reveal that P3 blend shows moderate mobilities (~10-5 cm2 V-1 s−1), hence its minimal solar cell performance is likely attributed to the very low-lying P3 LUMO (−4.28 eV), which is too deep for electron transfer from P3 to PC71BM. In comparison to P1 blend, P2 blend shows more balanced SCLC mobility, the degraded P2 cell performance is due to its coarse phase separation as revealed by film morphology characterization (vide infra). Figure 4b shows the external quantum efficiency (EQE) of the optimized P1, P2, and P3-based PSCs. P1 shows the highest EQE approaching 50% among the three devices and thus achieves the largest Jsc in the series. It is also noted that the energy loss in the P1 cells, defined as the difference between the optical bandgap and Voc, is 0.44 eV, which, to our best knowledge, is the lowest photon energy loss reported to date for the cells showing relatively high PCE (> 4%).43, 55 In comparison to that in phthalimide polymer-based PSCs,56-58 the small energy loss is attributed to the deep-lying LUMO levels of the BTZI-based polymers. In addition, tandem solar cells incorporating P3HT:ICBA as the front cell active layer and P1:PCBM as the back cell active layer were fabricated. The detailed device parameters of the tandem solar cells are included in supporting information (Table S5). It was found that the tandem solar cells show an enhanced PCE of 6.31% with a Jsc of 9.46 mA cm−2, a FF of 47.58%, and a Voc of 1.41 V. We admit the performance of the tandem cell is not state-of-the-art, the incorporation of the narrow bandgap polymer P1 leads to broad photoresponse (Figure S6b) and improved device performance than the front or real cells, demonstrating the advantage of using narrow bandgap polymers for improving solar



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absorption.

Figure 5. Tapping mode atomic force microscopy images of (5 × 5 µm) of (a) P1 film with annealing, RMS roughness = 2.35 nm; (b) P2 film with annealing, RMS roughness = 4.01 nm; and (c) P3 film with annealing, RMS roughness = 0.97 nm.(d) P1:PC71BM blend film, RMS roughness = 9.20 nm; (e) P2:PC71BM blend film, RMS roughness = 3.43 nm; (f) P3:PC71BM blend film, RMS roughness = 3.83 nm. TEM images of P1 (g), P2 (h), P3 (i) bulk heterojunction blend (polymer:PC71BM) films (scale bar: 500 nm). Film Morphologies and Their Correlations with Device Performance The morphology of active layers plays a crucial role in OTFT and PSC performance. A set of characterization techniques, including atomic force microscopy (AFM), 22 ACS Paragon Plus Environment

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transmission electron microscopy (TEM), and grazing incidence wide angle X-ray scattering (GIWAXS), were performed to study film morphology and microstructure. Under the optimal fabrication condition, the P1, P2, and P3 neat films (Figure 5a-5c) show structured morphology with the root-mean-square (RMS) roughness of 2.35, 4.01, and 0.97 nm, respectively. The structured surface morphology indicates that all BTZI-based polymers have a certain degree of ordering, which leads to substantial charge carrier mobility in OTFTs. Among all polymers, P2 shows the highest RMS roughness, which is likely attributed to its strongest aggregation as revealed by the UV-vis absorption of polymer solutions. P3 neat film shows a fibrillar ordered morphology with a much smoother surface, which is similar to phthalimide-based polymer analogues.50 Such morphology features greatly reduced domain boundaries, leading to substantial charge carrier mobility in OTFTs. The AFM images of polymers:PC71BM blend films are shown in Figure 5d-5f, and P2 and P3-based blend film exhibit large domain sizes with distinctive grain boundaries, which are likely due to the coarse phase separation of polymer donor and fullerene acceptor in the blend films, as reveled by TEM characterization (vide infra). TEM was also used to investigate polymer:PC71BM blend film morphologies. Figure 6b and 6c show distinctive structural features with large white and dark regions with sizes > 100 nm, corresponding polymer-rich and PC71BM-rich domains, respectively. Due to the limited exciton diffusion length, phase separation at hundred nanometer scale can lead to insufficient exciton dissociation, resulting in small Jscs.59 The coarse phase separation is likely attributed to the doubly intramolecular



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sulfur-oxygen interaction in polymer P2 and P3, similar to phthalimide-BTOR copolymers.51, 56 Among all polymer:PC71BM blends, a more uniform morphology with finer phase separation is observed in the P1:PC71BM film. In addition, the P1 blend features interconnected network, which provides continuous pathway for charge carriers to reach their respective electrodes. Therefore, among the PSCs, P1-based solar cells show the highest Jsc of 11.11 mA cm−2. The TEM results of the P1, P2, and P3 blend films are highly compatible with their photovoltaic performance.

Figure 6. GIWAXS images of (a-c) pristine film, (d-f) film after thermal annealing, and (g-i) polymer:PC71BM blend film of BTZI-based polymers. GIWAXS was carried out to examine polymer film microstructure and morphology of the as-cast neat polymer films, thermally annealed neat films, and polymer:PC71BM blend films.60 Figure 6 and Figure S7 show the two-dimensional images and the corresponding line cut profiles, respectively. On the basis of (010) diffraction patterns, all as-cast P1-P3 films adopt a face-on dominating bimodal 24 ACS Paragon Plus Environment

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orientation. It was found that the variation of the donor co-unit shows profound influence on the lamellar d-spacing distance of BTZI-based polymers, and the derived d-spacings in the qxy direction are 20.2 and 19.5 Å for P1 and P3 films, respectively, which are shorter than that (22.7 Å) of P2 film. The enlarged lamellar d-spacing of P2 is likely due to its most twisted backbone as revealed by the DFT calculation, resulting in a lower degree of side chain interdigitation. Polymer π-stacking diffraction peaks were observed at ~1.66 Å−1 in the qxy direction for P1 and qz direction for P2 and P3 polymers, corresponding a π-stacking distance of ~3.8 Å, which is slightly larger than for the phthalimide polymer analogue.50 After thermal annealing under the optimal OTFT fabrication condition, all BTZI-based polymers exhibit further improved film crystallinity, as revealed by the increased diffraction intensity and diffraction peaks progressing to higher orders. On the basis of (010) diffraction patterns of thermally annealed films, P1 shows an edge-on orientation, while P2 has a face-on dominant bimodal orientation, and P3 has an edge-on dominant bimodal orientation. Hence, thermal annealing leads to variation of polymer chain orientation, which has also been observed in other polymer semiconductors.61 Although the exact reason is unknown for such orientation change, it is likely induced by different intermolecular interaction between polymeric chains as well as interaction between polymer and substrate. The edge-on dominating polymer chain orientation of annealed P1 and P3 films likely accounts for their higher motilities than P2 film in OTFT devices. On the basis of the Scherrer equation, the crystal coherence length (CCL100) is



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Page 26 of 38

deduced from the full width at half maximum (FWHM) of the in-plan (100) lamellar peak. The derived CCLs of as cast-films were 4.9, 12.0, and 10.7 nm for P1, P2, and P3, respectively, and the CCLs of thermal annealed samples were increased to 10.5, 17.7, 17.5 nm for P1, P2, and P3, respectively. The results clearly indicate that thermal annealing improved long range crystalline order of BTZI-based polymers and hence benefitted carrier transport properties in OTFTs fabricated under optimized fabrication condition. For polymer:PC71BM blend film crystallinities (Figure 6. g-i), the addition of PC71BM leads to P1 and P2 with greatly reduced π-π packing diffraction peak, however P3:PC71BM blend film still maintain up to third order intermolecular diffraction peaks, suggesting less miscibility with PC71BM which leads to the coarse phase separation and partially accounts for its lower PCE in PSCs in comparison to P1 and P2. Conclusions In

summary,

a

highly

electron-deficient

2,1,3-benzothiadiazole-5,6-dicarboxylicimide

(BTZI),

was

acceptor synthesized

unit, and

incorporated into polymer backbone to afford three polymer semiconductors. The solubilizing chain on the N-imide group and the side chains on the head-to-head bithiophene afford the BTZI-based polymers with good solubility. The strong intramolecular charge transfer between the highly electron deficient BTZI and the electron-rich alkoxy-functionalized bithiophene (or thiazole) leads to the resulting polymers with narrow bandgaps of 1.13-0.92 eV and low-lying LUMOs of −3.96 to −4.28 eV. All polymers show broad absorption extended to the near-infrared region


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(beyond 1000 nm). The OTFT devices based on the three polymers show good hole mobilities of ~0.5 cm2 V-1 s-1. Furthermore, in addition to hole mobilities, P1 and P3 exhibit substantial electron mobilities up to 0.86 and 0.95 cm2 V-1 s-1, respectively, due to their lower LUMO levels by removing one oxygen and replacing electron-rich thiophene with electron-deficient thiazole, which result in high and balanced ambipolar transport in OTFTs. Despite the very narrow band gap (1.13 eV) of P1, the PCE of P1:PC71BM-based PSC devices reaches 4.1% with a remarkably small energy loss of 0.44 eV. The power conversion efficiency of P1 is among the highest reported to date with such small energy loss in polymer:fullerene solar cells. The results demonstrate that dicarboxylicimide and alkoxy-functionalized thiophene (or thiazole) are promising donor and acceptor pair for constructing high-performance ambipolar polymer semiconductors. ASSOCIATED CONTENT Supporting Information Supporting Information is available free of charge on the ACS Publications website. Experimental details, synthesis and characterization of monomers and polymers, optical spectra of copolymer solutions, details of OTFTs and PSC fabrication, GIWAXS measurements and data, TGA plots, DSC plots, DFT computational results, OTFT and PSC performance data, SCLC data. Notes The authors declare no competing financial interest. ACKNOWLEDGMENTS



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S.Z. acknowledges financial support from the National Key R&D Program of "Strategic Advanced Electronic Materials" (No.2016YFB0401100) and the National Natural Science Foundation of China (Grant No.61574077). X.G. thanks Shenzhen Peacock Plan Project (KQTD20140630110339343) and Shenzhen Basic Research Fund (JCYJ20160530185244662). G. X. acknowledges the financial support from the Science and Technology Development Fund of Macau SAR (FDCT-116/2016/A3) and Start-up Research Grant (SRG2016-00087-FST) from Research & Development Office at University of Macau. M. A. U. and H. Y. W. are grateful to the financial support

from

the

National

Research

Foundation

(NRF)

of

Korea

(2016M1A2A2940911, 20100020209). We are grateful to Prof. Tae Joo Shin for the useful discussion and helps for 2D GIWAXS image processing. We thank Dr. Byeongdu Lee for the GISAXShop program (by) for data conversion at https://sites.google.com/site/byeongdu/software. References 1.

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and

High-gain

Inverters

from

a

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Table of Contents Graphic


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2,1,3-Benzothiadiazole-5,6-dicarboxylicimide-Based Polymer

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