Small Molecule Acceptor and Polymer Donor Crystallinity and

Article pubs.acs.org/cm

Small Molecule Acceptor and Polymer Donor Crystallinity and Aggregation Effects on Microstructure Templating: Understanding Photovoltaic Response in Fullerene-Free Solar Cells Nicholas D. Eastham,† Alexander S. Dudnik,† Thomas J. Aldrich,† Eric F. Manley,†,§ Thomas J. Fauvell,†,§ Patrick E. Hartnett,† Michael R. Wasielewski,*,† Lin X. Chen,*,†,§ Ferdinand S. Melkonyan,*,† Antonio Facchetti,*,†,∥ Robert P. H. Chang,*,‡ and Tobin J. Marks*,†,‡ †

Department of Chemistry, the Materials Research Center, and the Argonne-Northwestern Solar Energy Research Center, Northwestern University, 2145 Sheridan Road, Evanston, Illinois 60208, United States ‡ Department of Materials Science and Engineering, the Materials Research Center, and the Argonne-Northwestern Solar Energy Research Center, Northwestern University, 2145 Sheridan Road, Evanston, Illinois 60208, United States § Chemical Sciences and Engineering Division, Argonne National Laboratory, 9700 South Cass Avenue, Lemont, Illinois 60439, United States ∥ Polyera Corporation, 8045 Lamon Avenue, Skokie, Illinois 60077, United States S Supporting Information *

ABSTRACT: Perylenediimide (PDI) small molecule acceptor (SMA) crystallinity and donor polymer aggregation and crystallinity effects on bulk-heterojunction microstructure and polymer solar cell (PSC) performance are systematically investigated. Two highperformance polymers, semicrystalline poly[5-(2-hexyldodecyl)-4Hthieno[3,4-c]pyrrole-4,6(5H)-dione-1,3-yl-alt-4,4″dodecyl-2,2′:5′,2″terthiophene-5,5″-diyl] (PTPD3T or D1) and amorphous poly{4,8bis(5-(2-ethylhexyl)thiophen-2-yl)benzo[1,2-b:4,5-b′]dithiophene2,6-diyl-alt-(4-(2-ethylhexyl)-3-fluorothieno[3,4-b]thiophene-2-carboxylate-2,6-diyl) (PBDTT-FTTE or D2), are paired with three PDI-based SMAs (A1−A3) of differing crystallinity (A1 is the most, A3 is the least crystalline). The resulting PSC performance trends are strikingly different from those of typical fullerene-based PSCs and are highly material-dependent. The present trends reflect synergistic aggregation propensities between the SMA and polymer components. Importantly, the active layer morphology is templated by the PDI in some blends and by the polymer in others, with the latter largely governed by the polymer aggregation. Thus, PTPD3T templating capacity increases as self-aggregation increases (greater Mn), optimizing PSC performance with A2, while A3-based cells exhibit an inverse relationship between polymer aggregation and performance, which is dramatically different from fullerene-based PSCs. For PBDTT-FTTE, A2-based cells again deliver the highest PCEs of ∼5%, but here both A2 and PBDTT-FTTE (medium Mn) template the morphology. Overall, the present results underscore the importance of nonfullerene acceptor aggregation for optimizing PSC performance and offer guidelines for pairing SMAs with acceptable donor polymers.

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INTRODUCTION

polymer-donor-based BHJ solar cell design and operation, and have enabled devices with impressive power conversion efficiencies (PCEs).12−21 Fullerene derivatives are the most commonly used molecular acceptors in BHJ devices due to their excellent electron transport properties, which afford high internal quantum efficiencies and PCEs.22−27 However, fullerenes have many drawbacks, such as limited optical absorption and relatively fixed energetic/band alignment.28−31 Furthermore, the high cost of fullerenes and their relative instability

Photovoltaic devices based on organic semiconductors offer promising technologies for inexpensive, lightweight, and scalable solar energy conversion.1−5 The most efficient polymer solar cell (PSC) devices to date are based on the bulk heterojunction (BHJ) architecture, which features an interpenetrating network of hole- and electron-conducting pathways formed upon blending electron-donating (donor) and electronaccepting (acceptor) semiconductors. Typical BHJ blends (active layer) consist of a semiconducting π-conjugated polymer donor (D) and a molecular acceptor material (A).6−11 Intense research efforts over the past years have provided deeper understanding of the phenomena governing © 2017 American Chemical Society

Received: March 8, 2017 Revised: May 2, 2017 Published: May 3, 2017 4432

DOI: 10.1021/acs.chemmater.7b00964 Chem. Mater. 2017, 29, 4432−4444

Article

Chemistry of Materials Scheme 1. Structures of Polymer Donors and Perylenediimide Small Molecule Acceptors

diisoquinoline-1,3,8,10(2H,9H)-tetraone (A1), [1,2:3,4]-bis[N,N 2 -bis-1−pentylhexyl-perylenediimide-1,12-yl]benzene (A2), and 9,9′-di(pentadecan-8-yl)-1H,1′H-[2,2′-bianthra[2,1,9-def:6,5,10-d′e′f ′]diisoquinolin]1,1′,3,3′,8,8′,10,10′(9H,9′H)-octaone (A3) shown in Scheme 1.42,43,67,88 Donor polymer samples of incrementally varied Mn values were prepared for the high-performance and wellcharacterized semicrystalline poly[5-(2-hexyldodecyl)-4Hthieno[3,4-c]pyrrole-4,6(5H)-dione-1,3-yl-alt-4,4″-dodecyl2,2′:5′,2″-terthiophene-5,5″-diyl] (PTPD3T, hereafter D1)15,69 and amorphous poly{4,8-bis(5-(2-ethylhexyl)-thiophen-2-yl)benzo[1,2-b:4,5-b′]dithiophene-2,6-diyl-alt-(4-(2-ethylhexyl)-3fluorothieno[3,4-b]thiophene-2-carboxylate-2,6-diyl) (PBDTTFTTE, hereafter D2)89 polymers (Scheme 1). The PDI acceptors A1−A3 are chosen to be structurally dissimilar, A1 having only a single PDI core and increasing intramolecular torsion between PDI units in the bis-PDI acceptor progression, determined by previously reported modeling, A2 → A3, with a moderate 32° twist in A2 and a 90° twist in A3.42,43,67,88 Previous studies have correlated increased intramolecular twisting with lower aggregation tendencies and reduced crystallinity, which suppresses BHJ blend phase separation and increases Jsc and PCE.39−42,67,90−92 All of the present polymer donor samples were characterized by high-temperature gel permeation chromatography (GPC), cyclic voltammetry (CV), and ultraviolet−visible (UV−vis) spectroscopy. When incorporated in PSC devices with the PDI acceptors, the D1derived solar cells exhibit three distinct performance responses with increasing polymer Mn, and the highest PCEs are obtained for the highest Mn for D1:A2 (3.42%) and, unexpectedly, the lowest Mn for D1:A3 (3.41%), while surprisingly, no specific dependence on Mn is observed for the D1:A1 devices (2.77%). However, a completely different trend is obtained for the D2 polymer donor blends, with the medium Mn polymer yielding the highest PCEs for the A2 (5.21%) and A3 (3.37%) acceptors and essentially no dependence observed for A1 (2.42%). Through detailed characterization of the photoactive layer blends with transmission electron microscopy (TEM), grazingincidence wide-angle X-ray scattering (GIWAXS), and charge transport measurements, it is established that the underlying morphological and electronic variations depend on the intrinsic aggregation propensities of both the polymer donor and PDI acceptor. BHJ morphology is templated by the PDI in one blend and by the polymer in another, in a manner dictated by the PDI crystallinity and the polymer aggregation, respectively. Here, a template refers to a structure preformed in solution or/

toward ambient O2 and H2O limit future large-scale deployment in PSC devices.32−35 Consequently, research efforts have shifted to developing alternative n-type materials for BHJ PSC application. Indeed, recent advances in developing nonfullerene small molecule acceptors (SMAs) have enabled remarkably high performing PSCs delivering PCEs comparable to fullerene-based devices and highly competitive for eventual largescale deployment.36−53 Among these promising fullerene alternatives, perylenediimides (PDIs) are attractive due to their low cost, highly tunable absorption properties, excellent solubility, intrinsic electron transporting ability, and favorable energy level alignments, affording high performance devices capable of competing with fullerenes.42,54−62 Current strategies to increase D:PDI device PCEs focus on optimizing BHJ film morphology to maximize D:A interfacial contact for improved charge generation and collection. This task is challenging due to the tendency of planar PDIs to self-stack, forming large crystalline domains that are inefficient at charge separation. To disrupt the aggregation of PDI acceptors, synthetic modification of the PDI core with bulky substituents as well as fusion of multiple PDI cores in twisted geometries have been successful in enhancing device PCEs and will likely see implementation as nonfullerene SMAs.27,42,56,63−68 Another promising strategy to enhance PSC performance involves tuning the donor polymer number-average molecular weight (Mn), which profoundly effects macromolecular aggregation and crystallinity, active layer thin film blend morphology, charge generation, recombination, and carrier mobility.12,69−74 The general conclusion is that increasing polymer Mn enhances PSC performance,75−85 although recent reports have questioned this paradigm.69,71,86,87 Importantly, the great majority of PSC donor polymer crystallinity/Mn effect studies have employed fullerene acceptors, and analogous studies with nonfullerene SMAs are currently lacking, which is surprising considering the generally accepted influence of polymer crystallinity/Mn on PSC performance. Considering that recent research to enhance fullerene-free device performance has yielded SMA materials rivaling fullerene performance, better understanding of the interplay between polymer donor and nonfullerene SMA is in its infancy, materials selection and matching principles are needed for advancing SMA-based PSCs. Here we report the first systematic study of polymer crystallinity and aggregation effects on the performance of D:SMA PSCs for the three representative PDI-based acceptors: 2,9-dioctyl-4,7,11,14-tetraphenylanthra[2,1,9-def:6,5,10-d′e′f ′]4433


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Chemistry of Materials and rapidly formed during film deposition morphological structure consisting of one BHJ blend component that directs the incorporation of a secondary component and defines domain sizes.75 The observed PSC performance trends are unprecedented and challenge the generally accepted notion that increasing the Mn is universally beneficial in SMA-based PSCs. The present results also show that such trends can be dramatically different from fullerene-based systems. Therefore, the SMA crystalline tendency must be a primary factor in PSC donor−acceptor pairing. The present study affords better understanding of D:PDI morphology−photovoltaic performance relationships and provides a guide for developing, testing, and appropriately pairing new nonfullerene acceptors with carefully selected polymer donors of optimal Mn.

polymerization reactions to obtain samples of D2 with Mn of ∼12 and ∼25 kg/mol (D2−12 and D2−25, respectively) as shown in Table 1. Table 1. Mn Values of D2 Polymers

a

polymer batch

excess 2 (mol %)

Mn (kg/mol)a

Đ

D2−12 D2−25 D2−36

10 3 0

12.4 25.3 36.2

1.67 2.06 2.23

Determined by GPC at 150 °C in 1,2,4-trichlorobenzene.

The predicted Mns from Carothers equation are found to be in good agreement with the experimental GPC Mns (Figure S1). In addition, five samples of donor polymer D1 with systematically varied Mns (8, 19, 30, 45, and 59 kg/mol abbreviated here as D1−8, D1−19, D1−30, D1−45, and D1− 59, respectively) were prepared by analogous methods reported previously.69 Polymer and PDI Physicochemical Properties. The optical and electrochemical properties of the D2 polymer samples were investigated by UV−vis both as thin films and in solution (Figure 1a, Tables 2 and S2), and CV (Figure S4 and Table S4). Note that the physicochemical properties of D1 with sequentially varied M n s were extensively characterized

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RESULTS Solution aggregation or “pre-aggregation” of both semicrystalline and amorphous polymers plays an important key role in the formation of solid state morphology of neat materials in the blends with SMAs.17,42,75,76,91 Here, we use a ratio of the lowest energy intramolecular charge transfer peak to the next higher one (A0−0/A0−1) as metrics for aggregation.80 Crystallinity of polymeric materials is estimated by differential scanning calorimetry (DSC) in bulk and/or by GIWAXS in thin films.69 For PDI acceptors, the twist angle between the core and the adjacent aromatic moiety is widely accepted as a predictor of aggregation with a larger angle resulting in less PDI aggregation and vice versa. PDI crystallinity is judged by GIWAXS and found to follow a general trend: the greater twist angle, the less crystalline the PDIs. Synthesis of Polymers with Defined Mns. We previously reported a synthetic strategy for the preparation of precisely controlled Mn samples of D1 which employs a modified version of Carothers equation to predict Mn from the degree of functional group conversion in a given polymerization reaction as well as from the stoichiometric imbalance of the comonomers. Using this approach, we synthesized D2 batches with Mns ranging from 12 to 36 kg/mol (Scheme 2). First, a Scheme 2. Synthesis of D2 Polymers

stoichiometrically balanced polymerization reaction of comonomers 1 and 2 followed by Soxhlet extraction purification yielded a D2 polymer with Mn = 36 kg/mol (a batch abbreviated hereafter as D2−36) with a molecular weight dispersity (Đ) of 2.23 by high-temperature GPC. From this polymerization, a degree of functional group conversion (p) for these polymerization reaction conditions of 0.9754 was calculated. Following the modified version of Carothers equation (eq S1), the amount of comonomer 2 was varied, while maintaining all other conditions, in two additional

Figure 1. (a) UV−vis absorption spectra of D2 polymers in CHCl3 (0.013 mg/mL). (b) UV−vis absorption spectra of PDI acceptors and polymer donors in thin films cast from CHCl3 (5.0 mg/mL). 4434


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Chemistry of Materials

results are summarized in Table S4. Overall, the HOMO energy of the lower Mn polymer D2−12 (−5.55 eV) lies slightly below those measured for D2−25 (−5.45 eV) and D2−36 (−5.44 eV), likely reflecting the higher fraction of electron-deficient fluorinated thieno[3,4-b]thiophene end groups in D2−12.94,95 The previously reported CV-derived HOMO levels of the D1 samples are independent of Mn and approximately −5.57 eV, which is ∼90 mV lower than those of the D2 samples. The PDI acceptor lowest unoccupied molecular orbital (LUMO) energies derived from the CV reduction onset are −3.79 eV for A1 and A2, whereas the LUMO of A3 is at −3.97 eV, similar to the previously reported data.42,43,67 Results are summarized in Table S5. Solar Cell Characterization. BHJ solar cells with inverted architectures96,97 were fabricated with donor polymers D1 and D2 and acceptors A1−A3. Glass substrates with cleaned, prepatterned indium tin oxide (ITO) electrodes were coated with a sol−gel-processed ZnO layer (∼25 nm) and annealed at 170 °C, and then the photoactive blend (∼90 nm) was spincoated on top. Optimal D1- and D2-based blend films were processed from chloroform as the primary solvent with 0.5% and 1.0% by volume of diphenyl ether (DPE) as a cosolvent, respectively. Vacuum deposition of MoOx (9 nm) and Ag layers (100 nm) completed the counter electrode. Further information on device fabrication can be found in the experimental section. Relevant photovoltaic data are summarized in Table 3 and Figure 2 for D1 and in Table 4 and Figure 3 for D2.

Table 2. Physicochemical Properties of D1 and D2 polymer batch D1−8f D1−30f D1−59f D2−12 D2−25 D2−36

λmax (nm)a 488, 582, and 631 582 and 629 583 and 629 692 and 627g 696 and 635g 694 and 637g

Egopt (eV)c

HOMO (eV)d

LUMO (eV)e

−

1.84

−5.57

−3.73

− − 1.10 1.29 1.32

1.85 1.85 1.61 1.61 1.61

−5.56 −5.57 −5.55 −5.45 −5.44

−3.71 −3.72 −3.94 −3.84 −3.83

A0−0/A0−1b

a

Solution absorption spectra (0.013 mg/mL in CHCl3). bAbsorbance maximum peak intensity relative to absorbance shoulder. cOptical energy gap estimated from absorption onset of the as-cast thin film (5.0 mg/mL in CHCl3). dElectrochemically determined vs Fc/Fc+; EHOMO = −(Eoxonset + 4.88). eCalculated according to ELUMO = EHOMO + Egopt. fAs previously reported by this lab in ref 69. gSecondary absorbance maxima or A0−1 band.

previously, and Table 2 collects only selected values.69 There is a clear dependence of the solution absorbance spectra on the D2 Mn (Figure 1a). All D2 polymer samples exhibit a solution absorbance maximum (λmax A0−0 band) located between 692 and 696 nm as well as a secondary vibronic shoulder (A0−1 band) between 627 and 637 nm. The ratio of the intensity between these two bands (A0−0/A0−1) indicates the degree to which the polymer is preaggregated and is known to depend on the Mn, Đ, and crystallinity.80 Generally, higher A0−0/A0−1 ratios for a series of the same polymer correlate with the higher device efficiencies in fullerene-based PSCs.75,84,93 By this metric, we observe that higher-Mn D2 samples display more preaggregation in solution (Table 2); however, the difference between D2−12 and D2−25 is much larger than that between D2−25 and D2−36. In thin films, the absorbance spectra of the D2 polymer samples display slightly red-shifted absorbance maxima relative to solution spectra, with λmax values ∼702 nm and a vibronic shoulder. A similar, but less pronounced trend in the aggregation behavior is observed for the D2 films, where D2− 12 is the least aggregated and D2−25 and D2−36 are found to be very similar (Figure S2), indicating primarily self-folding polymer aggregation.80 The D2 optical bandgaps calculated from the film absorption onsets are identical at 1.61 eV. Acceptors A1−A3 were synthesized according to literature procedures and characterized using the same conditions as for the polymers to provide comparable data.42,43,67 The thin film absorption spectra of D2−36 and D1−59 as well as those of A1, A2, and A3 are shown in Figure 1b. Note that the relatively low-bandgap D2 polymers absorb until ∼770 nm, whereas the higher optical gap (1.85 eV) D1 batches absorb until only ∼670 nm.69 As will be discussed later, the D2-derived PSCs exhibit enhanced photoresponse between 670 and 770 nm which contributes to the higher Jscs (Figure S7). Films of A1 exhibit the lowest-energy absorption onset of the three acceptor materials at ∼618 nm, the result of a significant red shift from solution to film due to aggregation. A2 films exhibit an absorption onset at ∼585 nm and pattern nearly identical to the solution spectrum. The vibronic structure of A3 is almost entirely retained in the film, with only a small red shift to an onset of ∼595 nm (Figure S3). The electrochemical properties of the D2 polymers and the PDI acceptors were next investigated as thin films by CV. The highest occupied molecular orbital (HOMO) energies of the D2 samples were extracted from the onset of polymer oxidation referenced to the ferrocene/ferrocenium redox couple, and

Table 3. PSC Performance of D1:A1−A3 Devices SMA

D1 Mn (kg/mol)

Jsc (mA/cm2)

Voc (V)

FF (%)

A1

8 19 30 45 59 8 19 30 45 59 8 19 30 45 59

4.48 4.38 4.68 5.53 5.28 1.81 5.62 6.43 6.68 7.00 7.59 7.14 7.34 7.51 5.62

0.954 0.958 0.939 0.959 0.945 0.986 0.965 0.971 1.00 0.981 0.861 0.856 0.846 0.834 0.849

61.9 62.4 63.0 52.2 54.7 35.7 35.4 38.7 50.6 49.9 52.2 54.2 51.9 43.0 47.6

A2

A3

PCE (%)a 2.64 2.62 2.77 2.77 2.73 0.64 1.92 2.41 3.38 3.42 3.41 3.31 3.22 2.69 2.27

(2.60) (2.55) (2.67) (2.69) (2.64) (0.53) (1.73) (2.36) (3.34) (3.36) (3.35) (3.23) (3.18) (2.61) (2.18)

a

Values in parentheses are the average PCE obtained from at least 8 devices.

Note that D1:A1 devices exhibit only small differences in performance as the aggregating ability of D1 is varied, with a maximum PCE = 2.77% for D1−45. As expected from the constant polymer energy levels, the open-circuit voltage (Voc) remains essentially constant (∼0.950 V) with Mn, while the short-circuit current density (Jsc) slightly increases from 4.38 mA/cm2 (D1−8) to 5.53 mA/cm2 (D1−45), in line with previous polymer Mn studies.12,69,75 In contrast, benzo-fused A2 exhibits a dramatic increase in performance with increasing aggregating ability of D1, achieving 0.64% PCE with D1−8 and 3.42% with D1−59, largely reflecting increases in both fill factor (FF) (from 35.7% to 49.9%) and Jsc (from 1.81 mA/cm2 to 7.00 mA/cm2) metrics. Surprisingly, an inverse trend is 4435


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Thus, an optimal PCE of 5.21% is found in the D2−25:A2 blend, whereas the PCEs of D2−12:A2 and D2−36:A2 are significantly lower at 3.33% and 3.38%, respectively. The superior performance of D2−25 reflects increases in Jsc (11.81 mA/cm2) and FF (47.7%). The A3-derived devices exhibit a similar, less pronounced trend with a small PCE increase to 3.37% for the D2−25:A3 blend relative to the lower and higher Mn blends (3.00% and 2.88%, respectively) reflecting an increased Jsc of 9.36 mA/cm2. Again, these results contrast the traditional assumption that higher polymer Mn means higher BHJ PSC performance. With regard to the present nonfullerene blends, the close relationship between blend morphology and performance suggests intriguing morphological trends as the SMA crystallinity and donor polymer aggregation is varied. Active Layer Morphology. TEM measurements next focused on each D:PDI acceptor film to investigate the role of the polymer Mn in nano- and microstructure formation (Figure 4). Previous studies of Mn effects in polymer−fullerene systems found that increasing Mn results in decreased polymer domain sizes and increasingly interpenetrated donor−acceptor networks, which enhance charge separation and reduce recombination.12,87 Additionally, increasing Mn is known to template the active layer morphology with both fullerenes and polymer acceptors due to solubility and crystallinity effects, schematized in Figure 5.69,75 TEM images of the D1:A1 blend films (Figure 4, panels a− e) show clear donor−acceptor phase separation for all Mns and a small increase in the polymer-rich domain sizes (darker region) from ∼10 nm in D1−8 to larger clusters of ∼20 nm in D1−59. The polymer network becomes slightly more pronounced and distinct from the acceptor phase as the D1 Mn is increased, typical of increasing aggregation evident in the UV−vis for the higher Mn polymers.69,80 Due to the observation of such a stable morphology, it is concluded that A1 templates the film formation process via its high degree of crystallinity. Overall, these data correlate well with the observed PSC performance: there is little impact on the overall device PCE as Mn is varied. D1:A2 blends exhibit increasing phase separation with increasing polymer Mn (Figure 4, panels f−j), with no distinguishable polymer domains in D1−8 (Figure 4f) and with larger 10−30 nm polymer-dominated domains in D1−59 (Figure 4j), reasonably responsible for the large current increase and concurrent PCE rise. Similarly, the A2 aggregates increase in density and size from ∼20 to ∼40 nm as D1 Mn is increased, due to the increasing ability of D1 to template the film morphology at high Mns. Films of D1:A3 exhibit very little phase separation in TEM images (Figure 4, panels k−o) in comparison to the other D1:PDI blends with no defined polymer network and only small ∼10 nm dark regions, reasonably assignable to polymer-rich domains. This argues that the more amorphous nature of A3 is more easily disrupted by D1, resulting in highly blended films at all Mns examined. Similar trends are also observed in the TEM images of the D2-derived films. Specifically, D2:A1 blend films (Figure S8, panels a−c) exhibit nanoscale phase separation with A1 aggregates of ∼30 nm that appear largely unaffected by polymer Mn changes. Notably, very crystalline A1 templates the morphology in the blend with amorphous D2, maintaining a nearly constant morphology regardless of the aggregating properties of the donor material. In contrast, for less crystalline A2 nanoscale phase separation with ∼40 nm PDI domains for D2−25:A2 films is observed (Figure S8e). Larger domains with

Figure 2. Dependence of power conversion efficiencies (PCEs) on polymer Mn for D1:A1-A3 devices showing increased PCE for A2 and decreased PCE for A3 with increasing D1 Mn.

Table 4. PSC Performance of D2:A1−A3 Devices SMA

D2 Mn (kg/mol)

Jsc (mA/cm2)

Voc (V)

FF (%)

A1

12 25 36 12 25 36 12 25 36

4.70 4.35 5.34 8.76 11.81 10.73 8.03 9.36 9.41

0.879 0.867 0.874 0.895 0.925 0.889 0.728 0.705 0.699

57.1 59.6 51.9 42.4 47.7 35.5 51.3 51.1 43.9

A2

A3

PCE (%)a 2.36 2.24 2.42 3.33 5.21 3.38 3.00 3.37 2.88

(2.20) (2.15) (2.21) (3.25) (5.10) (3.33) (2.92) (3.30) (2.78)

a

Values in parentheses are average PCE obtained from at least 8 devices.

Figure 3. Dependence of power conversion efficiency on polymer Mn for D2:A1−A3 devices showing increase in performance at D2−25 with A2 and A3 and the minimal change in efficiency with A1.

observed for the N-N-linked PDI dimer A3, where the performance decreases with increasing aggregation of D1, achieving 3.41% with D1−8 and 2.27% with D1−59 due to decreases in both Jsc (from 7.59 mA/cm2 to 5.62 mA/cm2) and FF (from 52.2% to 47.6%). These trends are unprecedented and opposite those in fullerene-based PSCs.12,75,98 Similar to the trend observed here for D1:A1 devices, the D2:A1 devices exhibit only minor performance variation (PCE ∼ 2.20%) upon increasing the D2 Mn (Table 4 and Figure 3). An increase in FF to 59.6% for the D2−25 blend relative to the higher and lower Mn blends is offset by a lower Jsc (4.35 mA/ cm2) and along with similar Vocs yields similar PCEs. PSCs fabricated with A2 show a far stronger dependence on D2 Mn. 4436


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Figure 4. TEM images of D1:A1−A3 blend films organized in rows with D1 Mn increasing from left to right from 8 kg/mol (D1−8) to 59 kg/mol (D1−59). Increased phase separation at greater polymer Mn is clearly visible for PDI acceptors A1 and A2 with minimal change in phase separation observed for A3.

cases, polymers prone to heavy preaggregation at higher Mns, such as D1, can exhibit significant chain entanglment that limits extended π-conjugation and depresses carrier mobility.81,101,102 Therefore, to examine charge transport in the vertical direction, space-charge limited current (SCLC) measurements were carried out on diode devices with a hole-only or electrononly architecture (see full fabrication details in the Supporting Information), and results are summarized in Tables 5 and S5− S6 and Figures S15−S17. Table 5. SCLC Hole Mobilities of D1 and D1:A1−A3 Filmsa D1 D1:A1 D1:A2 D1:A3

D1−8

D1−19

D1−30

D1−45

D1−59

341.00 1.84 0.04 2.00

47.00 0.21 0.11 1.03

6.11 4.55 0.99 1.26

3.48 4.98 1.87 2.20

1.78 5.90 23.80 2.16

Figure 5. Schematic illustration of the morphology templating process through either a PDI-directed or polymer-directed pathway.

a

D2−12 (Figure S8d) and D2−36 (Figure S8f) are visible in the form of bundled needles ∼50 nm in width resulting in much larger aggregates. The increased degree of phase separation and large A2 domains present in D2−12 and D2−36 are not conducive to high-performing devices, with more recombination likely occurring when large PDI acceptor aggregates are present.99,100 Clearly, the more homogeneous D2−25:A2 films, with domains on the nanometers scale, are much more likely to possess greater donor−acceptor interfacial area, which is consistent with the larger observed Jsc and PV performance of this blend. Similar to the D1:A3 films, the D2:A3 films again show no distinguishable phase separation at any Mn (Figure S8, panels g−i), likely reflecting the more amorphous nature of the A3. Active Layer Charge Transport. Mn is known to significantly affect the charge transport characteristics of polymeric semiconductors, with larger Mn values typically correlating with greater carrier mobility due to enhanced πstacking and interchain overlap.69,82−84,87 However, in some

Pristine D1 films exhibit a clear trend of decreasing hole mobility from 3.41 × 10−3 to 1.78 × 10−5 cm2/(V s) upon increasing Mn from 8 to 59 kg/mol, likely due to increased chain entanglement or/and grain boundary effects in the highMn crystalline polymers. These findings are contrary to those for the majority of PSC polymers which exhibit increasing mobilities with higher Mns.85 The overall variation in mobility from the lowest Mn to the highest is not as dramatic here as for previously reported polymer systems and subsequently enables high PSC performance at both high and low Mns (compare D1−59:A2 and D1−8:A3). Note that virtually amorphous D2 films exhibit the reverse trend, with hole mobilities increasing from 5.02 × 10−6 to 8.00 × 10−3 cm2/(V s) with Mn increasing from 12 to 36 kg/mol. Overall, D2 film charge transport increases with increasing Mn in accordance with low variation in aggregation between high and low Mn fractions of the polymer compared to D1 (evidenced by the solution UV−vis spectra, Figure 1a). Charge transport in pristine acceptor films was measured in electron-only devices of structure ITO/ZnO/A1-A3/LiF/Al. 4437

Mobilities (μh) shown in units of 10−5 × cm2/(V s).


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Chemistry of Materials Average SCLC electron mobilities of 4.65 × 10−5 cm2/(V s), 6.07 × 10−5 cm2/(V s), and 3.35 × 10−5 cm2/(V s) were measured for A1, A2, and A3, respectively (Table S7). A2 has the highest electron mobility, consistent with its observed higher PCEs. The slightly lower mobility of A1 may be due to its strong aggregating properties, expected to introduce increased grain boundary densities.103,104 The most amorphous PDI, A3, expectedly displays the lowest mobility with a structure designed to minimize self-stacking and aggregation. However, the difference in mobility between pristine A1, A2, and A3 films is small compared to the respective photovoltaic performance, indicating that pristine film mobility does not dictate the overall PSC performance and that blend film morphology plays an all-important role. To gain insight into the BHJ film transport properties, blend film hole SCLC mobilities were measured for D1:PDI and D2:PDI for all Mns (Tables 5 and S6 and Figures S15−16). A1 blend film mobilities display no dependence on D1 or D2 Mns and remain at ∼5 × 10−5 cm2/(V s), again supporting the highly templating nature of A1. In contrast, D1:A2 blend film hole mobilities monotonically increase from 3.89 × 10−7 to 2.38 × 10−4 cm2/(V s) by increasing D1 Mn from 8 to 59 kg/ mol, mirroring the PCE trends. In D2:A2 blend films, the hole mobility reaches a maximum of 2.24 × 10−6 cm2/(V s) for blends with D2−25 and decreases for both higher and lower Mn D2. This trend closely tracks the observed trend in device PCE. Finally, D1:A3 blends exhibit no trend in hole mobility, remaining at ∼2 × 10−5 cm2/(V s). However, D2:A3 blends achieve a maximum mobility value of 1.80 × 10−6 cm2/(V s) with D2−25, which falls for both higher and lower Mns, similar to the A2 case. Again, the observed trends in mobility closely track trends in device performance. GIWAXS. To elucidate the pristine and blend films morphologies, grazing incidence wide-angle X-ray scattering (GIWAXS) experiments were conducted. The GIWAXS measurements indicate that the crystalline polymer domains of both pristine and blend D1 and D2 polymer films exhibit an overall preference for π-face-on oriented polymer chains. Thus, those face-on domains are of greatest interest. Line cuts focus on the lamellar (100) peak (in-plane line cut) and the π−π stacking (010) peak (out-of-plane line cut) for the pristine D1 polymer and the pristine PDI acceptors and all of the resulting blends are shown in Figure 6. Full two-dimensional (2D) scattering images and line cuts are provided in the Supporting Information, as well as similarly focused face-on line cuts for the pristine D2 and D2:PDI blends (Figures S10 and S12). Figure 7 shows the d-spacings and correlation lengths of the polymer face-on domains in all films versus donor polymer Mn. Correlation lengths are calculated using the method of Smilgies to account for the broadening from a 2D detector.105 Pristine D1 films (Figure 6 (panels c and d) and 7a) exhibit a steady decline in (100) correlation length as Mn increases, from 224.9 Å for D1−8 to 90.7 Å for D1−59, indicating that the largest crystallites form in the lower Mn polymer films. D1−19, despite having slightly smaller domains than D1−8, exhibits the strongest peak intensity for both the lamellar (100) and π−π stacking (010) reflections, confirming that it has the highest crystallinity of the samples investigated. Note that this observation corroborates the previously reported DSC results for D1 polymers, which established that the maximum heat of fusion for D1 is for Mn = 19 kg/mol.69 Unlike D1, pristine D2 films show relatively constant (100) correlation lengths of ∼40 Å (Figure 7c) and are not influenced by Mn, although a gradual

Figure 6. GIWAXS line cuts focused on the face-on domains. The (100) peaks are result from in-plane (qxy) cuts while the (010) peaks are from out-of-plane (qz) line cuts.

increase in peak intensity as Mn is increased indicates higher crystallinity for higher Mn samples (Figure S12, panels a and b). A small increase in (100) d-spacing from 22.5 to 23.0 Å (Figure 7c) along with a similar increase in (010) d-spacing from 3.83 to 3.86 Å with a correlation length of ∼11 Å (Figure 7, panels g and h) is observed when the D2 Mn is increased from 12 to 36 kg/mol. The pristine PDI acceptor films (Figure 6, panels a and b) exhibit varying degrees of crystallinity, with A1 having obvious crystalline features visible in the 2D GIWAXS image (see Figure S11). Images of A2 show a significant decline in crystallinity relative to A1; however, they still retain some sharpness in the long-range (low q) diffraction peak. In contrast, A3 exhibits significant broadening of the long-range ordering diffraction peak and no visible features other than amorphous scattering in the short-range (higher q) scattering pattern. With regard to the BHJ films, note that the D1:A1 blend films (Figure 6, panels e and f) retain the crystalline features of A1 in all blends, although they are more pronounced in blends with D1−30 and D1−59. Notably, the D1−30:A1 blends exhibit significantly sharper A1 peaks than in the pristine A1 film. The polymer (100) correlation lengths maintain the same trend as the pristine D1 films, and the d-spacings remain constant at ∼25.5 Å (Figure 7, panels a and b). However, the (010) peak does appear to significantly shift and broaden at higher Mns. This is likely a result of the overlap between the 4438


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Figure 7. d-spacings and correlation lengths of (a−d) in-plane lamellar (100) and (e−h) out-of-plane π−π (010) peaks for the polymer domains in neat and blend D1 and D2 films. Correlation lengths are calculated using a modified Scherrer analysis that accounts for instrumental broadening from a 2D detector.105

the overall BHJ PCE trends with a systematic increase of donor polymer aggregation controlled by Mn. Acceptor A1 introduces negligible variations in device performance across all Mns tested for D1. The combination of TEM and GIWAXS (Figures 4, 6, 7, and S8−S14) trends indicate that A1 has a more crystalline nature than any of the D1 batches and largely dictates the film morphology by templating the surrounding polymer structure as illustrated with exaggeration in Figure 8a. The small variation in device performance, ranging from 2.62% to 2.77%, with only minimal variation in Jsc supports this assertion. Ring-fused and less crystalline PDI acceptor A2 affects the D1:A2 blend morphology very differently. The evolution in TEM images (Figure 4, panels f−j) from only PDI crystallites visible to blends with significant polymer tendrils indicates that A2 directs blend morphology at lower polymer Mn (Figure 8b) and shows no significant changes in TEM images from D1−8 to D1−30 Mns (Figure 8c). However, upon increasing Mn further, the polymer domain sizes increase as D1 becomes the directing component, reflecting increased pre-aggregation, and the blend morphology is then templated by the polymer. This is depicted in Figure 8d as the polymer structure begins to aggregate and define the blend morphology. The sharp increase in FF and the subsequent PCE increase from 2.41% to 3.38% is attributed to this evolution in domain size and templating capacity. These results are consistent with the GIWAXS analysis since disparate A2 and D1 (100) reflections are observed to merge as Mn is increased. The increased performance of the D1:A2 devices at high Mn reflects the increased hole mobility through the larger, polymer-directed domains. Indeed, the SCLC measurements on blend films of these materials directly correlate with device performance as the hole mobility maximizes at the highest Mn (D1−59). In contrast to the aforementioned A1 results, the D1:A3derived blend morphologies are almost completely polymerdriven. A3 is the most amorphous acceptor in this study as seen in the nearly featureless GIWAXS images and the lack of distinct PDI reflections in the blend GIWAXS data. Additionally, in blend films, the polymer domains exhibit properties almost identical to those evident in pristine polymer films, and TEM indicates a low degree of A3 phase separation from D1,

polymer and PDI reflections given the intensity and sharpness of the A1 peaks in the D1:A1 series. Note that the D1:A2 films (Figure 6, panels g and h) show a gradual intermixing of the polymer and PDI acceptor components with increasing D1 Mn. In D1−8:A2 and D1−19:A2, the A2 peak is readily identifiable as a significant shoulder on the polymer (100) feature. As Mn increases, this shoulder becomes less obvious and a shift in the (100) d-spacings from 25.3 to 24.5 Å is observed (Figure 7b), implying a gradual shift/contraction in the interaction between the polymer and PDI acceptor. The D1:A3 films (Figure 6, panels i and j) parallel the features of the pristine D1 in (100) correlation length and exhibit a constant but slightly larger dspacing (∼26 Å) (Figure 7b). In congruence with the amorphous nature of neat A3 films, no distinct A3 diffraction features are detected in the blend, only the polymer reflections. The D2:A1 blend films display sharp PDI diffraction peaks at 1.4 and 1.5 Å visible in the 2D images (Figure S10), suggesting crystalline features and significant phase separation between the blend components, in agreement with the TEM results (Figure S8), with marked decline in A1 crystallinity for D2−36. The D2:A2 blends exhibit additional features across polymer Mn, visible in both the 2D images and face-on line cuts (Figure S12, panels e and f), with some decline in sharpness of these features in D2−25 (Figure 7g), in accord with the TEM and PCE results. The crystallinity in each blend is relatively independent of polymer Mn, indicating that the large aggregates observed in the TEM have similar crystallinities. Similar to D1:A3 blends, the D2:A3 blends exhibit no A3 reflections (Figure S12, panels g and h) but instead exhibit very similar diffraction to neat D2. One notable exception is that the (100) d-spacings are ∼3 Å longer than in neat D2, suggesting that there is some PDI acceptor intercalation (Figure 7d).106 The absence of features in the A3 blends is consistent with the TEM data and with the amorphous character (Figure 4, panels k−o).

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DISCUSSION D1-Based PSCs. Three unique PCE trends are discerned here in PSCs fabricated from D1 with A1, A2, and A3. The combined optoelectronic and morphological measurements highlight the importance of acceptor crystallinity in governing 4439


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FF relative to blends D2−12 and D2−36. TEM images of the D2−25:A2 blend reveal an optimal small-scale phase separation with nanometer-sized A2 aggregates (Figure S8e), creating greater donor−acceptor interfacial area, increased current, and less recombination in the regime where neither component dominates the morphology (Figure 8h). The increase in Jsc is in part a direct result of this morphology enhancement as well as increased mobility of the blend film. Increasing the D2 Mn to 36 kg/mol significantly decreases the FF and slightly reduces Jsc. This blend contains large-scale A2 aggregates visible by TEM (Figure S8f). As D2 crystallinity increases, this observation is consistent with the polymer templating the active layer and excluding A2, which forms large self-aggregates, ultimately compromising the interpenetrating donor−acceptor morphology (Figure 8i). Reduced donor−acceptor interfacial area at the scale visible by TEM significantly increases recombination due to the greater distances excitons must travel before splitting, leading to lower Jsc and FF values. The blend film mobility data shown in Table S6 corroborate this model with a maximum hole mobility in the D2−25 blend and not D2−36, indicating increases in trap density and grain boundaries in the higher D2 Mn devices.107,108 Finally, the D2:A3 devices exhibit a trend similar to D2:A2. A peak PCE of 3.37% for D2−25 and lower PCE at higher and lower D2 Mns. The D2−25 blend performance increase reflects an enhanced Jsc. Interestingly, with higher D2 Mn, Jsc increases, but the FF falls; with lower D2 Mn, the Jsc falls while the FF rises. However, as for D1:A3, there is little variation in TEMvisible phase separation across this series, exemplifying the low templating ability of A3 (illustrated in Figure 8j).

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CONCLUSIONS BHJ PSCs employing three nonfullerene PDI-based acceptors, A1−A3, and two different donor polymers of varied crystallinity, PTPD3T (D1) and PBDTT-FTTE (D2), were fabricated and characterized. D1:PDI PSCs display three different and unique trends in PCE as Mn is increased. For the blends with the most crystalline A1, PCEs are insensitive to Mn while the cells based on less crystalline A2 and poorly crystalline A3 exhibit a monotonic increase and decrease, respectively, in the performance as Mn increases. Acceptor A1 dominates the morphology evolution of all blend films with minimal Mn effects. In the case of less crystalline A2, corresponding D1:A2 blends undergo a transition in which blend component templates the morphology, dominated by A2 at low D1 Mn and transitioning to polymer templating at high Mn, paralleled by a significant PCE increase. The most amorphous acceptor, A3, has no such transition, and the morphology is templated by the polymer at all Mns. Importantly, the PCE for corresponding D1:A3 devices decreases by ∼30% upon increasing polymer Mn, a hitherto unprecedented performance trend among systems implementing uniform fabrication conditions. In contrast, D2:PDI devices generally display the highest PCEs for the medium D2 Mns due to enhanced donor−acceptor blending and interfacial area for charge separation, evident in increased Jsc values. This trend is best exemplified in the D2−25:A2-based devices which achieve a maximum PCE of 5.2%. Overall, using three differently crystalline PDI-based SMAs, we demonstrated that BHJ morphology can be controlled by the aggregation properties of the donor polymer. The dynamic trends observed between these two variables (acceptor crystallinity and donor aggregation) offer guidelines in the

Figure 8. Illustrative cartoon of the relative aggregation tendencies and templating effects with resulting morphologies of donor polymers D1 and D2 with the indicated PDI acceptors.

regardless of the Mn (Figure 8e). The greater PSC performance of 3.41% in devices with lower Mn D1 likely reflects a combination of the uninterrupted low-Mn polymer domains forming larger crystallites, indicated by Scherrer analysis, as well as the relatively low templating capacity of the low-Mn polymer, thereby enabling some PDI domains to form and transport charges. The slight fall in blend film electron mobility when D1 Mn is increased (Table S7) accounts for the increased Jsc and FF at low D1 Mn values. D2-Based PSCs. Donor polymer D2 is less crystalline than D1 and thus less morphology-directing than D1 at all Mns (see Figure 8 for illustration). Unlike D1, where crystallinity maximizes at lower Mn, D2 crystallinity increases with increasing Mn. The PCE trends in D2-derived PSCs also differ significantly from those with D1, with each acceptor performing virtually equivalently in the D2−12 and D2−36 batches and medium Mn D2−25 providing the highest PCEs with the A2 (5.21%) and A3 (3.37%) acceptors. D2:A1 devices exhibit little dependence of performance (PCE ∼ 2.20%) on D2 Mn, which is not surprising from the aforementioned D1:A1 performance trends and relative D1 versus D2 crystallinities. A1 templates the morphology (Figure 8f) regardless of the polymer Mn, ultimately affording similar PSC performance. The D2:A2 devices display optimal PCEs of 5.21% with the intermediate Mn polymer D2−25, reflecting increased Jsc and 4440


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Design (CHiMaD) and E.F.M (X-ray characterization) by Qatar NPRP Grant 7-286-1-049.

development and fabrication of highly efficient PSCs with new nonfullerene materials and appropriately selected donor polymers.

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ASSOCIATED CONTENT

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S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b00964. Synthetic procedures and characterization data of all new compounds and details for all device characterizations (PDF)

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AUTHOR INFORMATION

Corresponding Authors

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REFERENCES

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ORCID

Michael R. Wasielewski: 0000-0003-2920-5440 Ferdinand S. Melkonyan: 0000-0001-8228-9247 Tobin J. Marks: 0000-0001-8771-0141 Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Research supported as part of the Argonne-Northwestern Solar Energy Research (ANSER) Center, an Energy Frontier Research Center funded by the U.S. Department of Energy (DOE), Office of Science, Basic Energy Sciences (BES) under Award Number DE-SC0001059 (N.D.E., R.P.H.C., and T.J.M. for device fabrication and characterization, T.J.F. and L.X.C. for materials and film characterization, P.E.H. and M.R.W for acceptor synthesis) and by AFOSR Grant FA9550-08-1-0331 (A.F. for project direction and data analysis). Use of the Advanced Photon Source, an Office of Science User Facility operated for the U.S. Department of Energy Office of Science by Argonne National Laboratory, was supported by the U.S. DOE under Contract DE-AC02-06CH11357. This work made use of the EPIC, Keck-II, and SPID facilities of Northwestern University’s NUANCE Center, which has received support from the Soft and Hybrid Nanotechnology Experimental (SHyNE) Resource (NSF ECCS-1542205); the MRSEC program (NSF DMR-1121262) at the Materials Research Center; the International Institute for Nanotechnology (IIN); the Keck Foundation; and the State of Illinois, through the IIN. A.S.D. (polymer synthesis and experimental design) thanks the Camille and Henry Dreyfus Postdoctoral Program in Environmental chemistry for a fellowship, and T.J.A. (polymer synthesis and characterization) thanks the NSF for a predoctoral fellowship. F.S.M. (project direction and data analysis) was supported by Award 70NANB14H012 from U.S. Department of Commerce, National Institute of Standards and Technology as part of the Center for Hierarchical Materials 4441


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