Molar Mass versus Polymer Solar Cell Performance: Highlighting the

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Molar Mass versus Polymer Solar Cell Performance: Highlighting the Role of Homocouplings Tim Vangerven, Pieter Verstappen, Jeroen Drijkoningen, Wouter Dierckx, Scott Himmelberger, Alberto Salleo, Dirk Vanderzande, Wouter Maes, and Jean V. Manca Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.5b00939 • Publication Date (Web): 24 Apr 2015 Downloaded from http://pubs.acs.org on May 3, 2015

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Chemistry of Materials

Molar Mass versus Polymer Solar Cell Performance: Highlighting the Role of Homocouplings Tim Vangerven†*, Pieter Verstappen‡, Jeroen Drijkoningen†, Wouter Dierckx†, Scott Himmelberger§, Alberto Salleo§, Dirk Vanderzande‡∇, Wouter Maes‡∇ and Jean V. Mancaф †

Material Physics Division, Institute for Materials Research (IMO-IMOMEC), Hasselt University, Universitaire Campus – Wetenschapspark 1, B-3590 Diepenbeek, Belgium

‡

Design & Synthesis of Organic Semiconductors (DSOS), Institute for Materials Research (IMO-IMOMEC), Hasselt University, Agoralaan 1 – Building D, B-3590 Diepenbeek, Belgium §

Department of Materials Science and Engineering, Stanford University, 476 Lomita Mall, Stanford, California 94305, United States ∇Imec

vzw, Division IMOMEC, Wetenschapspark 1, B-3590 Diepenbeek, Belgium

ф

X-LaB, Hasselt University, Universitaire Campus, Agoralaan 1, B-3590 Diepenbeek, Belgium

ABSTRACT: Although a strong link between the molar mass of conjugated polymers and the performance of the resulting polymer:fullerene bulk heterojunction organic solar cells has been established on numerous occasions, a clear understanding on the origin of this connection is still lacking. Moreover, the usual description of molar mass and polydispersity does not include the shape of the polymer distribution, although this can have a significant effect on the device properties. In this work, the effect of molar mass distribution on photovoltaic performance is investigated using a combination of structural and electro-optical techniques for the state-of-the-art low bandgap copolymer PTB7. Some of the studied commercial PTB7 batches exhibit a bimodal distribution, of which the low molar mass fraction contains multiple homocoupled oligomer species, as identified by MALDI-TOF analysis. This combination of low molar mass and homocoupling drastically reduces device performance, from 7.0 to 2.7%. High molar mass batches show improved charge carrier transport and extraction with much lower apparent recombination orders, as well as a more homogeneous surface morphology. These results emphasize the important effect of molar mass distributions and homocoupling defects on the operation of conjugated polymers in photovoltaic devices.

1. Introduction In recent years, we have witnessed considerable progress in the field of polymer solar cells. Power conversion efficiencies (PCE’s) above 10% for single junction and even 11.5% for triple junction devices have been realized.1-3 The application of low bandgap copolymers with broad absorption profiles and optimized interlayers has led to higher photon absorption and improved charge collection efficiency. Among the state-ofthe-art electron donor polymer materials, PTB7 (poly{[4,8bis(2-ethylhexyloxy)benzo[1,2-b:4,5-b′]dithiophene-2,6-diyl]alt-(3-fluoro-2-[(2-ethylhexyloxy)carbonyl]thieno[3,4b]thiophene-4,6-diyl)} is known to afford very high efficiencies for single-junction devices (7−9%).1,4-7 To reach optimal performance, the polymer is combined with a methanofullerene derivative (PC71BM) and the active layer blend is spincoated from an organic solvent (e.g. chlorobenzene) with the addition of 1,8-diiodooctane (DIO) to selectively dissolve PC71BM,7,8 leading to direct access of the polymer chains to pure PC71BM clusters, facilitating charge separation.9 Two parameters of high importance which are sometimes over-

looked or not clearly mentioned are the molar mass of the polymer and the accompanying polydispersity (D). A recent study by Cao et al. showed that the photovoltaic cells can exhibit drastically different characteristics when the numberaverage molar mass (Mn) of the PTB7 donor polymer is varied.10 They reached a PCE of 8.5% (compared to 5.4%) when increasing Mn from 21 kg mol−1 (D = 1.85) to 128 kg mol−1 (D = 1.12). Their devices were processed from orthodichlorobenzene (ODCB) without DIO and with fairly thick layers of 200 nm. The authors concluded that the increased performance is a consequence of larger phase separation (as visualized by transmission electron microscopy) and increased hole mobility, leading to higher short-circuit current densities. McNeill and co-workers varied the degree of fluorination and Mn of a series of PTB7-type copolymers and found that the effect of Mn is dominant over the degree of fluorination with respect to the final photovoltaic performance.11 They also discovered that the effectiveness of DIO reduces when the Mn decreases. Granted that their study revealed interesting insights on the effect of Mn for PTB7, they were not able to attribute the effect solely to Mn as both the degree of fluorina-

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tion and the molar mass were varied at the same time. Lu et al. very recently found that when PTB7 exhibits a large polydispersity, the performance drops from 7.5% (Mn = 47.6 kg mol−1, D = 2.1) to 2.5% (Mn = 24 kg mol−1, D = 4.3).12 They concluded that dispersity is the driving factor for the observed differences among their samples, although the Mn differs significantly as well. It is suggested by the authors that this high D is caused by (homocoupling) side reactions, which introduce structural defects. Whenever Mn and D values are disclosed in a polymer solar cell paper, the polymer distribution itself is often not shown, although this can have a severe impact on the properties of the devices. So et al. reported the impact of molar mass distribution for commercially purchased poly{N-9′-heptadecanyl-2,7-carbazole-alt-5,5-[4′,7′-(di-2thienyl)-2′,1′,3′-benzothiadiazole]} (PCDTBT) batches.13 They observed that the molar mass distribution was bimodal and noticed an improvement of the solar cell properties when the low molar mass peak was reduced. The origin of this low molar mass fraction was not discussed though. In this work, we report a detailed investigation of five commercially acquired batches of PTB7, used without any further purification, by a combination of different structural and electro-optical techniques. Thorough analysis of the molar mass distributions is performed by a combination of gel permeation chromatography (GPC) and matrix-assisted laser desorption ionization time of flight (MALDI-TOF) mass spectrometry. The effect of the molar mass distribution on the electro-optical properties of neat and blend films as well as their application in polymer solar cells is analyzed. To gain deeper understanding on transport limiting processes, the hole mobility as well as the apparent recombination order is investigated. Furthermore, a strong correlation between molar mass distribution and surface morphology is observed.

2. Experimental Section Materials: Five different batches of PTB7, obtained from Solarischem (SOL4700; P1−P3, P5) and Solarmer (ZP-002; P4), were used as received. The molar mass distributions of all polymers were determined by GPC using a Spectra Series P100 (Spectra Physics, Santa Clara, USA) pump equipped with two mixed-B columns (10 µm, 2 cm × 30 cm, Polymer Laboratories) and an Agilent 1100 diode array UV detector. The GPC was calibrated against linear (narrow) polystyrene standards. Chlorobenzene was used as the eluent at a flow rate of 1.0 mL min−1 at 60 °C. None of the batches showed visible formation of aggregates when dissolved in chlorobenzene. 1H NMR spectra (400 MHz) of batches P1 and P4 were measured in a mixture of CDCl3 and CS2 (1:3). MALDI-TOF mass spectra were recorded on a Bruker Daltonics Ultraflex II Tof/Tof. 1 µL of the matrix solution (4 mg mL−1 DTCB (trans-2-[3-(4tert-butylphenyl)-2-methyl-2-propenylidene]malononitrile) in CHCl3) was spotted onto an MTP Anchorchip 600/384 MALDI plate. The spot was allowed to dry and 1 µL of the analyte solution (0.5 mg mL−1 in CHCl3) was spotted on top of the matrix. Devices: Solar cells were manufactured on patterned ITO/glass substrates (Kintec, 20 Ω sq−1) which were ultrasonically cleaned with detergent, distilled water, acetone, and boiling isopropanol and dried with a nitrogen flow. Prior to spin-casting of a thin layer of PEDOT:PSS (Clevios Al 4083, 35 nm), the substrates were treated with UV-ozone for 15 minutes. The samples were transferred into a nitrogen filled glovebox and annealed at 130 °C for 10 minutes. The active

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layer contained a blend of PTB7:PC71BM (Solenne, 99%) in a 1:1.5 ratio with a total concentration of 25 mg mL−1 for all batches. The active materials were dissolved in chlorobenzene and stirred overnight at 50 °C. Before spin-casting of the active layer, 3 vol% DIO (Sigma Aldrich) was filtered and added to the solution at room temperature. All polymer solar cells were individually optimized for layer thickness. The samples were dried in a vacuum atmosphere (3 x 10-7 mbar) for 4 hours to eliminate all DIO. This was followed by spin-casting methanol on top of the active layer at 2500 rpm for 40 seconds. Methanol was completely evaporated before deposition of the metal contacts. The samples were completed by thermal evaporation of calcium (30 nm) and aluminum (120 nm) electrodes at a base pressure of 3 x 10-7 mbar through shadow masks, resulting in an active area of 3.0 mm². Transistor substrates were obtained from Philips Innovation Services (200 nm thick SiO2 dielectric on a highly n-doped Si wafer, with 100 nm gold on top of 10 nm titanium contacts, and coated with a HMDS layer). The patterning was done in such a way that the device length L and width W were 20 mm and 10 µm, respectively. The dielectric capacitance per unit area was C0 = 16.9 nF cm-2. The cleaning and spin casting steps for the transistors were the same as for the solar cell devices. Measurements: Photoactive layer thicknesses were recorded by a Bruker Veeco Dektak XT Profilometer. Ultravioletvisible (UV-vis) absorption spectra were recorded for films deposited on standard glass substrates using an Agilent Varian Cary 500. Current-voltage (I-V) characterization was performed using a class A solar simulator (Newport 91195A), calibrated with a silicon NIST reference solar cell to obtain 100 mW cm-² under AM1.5G conditions. External quantum efficiencies (EQE’s) were acquired by recording the monochromated (Newport Cornerstone 130 with sorting filters) output of a xenon lamp (100 W, Newport 6257) by a lock-in amplifier (Stanford Research Systems SR830). The light beam was mechanically chopped at 10 Hz. The recorded values were calibrated with a FDS-100 calibrated silicon photodiode. For the transient photovoltage (TPV) and transient photocurrent (TPC) measurements, a pulsed laser (Continuum minilite II, 532 nm) was used. The laser power was attenuated by neutral density filters to obtain a small perturbation which never exceeded 20 mV (∆V 100 kg mol−1) and three samples with low Mn (< 20 kg mol−1). Furthermore, the three low molar mass samples exhibit a bimodal distribution and consist out of a fraction of short chain oligomers (peak molar mass Mp ~ 3−4 kg mol−1) and a fraction of higher molar mass polymers (Mp ~ 20−30 kg mol−1). For these three samples, no large deviations of the Mp values of the different constituent fractions are observed, but the molar mass variation is rather due to a different ratio of the two molar mass fractions. The relative abundance (or the peak ratio) of both fractions was determined by taking the ratio of the integrals of the first (lowest molar mass) peak and the complete distribution from the GPC traces. In the lowest molar mass sample P1, the oligomer fraction accounts for approximately 28% of the total amount, while in the 20 kg mol−1 fraction P3, the contribution of the oligomer fraction decreases to approximately 8%. To visualize the peak ratios more clearly, linear molar mass plots (rather than the logarithmic plot in Figure 1) are added in Figure S1. Due to the fact that PTB7 batches P1−P3 are composed of two fractions, the polydispersities of these samples are slightly higher than those for P4 and P5.

Figure 1. Molar mass distributions of the as received PTB7 batches (as measured by GPC).

3.2. Electro-Optical Properties

Mp [kg mol-1] 21/3.3 25/3.7 28/3.9 141 238

Peak ratio [%] 28 22 8 / /

D 3.2 2.6 2.6 2.1 1.8

Polymers P1−P5 were blended with PC71BM and used as photoactive materials in polymer solar cells with a standard architecture.5 All devices were optimized according to layer thickness for each batch separately and processed under identical conditions. The typical photovoltaic parameters are shown in Table 2. The current-voltage curves of the optimized devices are depicted in Figure S2 together with the corresponding active layer thicknesses. There is a severe difference in performance between the different batches, yielding a maximum efficiency of 7.0% for P4 and P5. The short-circuit current density (Jsc) and fill factor (FF) are the main reasons for the improved performance. There is a small but significant decrease in open-circuit voltage (Voc) for P1−P3. When having a closer look to these low molar mass samples, it becomes evident that when the peak ratio decreases from 28 to 8% (Table 1), the PCE increases from 2.7 to 4.0%. Table 2. Overview of the photovoltaic parameters for the optimized polymer solar cells prepared in the standard device architecture ITO/PEDOT:PSS/PTB7:PC71BM/MeOH/Ca/Al.

Batch P1 P2 P3 P4 P5

Voc

Jsc JV

Jsc EQE

FF

η

[mV]

[mA cm-²]

[mA cm-²]

[%]

[%]

670 675 690 740 720

7.3 8.4 9.7 14.1 15.2

7.9 8.0 9.2 13.5 14.5

56 57 61 67 64

2.7 3.3 4.0 7.0 7.0

EQE measurements were used to examine whether the differences in Jsc can be allocated to the functioning of the donor polymer (Figure 2). Integration of the EQE spectra over the AM1.5G solar spectrum yields current densities which are within ~10% of the Jsc values as listed in Table 2. By increasing the chain length, the EQE values are boosted strongly in the range of 550 to 750 nm, which corresponds to the absorption window of PTB7,4,11 revealing improved photon absorption by the polymer and enhanced charge collection. The optical absorption of PTB7 samples P1−P5 was further investigated by UV-Vis spectroscopy for neat and blend films. Figure 3a shows the normalized optical absorption spectra of the neat PTB7 films. Polymers P4 and P5 exhibit similar profiles as reported in literature, with a strong absorption from 550 to 750 nm and a rather steep band edge.4,11 The vibronic peak at 677 nm is more pronounced and slightly red shifted for P5 as compared to P4 (667 nm). The presence of these vibronic peaks is used as an indication for the formation of polymer aggregates.14 PTB7 samples P1−P3, however, do not show vibronic features and, moreover, exhibit a broadened absorp-


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tion extending into the near-infrared (NIR) region. This bathochromically shifted absorption tail has previously tentatively been attributed to the presence of homocoupling defects in the polymer structure (vide infra).12,15 The optical absorption spectra of the PTB7:PC71BM blend films exhibit the same characteristics, with extra absorption in the 300 to 550 nm range caused by PC71BM (as shown in Figure 3b).

Figure 2. EQE spectra of the optimized devices prepared with the different PTB7 batches.

Figure 3. Optical absorption spectra of (a) PTB7 neat films, and (b) PTB7:PC71BM (1:1.5) blend films. In order to better understand the microstructure of the PTB7:PC71BM films, grazing incidence X-ray diffraction (XRD) measurements were performed. The PTB7 in the films showed a predominantly face-on packing orientation with the π-conjugated backbone parallel to the substrate. All pure and

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blend samples exhibited a π-π stacking distance of 3.9 Å, which is in agreement with literature data.16,17 Noriega et al. showed that in high molar mass semiconducting polymers, short range intermolecular aggregation is crucial for efficient long range charge transport.18 Quantification of the paracrystalline disorder did not yield noteworthy differences between the various molar mass fractions, with all samples exhibiting large amounts of disorder (g > 15%), but some aggregation in the π-π stacking direction. While not vital for charge transport, the alkyl stacking distance and coherence length, however, were significantly different between the high and low molar mass samples (Table S1, Figure S11). PTB7 batches P4 and P5 showed both a shorter spacing and coherence length in the alkyl stacking direction. We hypothesize that the greater occurrence of BDT (benzo[1,2-b:4,5-b′]dithiophene) homocoupling defects in the low molar mass fractions (vide infra) inhibits the close packing of adjacent lamellae due to the increase in monomers with side chains on both sides of the backbone. The lower coherence length in the high molar mass fractions is attributable to the increased difficulty for long, entangled chains to form long-range ordered structures. It appears that the molecular level packing is likely not a main contributor to the observed device performance differences.19 To gain more information on the compositional difference between PTB7 samples P1−P3 and P4−P5, the samples were subjected to MALDI-TOF mass spectrometry analysis. MALDI-TOF was not successful for P4 and P5 due to the absence of short chain oligomers. The mass spectra of P1−P3 (see Figure S3−S8) are quite similar and a complete peak assignment was done for P1. An important result obtained from MALDI-TOF is that homocoupling (i.e. the coupling of two identical building blocks rather than the expected crosscoupling) is observed for both the brominated TT (thieno[3,4b]thiophene) and the stannylated BDT monomer. These results confirm the recent work by Lu et al. that homocoupling defects can reside within the polymer distribution and have a noticeable effect on photovoltaic performance.12 Lu and coworkers provided evidence for homocoupling via model Stille reactions, spectroscopic studies and the synthesis of polymers with specifically designed dimeric TT repeating units,20 and an explanation based on energy transfer and dipolar effects was provided. From the intensity of the mass peaks it can be speculated that the stannylated BDT monomers undergo homocoupling more regularly in comparison with the brominated TT monomers. However, it is not possible to quantify the amount of homocoupling in the materials studied, especially since the mass spectra predominantly (only) show the oligomer fractions. Nevertheless, as the UV-Vis absorption spectra of these samples are nearly identical, no major difference in the total amount of homocoupling is expected for these samples. It is rather the relative abundance of the low molar mass fraction containing the homocoupling defects that differentiates these three PTB7 batches. Furthermore, it is worthwhile to mention that almost all polymer chains are terminated with methyl groups. This could indicate that a methyl shift during the transmetalation step of the Stille cross-coupling also occurs as a major side reaction, leading to termination of the polymer chains.21 In a recent study, Janssen et al. applied a systematic approach to incorporate intrachain homocoupling defects in diketopyrrolopyrrole (DPP) based push-pull copolymers.15 They noticed that these defects can increase the highest occupied molecular orbital (HOMO) and decrease the lowest unoccupied molecular orbital (LUMO), causing a lowering of the


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bandgap, resulting in the appearance of a low energy tail in the UV-Vis/NIR absorption spectra. Their findings for DPP-based systems are very similar to the results obtained for our PTB7based samples (Figure 3). Moreover, they observed a similar reduction in device performance when homocouplings were present. The decrease in Voc can be explained by the lower bandgap of P1−P3, as Voc is directly related to the energetic offset of the HOMO of the donor polymer and the LUMO of the electron acceptor in a bulk heterojunction solar cell.22

The small perturbation lifetime in the dark τ∆n0, n0 and are determined experimentally. The multiple trapping model can be used to explain the dependence of k on n. When traps are present, recombination is mainly controlled by free charges recombining with trapped charges. If an exponential trap distribution is assumed, R equals: ∝ ∙ !"#

Figure 4. FET hole mobility and Jsc in function of Mn.

3.3. Charge Transport and Recombination To further investigate the role of the observed homocoupling defects on the transport and recombination properties, several electrical transport techniques were used. Firstly, we studied the charge transport by determining the hole mobility for blend films in the in-plane direction by the field-effect transistor (FET) method, which corresponds with transport along the polymer backbone and π-π stacking. Although the FET mobility does not probe the mobility of the charge carriers in the direction of the solar cell operation mode, it can still be indicative for the device properties. The FET hole mobilities (µFET,hole) were extracted from the linear regime and are plotted together with Jsc in function of Mn in Figure 4. In general, the µFET,hole increases with increasing Mn. Looking into more detail, this is a similar behavior as observed for several other conjugated polymers.18,23,24 The highest achieved mobility was 1.6 x 10−4 cm2 V−1s−1 for Mn = 242 kg/mol (P5). The reduced mobility in the low molar mass fractions can be attributed to an increase in slow, intermolecular charge transport.25 Additionally, the observed homocouplings for P1−P3 can be expected to act as traps, due to the high nature of localization of the orbitals at the defect,15 severely reducing parameters like µ, Jsc and FF. The combination of TPV and TPC can be used to investigate the number of trap states present in a photovoltaic device, where we assume the charge generation rate to be field independent.26-29 The non-geminate recombination rate R is frequently described by a bimolecular Langevin model. However, when the recombination coefficient k is dependent on the charge carrier density n, R can be described by:26,30,31 ∙ ∝ ∙

(1)

wherein λ+1 represents the apparent recombination order, which is an empirical reaction order related to n, and k0 is the constant part of k which can be expressed as:

(2)

$% &' %

!"#

∝

(3)

wherein ET stands for the Urbach energy of the trap distribution, kB is the Boltzmann constant and T the temperature. The apparent recombination order itself should not be considered as a physical mechanism, but it provides information about the number of traps present in the photoactive layer. From the TPV and TPC measurements, the small perturbation lifetime τ∆n as a function of n can be determined. By taking the slope of the log-log plot of the small perturbation time versus the charge carrier density, the empirical reaction order -λ can be obtained and k can be calculated according to equation 2. When there is strictly bimolecular recombination, λ should be equal to 1 and the recombination coefficient is expected to be a constant. However, when traps are present, the recombination coefficient k is charge carrier concentration dependent, yielding slopes with λ > 1, where higher values for λ imply a higher trap concentration. Plots of the small perturbation lifetime and recombination coefficient as a function of the charge carrier density are gathered in Figure S12 and S13. The slopes validate the charge density dependence of k on n, implying the existence of trap states in the photoactive layer. From this data, the apparent recombination order can be calculated for the different polymer distributions. Kirchartz et al. showed that the reaction order can depend on the layer thickness of the device and that caution is needed when analyzing results with high reaction orders.32,33 It is often assumed that charge carriers are distributed uniformly while in reality the spatial variation can affect the charge carrier dependence of the recombination rate. The apparent recombination orders for the optimal devices are illustrated in Figure 5 by the black squares, while the open purple squares represent values for samples with thicker layers (> 170 nm). The exact layer thicknesses can be found in Figure S12 and S13. From Figure 5 it is clear that samples P4 and P5 show much lower apparent recombination orders compared to P1−P3, meaning that fewer traps are present for P4 and P5 (according to equation 3). It is not possible to effectively distinguish between P1−P3, probably due to the high level of structural defects caused by the homocouplings and short chain segments. It has to be pointed out that not all recombination orders for the corresponding batches are lowered with increasing layer thickness for PTB7:PC71BM cells, as proposed by Kirchartz et al.32 This effect could originate from the different morphologies for each batch, yielding different spatial distributions. It is widely accepted that the morphology of the photoactive layer plays a crucial role in the functioning of a polymer:fullerene solar cell.34,35 The phase separation is often controlled by using various solvents combined with additives.36 The degree of phase separation can be manipulated though by adjusting the Mn of the donor polymer.14


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properties, and hence high power conversion efficiency. It is also clear from our findings that discussing results purely based on the Mn, Mw and D values of the donor polymer is not sufficient. The shape of the polymer distribution should always be shown, as polymer samples with high D can possess a substantial amount of low molar mass components (possibly in a bimodal composition), which heavily influence the operation of polymer solar cells.

4. Conclusions

Figure 5. Apparent recombination order +1 for the different PTB7 batches.

3.4. Morphology To further investigate these features, the surface morphology of the PTB7:PC71BM blend films was recorded with PF-QNM (PeakForce Quantum Nano Mechanical) AFM (Figure S14S18). There is a noticeable difference between the surface topographies of the samples. Analysis of the PF-QNM AFM images revealed average roughnesses (Ra) of 4.3 nm (P1), 3.6 nm (P2), 3.4 nm (P3), 1.5 nm (P4) and 1.2 nm (P5). The surface morphologies of the P1−P3 blend films appear very similar to the eye, whereas those from P4 and especially P5 are much smoother with a finer size scaling. We believe this is an indication of better intermixing of the polymer and fullerene molecules in these high molar mass fractions, which provides, together with the hole mobility, an explanation for the improved Jsc and FF. The apparent recombination order also seems to be strongly linked to the active layer morphology, as the recombination order is much higher for a coarser morphology (P1−P3). We would like to highlight that all samples were processed with DIO, keeping PC71BM longer in solution than the polymer. McGehee et al. proposed,14 based on the findings of Russell et al.,17,37 that the solubility of the polymer is the main reason for the differences in morphology with varying molar mass when the solvent processing conditions are identical. Due to the lower solubility of high molar mass polymer, the formation of polymer aggregates is promoted in solution. These aggregates tend to fuse together during spin casting, creating a polymer network template where the residing locations of the, still solubilized, PC71BM molecules are limited, thereby preventing large scale phase separation. Low molar mass polymers show much faster kinetics of phase separation and are not able to form aggregate networks. The latter is possible for very long polymer chains because different parts of the same chain can get involved in many different aggregates. Shorter polymer chains form more localized aggregates where the chain is only involved in one single aggregate. A hypothesis resulting from this work is that the polymerization process is disturbed once side reactions, e.g. the formation of homocouplings, are initiated, thereby confining the final molar mass of the polymer. Low molar mass fractions have a higher solubility and do not form aggregated polymer networks which act as a template for PC71BM molecules to reside. This leads to a less favorable morphology with increased trap states, reducing the device parameters. The molar mass of the donor polymer seems to be a key driver in achieving an optimal morphology as well as favorable charge transport

In conclusion, we have found that a variation in polymer chain length and distribution translates directly to the photovoltaic device properties and we wish to highlight the particular role of homocouplings in this respect. Five commercial batches of the low bandgap copolymer PTB7 were investigated and polymer solar cell efficiencies of 7.0% were obtained when the molar mass was sufficiently high (Mn ≥ 141 kg mol−1). However, the low molar mass batches, exhibiting a bimodal molar mass distribution, showed a severely reduced photovoltaic performance, with efficiencies ranging from 4.0 to 2.7%. By MALDI-TOF analysis, various homocoupled species were detected in the samples with bimodal distributions. The presence of these homocoupling defects is reflected in several (device) parameters. A first effect resulting from the presence of homocouplings is the lowering of the polymer bandgap. This can be directly observed in the absorption spectra, due to the appearance of an additional absorption tail in the nearinfrared region, and it leads to a lowering of the Voc. Secondly, the combination of low molar mass and homocouplings results in lower hole mobilities and higher apparent recombination orders, which suggest a higher number of trap states in these materials, limiting the charge extraction and reducing the current. Thirdly, the various batches yielded a different surface morphology, with the finest size scaling for the highest molar mass batch, which points to a better intermixing of the polymer and the fullerene. Furthermore, the occurrence of homocoupled oligomers is a clear indication that the polymerization proceeded in an unintended manner and terminated at an earlier than expected stage, limiting the molar mass of the conjugated polymer. Current work is focused on the validation of these findings for other polymer:fullerene systems and the optimization of the polymerization protocols with respect to the occurrence of side reactions, e.g. homocouplings. As a general message, experimentalists should be aware that commercial batches of conjugated polymers may contain different distributions of chain lengths than they expect, which can strongly limit device performance. Inferior performances may in some cases be attributable to the presence of unintended oligomers, originating from poor polymerization.

ASSOCIATED CONTENT Supporting Information (GPC profiles on a linear scale, J-V curves for the optimized solar cells, MALDI-TOF mass spectra and analysis, 1H NMR spectra of P1 and P4, XRD data, TPV/TPC analysis, and AFM images) is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author * [email protected] Author Contributions All authors have given approval to the final version of the manuscript.


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Notes

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The authors declare no competing financial interest.

ACKNOWLEDGMENT The authors thank Hasselt University, the Research Foundation – Flanders (FWO) and the Science Policy Office of the Belgian Federal Government (BELSPO; IAP 7/05 project FS2) for continuing financial support. T. Vangerven and P. Verstappen acknowledge the Agency for Innovation by Science and Technology in Flanders (IWT) for their PhD grants. J. Drijkoningen acknowledges the support by the ‘Strategic Initiative Materials’ in Flanders (SIM) and the IWT under the Solution based Processing of Photovoltaic Modules (SoPPoM) program. Wouter Dierckx thanks Hasselt University for his PhD scholarship. S. Himmelberger would like to thank the National Science Foundation for support in the form of a Graduate Research Fellowship. The use of the Stanford Synchrotron Radiation Lightsource, SLAC National Accelerator Laboratory, is supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Contract No. DE-AC02-76SF00515.


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REFERENCES (1) Z. He, C. Zhong, S. Su, M. Xu, H. Wu, Y. Cao, Nat. Photon. 2012, 6, 591. (2) C.-C. Chen, W.-H. Chang, K. Yoshimura, K. Ohya, J. You, J. Gao, Z. Hong, Y. Yang, Adv. Mater. 2014, 26, 5670. (3) Y. Liu, J. Zhao, Z. Li, C. Mu, W. Ma, H. Hu, K. Jiang, H. Lin, H. Ade, H. Yan, Nat. Commun. 2014, 5, 5293. (4) Y. Liang, Z. Xu, J. Xia, S.-T. Tsai, Y. Wu, G. Li, C. Ray, L. Yu, Adv. Mater. 2010, 22, E135. (5) H. Zhou, Y. Zhang, J. Seifter, S. D. Collins, C. Luo, G. C. Bazan, T.-Q. Nguyen, A. J. Heeger, Adv. Mater. 2013, 25, 1646. (6) Z. He, C. Zhong, X. Huang, W.-Y. Wong, H. Wu, L. Chen, S. Su, Y. Cao, Adv. Mater. 2011, 23, 4636. (7) L. Lu, L. Yu, Adv. Mater. 2014, 26, 4413. (8) S. J. Lou, J. M. Szarko, T. Xu, L. Yu, T. J. Marks, L. X. Chen, J. Am. Chem. Soc. 2011, 133, 20661. (9) B. A. Collins, Z. Li, J. R. Tumbleston, E. Gann, C. R. McNeill, H. Ade, Adv. Energy Mater. 2013, 3, 65. (10) C. Liu, K. Wang, X. Hu, Y. Yang, C.-H. Hsu, W. Zhang, S. Xiao, X. Gong, Y. Cao, ACS Appl. Mater. & Interfaces 2013, 5, 12163. (11) X. He, S. Mukherjee, S. Watkins, M. Chen, T. Qin, L. Thomsen, H. Ade, C. R. McNeill, J. Phys. Chem. C 2014, 118, 9918. (12) L. Lu, T. Zheng, T. Xu, D. Zhao, L. Yu, Chem. Mater. 2015, 27, 537. (13) H. K. H. Lee, Z. Li, I. Constantinou, F. So, S. W. Tsang, S. K. So, Adv. Energy Mater. 2014, 4, 1400768. (14) J. A. Bartelt, J. D. Douglas, W. R. Mateker, A. E. Labban, C. J. Tassone, M. F. Toney, J. M. J. Fréchet, P. M. Beaujuge, M. D. McGehee, Adv. Energy Mater. 2014, 4, 1301733. (15) K. H. Hendriks, W. Li, G. H. L. Heintges, G. W. P. van Pruissen, M. M. Wienk, R. A. J. Janssen, J. Am. Chem. Soc. 2014, 136, 11128. (16) W. Chen, T. Xu, F. He, W. Wang, C. Wang, J. Strzalka, Y. Liu, J. Wen, D. J. Miller, J. Chen, K. Hong, L. Yu, S. B. Darling, Nano Lett. 2011, 11, 3707. (17) F. Liu, W. Zhao, J. R. Tumbleston, C. Wang, Y. Gu, D. Wang, A. L. Briseno, H. Ade, T. P. Russell, Adv. Energy Mater. 2014, 4, 1301377. (18) R. Noriega, J. Rivnay, K. Vandewal, F. P. V. Koch, N. Stingelin, P. Smith, M. F. Toney, A. Salleo, Nat. Mater. 2013, 12, 1038. (19) F. P. V. Koch, J. Rivnay, S. Foster, C. Müller, J. M. Downing, E. Buchaca-Domingo, P. Westacott, L. Yu, M. Yuan, M. Baklar, Z. Fei, C. Luscombe, M. A. McLachlan, M. Heeney, G. Rumbles, C. Silva, A. Salleo, J. Nelson, P. Smith, N. Stingelin, Prog. Polym. Sci. 2013, 38, 1978. (20) B. Carsten, J. M. Szarko, H. J. Son, W. Wang, L. Lu, F. He, B. S. Rolczynski, S. J. Lou, L. X. Chen, L. Yu, J. Am. Chem. Soc. 2011, 133, 20468. (21) F. Brouwer, J. Alma, H. Valkenier, T. P. Voortman, J. Hillebrand, R. C. Chiechi, J. C. Hummelen, J. Mater. Chem. 2011, 21, 1582. (22) K. Vandewal, K. Tvingstedt, A. Gadisa, O. Inganas, J. V. Manca, Nat. Mater. 2009, 8, 904. (23) R. J. Kline, M. D. McGehee, E. N. Kadnikova, J. Liu, J. M. J. Fréchet, M. F. Toney, Macromolecules 2005, 38, 3312. (24) J. J. Intemann, K. Yao, H.-L. Yip, Y.-X. Xu, Y.-X. Li, P.-W. Liang, F.-Z. Ding, X. Li, A. K. Y. Jen, Chem. Mater. 2013, 25, 3188. (25) S. Himmelberger, K. Vandewal, Z. Fei, M. Heeney, A. Salleo, Macromolecules 2014, 47, 7151. (26) A. Foertig, J. Kniepert, M. Gluecker, T. Brenner, V. Dyakonov, D. Neher, C. Deibel, Adv. Funct. Mater. 2014, 24, 1306. (27) C. G. Shuttle, B. O’Regan, A. M. Ballantyne, J. Nelson, D. D. C. Bradley, J. de Mello, J. R. Durrant, Appl. Phys. Lett. 2008, 92, 093311.


8

Page 9 of 10


(28) D. Spoltore, W. D. Oosterbaan, S. Khelifi, J. N. Clifford, A. Viterisi, E. Palomares, M. Burgelman, L. Lutsen, D. Vanderzande, J. Manca, Adv.

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Energy Mater. 2013, 3, 466. (29) T. M. Clarke, C. Lungenschmied, J. Peet, N. Drolet, A. J. Mozer, Adv. Energy Mater. 2015, 5, 1401345. (30) C. Deibel, D. Rauh, A. Foertig, Appl. Phys. Lett. 2013, 103, 043307. (31) C. G. Shuttle, B. O’Regan, A. M. Ballantyne, J. Nelson, D. D. C. Bradley, J. R. Durrant, Phys. Rev. B 2008, 78, 113201. (32) F. Deledalle, P. Shakya Tuladhar, J. Nelson, J. R. Durrant, T. Kirchartz, J. Phys. Chem. C 2014, 118, 8837. (33) T. Kirchartz, J. Nelson, Phys. Rev. B 2012, 86, 165201. (34) C. Shim, M. Kim, S.-G. Ihn, Y. S. Choi, Y. Kim, K. Cho, Chem. Commun. 2012, 48, 7206. (35) G. J. Hedley, A. J. Ward, A. Alekseev, C. T. Howells, E. R. Martins, L. A. Serrano, G. Cooke, A. Ruseckas, I. D. W. Samuel, Nat. Commun. 2013, 4, 2867. (36) H.-C. Liao, C.-C. Ho, C.-Y. Chang, M.-H. Jao, S. B. Darling, W.-F. Su, Mater. Today 2013, 16, 326. (37) Y. Gu, C. Wang, T. P. Russell, Adv. Energy Mater. 2012, 2, 683.


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Graphic entry for the Table of Contents (TOC)


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Molar Mass versus Polymer Solar Cell Performance: Highlighting the

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