Fullerene

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Relating Structure to Efficiency in Surfactant-free Polymer:Fullerene Nanoparticle-based Organic Solar Cells Stefan Gärtner, Andrew J. Clulow, Ian A. Howard, Elliot P. Gilbert, Paul L. Burn, Ian R. Gentle, and Alexander Colsmann ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b15601 • Publication Date (Web): 30 Oct 2017 Downloaded from http://pubs.acs.org on October 31, 2017

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ACS Applied Materials & Interfaces

Relating Structure to Efficiency in Surfactant-free Polymer:Fullerene Nanoparticle-based Organic Solar Cells Stefan Gärtner†, Andrew J. Clulow‡, Ian A. Howard§,∥, Elliot P. Gilbert⊥, Paul L. Burn‡,*, Ian R. Gentle‡ and Alexander Colsmann†,* †

Light Technology Institute, Karlsruhe Institute of Technology (KIT), Engesserstrasse 13, 76131

Karlsruhe, Germany ‡

Centre for Organic Photonics & Electronics, The University of Queensland, St Lucia, QLD

4072, Australia §

Institute of Microstructure Technology, Karlsruhe Institute of Technology (KIT), Hermann-

von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany ∥

HEiKA Heidelberg Karlsruhe Research Partnership (HEiKA), Heidelberg/Karlsruhe, Germany

⊥

Australian Centre for Neutron Scattering, Australian Nuclear Science and Technology

Organisation (ANSTO), Locked Bag 2001, Kirrawee DC, NSW 2232, Australia.

ACS Paragon Plus Environment

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KEYWORDS Organic nanoparticle, polymer solar cell, small-angle neutron scattering, transient absorption spectroscopy, printed electronics

ABSTRACT

Nanoparticle dispersions open up an eco-friendly route towards printable organic solar cells. They can be formed from a variety of organic semiconductors by using miniemulsions that employ surfactants to stabilize the nanoparticles in dispersion and to prevent aggregation. However, whenever surfactant-based nanoparticle dispersions were used to fabricate solar cells, the reported performance remained moderate. In contrast, solar cells from nanoparticle dispersions formed by precipitation (without surfactants) can exhibit power conversion efficiencies close to state-of-the-art solar cells processed from blend solutions using chlorinated solvents. In this work, we use small-angle neutron scattering (SANS) measurements and transient absorption spectroscopy (TAS) to investigate why surfactant-free nanoparticles give rise to efficient organic solar cells. We show that surfactant-free nanoparticles comprise an even distribution of small semiconductor domains, similar to bulk-heterojunction films formed using traditional solvent processing. This observation differs from surfactant-based miniemulsion nanoparticles that typically exhibit core-shell structures. Hence, the surfactant-free nanoparticles already possess the optimum morphology for efficient energy conversion before they are assembled into the photoactive layer of a solar cell. This structural property underpins the superior performance of the solar cells containing surfactant-free nanoparticles and is an important design criterion for future nanoparticle inks.


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INTRODUCTION The photoactive layers of the most efficient organic solar cells comprise blends of at least two organic semiconductors (a donor and an acceptor) to form a bulk-heterojunction. For highperformance, the bulk-heterojunction film must have the donors and acceptors distributed at a distance of no more than the exciton diffusion length to yield efficient exciton splitting. In addition, the donor and acceptor domains must provide percolation pathways for efficient charge carrier transport to the electrodes.1 Today, state-of-the-art solution processing methods use environmentally harmful solvents, with the morphologies required for optimized device performance achieved through manipulation of the wet-film drying processes, e.g., by using solvent annealing or high boiling point co-solvents.2-4 Subtle changes to the deposition conditions, e.g., solvent evaporation rates or temperature, can affect the photoactive layer morphology and hence the device performance. As a consequence, the transfer of laboratory scale fabrication methods to industrial roll-to-roll fabrication is not straightforward, and usually requires significant process modifications. Large-scale production for the commercialization of organic solar cells would benefit from the introduction of a simple, scalable and reproducible processing method as well as from environmentally benign solvents.

It has been suggested that employing polymer:fullerene nanoparticle dispersions instead of blend solutions solves two significant challenges for organic solar cell production. Firstly, solubility is no longer required for solution processing and therefore halogenated aromatic organic solvents that are typically used with organic semiconductor processing and that are eco-


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harming and/or toxic, can be avoided. In their place, eco-compatible aqueous or alcoholic dispersions can be employed.5-7 Secondly, nanoparticles may be used to prearrange material domains in order to instantaneously form the required photoactive layer morphology for good device performance without the need for complicated drying processes.8

Two principal methods have been reported for the formation of polymer:fullerene nanoparticles dispersions. The miniemulsion method of Landfester et al.9 involves intermixing a photoactive blend solution (the materials are in a good solvent, i.e., a solvent in which they are soluble) and an immiscible non-solvent (a solvent that the organic semiconductors are not soluble in) by ultrasonication in the presence of a surfactant. The surfactant is included to stabilize the interface between the solution of the materials and the non-solvent. During evaporation of the good solvent at elevated temperatures, nanoparticles of the photoactive blend form in the non-solvent. In such a process, dialysis is commonly required to reduce the amount of the surfactant "impurity" in the dispersion. Most studies of organic solar cells formed from nanoparticles have employed the miniemulsion method due to its simplicity and applicability to various materials.5,7,10-14 However, the downside of surfactant-containing nanoparticle dispersions is that the corresponding devices mostly yield only moderate power conversion efficiencies (PCEs), well below those of the standard solution-processed devices using the same materials. Recently it has been reported that polymer:fullerene nanoparticles formed by adding a good solvent containing the donor and acceptor to a miscible non-solvent can give rise to efficient organic solar cells whilst avoiding the use of surfactants.6,15 This nanoparticle precipitation is fast due to the rapidly decreased solubility of the polymer:fullerene blend in the good solvent/non-solvent mixture. From a practical perspective, a disadvantage of the latter


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process is that the surfactant-free nanoparticle dispersions can aggregate over time. For some material combinations, such as poly(3-n-hexylthiophene-2,5-diyl) (P3HT) and phenyl-C61butyric acid methyl ester (PCBM) precipitated from ethanol can aggregate rapidly, leading to solar cells with only moderate PCEs (1 %).15 However, by careful choice of material combinations, such as P3HT and indene-C60 bisadduct (ICBA), it is possible to form stable surfactant-free nanoparticles. Such nanoparticle dispersions can be used to form photoactive layers in organic solar cells yielding PCEs of 4 %, which is similar to the performance of devices containing bulk-heterojunction films of the same materials processed from blend-solution and comprising the same architecture.6,16,17 Notably, nanoparticulate P3HT:ICBA solar cells with photoactive layers prepared using the miniemulsion route were reported to have a significantly lower PCE of 2.5 %.5

In either case, the semiconductor distribution within the nanoparticles is critically important for the morphology of the final photoactive layer and hence solar cell performance. Dastoor et al. studied the internal morphologies of P3HT:PCBM and P3HT:ICBA nanoparticles formed via the miniemulsion route using Scanning Transmission X-Ray Microscopy (STXM), with both material combinations producing a core-shell morphology with fullerene-rich cores and P3HTrich shells.5,18,19 The core-shell morphology of P3HT:PCBM nanoparticles produced by the standard miniemulsion method was later confirmed by contrast variation combined with Small Angle Neutron Scattering (SANS) experiments.20 Although the core-shell structure of nanoparticles formed by the miniemulsion method is widely accepted, there has been a single contradictory report where analysis of Small Angle X-ray Scattering (SAXS) data was more consistent with nanoparticles containing pure P3HT and PCBM domains in a P3HT:PCBM


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matrix rather than by a core-shell architecture.21 The relationship between the morphology and performance of devices containing films formed from miniemulsion nanoparticles has provided the impetus to understand the structure of surfactant-free nanoparticles. Scanning Transmission X-Ray Microscopy (STXM) and Transient Absorption Spectroscopy (TAS) studies of films containing P3HT:PCBM formed from surfactant-free nanoparticles have suggested a more homogeneous distribution of the two materials than nanoparticles containing the same materials formed using the miniemulsion method.15,22,23 Yet, the internal structure of the as-formed surfactant-free nanoparticles has not been studied, nor has the evolution of morphology as they are processed into the photoactive film.

In this work, we close this knowledge gap by investigating the internal structure of surfactantfree P3HT:ICBA nanoparticles in alcoholic dispersion with SANS using contrast variation. We expand these morphology studies with the charge carrier generation dynamics obtained using TAS, which was measured for both the alcoholic nanoparticle dispersions and thin-films before and after annealing. We correlate the morphology studies with the corresponding solar cell performances and compare their properties to reference devices prepared from 1,2dichlorobenzene (o-DCB) blend solutions.

RESULTS AND DISCUSSION Nanoparticle formation and solar cell performance. Before discussing the nanoparticle structure and thin-film morphology, it is important to review the method for the preparation of the nanoparticle dispersions and the power conversion efficiencies of the solar cells containing thin-films formed from these dispersions. The


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production of nanoparticles by surfactand-free precipitation involves the rapid addition of the two organic semiconductors (donor and acceptor) dissolved in a good solvent to a miscible nonsolvent. The mixing leads to a rapid decrease of solubility of the organic semiconductors, causing the nanoparticles to form instantaneously. As shown in Figure 1a, the precipitation produces nanoparticles with radii between 50 nm and 150 nm (for details see method section), and no significant aggregate formation was observed. Given the poor solubility of P3HT and ICBA in alcohol and the rapidity of the precipitation, it was expected that the nanoparticles would contain P3HT and ICBA in approximately the same ratio as in the original solution. Furthermore, the rapidity of nanoparticle formation and the poor solubility of the organic semiconductors in alcohols would also be expected to preclude ripening, a process that is generally observed in the formation of nanoparticles from miniemulsions. In contrast, in the miniemulsion process, ultrasonication of the immiscible solvent/non-solvent system results in a complex heterogeneous system with droplets of the P3HT:fullerene solution (typically chloroform, CHCl3) in the non-solvent (often water) with the interface between the droplets and water being stabilized by amphiphilic surfactant molecules (e.g. sodium dodecyl sulfate, SDS). During CHCl3 evaporation and shrinking of the droplets, the difference between the P3HT and the fullerene surface energies5,18 or, more precisely, the difference in their affinity for the surfactant occupied interface20 leads to the formation of core-shell structures. Whilst the formation of the core-shell morphology during the miniemulsion process is well understood, the question that remains is whether the precipitation method also leads to a core-shell structure or whether the materials within the nanoparticles show a different spatial distribution.


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In this work we formed surfactant-free P3HT:ICBA (1:0.8 w/w) nanoparticles in the nonsolvents, methanol (MeOH) and ethanol (EtOH). The nanoparticle dispersions were processed to form the photoactive layers for organic solar cells with an inverted device architecture comprising an indium tin oxide (ITO)/zinc oxide (ZnO) cathode and a poly(3,4ethylenedioxythiophene):polystyrene sulfonate (PEDOT:PSS)/silver (Ag) anode as depicted in the inset of Figure 1b. For reference, solar cells incorporating P3HT:ICBA blend-films from a spin-coated o-DCB solution were also prepared. The films formed from the o-DCB solution were allowed to dry slowly to optimize the film morphology and to achieve the best reference cell performance. This slow drying procedure was not required for the formation of the nanoparticulate films, which dried rapidly after deposition. All of the P3HT:ICBA photoactive layers were annealed at 150 °C for 10 min. Annealing led to melding of the nanoparticles, a corresponding improvement in charge carrier transport to the electrodes and a reduced bimolecular charge carrier recombination as we reported earlier.6 The annealing time and temperature used were based on the previously optimized conditions for the best device PCE. The current density - voltage (J-V) curves of the solar cells are depicted in Figure 1b, with the key performance parameters summarized in Table 1.

From Table 1 it can be seen that the P3HT:ICBA devices with photoactive layers deposited from MeOH or EtOH nanoparticle dispersions yielded PCEs of 3.5 % and 3.9 %, respectively, which is consistent with the earlier reports.6,16,17 The P3HT:ICBA reference devices prepared from o-DCB solution exhibited a PCE of 4.8 %, which also matches reported PCEs of standard solution processed devices with an inverted architecture.24-27 The different PCEs of the nanoparticulate and solution-cast solar cells mainly reflect the somewhat different photoactive


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layer thicknesses of 160 nm and 220 nm, respectively. Knowing that the performance of organic solar cells is sensitive to morphology, the similar performances of nanoparticulate and state-ofthe-art solution-deposited devices would suggest that the BHJ morphologies of the photoactive P3HT:ICBA layers are similar, and this will be discussed in more detail in the following sections.

Internal nanoparticle morphology. To gain insight into the distribution of the P3HT and ICBA within the precipitated nanoparticles, we performed SANS experiments on P3HT:ICBA (1.0:0.8, 1.0:1.0 and 1.0:1.2 w/w) nanoparticle dispersions in MeOH/MeOH-d4 mixtures utilizing solvent contrast variation. The scattering length density (SLD) of the dispersant SLDD (i.e., the non-solvent in which the nanoparticles are suspended) was varied by changing the ratio of protonated and deuterated nonsolvent, i.e., MeOH:MeOH-d4. To measure the same nanoparticles at the same concentration with different dispersant SLDs, a stock dispersion was prepared in MeOH for each P3HT:ICBA blend ratio, and these were then diluted with MeOH:MeOH-d4 mixtures with different MeOH:MeOH-d4 ratios. We investigated five different MeOH:MeOH-d4 ratios to provide five different contrasts for each P3HT:ICBA blend ratio. The MeOH:MeOH-d4 ratios are provided in Figure 2a, and the SLDD are listed in Figure 2b. The dispersant SLDs ranged from −0.37 × 10-6 Å-2 to 4.58 × 10-6 Å-2, which afforded a wide range of contrast variation that extended beyond the SLDs of pure P3HT28 (SLD = 0.8 × 10-6 Å-2) and ICBA29 (SLD = 4.17 × 10-6 Å-2). The individual components of the blend nanoparticles could therefore be rendered invisible to neutrons by matching the SLDs of either P3HT or ICBA to that of the surrounding dispersant,


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thus providing information about the distribution of each of the semiconductors within the nanoparticle.

The scattered intensity (absolute scaling) versus the magnitude of the scattering vector Q is plotted for each dispersion and dispersant SLD combination in Figures 2a-c. The SANS profiles for each dispersion exhibited qualitatively the same shape with the intensity modulated by the dispersant SLD; the exponential decrease in intensity with Q2 at low Q (Guinier fitting) indicated that the particle radii were invariant with solvent contrast for all of the dispersions and that insignificant particle aggregation took place, which is consistent with dynamic light scattering (DLS) measurements on the same samples.

Initially, the intensity of the SANS peak in the low-Q plateau, a surrogate for forward scattering at Q = 0, was quantified. The scattering intensity in forward direction from particles of uniform SLD distribution with the same volume and concentration is proportional to the square of the difference between the SLDs of the particles and the dispersion medium.30 The background intensity of each measurement was subtracted from the intensity in the low Q plateau region (0.001 ≤ Q ≤ 0.002 Å-2). As shown in Figures 2d-f, the square root of the intensity at low Q follows a linear trend with SLDD for all blend ratios, which is thereby consistent with a uniform SLD distribution within the nanoparticles. A weak signal was observed in all cases for the dispersant mixture MeOH:MeOH-d4 = 60:40 (SLDD = 2.11 × 10-6 Å-2, green) due to poor contrast between the SLD of the nanoparticles and the SLD of the dispersant. That is, the nanoparticles become nearly invisible to the neutron beam as their SLDs almost match that of the solvent. If significant phase separation into core-shell or eccentric nanoparticles was occurring,


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we would expect a stronger signal, as the SLD of the MeOH:MeOH-d4 = 60:40 solvent mixture is significantly different to those of both P3HT28 (SLD = 0.8 × 10-6 Å-2) and ICBA29 (SLD = 4.17 × 10-6 Å-2) (see illustration in Figure 3). The average SLDs of the nanoparticles (SLDNP) with different polymer:fullerene weight ratios were calculated to be 2.05-2.38 × 10-6 Å-2 from the volume fraction weighted sum of the SLDs of pure P3HT and ICBA (Table 2), and these values were compared with the SLD contrast-matched points determined from the x-intercepts in Figures 2d-f. At the contrast-matched points, SLDD equals SLDNP and no scattered intensity would be observed. In both cases, SLDNP increases with increasing fullerene content with the SLDNP determined from the x-intercepts not deviating by more than 5% from the values calculated from the SLDs of neat P3HT and ICBA. These results reinforce the view that the nanoparticles formed by the precipitation method in MeOH do not have a core-shell structure but on average have the P3HT:ICBA homogeneously distributed for all blend ratios.

As a final confirmation of the SLD distribution within the nanoparticles, a Stuhrmann analysis was performed.31 The radii of gyration Rg of the nanoparticles were determined for each solvent contrast using the Guinier approximation.30 The Rg were obtained from the slope of ln(I) versus Q2 plots (not shown) in the region RgQ < 1.3 and the corresponding Guinier fits are shown in Figures 2a-c (black solid lines). According to previously reported atomic force microscopy (AFM) images of as-formed thin-films, the precipitated P3HT:ICBA nanoparticles are spherical.6 ହ

Thus their radii can be calculated from the radii of gyration using R = ටଷRg; these radii are shown in Figures 2g-i for the four different contrasts that provided sufficient intensity for accurate Guinier fitting (squares). The data for the MeOH:MeOH-d4 = 60:40 mixtures were not included due to their weak signal intensities and the resulting uncertainties in Rg. In addition to


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the radii determined from the Guinier approximation of the SANS data, the nanoparticle radii were measured by DLS and are shown in Figures 2g-i (triangles) for comparison. In most cases, the radii determined from SANS and DLS are in good agreement. A maximum deviation of 3 nm was observed between the SANS and DLS measurements, and both showed that no significant particle aggregation took place upon dilution of the dispersions with the different MeOH:MeOHd4 mixtures. The nanoparticle radii of 47 - 50 nm were also similar for the three P3HT:ICBA blend ratios and are consistent with the nanoparticle diameter distribution in Figure 1a. Most importantly, the radii Rg from the Guinier analyses for all of the P3HT:ICBA blend ratios were the same within experimental error for all of the dispersant SLDs investigated. In the Stuhrmann analysis, the invariance of Rg with solvent contrast indicates that P3HT and ICBA are homogeneously distributed within the nanoparticles.20,31 As such, the SANS data are only consistent with a uniform average distribution of P3HT and ICBA within the surfactant-free blend nanoparticles in MeOH for the 1.0:0.8, 1.0:1.0 and 1.0:1.2 blend ratios, and not a coreshell structure.

Although the SANS measurements are consistent with a homogenous distribution of P3HT and ICBA within the nanoparticles, it does not allow for resolving domains with radii smaller than of the order of Qmax-1, where Qmax is the Q value at which the intensity reaches the background level. Depending on the level of deuteration of the dispersant (with higher amounts of deuterated dispersant resulting in lower background intensities) we estimate a minimum resolvable domain size of around 10 nm diameter (Figures 2a-c). To gain insight into the P3HT donor and ICBA acceptor domains on smaller scale within the nanoparticles, we utilized optical spectroscopy.


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Photophysical properties of nanoparticle dispersions and films. At the simplest level, UV-Vis spectrometry of the nanoparticles provides important information. The UV-Vis absorption spectra of the nanoparticle dispersions in Figure 4 exhibit absorption peaks at 500 nm with shoulders at 550 nm and 600 nm, which are ascribed to absorption by the semicrystalline P3HT domains.32 These absorption features change only marginally during thin-film deposition and annealing. Thus, we conclude that the semicrystallinity of the P3HT domains remains relatively constant after nanoparticle precipitation. In contrast, thermal annealing significantly affects the P3HT semicrystallinity in thin-films processed from blend solution.33

We next used TAS to study the charge carrier generation dynamics within the nanoparticles in dispersion as well as in thin-films. Since charge carrier generation dynamics are very sensitive to the domain size and intermixing in BHJ morphologies,34 we can estimate the degree of phase separation within the nanoparticles and track the structural changes from when the nanoparticles were a dispersion to the thin-film before and after annealing. Using a time-resolved measurement in the ps-regime we determined the difference between the transmission of the sample in the excited and ground state divided by the ground state transmission (∆T/T). Photoluminescence from P3HT singlet excitons occurs at wavelengths between 700 nm and 800 nm. Hence, stimulated emission (giving a positive contribution to the ∆T/T signal) can be observed at these wavelengths if a P3HT singlet exciton population is present.35 In the case of P3HT:fullerene blends, the majority of excitons on P3HT are separated at a polymer-fullerene interface within a time-frame of 100 fs.36 However, some excitons require more time to diffuse to an interface before undergoing exciton splitting, which can be on the timescale of picoseconds. In general,


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the larger the domain in which the exciton is formed, the more time the average singlet exciton requires to reach an interface and, therefore, the longer the respective stimulated emission feature can be observed in TAS.37

The transient absorption spectra of P3HT:ICBA nanoparticle dispersions in MeOH and EtOH are depicted in Figures 5a and b. On top of a constant negative offset from the charge population that is created on the 100 fs timescale, and remains constant on the timescale up to 50 ps,35 a distinct stimulated emission signature of the P3HT singlet excitons for timescales up to 20 ps was observed. The presence of this long-lived stimulated emission is consistent with a degree of P3HT and ICBA phase separation within the nanoparticles in the MeOH and the EtOH dispersions. If the nanoparticles were a homogeneous blend of the two materials, then, based on the ratios of the two materials, all the P3HT excitons should have been quenched essentially instantaneously. Notably, the exciton dynamics previously reported for dispersions of P3HT:PCBM nanoparticles prepared by the precipitation method [tetrahydrofuran (THF) into water]23 showed a faster exciton decay. A faster decay is indicative of smaller P3HT domains within the nanoparticle blend. After depositing the P3HT:ICBA nanoparticles from MeOH or EtOH dispersions (Figures 5c and d) and thermal annealing of the nanoparticulate films (150 °C, 10 min, Figures 5e and f), the stimulated emission signatures of P3HT singlet excitons in the transient absorption spectra were still observed. After thermal annealing, the stimulated emission features were slightly more pronounced, indicating a minor increase in average domain size, but this effect was small. A much larger effect of annealing on the stimulated emission dynamics was observed in P3HT:fullerene thin-films cast from o-DCB blend solution.35 In this latter case, while no P3HT stimulated emission was observed in the as-cast films due to a very intimate


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mixing of P3HT and the fullerenes leading to an immediate exciton quenching, annealing the films led to phase separation of the P3HT and the fullerenes with a concomitant increase in the stimulated emission signal.35

Figure 6 provides a quantitative analysis of the decay kinetics of the stimulated emission. The average signal intensity between 725 nm and 775 nm has been normalized to the intensity of the photoinduced absorption of free charge carriers at longer timescales (200 ps) and is plotted versus the decay time for the dispersions as well as for the as-cast and the annealed films (data points). The characteristic decay times τ (to 1/e of the total decrease, see Methods for details) are summarized in Table 3. Both, the EtOH and the MeOH dispersions showed a very fast decrease of the stimulated emission intensity with decay times of 6 ps and 4 ps, respectively, which from previous studies is consistent with small P3HT domain sizes. After casting from EtOH, the nanoparticulate layer exhibited a very similar stimulated emission decay time (τ = 6 ps) as that observed in the dispersant. The as-cast nanoparticulate layer deposited from MeOH showed a slower decay (τ = 12 ps) than the nanoparticles dispersed in MeOH (see Table 3). This difference may originate from different amounts of dispersant retained in the as-cast layers due to the different dispersant boiling points and lipophilicities. However, after casting and annealing of the layers formed from either dispersant, the stimulated emission decay occurred at a slower rate in comparison to when in the dispersant, resulting in an increase of the decay times from 6 ps to 9 ps (EtOH) and from 4 ps to 12 ps (MeOH). Such an increase in decay times correlates to an increase of the domain sizes, and furthermore, is consistent with the surfactant-free nanoparticles not having a core-shell structure. It has been reported that core-shell P3HT:ICBA nanoparticulate films formed via miniemulsion become more homogenously blended after annealing.5 The


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transition from core-shell to more blended structures would yield shorter P3HT singlet exciton decay times in the thermally annealed layers compared to the dispersed nanoparticles.

Most importantly, the decay times in the annealed nanoparticulate films were very similar to the decay times obtained for state-of-the-art annealed P3HT:fullerene BHJs processed from blend solution (9 ps).35 Thus, the combination of SANS and TAS provides strong evidence that the precipitation method for forming surfactant-free nanoparticulate photoactive layers can give a similar film morphology to layers formed by deposition from blend-solutions and subsequent thermal annealing.

CONCLUSIONS We have investigated the semiconductor distribution within surfactant-free, precipitated P3HT:ICBA nanoparticles with diameters of around 100 nm in alcoholic dispersions on different length scales using SANS and TAS. The use of contrast variations in conjunction with SANS enabled elucidation of the fact that the P3HT and ICBA were evenly distributed throughout the nanoparticles. On smaller length scale, TAS experiments on the dispersions showed that there was some phase separation of the semiconductors into small domains. We hence conclude that the nanoparticles already contain optimized “nano-solar cells” that subsequently can be assembled into light-absorbing layers for efficient solar cells. After deposition from dispersion and layer formation, the TAS experiments indicated that this nanoscale structure and phase separation was largely retained. Annealing was found not to change the nanoscale structure of the blend significantly, but simply fuses the nanoparticles together, eliminating the voids within the nanoparticulate layer and hence enhances the long-range charge carrier percolation pathways.


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Altogether these observations explain the superior performance of solar cells fabricated from surfactant-free, precipitated nanoparticle dispersions when compared to those prepared by the miniemulsion method. Therefore, we expect future advancements on the deposition of organic semiconductor layers from eco-friendly dispersions to follow the surfactant-free precipitation rather than the currently prevalent miniemulsion route.

METHODS Nanoparticle precipitation. Poly(3-n-hexylthiophene-2,5-diyl) (P3HT, Rieke Metals MW = 57,000, ĐM = 2.4, RR = 91 %) and indene-C60 bisadduct (ICBA, Lumtec) were dissolved in chloroform (CHCl3) in different blend ratios (1.0:0.8, 1.0:1.0 and 1.0:1.2 w/w) with a total concentration of 10 mg/mL. The solution (50 °C) was added to the non-solvent ethanol (EtOH, 50 °C) or methanol (MeOH, 50 °C) in a volume ratio 1:4 under vigorous stirring in a beaker. To yield nanoparticle dispersions with concentrations of 10 mg/mL, CHCl3 was evaporated and the dispersant phase was reduced in volume at elevated temperatures (55°C for MeOH dispersions, 65°C for EtOH dispersions, 2-3 hours, no stirring). The same stock dispersions were used for all measurements except the SANS experiments, for which separate dispersions were prepared at the Australian Centre for Neutron Scattering (ACNS) of the Australian Nuclear Science and Technology Organisation (ANSTO) following the same procedure.

Device fabrication and characterization. Organic solar cells were fabricated according to the device architecture in Figure 1b. Patterned indium tin oxide (ITO) coated glass substrates (R□ = 13 Ω sq-1) were cleaned by sequential ultrasonication in acetone (10 min) and iso-propanol (10 min) followed by an oxygen plasma treatment (2 min). All subsequent fabrication steps were


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carried out under a nitrogen atmosphere. Zinc oxide (ZnO) layers were spin cast (4000 rpm, 30 s, 20 nm) from nanoparticle dispersions (Nanograde Ltd., N-10, 2.5 wt% in iso-propanol, diluted to 1 wt%) and thermally annealed (150°C, 10 min). The nanoparticulate P3HT:ICBA (1.0:0.8) photoactive layers were spin cast from MeOH dispersions 4 times or from EtOH dispersions 5 times to yield layer thicknesses of 160 nm. The reference photoactive layers were spin cast from 1,2-dichlorobenzene (o-DCB) solution (40 mg/mL, 800 rpm, 30 s, 220 nm) and solvent annealed (45 min). All photoactive layers were annealed at 150 °C for 10 min. Poly(3,4ethylenedioxythiophene):polystyrene sulfonate (PEDOT:PSS, HTL Solar, Heraeus, diluted with water 1:1 v/v) layers were spin cast (500 rpm, 5 s; 2000 rpm, 30 s; 35 nm) atop the photoactive layer and annealed on a hotplate (120°C, 10 min). The silver (Ag) top electrodes were then thermally evaporated (10-6 mbar, 100 nm) through a shadow mask, defining the photoactive area of the solar cells (3 × 3.5 mm²). Current density - voltage (J-V) curves were measured with a Keithley 238 source-meter unit under illumination from an Oriel solar simulator (300W, 1000 W/m2, AM 1.5), which was calibrated using a KG5 filtered silicon reference cell (PV Ref Cell System 91150-KG5, Newport Corporation). Unless stated otherwise, all measurement errors throughout the manuscript refer to 68% confidence intervals.

Dynamic Light Scattering (DLS). DLS measurements were performed using a Malvern Zetasizer Nano ZS yielding the intensity distribution depicted in Figure 1. It represents the fraction of the total scattering intensity scattered by particles of a given size, which affords the radius distribution. For the DLS measurements at ANSTO (Figure 2) a DynaPro Titan (Wyatt Technology) instrument was used. Each dispersion with a different SLD was measured twice, before and after the SANS measurements. The observed radii did not deviate significantly,


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suggesting that the particles were stable throughout the SANS measurement. The DLS radii of the dispersions (Figure 2) are the average radii of those two measurements. The DLS measurements of dispersions containing MeOH-d4 were corrected for its different viscosity by using the volume-weighted sum of the viscosities.

UV-Vis absorption spectrometry. UV-Vis absorption measurements of P3HT:ICBA (1.0:0.8) dispersions with a concentration of 0.05 mg/mL in MeOH and EtOH as well as of nanoparticulate P3HT:ICBA layers on glass were performed on an Agilent Cary 5000 spectrophotometer at room temperature.

Small Angle Neutron Scattering (SANS). For SANS experiments, 10 mg/mL P3HT:ICBA stock dispersions were prepared in MeOH for each investigated blend ratio (1.0:0.8, 1.0:1.0; 1.0:1.2) as described above. They were diluted to 2 mg/mL with MeOH/MeOH-d4 in different ratios to achieve the respective dispersant SLDs (Figure 2). MeOH-d4 (D, 99.8 %) was purchased from Novachem. SANS measurements were performed on the QUOKKA instrument at the OPAL reactor.38 Three instrument configurations were used, two with equal source-to-sample and sample-to-detector distances of 20 and 8 m, and the final configuration with a source-tosample distance of 4 m and a sample-to-detector distance of 2 m with a 300 mm lateral detector offset to increase the maximum Q. Source and sample aperture diameters of 50 mm and 12.5 mm, respectively, were used. At the 2 m and 8 m configurations, 5 Å neutrons were used (∆λ/λ = 10 %) and, at the 20 m configuration, 8.1 Å neutrons were used in combination with focusing lens optics.39 The focusing lens was required in the low Q configuration because only a small fraction of the Guinier region essential for the Stuhrmann analysis was measured in the standard


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configuration utilising 5 Å neutrons. The combined 20, 8 and 2 m configurations provided a continuous Q range of 0.001 to 0.548 Å-1 where Q is the magnitude of the scattering vector, defined by Q = (4π/λ)sin(θ), λ is the neutron wavelength, and 2θ is the scattering angle. All samples were enclosed in demountable cells with a 1 mm path length. The sample temperature was controlled by a Julabo thermostatted bath and held at 25 ºC throughout the measurements. All data were corrected for blocked beam measurements, normalized, radially averaged and placed on an absolute scale, following attenuated direct beam measurements, using a package of macros in Igor Pro software (Wavemetrics, Lake Oswego, OR, U.S.A.), and modified to accept HDF5 data files from QUOKKA.40 The radii of gyration of the nanoparticles were determined by linear fitting of ln(I) versus Q2 plots using the NCNR analysis macros built into the analysis software.

Transient Absorption Spectroscopy (TAS). P3HT:ICBA (1.0:0.8) dispersions with a concentration of 0.25 mg/mL were used for the TAS measurements. The film samples were prepared by spin coating 5 mg/mL P3HT:ICBA dispersions (1.0:0.8) 7 times on clean glass. The measurements were made on a home-built spectrometer based on a 1 kHz repetition rate titanium sapphire amplifier (Coherent Libra) pumping two optical parametric amplifiers (OPAs, Coherent OPerA Solo). One OPA (set to 1300 nm) was focused into a sapphire window to create the white-light supercontinuum probe, and the other tuned to 532 nm and chopped at 500 Hz to create the pump pulse. Individual laser pulses were read out using a diode array detector. To evaluate the decay kinetics the average signal intensity between 725 nm and 775 nm was normalized by dividing by ∆T/T at longer timescales (200 ps), determined by a fit. Therefore, the data was fitted by double exponential fits, beginning at the time ∆T/T starts decaying. The


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characteristic decay time τ is defined as the time needed for the signal to reach 1/e of the total decrease. The difference in ∆T/T (from the double exponential fits) between the time at ∆T/T starts decaying and 200 ps was used to determine this total decrease.


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FIGURES

Figure 1. (a) Size histogram of P3HT:ICBA (1.0:0.8) nanoparticles (NPs) in MeOH and EtOH measured by dynamic light scattering (DLS, straight lines show log-normal fits to data points). (b) Averaged J-V curves of the nanoparticulate P3HT:ICBA (1.0:0.8) devices in comparison with reference devices spun from o-DCB, all showing similar performance.


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Figure 2. (a-c) Neutron scattering intensity I versus the magnitude of the scattering vector Q for P3HT:ICBA nanoparticles with blend ratios 1.0:0.8, 1.0:1.0 and 1.0:1.2. The colored symbols represent data points at different MeOH:MeOH-d4 ratios (different SLDD) including error bars, whereas the black solid lines represent Guinier fitting curves performed in the region RgQ < 1.3. For comparison, all fits are plotted up to Q = 0.005 Å-1. (d-f) The background-subtracted intensities of the SANS peaks in the low Q plateau I(0) as a function of SLDD. (g-i) Stuhrmann


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plots of the nanoparticle radii determined from the Rg values of the Guinier fits (squares) in comparison with the DLS data (triangles).


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Figure 3. Cartoon of nanoparticles in dispersant of varying SLDs comparing nanoparticles with a homogeneous (top) and core-shell (bottom) structure. It can be seen that a nanoparticle with a uniform distribution is visible except when its average SLD matches that of the dispersant (top). In the case of the core-shell structure either the core or the shell can be independently observed at all dispersant contrasts.


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Figure 4. Normalized UV-Vis absorption spectra of the P3HT:ICBA (1.0:0.8) nanoparticles in (a) EtOH and (b) MeOH dispersions (black), of nanoparticulate as-cast thin-films (red) and of annealed thin-films at 150 °C for 10 min (blue).


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Figure 5. (a, b) Transient absorption spectra ∆T/T of P3HT:ICBA nanoparticles (1.0:0.8) in EtOH or MeOH, (c, d) of the corresponding thin-films and (e, f) of thin-films after annealing for 10 min at 150 °C. In the dispersant, a positive contribution to ∆T/T is observed, indicating some degree of phase separation (i.e., P3HT and/or ICBA domains). This feature is even more pronounced after annealing of the layers.


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Figure 6. Normalized average ∆T/T (725-775 nm) of the P3HT:ICBA nanoparticles in (a) EtOH and (b) MeOH dispersions (black squares), of thin-films (red circles) and of annealed thin-films at 150 °C for 10 min (blue triangles). The solid lines represent double exponential fits to the data. After casting and annealing of the films, the ∆T/T signal decays more slowly indicating a slight increase in the size of the P3HT domains.


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TABLES Table 1. Key performance parameters (fill factor FF, open-circuit voltage Voc, short-circuit current density Jsc) of nanoparticulate P3HT:ICBA (1.0:0.8) solar cells in comparison with reference devices spun from o-DCB. Data averaged over at least 6 devices. Photoactive layer

PCE

FF

Voc

Jsc

(%)

(%)

(mV)

(mA/cm²)

P3HT:ICBA NP, EtOH

3.9 ± 0.1

53 ± 1

797 ± 6

9.2 ± 0.2

P3HT:ICBA NP, MeOH

3.5 ± 0.1

51 ± 1

794 ± 1

8.7 ± 0.1

P3HT:ICBA ref., o-DCB

4.8 ± 0.2

56 ± 2

810 ± 5

10.6 ± 0.3


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Table 2. Measured and calculated scattering length densities of the investigated P3HT:ICBA nanoparticles (SLDNP) for the blend ratios tested.

P3HT:ICBA

SLDNP calculated for

SLDNP from x-intercept

blend ratio

uniformly mixed particles

of I(0)1/2 versus SLDD

1.0:0.8

2.05 × 10-6 Å-2

(2.17 ± 0.02) × 10-6 Å-2

1.0:1.0

2.23 × 10-6 Å-2

(2.29 ± 0.03) × 10-6 Å-2

1.0:1.2

2.38 × 10-6 Å-2

(2.44 ± 0.02) × 10-6 Å-2


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Table 3. Characteristic decay times τ of the P3HT singlet exciton emission between 725 and 775 nm, i.e., the time for the signal to reach 1/e of the total decrease in P3HT:ICBA (1.0:0.8) dispersions, thin-films and annealed thin-films (10 min, 150°C). The errors in this table are extrapolated from 95 % confidence intervals of the double exponential fits. τ (ps)

τ (ps)

τ (ps)

Dispersant

Dispersion

As-cast film

Annealed film

EtOH

6±1

6±1

9±1

MeOH

4±1

12 ± 2

12 ± 1


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AUTHOR INFORMATION Corresponding Author * Email: [email protected], [email protected] Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. ACKNOWLEDGMENT S.G. and A.C. acknowledge funding by the Federal Ministry for Education and Research (BMBF), Germany, under contract no 12N12295 (project UNICORN). I.A.H. acknowledges financial support from the Carl Zeiss and Baden-Württemberg Foundation alongside the HEiKA research partnership. This work was supported by the Australian Research Council (DP120101372). P.L.B is supported by University of Queensland Vice Chancellor’s Research Focussed Fellowship. A.J.C., I.R.G. and P.L.B. acknowledge funding from the University of Queensland (Strategic Initiative – Centre for Organic Photonics & Electronics). The authors acknowledge the support of the Australian Centre for Neutron Scattering (ACNS, formerly the Bragg Institute at the time of the measurements) and the Australian Nuclear Science and Technology Organisation (ANSTO) in providing the neutron research facilities used in this work and A. Rekas for her assistance with the DLS measurements at ANSTO. C. Feldmann and L. Brütsch (Institute of Inorganic Chemistry, KIT) kindly provided the DLS measurement setup at KIT. ZnO nanoparticles were provided by Nanograde Ltd. We thank E. Blasco (KIT) for conducting NMR experiments in preparation for the SANS experiments and D. Bahro (KIT) for scientific discussions.


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