Impact of Fullerene Intercalation on Structural and Thermal Properties

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Impact of Fullerene Intercalation on Structural and Thermal Properties of Organic Photovoltaic Blends Jessica Wade, Sebastian Wood, Elisa Collado Fregoso, Martin Heeney, James R. Durrant, and Ji-Seon Kim J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.7b05893 • Publication Date (Web): 05 Sep 2017 Downloaded from http://pubs.acs.org on September 12, 2017

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The Journal of Physical Chemistry

Impact of Fullerene Intercalation on Structural and Thermal Properties of Organic Photovoltaic 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

Blends Jessica Wade1, Sebastian Wood1,2, Elisa Collado-Fregoso3,4, Martin Heeney3, James Durrant3 and Ji-Seon Kim*1 1

Department of Physics and Centre for Plastic Electronics, Imperial College London, SW7 2AZ, 2National Physical Laboratory, Teddington, TW11 0LW, 3Department of Chemistry, Imperial College London, SW7 2AZ, 4University of Potsdam, Physik weicher Materie, Karl-Liebknecht-Straße, 14476

* [email protected]; 020 7594 7597

Abstract The performance of organic photovoltaic blend devices is critically dependent on the polymer:fullerene interface. These interfaces are expected to impact the structural and thermal properties of the polymer; with regards to the conjugated backbone planarity and transition temperatures during annealing/cooling processes. Here we report the impact of fullerene intercalation on structural and thermal properties of poly(2,5-bis(3-hexadecylthiophen-2-yl)thieno[3,2-b] (PBTTT), a highly stable material known to exhibit liquid crystalline behaviour. We undertake a detailed systematic study of the extent of intercalation in the PBTTT:fullerene blend, considering the use of four different fullerene derivatives and also varying the loading ratios.. Resonant Raman spectroscopy allows direct observation of the interface morphology in situ during controlled heating and cooling. We find that small fullerene molecules readily intercalate into PBTTT crystallites resulting in a planarization of the polymer backbone, but high fullerene loading ratios or larger fullerenes result in non-intercalated domains. During cooling from melt, non-intercalated blend films are found to return to their original morphology and reproduce all thermal transitions on cooling with minimal hysteresis. Intercalated blend films show significant hysteresis on cooling due to the crystallised fullerene attempting to re-intercalate. The strongest hysteresis is for intercalated blend films with excess fullerene loading ratio, which form a distinct nano-ribbon morphology and exhibit a reduced geminate recombination rate. These results reveal that careful consideration should be taken when during device fabrication, as post deposition thermal treatments significantly impact the charge generation and recombination dynamics. ACS Paragon Plus Environment

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Introduction Power conversion efficiencies of organic photovoltaic (OPV) devices have continued to increase and now rival conventional, inorganic devices in low light conditions and temperatures up to 120 °C.1,2 In order to commercialise the technology, research efforts must now focus on increasing the operational lifetime and ongoing increases in device efficiency. There have been substantial advances in structural design, but to truly overcome the limitations of the small exciton diffusion length and low charge carrier mobility of organics, a precise understanding of molecular packing at the polymer:fullerene interface must be realised. Recent research efforts have focussed on designing donor-acceptor copolymer systems, where chemical structure can be manipulated, localising the highest occupied molecular orbital (HOMO) to a donor unit and lowest occupied molecular orbital (LUMO) to an acceptor unit.3–5 However, the OPV architecture most likely to be commercialised is a bulk heterojunction consisting of a blended thin film of a simple electron donating polymer (such as poly3-hexylthiophene, P3HT) and an electron accepting small molecule (such as [6,6]-phenyl-C71-butyric acid methyl, PC70BM). When a thin film is excited in the donor absorption band, a tightly bound electron-hole pair is formed in the donor material. This bound pair is split into charge carriers at the donor-acceptor interfaces and travels through the neat materials before being collected at the electrodes. The efficiency of charge separation depends on several factors, such as crystallinity and effective conjugation length of donor, size of the donor and acceptor domains, relative orientation of donor and acceptor and charge carrier mobility in donor and acceptor molecules. The photogenerated charges can recombine before generating any useful photocurrent by geminate or nongeminate (bimolecular) processes. Geminate recombination occurs immediately after exciton dissociation, before charge separation, and can be identified by the dependence on electric field, as the charges have remained Coloumbically bound to one another. Non-geminate recombination occurs at longer time scales, when separated-charges recombine from different parts of the blend film.6–12 The properties of the polymer will play an important role in determining the final blend morphology through intermixing, diffusion and phase separation with fullerene molecules during device processing, eventually determining device performance. The complex morphology of the donor-acceptor interface is of crucial importance for the ACS Paragon Plus Environment

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mechanism of carrier recombination. A stable intercalated bimolecular crystal can result in efficient ultrafast 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

charge photogeneration, where pure phases support fast charge transport. The size of pure domains must be large enough for bound pairs to separate, with accessible pathways between donor and acceptor for the charges to move to the electrodes. The donor material considered is Poly[2,5-bis(3-dodecylthiophen-2-yl)thieno[3,2-b ]thiophene] (PBTTT), a polymer developed by McCulloch et al. to improve upon the oxidative stability of P3HT.13 The linear conjugated thieno[3,2-b]thiophene unit has a large resonance stabilisation energy- it does not favourably delocalise electrons. The position of the highest occupied molecular orbital (HOMO) is similar to that of P3HT (- 5.1 eV).13 It exhibits liquid crystalline behaviour, forming large domains on cooling from a liquidcrystalline phase. The optimised gas-phase geometry has an all anti-arrangement of sulphur across the shortaxis, which promotes lamellar packing. Alkyl side-chains are only attached to unfused thiophene rings, which permits the so-called ‘intercalation’ of small molecules between the extended chains. The extended πsystem ensures close intermolecular π-π distances, which results in excellent two-dimensional charge transport.13 The acceptor size and donor-acceptor blend ratio can be fine-tuned to control crystallinity and effective conjugation length of the PBTTT. The interfacial morphology can be dominated by intermixed (cocrystals of polymer and fullerene) or phase-separated domains.8–12,14 Further control of donor-acceptor domain sizes can be achieved via thermal annealing. Whilst the donor-acceptor interfacial structures described at the molecular scale, the molecular packing and dynamics of charge separation have been investigated using macroscopic techniques. Miller et al. evaluated the long-range molecular order of the blends using specular X-Ray Diffraction (XRD), an effective probe for long-range order of the polymer films, however a more sensitive local probe of polymer molecular order (e.g. backbone planarity) is urgently required.9 The lamellar spacing of neat PBTTT (21.5 Å) increases to 25.9, 30.0, 26.0 from 23.1 Å for 1:1 blends with bis-PC60BM, PC70BM, ICBA and ICTA, respectively.9 The increase in lamellar spacing upon blending indicates that appropriately sized acceptor molecules force a larger spacing, allowing the polymer side chains to extend perpendicular to the polymer backbone, which may or may not be beneficial for charge transport.

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Our study focuses on the distinct macroscopic photophysical processes at the various polymer/fullerene 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

interfaces. First, we directly probe the impact of fullerene intercalation in PBTTT blends on polymer molecular order (backbone planarity), second the impact on thermal properties of polymer during heating/cooling cycles, and third the impact on blend morphology. We discuss the distinctive structural and thermal properties of the polymer at the interfaces influenced by fullerene intercalation and their impact on photophysical and electrochemical processes in the polymer:fullerene blend film. The polymer:fullerene interface is controlled using: 1. Acceptor size and symmetry 2. Donor-acceptor blend ratio

We characterise these interfaces using a combination of UV-visible absorption spectroscopy, Transient absorption spectroscopy (TAS) and resonant Raman spectroscopy. In particular, we use in situ resonant Raman spectroscopy to evaluate the interactions at the molecular level between the donor and acceptor materials in the above blends. Results & Discussion Controlling polymer:fullerene interfacial morphology 1. Acceptor Size and Symmetry In Figure 1 (a) the absorption spectrum of PBTTT is shown for neat films and 1:1 blends with four acceptor molecules of varying diameters and symmetries. Here we consider the bis-adduct of [6,6]-phenyl-C60butryic acid methyl ester (hereafter bis-PC60BM), PC70BM, indene-C60 bis-adduct (ICBA) and indene-C60 tri-adduct (ICTA). The acceptor diameter varies from 15 Å (PC70BM) to 20 Å (bis-PC60BM).15 The neat polymer predominantly absorbs between 400 nm and 650 nm, with evidence of a shoulder at low energy (Table 1). The main fullerene absorption band lies at higher energy (wavelengths less than 350 nm). Considering first the largest (“bulky”) acceptors (bis-PC60BM, ICTA) in Figure 4(a), upon blending the polymer absorption peak broadens. The broadening affects the high-energy side more than the low energy side (∆abs max= - 2 nm, + 0.2 eV), indicating that this broadening is due to a change in the morphological


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distribution of the chromophore population, where the overall effect is a transition of chromophores from the 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

low-energy side (corresponding with planar backbone conformations) to the high-energy side (corresponding with twisted backbone conformations). Bis-adduct acceptor molecules are reported to be less miscible in organic polymers.15,16 The blend films are rougher than the neat material films, resulting in increased scattering at low energy. The well-resolved vibronic structure in the donor absorption band for 1:1 blends with PC70BM is typical of a highly planar polymer backbone.14 The red-shift of the absorption band maxima (∆abs max= + 14 nm, – 0.6 eV, Table 1) suggests an increase in the planarity and thus the effective conjugation lengths of the polymer chain. Blends with ICBA lie intermediate between the two extremes; a red-shift of the resolved fine structure (∆abs max= + 7 nm, – 0.3 eV, see Table 1) and broadening to high energy. Compared to ICTA, bis-PC60BM blends show a broad low energy shoulder, which can be combined with Miller’s observation of an increase in lamellar spacing (bisPC60BM - 25.9 Å, ICTA – 23.1 Å) to confirm an structure.9

intercalated

Figure 1: Normalized absorption (left) and resonant Raman spectra (right) for blends of PBTTT (1:1 by weight) with: (a) and (b) 4 small molecule acceptors (bis-PC60BM, PC70BM, ICBA and ICTA), (c) and (d) PC70BM at different loading ratios. Raman spectra are acquired with a 488 nm laser and normalized to the 1415 cm-1 mode.


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1:1 Blends

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Max Absorption Peak nm eV Neat PBTTT 540 2.30 PBTTT:bisPC60BM 538 2.31 PBTTT:PC70BM 554 2.24 PBTTT:ICBA 547 2.27 PBTTT:ICTA 535 2.32

C=C cm-1 1491.3 1490.9 1485.0 1487.8 1490.8

FWHM C=C cm-1 16.2 17.1 13.8 15.5 16.9

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∆C=C order/disorder vs. neat % 0 4.0 39.4 29.1 -0.5

Table 1: Absorption band maximum (nm), C=C shift (cm-1) and relative increase in planarity for PBTTT blends with respect to the neat material, extracted from Figure 1.

The Raman spectrum of PBTTT is described elsewhere.17 Our resources enabled us to perform DFT at a higher theory level and with longer oligomer segments, and found good agreement with the previous work. The assignment of Raman modes are discussed in the Supporting Information. The bulky acceptors (bisPC60BM and ICTA) have little impact on the position or shape of the Raman scattering peaks of PBTTT (see Figure 1 (b)). Interestingly, the situation is very different for the 1:1 blend with PC70BM, which shows a downshift of C=CTT peak from 1491 cm-1 for neat polymer to 1485 cm-1 and a narrowing of the FWHM from 16.2 cm-1 to 13.8 cm-1(see Table 1). As for the comparison of absorption spectra (Figure 1 (a)), blending with ICBA represents an intermediate case between the two ((C=C)TT peak at 1485 cm-1). We have previously demonstrated that the backbone planarity of P3HT can be interpreted from a simple linear combination of disordered (amorphous) and ordered (crystalline) components of the main C=C Raman scattering band.18–20 For PBTTT, our DFT calculations indicate that the unfused T units exhibit a dihedral angle between them ∠T-T = 18.5 ° and a larger dihedral angle with neighbouring TT units ∠T-TT = 34° due to partial positive charges on adjacent units (Figure SI 2). To evaluate the impact of inter-unit torsion in the PBTTT polymer chain, a series of polymer structures were optimized with frozen inter-unit dihedral angles ∠T-T and ∠T-TT ranging from 0 ° to 45 °, and their Raman spectra were calculated (Figure SI 3). For both dihedral angles, reducing backbone torsion shifts the (C=C)TT (1491 cm-1) peak position to lower wavenumber, narrows its FWHM and increases the relative intensity of the (C-C)T (1389 cm-1) peak, as found in P3HT. Whilst all the Raman peaks shift when the torsional angle is changed, the C=C is the most affected. The (C=C)TT mode can therefore be described as a superposition of planar and twisted phases of PBTTT (Figure SI 3). The degree of polymer order of the polymer chain is quantified by calculating the relative ratio of the planar vs. twisted (C= C)TT components (C=Corder/disorder). ACS Paragon Plus Environment

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This simple comparison (C=Corder/disorder, Table 1) provides evidence that 1:1 blends with the smallest 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

acceptor PC70BM result in the highest degree of polymer backbone planarity, which is consistent with our interpretation of the absorption spectra. As the largest shift is on the (C=C)TT mode, it is clear that intercalation directly influences torsion between the T and TT units. For small acceptor molecules, the relative polymer order increases on intercalation from 4 % up to 40 % for blends with PC70BM and to 29 % for blends with ICBA, whereas for bulky fullerenes the increase is negligible (4 % for bisPC60BM, 0.5 % for ICTA). The narrow FWHM on blending with PC70BM and ICBA further indicates a narrower range of different conformational states in blends vs. the neat PBTTT (Table 1).21 This result is surprising, as the PBTTT in blends with PC70BM and ICBA is more structurally ordered than the neat film. In summary: the bulky fullerenes (bis-PC60BM, and ICTA) have minimal effect on the torsion of the neat polymer, whereas smaller fullerenes (ICBA, and PC70BM) permit a planar polymer backbone. These results can be combined with the results of Miller et al. to establish the impact of total (PC70BM) and partial (ICBA) intercalation on the backbone planarity.22 Note that ‘partial’ intercalation results in a broadening of the C=CTT (1488 cm-1) Raman peak, which is evidence of both planar and twisted PBTTT chains, in good agreement with XRD studies.8

2.

Acceptor Loading Ratio

It has been shown that a fine-tuning of blend morphology can be achieved using the fullerene blend ratio.23 Figure 1 (c) illustrates the thin film absorption as the fraction of PC70BM is increased from 0 % to 80 % (neat PBTTT, 9:1, 1:1, 1:4). The main absorption band red shifts on increasing fullerene content from 540 nm for neat polymer to 554 nm for the 1:1 blend ratio and the low energy shoulder becomes more sharply resolved. However, for the 1:4 blend the absorption from the disordered polymer chains (the higher energy component of the polymer absorption band) dominates. Raman spectroscopy (Figure 1 (d)) indicates that there is an upper limit of fullerene (here 1:1) content which a polymer can incorporate, as the changes associated with intercalation saturate. Beyond this threshold, which also shows a dependence on fullerene size (see Figure SI-4), further fullerene loading results in a broadening of all the Raman bands towards higher energy. For example, in the PBTTT:PC70BM blend, increasing the acceptor loading ratio from 1:1 to 1:4 results in a shift of the C=C peak from 1484 cm-1 to 1489 cm-1, and a 30 % reduction of the ordered ACS Paragon Plus Environment

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phase vs disordered one (C=Corder/disorder). This indicates the formation of a multi-phase film, where the polymer backbone conformation is planar in the intercalated phase but excess fullerene results in a mixed phase containing polymer with a more twisted conjugated backbone. The FWHM of the (C=C)TT mode (1490 cm-1) can be used to interpret the extent of twisted polymer domains, which helps to refine the previous model. A comparison of Figures 1 (b) and (d) demonstrates the profoundly different polymeraccpetor interaction with PC70BM and bisPC60BM. Such different loading ratios for the two different acceptors (9:1, 1:1) result in in very similar morphologies of the polymer. A schematic of the 1:1 and 1:4 blends is illustrated in Figure 2.

Figure 2: Schematic of the donor:acceptor interaction in intercalated (1:1) and high fullerene loading (1:4) blends of PBTTT: PC70BM.

This model of the of the blend film morphology can be verified by considering the charge carrier generation and recombination dynamics. Transient Absorption Spectroscopy (TAS) is an established tool for probing charge generation in organic semiconductor blend films, and the specific application to PBTTT-based devices have been previously addressed in detail.11,12,14,24 Figures 3 & 4 show the transient decay signals of the polymer polaron band (1 µm wavelength) for the 1:1 and 1:4 PBTTT:PC71BM blend film ratios, respectively, comparing traces measured at different light intensities. We have explored the transient response of non-intercalated blends elsewhere.24


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Figure 3: 1:1 PBTTT/PC70BM blend film normalised polymer polaron kinetic traces probed at 1000 nm after excitation at 540 nm with different intensities. The red line corresponds to a tri-exponential fit of the data plus a power-law decay from ≈ 1 ns, showing a fast and field independent decay. All data were corrected for blend absorption at the excitation wavelength.

Figure 4: 1:4 PBTTT/PC70BM blend film normalised polymer polaron kinetic traces at 1020 nm after excitation at 540 nm with different intensities, showing a slow and clear intensity-dependent decay. All data was corrected for blend absorption at the excitation wavelength.

The transient decay signals in Figure 3 show no variation in lifetime for the different intensities plotted, whereas in Figure 4, there is a clear trend of decreasing polaron lifetime as the excitation intensity increases. Intensity-dependent polaron recombination dynamics are indicative of bimolecular recombination (nongeminate) in the 1:4 ratio film, whereas the intensity-independent dynamics of the 1:1 ratio (Figure 3) film indicate that this sample is limited by geminate recombination. The intercalated 1:1 blends have an intimate donor- acceptor interface, and so we expect charges generated at these interfaces to recombine geminately. In 1:4 blends (Figure 4), large domains of pure fullerene disrupt the packing of polymer chains and charge dissociation is less field dependent, as described in the work of Zusan et al.23 The microstructure of these blends exhibits some intercalation accompanied by large regions of disordered polymer material, which is consistent with the broad FWHM of the C=CTT mode encompassing both ordered and disordered polymer conformations (Figure 1). In non-intercalated systems it is not possible to form free charges in high mobility ACS Paragon Plus Environment

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domains that favour charge extraction.24 These results confirm that whilst intercalation is beneficial for 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

charge generation, there must be pure domains of an appropriate size to avoid geminate recombination.24–26

Thermal properties of blends The dependence of the dynamics of charge separation on domain size offers promise for improving device lifetime through careful consideration of the blend ratio. A clear understanding of the phase behaviour of the blend film on heating/cooling could further assist our understanding of the donor-acceptor interaction at molecular heterointerfaces. Previous attempts to optimise the molecular packing via thermal annealing have combined optical microscopy, differential scanning calorimetry (DSC), PL, XRD and cyclic voltammetry (CV) to determine the impact of thermal treatment on PBTTT-fullerene blend films.8,27 Large, pure, crystalline domains of PC70BM are reported to increase the electron affinity of the acceptor, which in-turn increases the polymer- fullerene energy level offset and creates spatially separated long-lived free charges.27 However, there have been no direct in situ studies of molecular structures at the polymer-fullerene interfaces revealing the impact of interfacial molecular interactions on conformational changes of the polymer backbone during heating/cooling treatments, and no reported improvement in device efficiency.9 In order to evaluate such changes a series of in situ resonant Raman measurements were performed. The (C=C)TT peak shift and (C=C)order/disorder were extracted from Gaussian fits to the experimental spectra acquired at every 10 °C (10 °C per minute heating rate). We note that these two parameters are closely related to each other and so we expect them to show the same trends, however, in this case the Raman peak position (Figure 5) shows the transitions a little more clearly than the peak intensity ratio (see Figure SI 5). On heating the neat polymer (Figure 5 (a)), the relative intensity of (C=C)order/disorder increases and reaches a maximum between 100 °C – 110 °C, which is similar to the side-chain melting temperature of PBTTT reported by Miller et al.8 Beyond this temperature the mobile side chains no longer maintain the regular packing of adjacent polymer chains and disorder increases. The most significant polymer phase transition occurs between 200 °C – 225 °C, consistent with the backbone melting temperature for PBTTT, demonstrated here as a dramatic decrease in (C=C)order/disorder and shift of (C=C)TT to a higher frequency. On cooling, there is no hysteresis in the return of (C=C)TT to lower wavenumbers. Interestingly, there is a sudden increase in the polymer backbone

order

once

the

side

chains

start


to

recrystallize

(115

°C). 10

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The Journal of Physical Chemistry Least intercalation (1:1 ICTA)

Most intercalation (1:1 PC70BM)

Figure 5: (upper) in situ Resonant Raman spectra acquired on heating at 10 °C/minute for the four fullerenes discussed in the text, normalised to C=CT. The peak position of the C=CTT Raman mode (≈ 1490 cm-1) at room temperature can be used as a simple measure of the extent of intercalation.The most intercalated (PC70BM) and least intercalated (ICTA) samples are indicated. (lower) The extracted peak shift of the C=CT Raman active vibrational mode for (a) Neat PBTTT, 1:1 blends of PBTTT with (b) ICTA, (c) bisPC60BM, (d) ICBA, (e) PC70BM during heating (red) and cooling (blue). The labels 3 – 6 are discussed in the text.

1. Impact of Acceptor Size We observe the strikingly different temperature dependent features for the different fullerene species (Figure 5 (b-e)). For 1:1 blends with large fullerenes that are too bulky to penetrate the polymer side chains (bisPC60BM, ICTA), thermal energy causes only gradual structural changes in the molecular packing, until a temperature above the polymer melt (200 °C - 225 °C) is reached, where a sharp transition occurs, which is the same as for neat polymer. The lack of a clear transition relating to the side-chain melt on heating or cooling could either indicate a slight disruption of side-chain packing in these blend films, or that without intercalation the movement of the side-chains has little impact on backbone conformation. For ICTA, the C=C shift and evolution of (C=C)order/disorder is completely reversible on heating with negligible hysteresis, equivalent to the neat polymer (Figure SI 5). The hysteresis on cooling for bisPC60BM blends could indicate a stronger interaction with the polymer backbone. This observation supports our previous observations of the partially resolved vibronic structure in the absorption spectrum and slight increase in (C=C)order/disorder (only 4 %) in bis-PC60BM blend.


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In the case of the intercalated blend (PC70BM), as temperature increases beyond the onset of side-chains 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

becoming more flexible (110 °C) the overall backbone order increases (down shift from 1485 cm-1 to 1482 cm-1) as the polymer chains become more planar (up to 150 °C). We have previously demonstrated that this temperature is also the onset for PC70BM to start to diffuse out of a P3HT matrix.18 The temperature at which PBTTT order is maximised (150 °C) should indicate the annealing temperature to optimise polymer fullerene phase separation at these blend ratios. This is higher than typical annealing temperatures used for this system, and perhaps why previous authors have not been able to measure an improved device performance upon annealing.10,28,29 Beyond this temperature (150 °C – 190 °C) both the polymer and fullerene become more liquid-like, which is consistent with the C=C peak shifting to a higher energy and (C=C)order/disorder increasing as PC70BM separates out from the polymer side chains and crystallises.22 The disorder increases until some fullerene molecules have disentangled from the polymer matrix and formed macroscopic crystalline domains, visible in optical microscope images, completing by 230 °C (Figure SI 4). The film now represents a multi-phase system; with domains of fullerene, liquid polymer and a partially intercalated blend phase. There is noticeable hysteresis on cooling from the liquid polymer: crystal PC70BM system, implying the now crystalline fullerene does not effectively re-enter the polymer side chains.27 This hysteresis correlates well with the extent of intercalation. The polymer becomes less flexible once the noncrystalline fullerenes become immobile (150 °C) but recovers the original C=C peak position. The macroscopic fullerene domains remain at room temperature and their chemical composition in phase separated domains is probed using resonant Raman spectroscopy and PL spectra in regions of annealed film and aggregate crystal (Figure SI 4). The cartoon in Figure 8 can be used to interpret the interactions in 1:1 blends with PC70BM and parts 3 – 7 correlate with the numbers on Figure 5. This result supports the decrease in lamellar spacing recorded using specular XRD with in situ annealing (≈ 30 to 23 Å) reported by Miller et al., attributed to thermal expansion and side chain melting.22 The total recovery of the peak position and (C=C)order/disorder during slow cooling emphasises the need for such in situ analysis (Figure SI 5). A limited understanding of the impact of the thermal cycle on polymer backbone, or fast cooling, can result in a reduction of intercalation.


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Figure 6: (upper) in situ Resonant Raman spectra acquired on heating at 10 °C/minute for four PC70BM loading ratios (0, 9:1, 1:1, 1:4 by weight) discussed in the text. The extracted peak shift of the C=CTT Raman active vibrational mode during heating (red) and cooling (blue). (lower) Relative intensity of C=Corder/disorder, corresponding to the proportion of planar/twisted polymer chains.

Figure 7: AFM surface height images of the (upper) as-prepared spin-coated PBTTT:PC70BM blend films at the following loading ratios: (left to right) neat, 9:1, 1:1, 1:4. (lower) AFM surface height images of the blend films after annealing to 270 °C and slow cooling to room temperature.

Whilst ICBA causes some planarization of the polymer chains as similar to PC70BM, the phase behaviour of these blends is remarkably different (Figure 5). Instead of polymer order improving after side chains become more flexible, the disorder gradually increases until the polymer melt temperature. This could be attributed to the indene units penetrating the PBTTT side chains.22,30 Overall, heating 1:1 PBTTT:ICBA blends ACS Paragon Plus Environment

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gradually decreases backbone order as ICBA separates from the partially intercalated matrix. Macroscopic 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

circular domains of micrometre-wide domains of ICBA and disordered PBTTT are visible in optical microscope images (Figure SI 4). ICBA starts to become mobile at slightly lower temperature (140 °C) than PC70BM, completing by 240 °C. There is a noticeable loss of polymer order once these films are cooled to room temperature, with the maximum C=C peak shift increasing from 1488 cm-1 to 1490 cm-1. In the case of 1:1 blends with PC70BM, heating to 270 °C and subsequent slow cooling allows partial disentanglement of the fullerene molecules from the intercalated sections, which permits the remaining polymer to re-arrange. It is interesting to compare the local changes in polymer backbone order detected with resonant Raman spectroscopy to the macroscopic changes detected using absorption (Figure SI 6). For the 1:1 blends, the distinct fullerene absorption band (< 350 nm) represents a less intimately mixed donoracceptor material. The absorption maxima red-shifts (540 nm to 555 nm) and broadens at lower energy, with resolved fine structure from ordered polymer chains becoming increasingly apparent. Resonant Raman spectroscopy at λEx = 488 nm is particularly sensitive to this effect since it resonantly excites the highenergy, shorter conjugation length portions of the polymers, so the increase in disorder is emphasised. The driving force for polymer chains to rearrange is from the thermal energy released when PC70BM crystallises out of polymer side-chains. 2. Impact of loading ratio For the intercalated blends, the impact of fullerene loading on the thermal phase behaviour is further investigated (Figure 6). For small amounts of PC70BM (9:1 blends) the disrupting effect of the polymer backbone melting is somewhat softened and completed by 240 °C. On cooling there is minimal hysteresis, and the re-organisation of the mobile side-chains results in an increase in returning to room temperature. For blends with excess PC70BM (1:4) the polymer steadily becomes more ordered (C=CT max shifts from 1489 cm-1 to 1484 cm-1) until the intercalated acceptors start to mobilise (140 °C - 150 °C). The progression to a liquid polymer and solid PC70BM system is completed by 220 °C. Increased PC70BM content develops our proposed model as follows (Figure 8): the liquid polymer chains (≈ 100 °C) allow the disordered polymer backbones close to large fullerene domains to re-assemble, ACS Paragon Plus Environment

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promoting polymer backbone order (step 3). The high fullerene loading results in a more significant loss of 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

backbone order when these start to mobilise (step 4). The large hysteresis on cooling represents an increase in the proportion of solid fullerene domains that cannot re-intercalate into the polymer side chains (step 7). The process is not reversible: after the PC70BM becomes immobile (150 °C) the backbone is frozen into a more twisted structure than the pristine film. The (C=C)TT peak position and (C=C)order/disorder (Figure SI 7) is now comparable to the neat PBTTT, indicating the partial intercalation that was present in pristine 1:4 blends is eliminated. These results support the loss of lamellar spacing observed after heating these blends above 250 °C.22 It is also important to consider the evolution of thin film surface microstructure with increasing PC70BM content. Figure 7 displays atomic force microscope (AFM) images (2 µm × 2 µm) for the pristine (nonannealed) and annealed films discussed above. For the pristine films, increasing fractions of PC70BM decreases the surface roughness (rq(pristine) = 2.46 nm, rq(TA) = 0.50 nm, Table SI 1). After annealing the samples to 270 °C and the subsequent slow-cooling to room temperature, high aspect ratio ribbons of polymers (1-2 nm high, 20 nm wide, 100s of nm long) exist between the crystalline fullerene domains. The combination of AFM with absorption and in situ Raman spectroscopy provides a detailed picture of molecular packing throughout the whole film thickness. The broad polymer absorption with increased absorption from the low-energy shoulder, but reversible changes detected by in situ Raman analysis, indicate that for blends with low PC70BM content annealing forms surface terraces of polymer domains (Figure 7 (b)) but has minimal effect on the film below.31 The terraced morphology for pristine and 10 % PC70BM blends is typical of annealed PBTTT, with the height of each terrace (11 nm) representing the distance between in-plane π-stacked polymer backbones.31 For high fullerene content, annealing and subsequent removal of fullerenes from the polymer side-chains results in ordered ribbons of material that float to the surface. The nano-ribbon morphology has been described before for films of PBTTT heated beyond their second phase transition (T = 275 °C) and is characteristic of fully extended PBTTT chains.31 The presence of higher loading ratios of acceptor molecules appear to align the nano-ribbons of PBTTT (1:1, 1:4), similar to those obtained by zone-casting and dip-coating.31,32 In particular, these higher loading ratios support wider nanoribbons (50 nm – 60 nm) (Figure 7 (c, d)). ACS Paragon Plus Environment

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These results identify how the annealing process affects the molecular packing of each component within a 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

polymer:fullerene

blend

and

the

ultimate

phase

segregated

domains.

Figure 8: The mechanism of heating and cooling an intercalated blend of 1:1 PBTTT:PC70BM.


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Conclusion 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

The conjugated backbone modes of PBTTT have been identified and the peak positions, relative intensities and shapes studied to assess the molecular conformation in intercalated and non-intercalated blends. Intercalation with small fullerenes (PC70BM, ICBA) is found to planarise the polymer backbone. High loading ratios are found to decrease the co-planarity of adjacent units as polymer chains twist around large fullerene domains. On the other hand, large fullerene molecules (bisPC60BM, ICTA) have minimal influence on the backbone packing. The phase-behaviour of the intercalated and non-intercalated blends is explored using in situ Raman spectroscopy. For both intercalated and non-intercalated blends with PC70BM, heating beyond the side-chain melt is found to improve backbone planarity. Heating to the second phase transition increases the relative proportion of disordered polymer units. For the non-intercalated blends, the transitions are found to be totally reversible with minimal hysteresis. For intercalated blends the observed hysteresis on cooling rises due to the now-crystalline fullerene attempting to re-intercalate. The topography of the annealed intercalated films reveals a surface network of polymer nano-ribbons hundreds of nanometres long. At high blending ratios an interesting microstructure results from annealing beyond the second phase transition: increasingly narrow nano-ribbons (tens of nanometres across) of highly ordered material, where there is a strong electronic interaction between the polymer and fullerene. For the systems we have investigated, it is only this blend comprising both intercalated and phase-separated regions, that exhibits non-geminate recombination. This understanding of polymer:fullerene interactions can be used to optimise the postdeposition processing conditions for PBTTT:PC70BM blends, and is applicable to a range of acceptor molecules where thermal transitions can be exploited to optimise the processing conditions for high performance organic photovoltaic devices. Methods Sample preparation: For PBTTT and 1:1 associated fullerene blends: PBTTT was provided by Bob Schroeder. 30 mg/ml solutions were dissolved in chlorobenzene, left to stir overnight at 100 °C and clarified through a 0.45 µm ACS Paragon Plus Environment

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filter. The glass substrates were cleaned (detergent, deionised water, acetone, IPA (twice), plasma ashing 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

for 3 minutes at 80 W) prior to spinning the polymer film at 1500 rpm for 60 s. For PBTTT and mixed weight PC70BM blends: the same processing steps were carried out but instead of glass the films were deposited on clean quartz. Spectroscopic Measurements: Transmittance spectra were measured using a Shimadzu UV-2550 UV-visible spectrophotometer. Absorbance was calculated assuming no scattering or reflection correction based on the natural logarithm and the substrate contribution simply subtracted.

Raman spectra were measured using a Renishaw inVia Raman microscope in backscattering configuration. Thin film samples were measured in a nitrogen-purged sample chamber to minimize photooxidation of the samples during the measurement. Raman spectra was acquired at 785 nm and 488 nm. The excitation conditions were: 785 nm (130 mW) acquisition time: 20s and 488 (9.0 mW) 20 s, , in each case the beam was focused to an diameter of roughly 1 µm. The spectrometer was calibrated for frequency using a silicon reference sample. Raman spectra of PBTTT were calculated using 6-31 and 6-311 G (d,p) basis sets at the B3LYP functional. 6-311 is found to be a more appropriate approximation. The additional ‘1’ adds more gaussians to describe valence orbitals, which is more computationally expensive than 6-31 but results in a better approximation of the molecular orbitals along the chain. In each case, the ground state geometries of the molecules were optimized in the gas phase (tetradecylthiophene-2-yl chains were replaced with methyl side groups to reduce computation time). The geometry optimisation results in an all trans arrangement of sulphur atoms along backbone short-axis. In situ thermal studies: Raman spectra were recorded at a heating rate of 10 °C/ minute and measurements taken every 10°C. Spectra were acquired in non-resonant (785 nm excitation, neat polymer) and resonant conditions (488 nm, blends) conditions in a nitrogen environment to prevent degradation.


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Transient absorption spectroscopy was carried out on thin film samples under nitrogen with a commercially 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

available setup that comprises a 1 kHz Solstice (Newport Corporation) Ti:sapphire regenerative amplifier with 800 nm, 90 fs pulses. The output of this laser was split into two parts that ultimately generate the pump and the probe pulses. The tunable pump pulse was generated in a TOPAS-Prime (Light conversion) optical parametric amplifier and used to excite the sample. The probe light was used to generate a Near-IR continuum (900 nm - 1450 nm) in a sapphire crystal. A HELIOS transient absorption spectrometer (Ultrafast systems) was used for collecting transient absorption spectra (900 nm - 1450 nm) and decays up to 6 ns.

Acknowledgments The authors gratefully acknowledge funding provided by the EPSRC via a DTA studentship and the Centre for Doctoral Training (EP/G037515/1). We also acknowledge EPSRC support from EP/G060738/1, EP/I019278/1 and EP/K029843/1.

The Supporting Informaton includes: Quantum chemical simulations of PBTTT oligomers with Mulliken charges and Raman mode assignments (Figure SI 1), also considering varying dihedral angles within the polymer backbone (Figure SI 3); Raman and photoluminescence spectra for aggregate regions in blend films (Figure SI 4); Temperature dependence of Raman spectra and fitted peak intensities for blend films with different acceptors (Figure SI 5); Absorption spectra for pristine and annealed films of PBTTT blended with PC70BM and ICBA (Figure SI 6); Surface roughness (rq) for PBTTT : PC70BM blends pre- and postannealing (Table SI 1); Temperature dependence of Raman spectra and fitted peak intensities for blends with different loadings of PC70BM (Figure SI 7).

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Graphic:

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Impact of Fullerene Intercalation on Structural and Thermal Properties

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