Nanoscale Chain Alignment and Morphology in All-Polymer Blends

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Nanoscale Chain Alignment and Morphology in All-Polymer Blends Visualized Using 2D Polarization Fluorescence Imaging: Correlation to Power Conversion Efficiencies in Solar Cells Daniela Täuber, Yuxi Tian, Yuxin Xia, Olle Inganäs, and Ivan G. Scheblykin J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.7b05244 • Publication Date (Web): 19 Sep 2017 Downloaded from http://pubs.acs.org on September 19, 2017

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The Journal of Physical Chemistry

Nanoscale Chain Alignment and Morphology in All-polymer Blends Visualized Using 2D Polarization Fluorescence Imaging: Correlation to Power Conversion Efficiencies in Solar Cells Daniela Täuber,1 Yuxi Tian,1 Yuxin Xia,2 Olle Inganäs,2§ and Ivan G. Scheblykin1* 1

Chemical Physics, Lund University, P.O. Box 124, SE-22100, Lund, Sweden

2

Biomolecular and Organic Electronics, IFM, Linköping University, SE-58183 Linköping,

Sweden *

Email: [email protected] § Email: [email protected]

Abstract All-polymer blends are promising materials for organic electronics. Their performance critically depends on the quality of mixing of the electron donor and acceptor polymers and on the local chain organization. We investigated spatially-resolved photoluminescence properties of as prepared and annealed blends of poly[2,3-bis(3-octyloxyphenyl)quinoxaline-5,8-diyl-altthiophene-2,5-diyl] (TQ1) and poly(N,N-bis-2-octyldodecyl-naphtalene-1,4,5,8-bisdicarboximide-2,6-diyl-alt-5,5-2,2-bithiophene) (N2200) using two-dimensional polarization imaging (2D POLIM). N2200 is known to aggregate into fiber-like morphologies with a few 100 nanometers lateral extensions. Our findings suggest a highly parallel chain organization within individual domains. Comparing blends differing in the batch of the N2200 component, we could relate decreased power conversion efficiencies of the corresponding devices to aggregation of N2200 in tens of micrometer sized elongated structures. TQ1 showed less sensitivity to preparation conditions. Other than N2200, TQ1 is liquid crystalline and its side chain structure hinders aggregation. It thus might be feasible to consider similar properties for the design of acceptor polymers as well.

1

ACS Paragon Plus Environment


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INTRODUCTION Low cost fabrication from solution in combination with large area coverage as well as the advantage of lightweight and mechanically flexible materials for renewable energy harvesting have drawn increasing interest towards organic photovoltaics (OPV).1–4 Polymer solar cells utilizing fullerenes as electron acceptor have reached power conversion efficiencies (PCE), exceeding 10%.3,5,6 The drawback of fullerenes, however, is their costly fabrication and a potential lack of morphological stability,7 which promotes a turn towards fullerene free allpolymer solar cells, using conjugated polymers not only as electron donor, but also as acceptor materials. Meanwhile a PCE of 8% has been reached from such all-polymer solar cells.8,9 Although also OPVs employing sequential stacking of thin layers of donor and acceptor materials are investigated, the most commonly used configuration of the active material in OPVs is the bulk heterojunction (BHJ), which is obtained from spincasting or even printing a mixture of donor and acceptor materials dissolved in a suitable solvent.2,10–12 The purpose for using BHJ configurations is to enable short diffusion paths for photo-generated excitons towards the donoracceptor interface, thus decreasing the probability for geminate recombination and enhancing charge separation.1,2,13 The requested short distances towards the donor-acceptor interface, of course, put constraints on the morphology of the blend, which has lead to considerable research effort aiming to control the morphology of BHJ cells and studying its impact on solar cell performance.6,13–18 On the one hand, good mixing of the components is required to provide close proximity of the donor-acceptor interface to all photo-generated excitons. On the other hand, improved solar cell performance was reported after annealing routines were applied to the spincast blend, 19–23 Such annealing induces partial phase separation leading to increased domain sizes, which enhance charge delocalization.16,24 In all-polymer solar cells, the nanoscale polymer chain organization is crucial for the delocalization of photo-generated excitons over interchain and intrachain chromophoric units.25–27 Enhanced delocalization reduces the likelihood of geminate charge recombination,25,28 improving solar cell performance. Nevertheless, some of the photo-generated excitons will recombine, whereby improved delocalization will lead to enhanced excitation energy transfer (EET) from absorbing to nearby emitting sites. This EET 2


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does not preserve the polarization phase between absorbed and emitted light and thus can be studied by polarization microscopy.26,27,29 In general, morphology studies of OPV materials are not straightforward due to the small chemical difference of the components, which in particular puts constraints to an investigation by conventional AFM measurements. The differing electronic properties of donor and acceptor materials can be used for applying conducting AFM measurements.23,30,31 By X-ray investigation, structural information on domain sizes can be achieved,23,32 however, X-ray techniques in general yield only average domain sizes in the investigated area and cannot provide a localized visualization of the structure. The recently developed two-dimensional polarization imaging (2D POLIM)33–35 has been proven a powerful tool for the investigation of EET and local organization within differently prepared

films

of

the

polymer

poly[2,3-bis-(3-octyloxyphenyl)quinoxaline-5,8-diyl-alt-

thiophene-2,5-diyl] (TQ1) blended with a fullerene33 and pristine TQ1.36 2D POLIM is based on fluorescence microscopy. Thereby, the polarization phase angle φex of the linearly polarized excitation is controlled by rotation of a λ/2 plate, while a polarization analyzer is used to control the polarization phase angle φem of the detected fluorescence from the sample; see Figure 1a. Fluorescence images are acquired for several combinations of φex and φem. By this, twodimensional intensity maps I(φex, φem) called “polarization portraits” for each pixel on the image are achieved,33,35 an example is shown in Figure 1b. The portrait contains information on the orientation of light absorbing and light emitting transition dipole moments in the sample plane and on the EET between them.33

3



le Samp

plane

M=

100 80 60 40

Imin

20 0

2

Imax − Imin Imax + Imin

(c)

180

x

120

160

y

θem , Imax

I (φex ,φem)

140

z

140

120

Objective

160

80

φex

0 180

100

Laser

Polarization analyzer

1

60

φem

2

Integration of I (φex, φem ) over φex

40

Camera

Polarization controller (λ/2 plate)

I (φem )

θex Excitation polarization angle, φex

Integration of I (φex, φem ) over φem

(b)

20

(a)

Emission polarization angle, φem

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I (φex )

1 0

(d) C8H17O

N2200

C8H17O

TQ1

Figure 1. (a) Scheme of 2D POLIM microscope setup. (b) Example of a polarization portrait I(φex, φem) obtained from a BHJ blend film studied in this paper (TQ1/N2200 1:1), and derivation of the polarization parameters. Modulation depths of PL excitation (Mex) and emission (Mem) are obtained after integration of I(φex, φex) over φem and φex, respectively. (c,d) Chemical structures of (c) TQ1 and (d) N2200. Modified from Täuber et al.37

From each polarization portrait, the degree of photoluminescence (PL) polarization is retrieved as modulation depths in excitation (Mex) and emission (Mem), which are calculated from I(φex) and I(φem), respectively, as shown in Figure 1b.34 Thereby, the integration over one of the angles reduces I(φex, φem) to a function in the form of Malus’ law, I(φi) = I0(1 + Mi cos[2(φi − θi)]), where Mi is the modulation depth and θi is the phase angle with the index i = “ex” and i = 4


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“em” when integrated over φex and φem, respectively.37 M reports on the degree of orientation of the transition dipole moments from isotropic (M=0) to uniaxial (M=1) within the sample plane. In conjugated polymers the orientation of the optical transition dipole usually is approximated to be parallel to the extension of the π-conjugation, as the effect of bends and kinks of the conjugated system on the orientation of the transition dipole quite often can be neglected.38 Assuming this approximation to be valid, M also reports on the order parameter of polymer chain alignment. Besides revealing the orientation and the degree of PL polarization in excitation and in emission, the polarization portrait also contains information about EET. To quantify EET, the full polarization portrait is analyzed via the so-called single funnel approximation.34,35 In this approach, the EET from each excited dipole in the system towards a single emitter is approximated by the energy funneling efficiency parameter ε, which ranges from zero (no EET) to one (all excitation energy is transferred to a single emitter). The single emitter is not necessarily a single chromophore, it can be a group of coupled chromophores thus potentially having any emission polarization character from pure dipolar to isotropic.34 Information on EET may also be obtained using fluorescence anisotropy measurements.39,40 However, the latter can only be applied for the investigation of isotropic samples, while ε can also be used with anisotropic samples. Here we employed 2D POLIM to study thin films of a polymer/polymer blend, TQ1 and poly(N,N-bis-2-octyldodecyl-naphtalene-1,4,5,8-bisdicarbox-imide-2,6-diyl-alt-5,5-2,2bithiophene) (N2200, also known as P(NDI2OD-T2)); for the chemical structures see Figure 1. Both polymers have been used for OPV together with other semiconductors showing good performances.41–44 N2200 is one of the few acceptors with high charge mobility.45 We investigated pristine TQ1 and N2200 films as well as their blends (concentration ratios 1:3, 1:1 and 3:1) before and after annealing from three different series differing in the batch of N2200. Thereby we found considerable variations in micromorphology, which were compared to the PCEs obtained in devices prepared from these series of polymers. Further, we were able to see changes in EET for as prepared and annealed films, thus, probing the nanoscale polymer chain 5



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organization. We compare our results from 2D POLIM with PL lifetime measurements and PL spectra obtained from the same samples. RESULTS AND DISCUSSION We investigated BJH blends of TQ1 and N2200 differing in the batch of the N2200 component (labeled A and B) and in the age of the polymer powder used for preparation, whereby subscripts denote the differing age of the used powder of batch A (Af, f = fresh, and Ao, o = old). Batch B had a slightly smaller molecular weight and polydispersity, see Experimental details in the Supporting Information. For solar cells produced from those blends, considerable changes in PCE were found. Xia et al. report PCE for cells from as prepared blends of the series Af to be 2.4%, while the PCE could be increased to 4.4% upon annealing.23 However, six months later, for similar cells, only about 2% PCE could be reached upon annealing (our series Ao). For cells made from a new batch of N2200 (our series B), only a PCE of about 2% could be reached upon annealing also. All films were prepared using the same batch of TQ1. Here we first present the differing local micromorphologies obtained from the investigation of spatially resolved photoluminescence properties of the three sample series before and after annealing. In a second part we then discuss the obtained energy funneling efficiencies ε. Differing micromorphology visualized by imaging polarization parameters Figure 2 compares images of PL intensity, modulation depth in emission Mem and polarization

phase in emission θem obtained from 1:1 blends of all the three series. The pattern seen in the intensity images in Figure 2 (top row) is a sample-independent artifact related to the excitation beam profile. Comparing intensities from as prepared to annealed films, a slight decrease can be seen for the annealed ones, which reports on enhanced PL quenching in the annealed films. While the films appeared homogeneous in the intensity images, Mem as well as θem images reveal varied micromorphologies, which will be discussed in the following.

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4.104

intensity 3.104 (a.u.)

N2200 fresh batch A as prepared annealed

N2200 old batch A as prepared

annealed

N2200 batch B as prepared

annealed

2.104

50 µm 1.104 0.20 0

0.16

Mem

0.12 0.08 0.04

180 150

0.00

120

θ = 0°

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


90 60 30 0

θem (°)

PCE 2.4%

PCE 4.4%

θ = 90°

PCE ≈ 2%

PCE ≈ 2%

Figure 2. Images of PL intensity (top), Mem (middle) and θem (bottom) obtained by 2D POLIM from 1:1 blends of series Af (left), Ao (middle) and B (right). For each series as prepared (left) as well as annealed (right) films are shown. Excitation at 488 nm, detected emission range 695 - 1000 nm. The shown films of the old batch A were measured at differing sample orientations, thus, yielding different phase angles θem.

All blends showed a rather low modulation depth, both in excitation and in emission (Mem < 0.2, see Figure 2 middle row). Investigation of the spincast pristine polymer films revealed the PL of TQ1 to be almost isotropic while that of N2200 was slightly anisotropic, see Figure S2 This feature is caused by the tendency of N2200 to form fiber-like morphologies45 yielding a modulation depth of up to 0.2, while TQ1 chains appeared to be isotropically distributed in the spincast films. Thus, locally increased values of Mem pointing to a slightly higher degree of polarization in the blends in Figure 2 (middle row) can be associated with enhanced contributions 7



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from highly oriented domains of N2200. Thereby, we have to keep in mind that PL intensities recorded by 2D POLIM are averages over the film thickness of the samples and that the investigation of lateral structures is restricted by the spatial resolution of the microscope (≈ 1 µm). Thus, highly polarized PL emission from nanoscale domains of N2200 as well as from nanoscale domains of ordered TQ136 are averaged out over the sampling area, and the actual modulation depth of the nanodomains can be considerably higher than the values obtained using 2D POLIM. For TQ1-N2200 blends, emission from a charge transfer (CT) state in the near infrared was reported from electroluminescence measurements.23 However, the PL emission obtained from the 1:1 blend is nicely matched by a linear combination of the PL spectra from pristine TQ1 and N2200 (see Figure S5). Thus, luminescence from the CT state can be regarded as very weak adding a negligible contribution to the PL signals probed using 2D POLIM. Nevertheless, there is an indirect contribution from the CT state, as it causes a considerable quenching of the PL from TQ1 in the blends. PL from pristine TQ1 is about 100 times stronger than from pristine N2200, whereas in the 1:1 blends the contributions from the TQ1 band and the N2200 band to the PL signal are of similar order of magnitude; see Figure 3b. The micromorphology seen in Mem and θem reveals micrometer sized domains for series Af (Figure 2, left block, bottom and middle row), for which the solar cells yielded higher PCE upon annealing (4.4%) than the cells made from as prepared films (2.4%).23 Comparing the left (as prepared) to the right (annealed) column in this block, we see a slight increase of the domain size upon annealing, from ≈ 2 µm to ≈ 5 µm on average. This points to partial phase separation of the two components upon annealing in agreement with the reports from conducting AFM measurements and X-ray investigation.23 The former revealed an entangled network of N2200 polymer chains, which spanned the film thickness and had micrometer-sized lateral extensions for the annealed films,23 while results from X-ray were interpreted to point to planarization of TQ1.23 Thereby, the alignment within each polymer domain is considerably underestimated by the observed modulation depths, because they are obtained from the fluorescence signal averaged over the film thickness and the diffraction limit of the microscope. Nevertheless, the 8


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phase separation is only partial, leaving still well enough mixing for an efficient charge transfer to occur. In conjugated polymers conformational subunits are formed by conjugation breaks and disorder.26 The local alignment of the polymer chains determines to which extent such subunits electronically couple to neighboring units.26 The slightly increased values of Mem observed from some structures in the annealed film point to improved local alignment of the polymer chains of N2200. Further evidence on improved local polymer chain alignment of both components was found from investigations of the energy funneling efficiencies as will be discussed in the next section. The improved local alignment of each of the two components observed upon annealing, allows for an improved charge transport to the electrodes23 as well as for enhanced delocalization of the excitation energy. The latter leads to lower energy of the emitting state, while the improved transport enhances the probability of reaching the lowest states. Thus, the redshift upon annealing observed from absorption and PL spectra (see Figure 3) agrees with the improved local alignment found from comparing the Mem images obtained by 2D POLIM. This redshift is more pronounced for TQ1, which may point to a stronger change in the reorganization of this component. However, the sensitivity of the CCD used for detection considerably decreases in the range > 850 nm which might lead to an underestimation in the redshift of the PL from N2200. In accordance with Xia et al.,23 we observed a blueshift of the emission of the blend at the TQ1 band in respect to pristine TQ1 for as prepared as well as for annealed films (see Figure 3b). It can be explained by less favorable polymer chain organization due to the mixing with N2200. Upon annealing the polymer chains of TQ1 reorganize and thus the emission stems from chromophoric subunits of larger extension than in the as prepared films. This is also observed for pristine TQ1, however, in the blend, the N2200 still restricts the formation of extended chromophoric subunits even upon annealing, so its emission remains blueshifted in respect to annealed pristine TQ1. In contrast to emission, in absorption there is an overlap of the N2200 emission band into the spectral range of the TQ1 band (see Figure 3a). This causes a redshift of the absorption of the blend in respect to pristine TQ1. For resolving the effect from 9



conformational changes, the relative amplitude of the absorption of N2200 in the blend needs to be determined first.

Normalized absorption

(a)

Wavelength (nm) 500 600

400

TQ1 band

1.0

700 N2200 band

0.8 0.6 0.4 0.2 0.0

3.0

(b) 700

2.5 Energy (eV)

2.0

Wavelength (nm) 800 900

TQ1 band

1.5

1000

N2200 band

1.0

Normalized PL

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

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0.8 0.6 0.4 0.2 0.0 1.9

1.8

1.7

1.6 1.5 Energy (eV)

1.4

1.3

1.2

Figure 3. (a) Absorption spectra and (b) PL spectra obtained upon excitation at 488 nm (2.54 eV) for pristine TQ1 (blue), 1:1 blend (pink) and pristine N2200 (green), as prepared (solid) and annealed (dashed) films from series B. All films show a redshift upon annealing in absorption as well as in PL.

Compared to the 1:1 blends of series Af, those of series Ao (Figure 2, middle block) exhibit higher Mem and a spatially uniform value for θem within each film. The values of θem measured at several positions of each film showed a correlation with the orientation of the sample upon its rotation on the sample stage (Figure 2 exemplarily shows two differing sample orientations of the films of the old batch A). Such a correlation was absent in series Af,. The dependence of θem on 10


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the sample orientation while being uniform within the field of view means that the polymer chains in the whole investigated region of the sample possess some predominant orientation. Mem as high as 0.7 was reported from differently prepared self-assembled TQ1 films.36 However, the here used spincasting procedure does not allow for TQ1 chain alignment exceeding the spatial resolution of our microscope in the range of 1 µm, as can be seen from the investigation of spincast pristine TQ1 films which yielded Mem ≈ 0.04 (see Figure S2). By contrast, the pristine N2200 films of this series Ao yielded even higher Mem (up to 0.2) than the blends and were oriented in a similar way, which points to alignment of N2200 throughout the whole film. The preparation of polymer blends by spincasting starts from an initial solution, where the solvent solubilizes both polymer components, which means that the system is in the intermixed region of the solution phase diagram. Evaporation of the solvent decreases its overall concentration and at some stage the system will become either metastable or unstable, whereby the manner in which a system phase separates is very sensitive to the experimental conditions.46 In particular, the substrate/film and film/air surfaces break the symmetry and can impose a direction on phase separation.46 The rather homogeneous structure seen throughout the blend films of series Ao, might thus result from vertical phase separation of the polymer components, which could be caused by degradation of N2200 inducing a shift in the solution phase diagram towards a narrower metastable region for the mixing in the blend, enhancing the likelihood for the occurrence of large-scale phase separation.46 Comparing polarization parameters observed from as prepared and annealed blends, we see no change upon annealing in series Ao contrary to series Af, which further supports the interpretation with a vertical phase separation and large scale alignment already formed during film preparation. In comparison to Af and Ao, series B revealed some larger irregular structure appearing in the Mem as well as in the θem images, whereby stripes of similar orientation extend over up to 100 µm having a width of about 10 µm, see Figure 2 right block). The pattern varied between different positions investigated by 2D POLIM on the same film, rendering a discrimination of 11



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structural changes upon annealing not feasible. Similar to series Ao, the pristine N2200 films from series B revealed a predominant average orientation in sample areas exceeding tens of micrometers, see Figure S2d. However, the Mem was considerably lower than in pristine N2200 films from batch A showing a smaller degree of local alignment of the polymer chains in series B. Batch B had a slightly smaller Mw and a smaller polydispersity (see Experimental details in the SI) than batch A, which could be the reason for the smaller Mem. Yet, the larger irregular structure appearing in the blends of series B can be related to the tendency of N2200 to form large-scale oriented structures found also for batch B. Absorption spectra of the 1:1 blends and pristine N2200 films obtained from series Ao, appear redshifted compared to those of series B (see Figure S4), which agrees with enhanced local alignment of N2200 for series Ao. The control experiment comparing the pristine TQ1 films from series Ao and B showed no shift in the absorption spectra (see Figure S4), which means there was no sign of degradation of the TQ1 powder used for preparing the films. A slight increase in domain size upon annealing, which we interpret as a local phase separation yielding µm-sized domains of each polymer, can be discriminated for series Af only. For the other two series, local changes upon annealing are hidden by the irregular micrometer sized structure of N2200, which also prevents improvement of PCE upon annealing in those series. Energy funneling efficiency reports on variation of nanoscale polymer organization In the previous section we discussed the differing local micromorphology of the blends fabricated from different batches of N2200. In particular, we found macroscopic structures exceeding tens of micrometers in those blends, which yielded low values of PCE. However, the information obtained from imaging the polarization order parameters using 2D POLIM is limited by the optical resolution of our microscope, which was about 1 µm in this case, whereas EET occurs on a nanometer scale. Additional information on the nanoscale polymer organization can be provided by studying and comparing energy funneling efficiencies ε obtained from the various polymer films. 12


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(a)

1.0

as prepared

25 µm 0.9

ε ≈ 0.99

ε ≈ 0.83

ε ≈ 0.75

ε ≈ 0.70

ε ≈ 0.69

0.8

annealed

0.7

0.6

ε ≈ 0.96

ε ≈ 0.75

ε ≈ 0.69

ε ≈ 0.64

ε ≈ 0.67 0.5

pristine

(b) Energy funneling efficiency ε

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


TQ1

1:1

N2200

pristine

1.0 0.9 0.8 0.7 0.6 0.5 0.4

1.00 0.75 0.50 0.25 0.00 mass fraction of TQ1 in the films

Figure 4. (a) Energy funneling efficiencies ε obtained by 2D POLIM for as prepared (top) and annealed films (bottom) from series B, with TQ1 fraction in the films decreasing from left to right. (b) Overview of ε from all sample series: series Af (p blue), series Ao (ugreen) and series B (l pink), where filled symbols represent as prepared and open symbols annealed films. Variation of ε obtained at three different positions in each film was < 0.02 (smaller than the symbols). Excitation at 488 nm, detected emission range 695 - 1000 nm. Figure 4a exemplarily shows energy funneling efficiencies ε from as prepared (top) and

annealed (bottom) films from series B. The shown values are averages obtained at three different positions of the films, however, the range of variation of ε was smaller than 0.02 in all investigated films from all series. In general, ε decreased with increasing concentration of N2200, as well as upon annealing. This trend was seen in all the three sample series independently on the 13



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batch of the polymers. In all cases, pristine TQ1 films yielded ε close to 1, however, the absolute values obtained for pristine N2200 differed considerably for different batches (largest value for B, smallest for Af); see Figure 4b. In conjugated polymers, EET is due to long-range dipole-dipole interaction 27,29,47 leading to exciton diffusion.27,47 The typical average exciton diffusion length in conjugated polymer films is of the order of 10 nm and depending on the nanoscale organization of the polymer chains. The variations in ε observed for the different batches of N2200, therefore, report on differing nanoscale polymer organization in those films. In order to discuss the cause of the variations in ε, let us go a bit deeper to the meaning of the energy funneling parameter ε. This coefficient tells the position of the system between the two limiting cases: i) non-rotating chromophores without energy transfer between them (ε = 0) and ii) complete energy funneling to an emitter with fixed polarization properties which are independent of the excitation conditions (ε = 1). Note that ε = 1 even in the case of randomly oriented chromophores (Mex = Mem = 0) if they either rotate or exchange energy.33 In this case the effective funnel is simply an unpolarized emitter. We mention this to stress that the parameter design ensures that its value does not depend, within limits of experimental accuracy, on the degree of orientation of chromophores, contrary to commonly used fluorescence anisotropy.33,35 However, as any parameter based on light polarization, it is not very useful for systems where all chromophores are very well aligned, because energy transfer between chromophores of the same orientation does not change any polarization property. Due to experimental accuracy, ε will not be able to adequately reflect the energy transfer in systems with M > 0.9. So, if two polymer chain segments are aligned parallel, their optical transition dipoles cannot be discriminated by 2D POLIM and they yield ε close to 0 no matter, which amount of EET actually occurred between the two chains. Thus, the energy funneling parameter may underestimate the extent of energy transfer in films containing regions where polymer chains are oriented parallel to each other. 14


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N2200 tends to form fiber-like structures,45 i.e. domains of highly aligned polymer chains. Thus, the ε obtained from N2200 films is limited to values considerably smaller than 1. The large difference in ε of about 0.2 between batch A and batch B shows the high sensitivity of this parameter to the nanoscale polymer chain organization. Obviously, for N2200 the latter is quite sensitive to preparation conditions and to the polymer batch and its age. The higher ε obtained from series B agrees with the lower Mem found from these series (see Figure S2) showing less aligned polymer chains than in the N2200 films from batch A. The lowest ε was obtained for series Ao, for which the investigation of the micromorphology of the blends showed the highest modulation depth (see Figure 2 middle columns). In pristine N2200 of series Ao, ε did not change upon annealing (see Figure 4b), which further supports the assumption that the polymer chains already align during spincoating. The chain architecture of TQ1 was designed to prevent aggregation in solution, and the formation of π-stacked aggregates in spincast films is limited to a few nanometers,48 far below the optical resolution of our microscope. Nevertheless, it is liquid-crystalline at room temperature,48 and the formation of millimeter-sized ordered domains, exhibiting Mem up to 0.7, was reported from TQ1 films prepared by a floating film transfer method.36 However, spincast TQ1 was found to be almost isotropic on the length scale of our instrumental resolution; see Figure S2. The high ε ≈ 1 observed for the as prepared pristine TQ1 (Figure 4a top left), can be explained by the ultrafast delocalization of the photo-generated excitons over interchain (ps to tens of ps) and intrachain (tens to hundreds of ps) chromophoric units succeeded by energy funneling (tens to hundreds of ps) towards emitting sites and radiative exciton relaxation.25–27,49 The excited state lifetime of pristine TQ1 is in the range of one nanosecond (see Figure 5) and thus long enough for the various EET processes to take place.27 Upon annealing a slight decrease of ε was found for pristine TQ1 films (Figure 4 left column). The preparation by spincasting from polymer solution in chloroform does not provide sufficient time for the polymer chains to reach thermal equilibrium before the solvent has evaporated, resulting in unfavorable chain conformations.50 Moreover, TQ1 has thermotropic liquid-crystalline properties,48 and annealing it above its Tg (≈ 100°C),51 therefore, is likely to 15



induce structural changes which due to long-range interactions with the substrate52–54 and confinement in the thin film geometry lead to enhancement of the local order parameter in respect to the as prepared films. The decrease of ε thus points to an increased parallelization of TQ1 chains, in accordance with the observed redshift in the spectra (Figure 3) and the planarization of TQ1 chains suggested from X-ray investigations.23 Nevertheless, the effect on ε is limited as the side-chain architecture of TQ1 restricts highly parallel π-stacking of TQ1 chains to the range of few nanometers.48

100

IRF TQ1 N2200 3:1 1:1 1:3

Normalized intensity

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10–1

10–2

0.0

0.5

1.0 1.5 2.0 Decay time (ns)

2.5

3.0

Figure 5. Normalized PL decay curves obtained upon excitation at 640 nm (1.94 eV) comparing pristine polymers and blends. Pristine TQ1 (blue) shows a fast PL decay of < 1 ns, while N2200 (green) has a much shorter PL lifetime, close to the experimental time resolution of 0.1 ns shown by the IRF (black). The very fast PL decays of the blends cannot be resolved by our instrument, and their decays resemble the IRF.

Having discussed ε obtained from pristine polymer films, we now can turn to the polymer blends to see how ε is related to the PCE observed from corresponding solar cells. Fluorescence lifetimes from the blends show fast decays in the range of our instrumental resolution of 100 picoseconds or even shorter (Figure 5). As explained above, this is related to the non-emissive CT state. Initial exciton relaxation, thus, occurs very fast, pointing to ultrafast exciton localization (≈ 100 fs)25–27 as its main source. The dependence of ε on the concentration ratios of the polymer components, thus, can be simply explained by the differing ε values of the pristine 16


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components, and the decrease of ε upon annealing again can be understood by improved polymer chain alignment. Blends containing batch B yielded larger ε than those of batch A (see Figure 4b) for which PCE = 4.4% could be reached upon annealing.23 As high ε is associated with low nanoscale polymer chain alignment, this shows that the poor performance of the corresponding solar cells in case of batch B is not only related to the appearance of macroscopic structures in the BJH films as seen from the polarization parameters (see Figure 2), but is also caused by a lower degree of chain alignment within the N2200 domains than in films containing batch A.

CONCLUSION Using 2D POLIM we were able to reveal changes in the micromorphology and the nanoscale chain organization of bulk-heterojunction (BHJ) all-polymer films, thus complementing wellestablished investigation methods as atomic force microscopy, conductive AFM and grazing incidence wide-angle X-ray scattering. Improved power conversion efficiency (PCE) upon annealing was found to be related to an increase of micrometer sized domains caused by partial phase separation of the two components. Changes in energy funneling efficiency report on improvement of nanoscale alignment of the polymer chains upon annealing above the Tg of TQ1. Investigating BHJ blends differing in the batch and in the age of the acceptor component N2200, we found macroscopic structures exceeding tens of micrometers for those blends, which yielded only low PCE. The different batches of N2200 differed slightly, yet systematically in the energy funneling efficiency ε obtained using 2D POLIM, which is sensitive to the nanoscale polymer chain organization in respect of chain parallelization. By contrast, the liquid-crystalline TQ1 showed no sign of degradation in form of a dependence on the age of the powder used for film fabrication. TQ1 was designed in such a way as to restrict π-stacking of its chains and to avoid aggregation in solution. These findings emphasize the importance of the structural design of conjugated polymers used for BJH solar cells. On the one hand, large-scale aggregation, which hinders mixing of the donor and acceptor materials has to be avoided. On the other hand, the solar cell performance is 17



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improved by nanoscale phase separation and chain alignment. The properties of TQ1 causing its liquid-crystalline behavior are well-suited to support chain alignment when annealed above its Tg, while its side chain structure successfully hinders large-scale aggregation during preparation by spincasting. Our findings suggest to consider such properties also for the design of the acceptor polymer material.

SUPPORTING INFORMATION Images of PL intensity, Mem and θem obtained from 2D POLIM for pristine TQ1 and N2200 films. PL and absorption spectra of all ten films from series B. Comparison of absorption spectra from pristine polymers and the 1:1 blends of series Af and B. Comparison of PL spectrum of 1:1 blend with a linear combination of PL spectra from pristine polymers (pdf).

ACKNOWLEDGEMENTS DT acknowledges a research grant DFG-TA 1049/1-1 provided by the German Research Foundation. OI acknowledges a Wallenberg Scholar grant provided by the Knut & Alice Wallenberg foundation. YX acknowledges support by the China Scholarship Council (CSC). IGS acknowledges grant KAW 2016.0059, "Mastering Morphology for solution-borne electronics" by the Knut & Alice Wallenberg foundation. This work was further supported by the Swedish Research Council, NSFC (21504006, 21534003), the Crafoord Foundation and by LASERLAB-EUROPE (grant agreement no. 654148, Horizon 2020).

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Nanoscale Chain Alignment and Morphology in All-Polymer Blends

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