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DOI: 10.1002/((please add manuscript number))

Article type: Full Paper

Impact of initial bulk-heterojunction morphology on operational stability of

polymer:fullerene photovoltaic cells

Saurav Limbu, Sebastian Pont, Alexander Doust, Sooncheol Kwon, Patrick Fuller, Ellasia Tan,

James R. Durrant, and Ji-Seon Kim*

S. Limbu, P. Fuller, E. Tan, Prof. J. S. Kim

Department of Physics, Imperial College London, SW7 2AZ, United Kingdom

E-mail: ji-seon.kim@imperial.ac.uk

S. Pont, Prof. J. R. Durrant

Department of Chemistry, Imperial College London, SW7 2AZ, United Kingdom

Dr. A. Doust

Cambridge Display Technology, Unit 12 Cardinal Business Park, Cardinal Way, Godmanchester

PE29 2XG

Dr. S. Kwon

Research Institute for Solar and Sustainable Energies (RISE), Gwangju Institute of Science and

Technology (GIST), Gwangju 500-712, Republic of Korea

Keywords: Organic solar cells, Operational stability, Bulk-heterojunction morphology, Interface

stability, Phase-segregation temperature

Controlling initial bulk-heterojunction (BHJ) morphology is critical for device performance of

organic photovoltaic (OPV) cells. However, its impact on performance, specifically long-term

operational stability is still poorly understood. This is mainly due to limitations in direct

measurements enabling in-situ monitoring of devices at a molecular level. Here, we utilize thermal

annealing preconditioning step to tune initial morphology of model polymer:fullerene BHJ OPV

devices and molecular resonant vibrational spectroscopy to identify in-situ degradation pathways.

We report direct spectroscopic evidence for molecular-scale phase segregation temperature (TPS)

which critically determines a boundary in high efficiency and long operational stability. Under

operation, initially well-mixed blend morphology (no annealing) shows interface instability related

to the hole-extracting PEDOT:PSS layer via de-doping. Likewise, initially phase-segregated

morphology at a molecular level (annealed above TPS) shows instability in the photoactive layer via

continuous phase segregation between polymer and fullerenes in macroscales, coupled with further

fullerene photodegradation. Our results confirm that a thermal annealing preconditioning step is

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essential to stabilize the BHJ morphology; in particular annealing below TPS is critical for improved

operational stability whilst maintaining high efficiency.

1. Introduction

Bulk-heterojunction (BHJ) blend of electron donors and acceptors as a photo-harvesting layer is

state-of-the-art structure employed in organic photovoltaic (OPV) devices primarily due to low cost,

solution processing and freedom in scalability.[1] BHJ has huge practicality because functional

devices can be fabricated from a simple mixture of organic semiconductor molecules in a selection

of compatible solvents. However, complexity in the control of BHJ thin film morphology at a

molecular level is still a major bottleneck towards technological maturity mainly due to intricate

intermolecular interactions[2] and large selection of organic molecules.[3] Desired BHJ morphology

should possess a well-mixed nanoscale domain of moieties (for efficient dissociation of

photogenerated excitons) and pure domains of each moiety (for efficient charge transport to

respective electrodes) forming a bicontinuous pathway throughout the layer.[4] Difficulty lies in the

optimization of each of these domains for maximum device performance. Choices of compatible

molecules, blend ratio, substrate and interlayers together with solvent formulations, additives,

thermal/solvent annealing and film drying techniques are many different variables and processes

that can be utilized to create desired BHJ morphology for high efficiencies.[4-5]

Additional processing steps (thermal/solvent annealing or additives) during device

fabrication for morphology optimizations is known to be beneficial to maximize efficiency.[6]

However, such additional processing steps can have major influence in device operational stability.

Likewise, BHJ phase morphology is vulnerable to external stresses including temperature, light,

humidity and oxygen.[7] Moreover, molecular incompatibility is one of the limiting factors of

stability due to the intrinsic nature of organic materials to de-mix. In fact, spinodal de-mixing[8] has

been shown to be responsible for the complex initial burn-in[9] degradation of OPV devices. Other

studies have also identified different origins of burn-in degradation including polymer molecular

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weight[10],crystallinity[11], fullerene dimerization[12] and degradation[13], and interlayer interfaces.[14]

Thus, a comprehensive study of BHJ unifying both efficiency and operational stability is of utmost

importance.

In this study, we utilize thermal annealing preconditioning as a tool to tune BHJ morphology

where thin films formed kinetically via spin coating are allowed to reorganize thermodynamically

by treating at various temperatures before metal electrode deposition (whilst keeping all the other

variables contributing to morphology tuning the same). Studies have identified how different states

of a typical BHJ morphology can be achieved by a simple yet effective thermal annealing process

with direct influence on device efficiency.[4, 5b] Novelty of this study lies in unveiling how such

subtle variation in BHJ morphology influences operational stability of devices. For this, we utilize

in-situ molecular resonant vibrational spectroscopy to probe the degradation pathways of devices.

We identify the correlation between efficiency and operational stability tuned by the thermal

annealing step. A donor-acceptor type copolymer F12TBT, poly[9,9-didodecylfluorene-4,7-

di(thiophen-2-yl)benzo[1,2,5]thiadiazole and PC70BM, phenyl-C71-butyric acid methyl ester, are

selected to represent a model BHJ system. Fullerene-based BHJ system is chosen because fullerene

as an electron acceptor are particularly relevant due to their universality matching with all kinds of

polymers and their tunability in terms of solubility and stability.[15] Likewise, F12TBT is a typical

example of a classic push-pull copolymer. F12TBT used here is an industrial grade research

polymer with well optimized synthetic formulations and good potential for cost-effective scale-up.

Thin film morphology in blends in terms of molecular vibrations and optical properties; their

evolution with thermal annealing and device operation is probed with combined resonant Raman

and photoluminescence spectroscopy.[16] Complete devices (fresh and aged) are analysed to identify

molecular origins of device degradation in these BHJ devices.

2. Results and Discussion

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2.1. Photoactive layer morphology. In order to study the morphological changes (both macro and

molecular scales) of photoactive layer BHJ blends upon thermal annealing, we prepared pristine and

blend films of F12TBT and PC70BM. F12TBT film (Figure 1a) shows the characteristic “camel-

back” absorption spectrum of a typical donor-acceptor type copolymer.[17] Inclusion of PCBM

complements the absorption gap between 400-500 nm. To study the morphology variation

(specifically molecular packing of F12TBT polymer chains) upon blending and thermal annealing,

we employed resonant Raman spectroscopy.

Figure 1b shows Raman spectra of as-cast F12TBT and its blend with PC70BM.

Assignments of various vibrational modes of polymer are performed via DFT simulations and

supported by literature[18] (Figure S1). Vibrational modes are distinguished into 2 groups (Figure

1b): single-unit vs coupled vibrations. Single-unit vibrations represent vibrational modes primarily

localized in individual units, nominally Peak-1 (1606 cm-1), 2 (1542 cm-1), 3 (1445 cm-1) and 6

(1272 cm-1) localized to ring stretching modes of Fluorene (F12), Benzothiadiazole (BT),

Thiophene (T) and Fluorene (F12) units respectively. Coupled vibrations represent vibrational

modes delocalized over several different units, Peak-4 (1373 cm-1) and 5 (1349 cm-1) which are

antisymmetric and symmetric forms of C=C mode of BT unit coupled to C-C mode of T unit

respectively. For reference, laser excitation wavelength dependent Raman spectra of F12TBT

polymer is shown in Figure S2.

Upon loading PCBM on pristine F12TBT polymers, T-BT-T coupled vibrations (Peak-4 and

5) are selectively quenched with respect to the donor F12 (Peak-1) highlighting a preferential

conformational change on polymer π-conjugated backbone (Figure 1c). Such spectral changes have

been observed previously on a similar polymer (F8BT) containing BT units, where the decreased

intensity of benzothiadiazole with respect to fluorene ring stretching modes is identified as an

indicator for increased inter-unit dihedral torsion angle.[19] Hence, for the case of F12TBT:PCBM

blend, the backbone of polymer becomes more twisted than its original conformation (in pristine

polymer films) upon PCBM loading. It is understood that PCBM molecules find spaces near T-BT-

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T unit of the polymer (where no side chains exist) to diffuse and intercalate, resulting in an increase

in torsion angle between F12 and T-BT-T units. On the other hand, thermal annealing on pristine

F12TBT (Figure 1c) show selective enhancement of T-BT-T coupled vibration with respect to F12,

indicating improved backbone planarity which is opposite to the case of PCBM loading. These

subtle changes in intra-molecular planarity of F12TBT polymer upon PCBM loading and thermal

annealing are expected to have major consequences in inter-molecular packing of the polymer and

thus its charge transport properties, as establised in similar polymer (F8BT).[20]

Figure 1d summarises the effects of PCBM loading (blue circles) and thermal annealing (red

triangles) on F12TBT polymer conformation (black squares) in particular, F12 and T-BT-T unit

torsion angle, where higher Raman peak intensity ratio (Peak-5/Peak-1) represents an increase in

the polymer backbone planarity. Maximum planarity is observed in thermally annealed films and

maximum twists in blends (at 50 % wt PCBM loading). Twisting of F12TBT polymer upon PCBM

addition is already saturated at 25 % wt PCBM, indicating the upper limit of PCBM loading that

can physically interact and alter the polymer conformation. The morphology in molecular scale for

such blends (> 25 % wt PCBM) is therefore a well-mixed domain with saturated intercalation of

PCBM near T-BT-T unit between bulky F12 units of polymer. Intensity ratio is completely

recovered in the case of 75 % wt PCBM (8.7 compared to 8.6 for pristine annealed) upon thermal

annealing, implying that thermal annealing of blend improves polymer conformation with no

hindrance by the large amount of co-existing PCBM molecules. At this stage, we observe large

PCBM microaggregates (Figure S3) indicating significant phase segregation of PCBM from

polymer matrix.

More detailed study of nanoscale phase morphology of the BHJ blend thin films upon a

thermal annealing cycle is investigated by in-situ annealing of samples. Blend weight ratio of 1:3

(polymer:fullerene i.e. 75 % wt PCBM loading) is selected to incorporate in both thin films and

devices hereafter. Figure 2a shows the gradual progression of Raman spectrum in heating phase

acquired by 514 nm laser excitation which is resonant with the low-energy absorption band of

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F12TBT (Figure 1a). Thus, background of Raman spectrum is originated from photoluminescence

(PL) of the polymer in the range of 530-570 nm (Figure S4). Figure 2b shows the background

corrected resonant Raman spectra, all normalized to polymer F12 peak (1606 cm-1). Relative

intensity of polymer peaks with respect of PCBM peaks (denoted by asterisk) is a direct measure of

relative chemical compositions of the two moieties within the laser focus (≈ 10x10 µm2 focal area)

used to acquire Raman spectrum. Prominent polymer peaks are more than 2.5 times stronger than

that of the maximum PCBM peak (at 1567 cm-1) although the chemical composition of PCBM is

three times more than polymer due to stronger Raman cross-section of polymer at 514 nm resonant

excitation. Probing multiple areas of the as-cast blend film shows good homogeneity of chemical

composition throughout the whole sample (Figure S4).

Upon heating (Figure 2a and b) we observe an increase of PL which is strongly coupled to

the decrease of PCBM peak (1567 cm-1). Monotonic downshift of all Raman peaks’ position to

lower wavenumbers (by 2.5 cm-1 from 30 to 180 °C) is a reversible thermal effect and recovers

completely upon cooling (shown by 30 °C (Cooled) spectrum in Figure 2b). Irreversible thermal

effects after the heating/cooling cycle are PL and PCBM peak intensity ratio with respect to

polymer peak (Figure 2c and d). Figure 2c shows a gradual increase of PL with temperature from

25% (at 30 °C) to 45% (at 130 °C) of the maximum PL, with a small kink at 70 °C. Interestingly,

beyond 130 °C, the gradient of PL increase steepens and achieves maximum value at 180 °C clearly

indicating a real transition point. Likewise, PCBM peak intensity with respect to polymer peak F12

(Figure 2d) show almost unity until 140 °C, beyond which the intensity ratio drops significantly to

0.5 at 180 °C annealing temperature. Thus, the inflection point ≈130 °C indicates a critical

transition temperature for phase separation (TPS) at a molecular level of F12TBT:PC70BM (1:3)

BHJ blend.

Meaning of the observed critical transition temperature is explained clearly by micrograph

images taken in-situ during the thermal experiment (Figure 2e). Blend film appears homogeneous

until 180 °C, only after which microscale aggregates of PCBM begin to develop (Refer Figure S5

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that identifies the microaggregates to be PCBM). By 180 °C, PCBM diffuses out from a well-mixed

domain to form phase-separated and disconnected pure PCBM domains in microscales. The

initiation of nucleation sites (in a molecular scale) for disconnected phase separation should start at

a temperature below 180 °C which cannot be identified by in-situ optical micrograph images (or ex-

situ Atomic Force Microscopy images, Figure S6). In this respect, our in-situ resonant Raman

results provide very important information about this critical transition temperature revealed to be

≈130 °C, which is much lower than 180 °C (where a distinctive microscale phase-separated

morphology is observed in blends optically). Thus, we define this observed transition temperature

(TPS) to be the temperature required for initial phase segregation to occur at a molecular scale in

F12TBT:PC70BM (1:3) BHJ blend.

Gradual increase of PL of the polymer up to 45% at 130 °C indicates minor morphology

optimizations in nanoscale, where the kink at 70 °C might indicate the beginning of this process

(Figure 2c). This is a minor thermal effect on the polymer upon annealing, which contributes to a

small increase of PL (Figure S7). However, there is a sudden steep increase in PL at a transition

point (TPS). This strong increase in PL of the polymer indicates a phase separation between the

polymer and PCBM starting at this temperature, i.e. PCBM diffusing out from well-mixed

polymer:fullerene domains to form PCBM aggregates, thus reducing quenching of the polymer PL.

Here, the observed transition temperature is slightly higher than PL observation. It can be

understood from the fact that PL is slightly more sensitive, as significant amount of PCBM has to

physically move out to show measurable changes in relative chemical composition of PCBM with

respect to polymer. Phase segregation between PCBM and the polymer is maintained throughout

(and end) of the cooling cycle as shown by the hysteresis of PL (i.e. 50% of max PL after cooling

compared to 25% before heating) and PCBM peak intensity ratio (i.e. 0.55 after cooling compared

to 1 before heating). 2D Raman mapping of phase segregated areas reveal that as PCBM domains

develop from a nucleation site even at a molecular level, PCBM diffuse out from a well-mixed area

to create larger PCBM domain such that distinctive areas of PCBM-deprived regions form around

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the large PCBM aggregates (Figure S5). Morphology variations (well-mixed to phase-segregated)

determine charge extraction efficiency where both extremes: well-mixed and phase-segregated

morphologies suffer from immediate geminate recombination and the solution is have a balance of

both phases.[21] Thus, the macroscopic parameters such as current-voltage characteristics of OPV

devices must be investigated together to understand the influence of different morphologies on

device performance and stability thereafter.

2.2. OPV device efficiency and operational stability. Between the two extremes of well-mixed

and phase-segregated morphology tuned by thermal annealing process, the effects on efficiency and

operational stability are studied in conventional BHJ OPV architecture (Figure S8): ITO/

PEDOT:PSS (30 nm)/ F12TBT:PC70BM (1:3 by wt, 80 nm)/ Ca (20 nm)/ Al (100 nm). Chosen

temperatures: Room Temperature (RT, no annealing, well-mixed), 70 and 100 °C which represent

mild-thermal annealing (T < TPS, minor morphology optimizations but still well-mixed) and 140 °C

which represent high-thermal annealing (T > TPS, molecular scale phase-segregation but no

optically or topologically measurable PCBM aggregates).

Figure 3a and Table 1 show the impact of thermal annealing preconditioning of photoactive

layers on current-voltage characteristics of the OPV devices. Initial power conversion efficiency,

PCE (at RT) is relatively maintained at higher thermal annealing temperatures until 140 °C where it

deteriorates considerably by ≈ 40 % from average value of 3.34 to 1.95 %. So, the trend of PCE

against thermal annealing temperatures reveal a transition point below 140 °C beyond which PCE

deteriorates significantly. Dominant parameters for the loss of PCE are short-circuit current density

(JSC) decreasing by ≈ 30 % from 6.4 (at 22 °C) to 4.6 mA cm-2 (at 140 °C) and Fill Factor (FF)

decreasing by ≈ 20 % from 0.5 (at 22 °C) to 0.4 (at 140 °C). However, interestingly, open circuit

voltage (VOC) show a very little change, ≈ 2 % increase from average value of 1.01 (at RT) to 1.03

V (at 140 °C). Major losses in JSC and FF upon thermal annealing observed here imply that the main

cause of PCE decrease is due to the changes in nanoscale phase-morphology of BHJ photoactive

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layer which is not optically or topologically visible (even at ~100 nm scale), but already sternly

influences the efficiency and transport properties of photogenerated charge carriers. By correlating

device efficiencies with BHJ thin film morphology (Figure 2), we can again identify the transition

temperature (≈130 °C) (Figure 3b) for device PCE trend against annealing temperature. As unveiled

by Raman and PL results, 140 °C represent temperature above this transition temperature where the

nucleation points for PCBM segregation is already initiated albeit in nanoscales. Such phase

segregated BHJ morphology is not ideal, in particular for effective charge generation due to reduced

number of heterointerfaces formed between polymer and PCBM, so directly reduces photocurrent

as observed in our devices.

Figure 3b, c and Table 1 show operational stability of the OPV devices under continuous

white LED, 1 sun illumination (JSC matched with solar simulator) inside temperature and nitrogen

controlled atmospheric chamber. We observe a clear effect of thermal annealing on the operational

stability of devices (Figure 3c and d). Mild-thermal annealing (70 and 100 °C devices) show 5-7 %

PCE degradation via < 5 % JSC reduction compared to non-annealed RT (> 20 % PCE degradation

via > 15 % JSC reduction). The stability improvement is particularly in the initial burn-in phase (first

1000 mins) showing around 50 % improvement. This is a clear indication that thermal annealing

has alleviated some specific factors contributing to initial burn-in phase, which will be discussed

later. Interestingly, stability worsens for high temperature (140 °C) annealed devices. Mild-thermal

annealing improves operational stability, maintaining relatively high PCE (only 5-7 % drop) under

continuous illumination such that a small sacrifice in initial fresh device PCE specifically

comparing RT (3.46 %) with 100 °C (3.30 %) is more effective in long-term where RT degrades to

2.66 % compared to 3.02 % for 100 °C case. Losses in VOC (≈ 2 %) and FF (≈ 4 %) during

operation are comparatively minimal for all annealing cases (Figure S10). The improved stability

shown by mild-thermal annealing cases is mainly through the stabilization of JSC (≈12 %). This

result highlights an important role of the mild-thermal annealing on stabilising BHJ morphology

and hence improving device operational stability for polymer:fullerene solar cells.

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To identify the origins of loss of JSC and PCE during device operation, we employed

resonant vibrational spectroscopy on the complete multi-layered devices and compared fresh and

aged devices (i.e. before and after LED degradation). Figure S11 shows the reference spectrum of

each Raman active component present in the complete device. Raman spectrum of devices is

convolution of all these components (PEDOT:PSS, F12TBT and PC70BM) along with additional

signatures induced by degradation under operation. Figure 4a shows representative Raman spectra

of Fresh and Aged devices for the two unstable cases: RT and 140 °C which are morphologically

the two extremes of the variable under investigation i.e. no thermal annealing (well-mixed initially)

and thermal annealing above the transition temperature (TPS) (molecular scale phase-segregation

initially). Examining RT case, we observe that PCBM peaks (denoted by asterisk) are unchanged

after degradation showing that the relative chemical compositions in blend film (or mixed

morphology) is unchanged. In contrast, the peaks between 1400-1500 cm-1 are selectively

enhanced, which is a convolution of polymer Peak-3 (T, 1445 cm-1), PCBM and PEDOT:PSS peak

(1447 cm-1). Raman difference spectrum (Figure 4b) for RT device obtained by the difference

between aged and fresh spectra (LED Aged - Fresh) show the emergence of the peak at 1447 cm-1

upon device degradation. This peak is from PEDOT:PSS (Figure S11) and previously identified to

be symmetric C=C PEDOT intraunit stretching mode.[22] Importantly this peak is strongly related to

doping level of the PEDOT:PSS layer and its intensity becomes stronger when the doping level of

PEDOT:PSS is reduced (de-doped PEDOT:PSS). Such Raman intensity changes of PEDOT:PSS

layer is originated from a shift of the absorption band of PEDOT:PSS from near IR (> 700 nm) to

visible range (≈ 630 nm) upon de-doping, which is now resonant with the visible excitation

wavelength used for Raman measurement.[23] For long-term operated RT device, there is a stronger

selective enhancement of this PEDOT:PSS peak, indicating a change in the electronic property of

the PEDOT:PSS layer during device operation, i.e. PEDOT:PSS becomes less doped.[24] Such

change in the PEDOT:PSS doping level is expected to influence its capability acting as a hole

extraction layer, lowing JSC and PCE of the device.

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Examining the 140 °C case in Figure 4a, we do not observe such strong selective

enhancement of the PEDOT:PSS Raman peaks in the range of 1400-1500 cm-1, indicating no

significant change of PEDOT:PSS electronic property upon device operation. However,

surprisingly the intensities of all PCBM peaks (asterisked in Figure 4c) are markedly decreased with

respect to the intensities of polymer peaks. Such reduction of all PCBM peaks clearly indicates a

further phase separation between polymer and PCBM under continuous illumination, which alter

relative chemical compositions in a probed area. The micrograph of 140 °C devices in Figure 4c

clarifies the speculation raised; it shows appearance of macroscale PCBM aggregates after

degradation. Raman difference spectrum (Figure 4b) for 140 °C show the quenching of PCBM

peaks, however the quenching is not identical for all PCBM peaks. Such peak selectivity in

quenching of PCBM peaks after device aging is indicative of PCBM photodegradation.[13]

On another note, a common feature for both RT and 140 °C cases is that Raman peaks of

polymer specifically Peak-2 (BT, 1542 cm-1), 4 (T+BT, 1373 cm-1), 5 (T+BT, 1349 cm-1) and 6 (F,

1272 cm-1) show no differences between fresh and aged devices, suggesting minimal polymer

specific photodegradations. Increase in Peak 3 (T, 1445 cm-1) is extrinsic and due to other

component in device specifically PEDOT:PSS interlayer. We also observe the appearance of a new

peak (at 1624 cm-1) signifying degradation. The vibrational origin of this peak is not obvious as it is

not observed in any of the reference Raman spectra of the single layered Raman active materials

(Figure S11). But, it appears in BHJ/PEDOT:PSS multilayer (Figure S12) indicating a chemical

interaction between photoactive BHJ and hole extracting PEDOT:PSS layers which enhances upon

degradation. Nevertheless, this degradation signature is a common degradation route for both

unstable RT and high temperature annealed devices, not specific to thermal annealing control.

Based on what is observed in blend films and devices, we propose two different degradation

mechanism of BHJ devices during operation depending on the thermal annealing of photoactive

layer as shown in Figure 4c. Non-annealed morphology (RT devices), which are exclusively

kinetically developed via spin coating, comprises of optimum bicontinuous mixed domain to

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produce optimum efficiency (high JSC and PCE). However, the failure in operational stability

originates mainly from the unstable interface between the photoactive blend layer and the

PEDOT:PSS hole extraction layer leading to de-doping of PEDOT:PSS. This interfacial

degradation is responsible for observed reduction of JSC during operation (Figure 3d). We speculate

the reason of interlayer degradation to be due to inability in forming ideal PEDOT:PSS/photoactive

layer interfaces. Remnants of trapped solvents (due to no thermal treatment) together with ingress

of H2O into the interlayer (due to non-ideal interfaces) can act as the medium of permanent

PEDOT:PSS de-doping affecting specifically the initial phase (first 1000 mins) of device

degradation. Thus, the degradation is not due to initial photoactive layer morphology, rather due to

thermal anneal process itself.

On the other hand, mild-thermal annealing (< TPS) improves device stability mainly via the

stabilization of JSC. So, mild-thermal annealing of the photoactive layer is critical to improve

interface quality and partially remove trapped solvents whilst still maintaining the mixed blend BHJ

morphology desirable for high JSC during device operation. The upper limit of annealing

temperature to maintain improved stability is given by the phase segregation temperature (TPS) of

the bulk heterojunction blend. Likewise, high-thermal annealing (> TPS) creates non-ideal

microscale phase morphology where PCBM further diffuse out from mixed domains. In this case,

trapped solvents are removed and the multilayer is allowed to thermodynamically settle to create

better interfacial contacts protecting the PEDOT:PSS interlayer degradation (thus alleviating the

instability originated in PEDOT:PSS layer). However, nucleation sites for phase segregation in the

BHJ photoactive layer is already present due to initial high temperature annealing as established

previously in phase morphology study (Figure 2). This promotes further phase segregation between

the polymer and PCBM (demixing) during device operation leading to further loss of JSC and PCE,

generating macroscale PCBM aggregates (optically visible, Figure 4c). Further phase segregation

occurs during the continuous operation of devices in room temperature controlled environment.

PCBM in this phase is now susceptible to photodegradation. Further phase segregation and PCBM

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degradation are responsible for reduced JSC. Although similar reductions in JSC upon operation is

observed in both non-annealed (RT) and 140 °C annealed devices, the mechanisms for such

reductions are very different.

3. Conclusion

Thus, phase morphology tuning in polymer:fullerene BHJ OPV devices is vital not only for

high efficiency but also for long-term operational stability. With thermal preconditioning as a tool

to tune BHJ morphology of a model F12TBT:PC70BM BHJ blend system, a critical transition

temperature (TPS) is identified to be around 130 °C above which molecular-scale phase-segregation

occurs. Under operation of devices, stability clearly shows dependence on the state of initial BHJ

morphology. Initially phase-segregated morphology at a molecular level (annealed above TPS)

shows instability intrinsic to the photoactive layer via continuous phase-segregation between

polymer and fullerenes in macroscales, coupled with further fullerene photodegradation. On the

other hand, initially well-mixed blend morphology (no annealing) shows the interface instability

related to the hole transport PEDOT:PSS layer via de-doping. Based on these results, we conclude

that thermal annealing preconditioning of the photoactive layer is essential for improved device

operational stability; in particular annealing below TPS is critical in order to maintain a well-mixed

morphology with a stable interface desirable for high efficiency and long-lived devices.

4. Experimental Section

Materials. High purity F12TBT polymer was synthesized by Sumitomo Chemical Co., Ltd

with Mn of 10.5 kg/mol, Mw of 24.7 kg/mol and polydispersity of 2.35. PC70BM with 99 % purity

rating was obtained from Solenne BV products.

OPV device fabrication. Bulk heterojunction OPV were fabricated in conventional

geometry: Glass/ITO/PEDOT:PSS/F12TBT:PC70BM(1:3)/Ca/Al. 12x12 mm2 substrate of ITO

patterned in glass (Psiotec Ltd.) was cleaned sequentially by ultrasonication in acetone (5 mins),

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detergent (20 mins), DI water (5 mins, 3 times), acetone (10 mins), IPA (10 mins) followed by a

blow dry with N2 and oxygen plasma treatment at 100 W for 8 mins. Filtered PEDOT:PSS

(Heraeus) solution was spun at 3500 rpm for 35 s at acceleration of 10000 rpm/s and baked at 140

°C for 20 mins to obtain ≈30 nm films. 26.5 mg/mL total concentration in 1,2-Dichlorobenzene

solvent was used for all pristine and blend solutions of F12TBT and PC70BM. F12TBT:PC70BM

(1:3) was spun over ITO/PEDOT:PSS at 1500 rpm for 60 s at 10000 rpm/s to obtain thickness of

≈80 nm. Thermal annealing preconditioning was performed at this stage using a hot plate for 20

mins and allowed to cool down normally back to room temperature before metal electrode

deposition. Ca (20 nm) and Al (100 nm) metal evaporation deposition were performed through a

shadow mask (yielding 6 individual devices with active area of 0.045 cm2) inside an evaporator

chamber under high vacuum (< 5×10-6 mbar) with quartz crystal thickness monitor.

Device Characterizations. Xenon arc lamp was used for AM1.5G solar simulation, distance

calibrated with a Si-photodiode using Keithley 2400 source meter. Light degradation experiments

were performed inside nitrogen filled atmospheric chamber, temperature controlled at 25 °C by a

heat pump. All the devices were degraded together simultaneously to ensure the environmental

conditions are identical for all cases. White LED array (Bridgelux RS Array Series) was used for

continuous illumination and periodic in-situ J-V acquisition. 1 sun calibration for white LED was

performed by matching JSC of OPV sample from solar simulator.

Thin Film Characterizations. All thin film samples were fabricated by spin cast on top of

cleaned Quartz (Spectrosil 2000, UQG Ltd.) substrates. UV-Vis absorption spectra were acquired

from Shimadzu UV-2550 spectrophotometer in a transmission mode. Atomic Force Microscopy

was performed by NX-10 (Park Systems) in non-contact mode. Raman and Photoluminescence

spectroscopy were acquired from Renishaw inVia Raman Microscope in a backscattering

configuration with Argon-ion Laser pump for 457 and 514 nm excitation sources. The filter and

grating calibration was performed against a well-defined peak at 520 cm-1 of a Si reference. All

measurements were performed inside Linkam THMS600 temperature stage under continuous

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nitrogen purging. The temperatures were accurate to 0.1 °C. A constant heating/cooling rate at 10

°C/min was maintained in in-situ thermal experiments whilst the samples were allowed to settle for

5 mins at a given temperature before taking a scan to ensure thermal equilibrium. Around 10 µm

defocussed laser with 0.5 mW power and ≈30 s exposure time was used in order to avoid laser

induced degradations of organic materials. 5 different points of a sample was used to test

homogeneity of films and reproducibility of Raman scans.

Raman spectroscopy of whole device. Raman spectrum of the whole intact devices were

performed through the transparent ITO side, where the backscattered light from the opposing metal

electrode (Al) is enhanced due to high reflectivity of the metal hence increasing the backscattered

photon counts for good signal/noise ratio. LED-aged devices themselves and the Duplicates of

Fresh devices were transported into Linkam Stage to perform Raman scans under Nitrogen purged

environments.

DFT simulations. Symmetric F12TBT oligomers were built in Gaussian 09 with methyl

group replacing long alkyl side chains. All energy and frequency optimizations were performed

using B3LYP hybrid functional with 6-31G(d,p) basis set.

Supporting Information

Supporting Information is available from the Wiley Online Library or from the author.

Acknowledgements

Authors acknowledge the UK ESPRC for the Plastic Electronics Centre for Doctoral

Training (EP/L016702/1) funding and Cambridge Display Technology Ltd for a CASE studentship.

Received:

Revised:

Published online:

16

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19

Figure 1. (a) Normalized absorption spectra of F12TBT and F12TBT:PC70BM (1:3) films. (Inset)

Chemical structures of F12TBT polymer (electron donor) and PC70BM (electron acceptor). (b)

Resonant Raman spectra of the same films acquired by 514 nm excitation. Vibrational assignment

of F12TBT polymer is shown, derived from DFT simulation (Figure S1). Numbered peaks (1-6)

belong to polymer and asterisked peaks belong to PC70BM. (c) Normalized resonant Raman

spectrum of F12TBT acquired by 457 nm excitation, comparing pristine as-cast with thermally

annealed and PC70BM loaded films. (d) Raman peaks intensity ratio: Peak-5 (1349 cm-1)/Peak-1

(1606 cm-1), of as-cast F12TBT films against PC70BM content. Thermally annealed cases for 0 and

75 % wt of PC70BM content is shown by red triangles.

20

Figure 2. Resonant Raman spectra (via 514 nm excitation) of F12TBT:PC70BM (1:3) films

acquired in-situ during heating cycle with (a) showing absolute Raman spectra and (b) showing

baseline corrected Raman spectra normalized to F12 peak of polymer F12TBT. (c) Integrated

Photoluminescence (PL) intensities of F12TBT:PC70BM (1:3) films during the heating cycle. (d)

Raman peak intensity of PC70BM (1567 cm-1) with respect to F12 (1606 cm-1) during the heating

cycle. (e) Micrographs of F12TBT:PC70BM (1:3) film obtained in-situ at specified temperatures

during the heating/cooling experiment. Scale in micrograph images symbolized by black dash is 10

μm.

21

Figure 3. (a) J-V characteristics of conventional F12TBT: PC70BM (1:3) OPV devices fabricated at

different thermal annealing preconditioning temperatures. (b) Power conversion efficiency (PCE) of

the devices as a function of thermal annealing preconditioning temperatures. Additionally, it shows

PCE of these devices before (Fresh) and after (Aged) continuous white LED degradation under

operation for ≈3500 mins. LED degradation profiles of the devices with (c) showing normalized

PCE and (d) showing normalized JSC degradations.

Samples

Precondition

Temperature

VOC

(V)

FF JSC

(mA cm-2)

PCE

(%)

VOC

(V)

Fresh

(Aged)

FF

Fresh

(Aged)

JSC

(mA cm-2)

Fresh

(Aged)

PCE

(%)

Fresh

(Aged)

RT (22 °C) 1.01

±0.01

0.52

±0.02

6.44

±0.27

3.34

±0.22

0.99

(0.96)

0.54

(0.52)

6.47

(5.33)

3.46

(2.66)

70 °C 1.01

±0.01

0.51

±0.01

6.62

±0.25

3.41

±0.15

1.00

(0.98)

0.52

(0.51)

6.72

(6.46)

3.50

(3.22)

100 °C 1.02

±0.00

0.50

±0.02

5.98

±0.21

3.08

±0.15

1.01

(0.99)

0.50

(0.48)

6.53

(6.30)

3.30

(3.02)

140 °C 1.03

±0.02

0.41

±0.01

4.63

±0.17

1.95

±0.10

1.01

(0.97)

0.41

(0.38)

5.11

(4.80)

2.11

(1.79)

Table 1. Photovoltaic parameters of F12TBT: PC70BM (1:3) devices fabricated at different thermal

annealing preconditioning temperatures along with values of Before (Fresh) and After (Aged) LED

degradation as shown in Figure 3. Parameters were averaged over 12 devices.

22

Figure 4. (a) Raman spectra (via 457 nm excitation) of whole multi-layered devices comparing

Fresh and LED Aged of two unstable cases: Room Temperature (RT) and 140 oC annealed devices.

Asterisked peaks belong to PC70BM and numbered peaks (1-6) belong to F12TBT polymer. (b)

Raman difference spectrum (LED Aged – Fresh) of the two aforementioned cases. (c) Schematic

Illustration of proposed degradation mechanism of a model polymer:fullerene bulk-heterojunction

OPV devices dependent on its history of photoactive layer annealing during device fabrication.

23

Copyright WILEY-VCH Verlag GmbH & Co. KGaA, 69469 Weinheim, Germany, 2016.

Supporting Information

Impact of initial bulk-heterojunction morphology on operational stability of

polymer:fullerene photovoltaic cells

Saurav Limbu, Sebastian Pont, Alexander Doust, Sooncheol Kwon, Patrick Fuller, Ellasia Tan,

James R. Durrant, and Ji-Seon Kim*

Figure S1. (a) Absorption spectrum of F12TBT thin films marked with 514 nm (resonant with low-

energy absorption band) and 785 nm (non-resonant). (b) DFT simulated Raman spectrum compared

with experimentally obtained spectrum at 514 nm and 785 nm. (c) Structure of F12TBT oligomer

used in DFT calculation of Raman peaks. Also shown, is vibrational assignments of the numbered

peaks (1-6) to different units of the F12TBT monomer.

24

Figure S2. (a) Absorption spectrum of F12TBT thin film labelled with different wavelengths (457,

488, 514 and 765 nm) of laser excitation used for Raman acquisition. (b) Normalized Raman

spectra of F12TBT thin film using aforementioned laser excitations.

In Figure S2, we show the wavelength-dependent Raman spectrum of F12TBT. At non-

resonance excitation (785nm), different Raman peak intensities are determined by their Raman

cross-sections which are intrinsic property of the material. Changing to shorter wavelengths of

excitation (< 700nm), the excitation becomes resonant with the optical absorption transitions. At the

shorter wavelengths used here (457, 488, and 514 nm), the excitation is now strongly resonant with

the low-energy absorption transition which has a charge-transfer (CT) character. When the

excitation is resonant with this transition, there is a strong increase in Raman peak intensities

originated from the acceptor T-BT-T units.

25

Figure S3. Micrographs of as-cast and post-annealed (180 °C) F12TBT:PC70BM (1:3) thin films,

from which Raman spectra were acquired to identify the conformational variations of F12TBT upon

blending with PC70BM vs thermal annealing.

Figure S4. (a) Raw Raman spectrum of thin films between 800-1700 cm-1 Raman shift acquired by

514 nm laser excitation. (b) Raman spectra taken at different positions of a sample showing a good

level of homogeineity, acquired by 457 nm laser excitation.

514 nm excitation is resonant with low-energy absorption band of F12TBT polymer (Figure

S1a). For 514 nm excitation source used, the Raman shift converted to photon wavelength gives

range between 530-570 nm. Higher background obtained from F12TBT film is due its

photoluminescence (PL). There is no PL observed for pristine PC70BM at this range. PL is almost

completely quenched for F12TBT:PC70BM (1:3) blend film indicating charge transfer from

polymer (electron donor) to fullerene (electron acceptor) at heterointerfaces.

26

Figure S5. (a) Micrograph image of a phase segregated state of F12TBT:PC70BM(1:3) blend film

obtained post thermal annealing at 180 oC. Scale marked by black dash is 10 µm. (b) 2D Raman

map of the microaggregate. (c) Typical Raman spectrum obtained during 2D scan, which is used to

produce the whole 2D Raman map. Asterisk denote the peak due to PC70BM.

PC70BM peak (1567 cm-1) with respect to polymer peak-2 (BT, at 1542 cm-1) was traked in

X-Y plane to obtain the 2D Raman map of the micrograph image. The scan confirms that the

microaggregate is a PCBM aggregate. Furthermore, it also shows that the aggregate forms by

PCBM diffusing out from surrounding well-mixed domains, thus leaving out distinctive PCBM

deprived regions (blue shades in Figure S5) surrounding the microaggragate.

Figure S6. Atomic Force Microscopy (AFM) scans of Room Temperature RT (Rq = 1.0 nm), 100 oC (Rq =1.1 nm) and 140 oC (Rq = 0.7 nm) post-annealed cases for F12TBT:PC70BM(1:3) blend

films.

27

Figure S7. In-situ integrated photoluminescence (PL) intensities of F12TBT:PC70BM blend

compared to neat F12TBT film.

Figure S8. Energy level landscape of the OPV device structure used in this project which is a

conventional architecture utilizing F12TBT:PC70BM(1:3) as the photo harvesting layer. Air

photoelectron spectroscopy (APS) of F12TBT thin films used to calculate its HOMO level. The

LUMO level was approximated by adding the band-gap value from absorption on-set (Figure S1 a).

28

Figure S9. Absolute decay profiles of F12TBT:PC70BM(1:3) OPV devices fabricated at different

thermal annealing preconditioning temperatures showing decay of (a) PCE, (b) JSC, (c) VOC and (d)

FF. The degradation was performed all together in atmospheric chamber filled with N2 and

temperature controlled at room temperature (25 oC). Continuous white LED illumination was used

for degradation.

Figure S10. Normalized decay profiles of (a) VOC and (b) FF, obtained by normalizing absolute

values of Figure S9c and d.

29

Figure S11. Reference spectrum of single-layered films on quartz substrate acquired by 457 nm

laser excitation.

Figure S12. (a) Observation of 1624 cm-1 peak only in PEDOT:PSS/F12TBT:PC70BM(1:3) multi-

layered case. This peak is not present in single-layered films. (b) and (c) enhancement of 1624 cm-1

peak as a common degradation signature for both extremes of BHJ states (RT: well mixed

morphology and 140 oC: molecular scale phase segregation).