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Constructing nanostructured donor-acceptor bulk heterojunction via interfacial templates for efficient organic photovoltaics Zhiping Wang, Ying Zhou, Tetsuhiko Miyadera, Masayuki Chikamatsu, and Yuji Yoshida ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b13989 • Publication Date (Web): 27 Nov 2017 Downloaded from http://pubs.acs.org on November 27, 2017

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ACS Applied Materials & Interfaces

Constructing nanostructured donor-acceptor bulk heterojunction via interfacial templates for efficient organic photovoltaics Zhiping Wang,a,† Ying Zhou,b,* Tetsuhiko Miyadera,a Masayuki Chikamatsu,a Yuji Yoshidaa a

Research Center for Photovoltaic Technologies, National Institute of Advanced Industrial

Science and Technology (AIST), AIST Tsukuba Central 5, 1-1-1 Higashi, 305-8565, Tsukuba, Japan b

Electronics and Photonics Research Institute, National Institute of Advanced Industrial Science

and Technology (AIST), AIST Tsukuba Central 5, 1-1-1 Higashi, 305-8565, Tsukuba, Japan

KEYWORDS: organic solar cells, nanostructured bulkheterojunction, vacuum evaporation, molecular growth, structural template, cascade energy transfer.

ABSTRACT

We

demonstrate

that

a

poly(3,4-ethylenedioxythiophene):poly(styrene

sulfonate)/diindenoperylene (PEDOT:PSS/DIP) interfacial bilayer could serve as a structural

ACS Paragon Plus Environment

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template to enable the morphological control of bulkhetejunction by co-evaporation of tetraphenyldibenzoperiflanthene:fullerene(DBP:C60), which

greatly improves

the device

performances. Especially, we show that isolated crystalline domains of C60 can be well controlled at nanoscale during the co-evaporation. Photoluminescence spectra indicate the realization of DIP/DBP cascade energy architecture, which significantly facilitates both of the energy transfer and photocurrent generation. In addition, with bias-dependent external quantum efficiency analysis, we reveal that such cascade energy device architecture greatly suppresses the energy recombination in both of carrier and exciton transfer, resulting in a high open-circuit voltage and a high fill factor. By carefully optimizing the interfacial and BHJ layers, we achieved a high-performance OPV cell with a power conversion efficiency of 5.0±0.3%.

1. INTRODUCTION Organic photovoltaics (OPVs) have shown great potential as a new generation of low-cost, light-weight solar cells based on eco-friendly and elementally abundant materials.1–5 Significant breakthroughs in material design and device architectures have pushed OPV closer to commercial viability as the power conversion efficiencies (PCE) over 12 % have been achieved in both of single-junction and tandem cells.6,7 Among various device architectures currently being explored, bulk heterojunction (BHJ) comprising a blend of electron-donating and electronaccepting organic molecules typically exhibits more efficient exciton dissociation over its planar heterojunction (PHJ) counterpart, due to the widely dispersed donor/acceptor (D/A) interfaces.8,9 Achieving a phase-separated morphology at the nanoscale in a BHJ film is critical to the device performance, which is closely associated with the photoactive layer composites, processing


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conditions and the interfacial layer between active layer and electrode.10,11 Fine-tuning of the D/A interfaces in a BHJ photoactive layer at the molecular level for optimal exciton dissociation and charge transport guarantees the further improvement of device efficiencies as well as device reproducibility. For solution-processed OPV devices, thermal or solvent annealing is generally applied to the BHJ layer to achieve nanoscale interpenetrating D/A networks.12,13 For vacuum-processed OPV devices, morphological control of co-evaporated films is relatively complicated and elusive. Adopting substrate heating, alternative deposition, adjusting the co-evaporation rates of donor and acceptor molecules and introducing buffer layers or structural templates have been used to optimize the BHJ morphology.14–21 We have previously demonstrated that an introduction of nanostructured templating layers enables fine regulation of molecular growth in zinc phthalocyanine (ZnPc) and fullerene (C60) co-evaporated films, which enables high-performance OPV device with a PCE of over 4% due to the construction of interpenetrating D/A networks.22– 24

Despite remarkably increased short-circuit current (Jsc), the device performance is greatly

limited due to the low open-circuit voltage (Voc) owing to the shallow highest occupied molecular orbitals (HOMO) level of the ZnPc molecule (5.2 eV)25. Various molecules, such as squarine26 and tetraphenyldibenzoperiflanthene (DBP),27 that having deeper HOMO levels have been developed to enhance the Voc. In contrast to the planar molecules like ZnPc, these novel molecules typically possess complicated three-dimensional structure (i.e. non-planar molecules), which brings great difficulties in controlling the growth of these non-planar molecules due to the weak intermolecular interaction. Therefore, an amorphous nature is always observed in the final films. Further mixing the non-planar molecule with acceptor molecules (i.e. fullerene C60) forming a BHJ film makes it extremely challenging to achieve a well-defined nanostructure.28,29


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Corresponding solutions are highly desired to fully realize the potentials of these molecules towards higher device efficiencies. On the other hand, energy cascade concept has been proved to be effective in improving cell performances, especially in photocurrent generation, due to the exciton confinement effect.30–32 By employing multi-donors that have bandgaps monotonically decrease from donor side to acceptor side, the energy singlet exciton Förster transfer from the wide bandgap layer to low bandgap layer can be realized. However, such energy transfer effect has merely proven to be effective in PHJ configuration, and has not been applied to the more efficient BHJ architecture. In this paper, we demonstrate an effective method to control the growth of non-planar DBP molecule

in

co-evaporated

DBP:C60

BHJ

film.

A

ternary

poly(3,4-

ethylenedioxythiophene):poly(styrene sulfonate) (PEDOT:PSS)/ diindenoperylene (DIP)/DBP layer has been introduced between indium-tin-oxide (ITO) and DBP:C60 BHJ layer as a multifunctional template, which not only enables morphological control at nanoscale for efficient charge separation and transport, but also acts as an energy cascade buffer that improves photocurrent generation and ensures a high fill factor (FF) and Voc. Benefiting from these advantages, a champion PCE of 5.34 % is achieved. Our results highlight the strategy to control the growth of non-planar molecules in BHJ films and the importance of optimizing energy transfer process in BHJ-type OPV devices.

2. EXPERIMENTAL SECTION All photovoltaic cells were fabricated on commercially available ITO patterned on glass substrates with a sheet resistance of 10 Ω/square that was pre-treated in oxygen plasma for 30 min. A 40-nm-thick PEDOT: PSS (Clevios P VP Al 4083, Heraeus) film was spin-coated on ITO substrate, and then, annealed at 135°C on a hotplate for 30 min under ambient atmosphere. Prior


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to use, DIP (Lumtec) and DBP (Lumtec) were purified two times by vacuum gradient sublimation, while C60 (Frontier Carbon Corporation, 99.9%) and bathocuproine (BCP, Dojindo, 98%) were used as received. These materials were vacuum evaporated with a pressure less than 2 × 10-5 Pa at room temperature. OPV devices were fabricated with p-i-n structure that is ITO/PEDOT:PSS (0 or 40 nm)/DIP (0 - 15 nm)/DBP (5 nm)/i-layer (DBP:C60 blend: 20 - 60 nm)/n-layer (C60: 20 nm)/ BCP (3 nm)/Al (100 nm). The active device area was 4 mm2. At least 8 OPV devices were fabricated to obtain the standard deviations. The current-density versus voltage (J-V) characteristics of the cells were measured under dark conditions and simulated solar illumination (AM 1.5G) with a digital source meter (Keithley 2400). Incident power was calibrated using a standard silicon photovoltaic to match 1-sun intensity (100 mW/cm2). External quantum efficiency (EQE) spectra were collected using a Xe lamp that was integrated with a computer-controlled monochromator (CEP-3000 Spectral Response Measurement System, Bunkoukeiki.co). The irradiation intensity of monochromatic light was set to 5 mW/cm2 being independent on wavelength. For bias-dependent EQE measurement, the EQE spectra were collected by further applying a DC bias voltage between the electrodes from 0 to the VOC. Surface morphology was investigated using atomic force microscopy (AFM; Nanonavi/E-sweep, SII nanotechnology) in dynamic mode. Crystallinity was investigated by X-ray diffraction (XRD) analysis using the Bragg-Brentano configuration with a 9 kW rotating anode generator (Rigaku, Smartlab). Photoluminescence (PL) measurements were performed with a spectrofluorometer (JASCO) at 45o incident and detection angles. Excitation spectra of the films were measured with an emission wavelength of 690 nm. Absorption spectra were collected using an ultraviolet-visible–near-infrared (UV–vis–NIR) spectrophotometer


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(Shimadzu UV-3600). The samples for AFM and XRD were prepared using the same process as that used for fabricating OPV cells.

3. RESULTS AND DISCUSSION

Figure 1. AFM height images of co-evaporated 20-nm-thick DBP:C60 (1:4) films on (a) ITO; (b) ITO/PEDOT:PSS; (c) ITO/DIP; (d) ITO/PEDOT:PSS/DIP;

the corresponding AFM phase

images are in (e), (f),(g), (h). Note that a 5-nm-thick DBP layer was deposited between the coevaporated layer and the substrate layer, is consistent with the actual device structure. The scale bars are 100 nm.

Figure 1 shows height- and phase-contrast AFM images of DBP:C60 (20 nm) films with an evaporation ratio of 1:4. All the films show similar morphology with a mixture of small and large isolated domains. Their corresponding phase images (see Figures 1e-1f), which are


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sensitive to the variations in material characteristics such as adhesion force and viscoelasticity clear show the structural differences. Phase images indicate very few structural defects and the significant contrast between center and boundary in large domains, which are consistent with the characteristics of single crystals. On the other hand, the phase contrast for small domains is very small, possibly corresponding to amorphous or micro-crystallized structures. It can be considered that high concentration of C60 during co-evaporation results in a phase separated morphology, and the existence of amorphous DBP suppresses the crystalline growth of C60. Therefore, the small and large domains possibly correspond to the DBP:C60 blend and pure C60 areas, respectively. Moreover, it seems that the interfacial layers lead to structural variations of these domains. On bare ITO substrate (Figure 1a), the blended film displays round-grain morphology with a mixture of small domains (20 – 50 nm) and large isolated domains (100 – 120 nm). Its phase image (Figure 1e) clearly shows the different structural characteristics of these domains. When the ITO substrate is smoothened with PEDOT:PSS, the co-evaporated film shows increased C60 domains (150-200 nm) and correspondingly reduced density. It indicates that the crystalline growth of C60 during the evaporation is strongly dependent on the surface morphology. Compared with bare ITO substrate, similar size and distribution of C60 domains are observed with introducing 10-nm-thick DIP film on ITO, as shown in Figure 1c. On the other hand, the small domains disappear (in Figure 1g), suggesting the formation of uniform DBP:C60 domains. Further employing a PEDOT:PSS/DIP bilayer as a structural template, the C60 domains grow similarly with those on PEDOT:PSS, simultaneously, very few of small domains can be observed around C60 domains (Figure 1d and 1f). Apparently, C60-rich condition during the coevaporation leads to three-dimensional island growth. As a result, the height of these C60 crystals is significantly higher than DBP:C60 film. Considering that these C60 crystals play an important


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role in efficient photon-electron conversion including exciton diffusion, exciton dissociation and charge transfer, the morphological control of C60 is crucial here. It has been widely investigated that PEDOT:PSS with a high resistance not only improves the hole extracting efficiency due to the deep workfunction but also smoothens the rough ITO surface to enhance the crystalline growth of the organic film.22,23,25 Generally, organic molecules have much more chance to diffuse and form larger domains on PEDOT:PSS than on bare ITO. On the other hand, owing to their structural and electronic similarities (Figure S1), DIP and DBP molecules should have relatively strong intermolecular interactions between each other. We have previously demonstrated that DBP molecules prefer to grow on DIP nanocrystals, leading to the formation of nanostructured DBP film.30 Here, DIP layer also acts a structural template to control the structure of DBP and the following co-evaporated film, resulting in the formation of uniform DBP:C60 blend film. Figure S3 indicates that C60 domains are randomly distributed on a larger scale.


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Figure 2. (a) XRD patterns of DBP/DBP:C60 films on with/without interfacial layers of PEDOT: PSS or /and DIP. The AFM height images of 10-nm-thick DIP films on (b) bare ITO and (c) ITO/PEDOT:PSS. The scale bars are 200 nm. Figure 2a shows the XRD patterns of DBP:C60 blended films on various substrates. The peaks at 2θ = 21.5o and 30.5o come from the ITO substrate (polycrystalline In2O3).33 Distinctively, the sample prepared on ITO/PEDOT:PSS/DIP shows a peak at 2θ = 5.1°, which can be assigned to DIP crystals with standing-up order.30 In contrast, no DIP peaks are found on bare ITO substrate. Note that, here, DIP was evaporated by a general vacuum system with a deposition pressure of 2 × 10-5 Pa, while our previous work30 used a molecular beam deposition (MBE) system to grow highly crystallized organic film. To verify the growth of DIP, the surface morphology was investigated by AFM, as shown in Figures 2b and 2c. On bare ITO, 10-nm-thick DIP film


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exhibits poly-crystallites with grain sizes ranging from 10 to 50 nm. Apparently, the rough polycrystalline surface of ITO has a strong influence on the molecular growth of DIP. These grains are greatly increased to 50-100 nm when DIP grows on ITO/PEDOT:PSS. Similar surface morphology of DIP grown on PEDOT:PSS has been reported.34 The increased grain size should be consistent with the appearance of diffraction peak for DIP. Thus, it can be concluded that the smooth PEDOT:PSS surface enhances the crystalline growth of the DIP molecules because PEDOT:PSS-induced variations in the interaction between molecules and substrate should be small. Unfortunately, no crystalline DBP phase is identified, indicating the amorphous state of DBP film. Previous research suggests that the growth of DBP molecules preferentially follow the facet of DIP crystals due to their strong intermolecular interaction, which may responsible for the no long-range DBP crystalline order in our BHJ films.30 On the other hand, all the samples exhibit similar peaks at around 2θ = 27.3 and 29°, corresponding to the (311) and (420) diffractions of C60, respectively.35 The XRD results suggest that the DIP or PEDOT:PSS has negligible influence in the crystalline growth of C60.


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Figure 3. (a) Absorption and (b) photoluminescence excitation spectra of DBP/DBP:C60 films with/without interfacial layers of PEDOT: PSS or/and DIP.

Figure 3a shows the absorption spectra of DBP/DBP:C60 films with/without interfacial layers of PEDOT: PSS or DIP. All the films exhibit similar optical absorption for wavelengths ranging from 500 to 700 nm, suggesting that introduction of interfacial layer cannot modify the absorbance of the DBP. Note that introduction of PEDOT:PSS clearly increases the absorbance for wavelengths ranging from 380 to 480 nm, corresponding to optical absorption of C60 or PEDOT:PSS. However, it has been widely investigated that PEDOT:PSS shows typical absorption peaks at ultraviolet region between180 and 250 nm, and slightly increased absorbance with wavelength from 450 nm to near infrared region.36–41 Therefore, such increase may be attributed to the enhanced crystallinity of C60. Especially, the films on ITO/PEDOT:PSS exhibit a weak peak at a wavelength of around 440 nm, which are usually observed in crystalline C60.42– 44

Similar increase has been also observed in the previous literature,45 which introduced a co-

evaporant component to enhance the crystalline growth of C60 during co-evaporation. The physical origin of such increase remains unclear, but some reports have suggested that when C60 crystallite size becomes larger, the optical polarizability would be increased, leading to more efficient optical absorption.46 In spite of very similar XRD patterns, both of the AFM images and optical absorbances identify the enhanced crystallinity of C60. To understand the exciton transport in the DBP/DBP:C60 structure, we performed PL measurement. Figure 3b shows the PL excitation spectra with an emission wavelength of 690 nm. Here, since DIP and C60 have much smaller absorption coefficients or smaller PL efficiencies than DBP, the PL signals corresponding to the optical absorption of DIP and C60 layers are negligible. The PL intensity for


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the DBP absorption region is significantly enhanced by inserting a DIP layer. Apparently, the isolated C60 domains in co-evaporated DBP:C60 layer cannot cover DBP domains, and quench the excitons in DBP absorption region. It also indicates a phase-separated morphology of DBP:C60. It has been discussed that such increase may be ascribed to the exciton confinement effects.47,48 Here, DIP layer can prevent excitons from being quenched on the ITO or PEDOT:PSS. On the base of Förster resonance energy transfer model,49 singlet exciton transport can be considered as a hopping process of energy. The availability of empty states in the density of state determines the exciton transfers in an organic layer. Excitons will randomly move to the nearest neighbor hopping sites until annihilation by recombination. However, there is a very small chance of the excitons transferring from lower to higher lying-energy states. Apparently, excitons generated in the DBP domains cannot move to the DIP domains because of the large energy barrier (0.8 eV) between their lowest unoccupied molecular orbital (LUMO) levels. As a consequence, excitons generated in DBP domains can be concentrated at the DIP/DBP interface, which leads to energetically ordered hopping sites toward to the DBP/C60 interface. Therefore, a DIP/DBP cascade structure pushes more excitons to the DBP/C60 interface where dissociation of excitons occurs. Here, further deposition of C60 continuous layer is required to lead such enhancement directly contribute to photocurrent generation.


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Figure 4. (a) J-V characteristics and (b) EQE spectra of OPV cells with different interfacial layers. The device structures are ITO/PEDOT:PSS (0 or 40 nm)/DIP (0 or 10 nm)/DBP (5 nm)/DBP:C60 (20 nm)/C60 (20 nm)/BCP/Al.


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Table 1. Device performance of OPV cells. The errors represent standard deviation. Device architecture Interfacial layer

a)

Device characteristics DBP:C60 Thick. (nm)

PCE (%)

Jsc mA/cm-2

Voc (V)

FF

PEDOT:PSS

DIP (nm)

Blend ratio

None

0

1:4

20

3.3±0.1

5.9±0.2

0.82±0.01

0.70±0.01

None

10

1:4

20

3.6±0.1

6.7±0.2

0.91±0.01

0.60±0.02

PEDOT:PSS

0

1:4

20

3.7±0.2

6.5±0.1

0.90±0.01

0.63±0.02

PEDOT:PSS

10

1:4

20

4.1±0.3

7.1±0.2

0.90±0.01

0.64±0.02

PEDOT:PSS

5

1:4

40

3.1±0.5

7.9±0.5

0.88±0.03

0.44±0.04

PEDOT:PSS

10

1:4

40

5.0±0.3

9.4±0.2

0.91±0.01

0.58±0.03

9.54

0.92

0.61

a)

PEDOT:PSS

10

1:4

40

5.34

PEDOT:PSS

15

1:4

40

2.2±0.3

7.1±0.3

0.86±0.04

0.36±0.04

PEDOT:PSS

10

1:1

40

2.6±0.4

7.3±0.3

0.90±0.01

0.39±0.05

PEDOT:PSS

10

1:2

40

3.9±0.3

9.0±0.3

0.91±0.01

0.48±0.03

PEDOT:PSS

10

1:6

40

2.9±0.6

6.7±0.8

0.90±0.02

0.48±0.07

PEDOT:PSS

10

1:4

60

3.0±0.4

8.9±0.4

0.90±0.02

0.38±0.05

Best device.

To study the effects of the interfacial layer, OPV devices with a structure of DBP (5 nm)/DBP:C60 (20 nm)/C60 (20 nm)/BCP (10 nm)/Al (100 nm) on various substrates were prepared and their device performances are evaluated. Figure 4a shows the J-V characteristics, and Table 1 summarizes the detail parameters. On bare ITO film, the OPV device exhibits a small Jsc of 6.02±0.10 mA/cm2, a small Voc of 0.82 ± 0.01 V, and a high FF of 0.70 ± 0.01. The introduction of a PEDOT:PSS layer between ITO and DBP enhances both of the Jsc and the Voc. When replacing PEDOT:PSS with DIP forming a DIP/DBP cascade buffer layer, the OPV cells show a substantial improvement in Jsc from 6.46 ± 0.15 to 6.76 ± 0.12 mA/cm2 while retaining a high Voc of 0.91 ± 0.01 V. Inserting a bilayer of PEDOT:PSS/DIP layer between ITO and DBP further boost the Jsc up to 7.21 ± 0.10 mA/cm2, without sacrificing the Voc and FF as comparing


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to the OPV devices with other types of interfacial layers. Such enhancements result in a PCE of 4.1 ± 0.3 %, being 125% higher than that of OPV device on bare ITO. To understand the effect of the interfacial layer on Jsc, we measured the EQE spectra in Figure 4b, where EQE peaks at wavelengths of λ = 560 nm and λ = 605 nm correspond to the DBP absorption region, and the peak at around λ = 440 nm corresponds to C60. Inserting the interfacial layers leads to an increase in the EQE for the DBP absorption region. PL results show that exciton confinement effects due to the DIP increase the exciton transfer efficiency in DBP, which directly contributed to Jsc. On the other hand, only the devices with PEDOT:PSS shows an increase in EQE for C60 absorption region. Generally, PEDOT:PSS as a hole extraction layer will improve the photocurrent generation efficiency independently of the spectra. Thus, we attribute such increase to the structural improvement of co-evaporated DBP:C60, especially the enhanced crystalline growth of C60. Figure 1 clearly shows that C60 crystalline domains grow larger on PEDOT:PSS. Figure 3a identifies the increase in optical absorbance of C60 due to the crystalline growth. In addition, the well-crystallized C60 domains would be beneficial to the exciton diffusion efficiency as well as the charge transfer efficiency in C60, and thus enhance the photocurrent generation in both of the DBP and C60 absorption region. As a result, the device on ITO/PEDOT:PSS/DIP exhibit a higher EQE for DBP absorption region than that on ITO/DIP. Consequently, the OPV device with PEDOT:PSS shows an overall enhancement of the entire EQE spectra. In contrast to the device on the bare ITO, the device on PEDOT:PSS/DIP bilayer shows a 150% increase in the EQE peak at λ = 605 nm.


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Figure 5. (a) Bias dependent EQE spectra of the OPV cells with different interfacial layers, and the reduction rate of EQE spectra with a bias voltage of (b) VM and (c) VM + 0.1 V.

A few of models has been well established to understand the physics governing JSC50 and VOC51 of OPV devices, while thorough physical understanding is still lacking for FF. FF characterizes maximum power (JM × VM) by Jsc and Voc, where JM and VM are the current density and voltage at the maximum power point, respectively. Apparently, charge recombination should dominate such energy-loss phenomena induced by an electrical field. There has been an increasing number of publications52–60 that correlate the changes in FF to some parameters including mobilities of holes and electrons, interfacial layer, light intensity, temperature, and so on. Especially,


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monomolecular and bimolecular recombination can be distinguished depending on whether the electron and hole are bound as a geminate pair or dissociated into free carriers. These works make it possible to understand the dynamical behavior of charge recombination, and its correlation with FF. For example, Chen et.al28 have systematically characterized light-intensity dependent DBP:C60 based OPV devices, which indicate that biomolecular recombination is dominate for C60-riched BHJ. With considering the device structure and performance here, similar biomolecular recombination is possibly dominated for all of the devices. On the other hand, very few paper investigate the recombination of execution, which should be one of prominent recombination mechanisms in organic semiconductors. We have previously demonstrated that bias-induced EQE reductions may origin from a bias-induced leakage current or energy recombination in exciton transfer, exciton dissociation, charge transfer, and charge extraction, which is directly related to FF.30 Especially, the correlation between exciton recombination and FF has been demonstrated. Here, bias-dependent EQE spectra are utilized to understand the effects of interfacial layers on Voc and FF, as shown in Figure 5. Apparently, smaller bias-dependent EQE leads to higher FF and Voc. Note that all the cells show very similar VM, despite the fact that the cell on bare ITO has a smaller Voc. As shown in Figure 5b, for a cell on bare ITO, EQE spectra ranging from 350 to 630 nm decrease proportionally with increasing the bias to VM of 0.67 V, which leads to a 30% reduction in overall EQE. With a PEDOT:PSS film, the EQE in C60 absorption region seems to be more sensitive to bias, resulting in a 35% reduction in overall EQE. For the devices with DIP or PEDOT:PSS/DIP, the reduction of EQE spectra in DBP absorption region is further suppressed, especially for λ > 630 nm, while those in C60 region are not improved. DIP crystalline film could enhance the hole transfer or suppress the recombination for efficient hole extraction at ITO electrode due to the wide bandgap. Generally,


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the recombination in carrier transfer in either C60 or DBP will lead to a variation in the magnitude of the EQE spectrum, but not a variation in the form. However, for the devices with DIP, the EQE spectra for C60 absorption range show a larger bias-induced reduction than those for DBP absorption range. These phenomena can be attributed to polaron induced exciton quenching, where an exciton (regarded as a coulomb-bounded electron-hole pair) recombines with a third charge carrier (polaron) generated by applied bias.61,62 Clearly, as compared to DBP, exciton recombination is more serious in C60. PL results (in Figure 3b) imply that DIP/DBP cascade structure can assist more excitons to transfer from DBP to DBP/C60 interface. It seems that such cascade structure also suppresses the polaron induced recombination. Furthermore, these phenomena are more apparent when the bias voltage is increased to VM + 0.1 V of 0.77 or 0.78 V, as shown in Figure 5c. Such bias voltage is close to (or larger than) the threshold of the diode, which is mainly determined by the HOMO level of DBP and LUMO level of C60, resulting in a larger leakage current. Significant reductions in both DBP and C60 absorption region are observed in all of the devices. For the device on bare ITO, the reduction in C60 absorption region is twice high than that for DBP, indicating more serious recombination. On the other hand, inserting a PEDOT:PSS film clearly suppresses the reduction in C60, while those in DBP absorption region remains similar. For the devices on bare ITO, the relatively smaller VOC of 0.82 should be attributed to larger bias-dependence of EQE. Especially, when the bias is increased to 0.77 V, significant exciton recombination for C60 occurs. As well known, the efficiency of exciton dissociation of C60 is strongly affected by the crystallinity, morphology, and thickness of the film. Here, the effects of electron extraction layer can be ignored, since all the devices use BCP/Al to collect electrons. In spite of the factor that no clear difference can be observed in XRD in Figure 2, PEDOT:PSS improves the crystallinity of C60 domains during the


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co-evaporation, and suppress the polaron induced recombination. Furthermore, the introduction of a PEDOT:PSS/DIP bilayer greatly suppresses the reductions in both of DBP and C60 absorption region. Especially, the ratio for DBP remains as high as 0.7. Therefore, PEDOT:PSS/DIP not only leads to a cascade architecture for efficient energy harvesting but also enables structural control of DBP:C60 during the co-evaporation as a structural template. As a result, the bias-induced energy recombination is suppressed, and a high FF of 0.64±0.02 and a high VOC of 0.90±0.01 are obtained. These results also indicated that constructing cascade devices architecture with the desired nanostructure at the nanoscale should be crucial to achieve high FF and VOC for BHJ, no matter whether being vacuum evaporated or solution-processed.

Figure 6. J-V characteristics of the OPV devices with structures of (a) ITO/PEDOT:PSS/DIP (10

nm)/DBP

(5

nm)/DBP:C60

(various

thickness)/C60

(20

nm)/BCP/Al,

(b)

ITO/PEDOT:PSS/DIP (various thickness)/DBP (5 nm)/DBP:C60 (40 nm)/C60 (20 nm)/BCP/Al,


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(c) ITO/PEDOT:PSS/DIP (10 nm)/DBP (5 nm)/DBP:C60 (various co-evaporation ratio)/C60 (20 nm)/BCP/Al. Their corresponding EQE spectra are (d), (e) and (f), respectively.

We show that a bilayer interfacial layer of PEDOT:PSS/DIP is effective to improve the device performance. In addition, the dependencies of thickness and the co-evaporation ratio of the coevaporated film, and the thickness of DIP film are investigated in detail, as shown in Figure 6. Figure 6a shows J-V curves of OPV devices, where the thickness of co-evaporated DBP:C60 is varied from 20 to 60 nm. Increasing the thickness from 20 to 40 nm significantly increases Jsc from 7.1 to 9.4 mA/cm2. Both of FF and Voc remain high, resulting in a high PCE of 5.1%. However, further increasing the thickness to 60 nm leads to a large degradation in FF. The amorphous DBP film has a poor carrier mobility, so that, the recombination of charge due to the applied bias could be significant for a thick DBP film or large DBP domains. The bias-dependent EQE spectra are also studied, as shown in Figure S5. Apparently, the device with 40 nm-thick DBP:C60 shows a similar bias-dependent behavior with 20 nm-thick DBP:C60. While, for 60 nmthick DBP:C60, overall of the EQE spectrum with an applied bias of 0.52 V becomes 50% smaller compared to the devices with 40 nm-thick DBP:C60. As a result, it shows a poor FF of 0.38. On the other hand, the J-V characteristics are strongly influenced by the co-evaporation ratio, as shown in Figure 6b. It shows that the highest Jsc can be achieved for co-evaporation ratios in a range from 1:2 to 1:4. The corresponding EQE spectra in Figure. 6e indicates that the photocurrent generation efficiencies are lower in C60 and DBP absorption region for coevaporation ration of 1:1 and 1:6, respectively. Furthermore, in Figure S6, their bias-dependent EQE spectra indicate that with increasing C60 composition, the reduction of EQE spectra due to the applied bias becomes smaller, leading to a higher FF. It is possibly attributed to the


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crystalline growth of C60, which is preferred for a higher evaporation rate during the coevaporation of DBP:C60. As a result, co-evaporated DBP:C60 with a ratio of 1:4 should be the best BHJ structure, which can not only balance the exciton dissociation and charge transport processes of DBP and C60 but also suppress the exciton and carrier recombination due to the applied bias. Finally, Figures 6c and 6f suggest that the thickness of DIP film mainly influences FF, and optimized thickness for the DIP is 10 nm.

4. CONCLUSION In summary, we have demonstrated that constructing an interfacial bilayer of PEDOT:PSS/DIP as a template can greatly improve the performance of OPV device based on co-evaporated DBP:C60 BHJ. PEDOT:PSS can smooth the ITO surface, which enhances the crystalline growth of C60 during the co-evaporation. These three-dimensional crystalline C60 domains greatly improve the photocurrent generation for C60 absorption region. On the other hand, DIP/DBP cascade architecture enables efficient exciton transfer in DBP domains, which significantly facilitates the energy transfer and photocurrent generation. In addition, the study on the biasdependent photocurrent generation behavior by EQE indicates that the cascade architecture can suppress energy recombination in exciton and carrier transfer under a forward bias. The optimized OPV device shows a power conversion efficiency of 5.0±0.3%.

ASSOCIATED CONTENT Supporting Information.


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Schematic energy diagram, additional AFM images, J-V characteristics and EQE spectra of the OPV devices. This material is available free of charge via the Internet at http://pub.acs.org AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] Present Addresses †Present address: Clarendon Laboratory, Department of Physics, University of Oxford, Parks Road, Oxford, OX1 3PU, United Kingdom Notes There are no conflicts to declare ACKNOWLEDGMENT

This work was partially supported by JSPS Grant-in-Aid for Research Activity Start-up and Challenging Exploratory Research.

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