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Materials Chemistry and Physics 132 (2012) 131– 137

Contents lists available at SciVerse ScienceDirect

Materials Chemistry and Physics

j ourna l ho me pag e: www.elsev ier .com/ locate /matchemphys

ffects of annealing on the polymer solar cells based on CdSe–PVK electroncceptor

zong-Liu Wanga,∗, Chien-Hsin Yanga, Yeong-Tarng Shieha, An-Chi Yehb, Chin-Hsiang Chenc,sung-Han Hod

Department of Chemical and Materials Engineering, National University of Kaohsiung, Kaohsiung 811, Taiwan, ROCDepartment of Chemical and Materials Engineering, Cheng Shiu University, Kaohsiung 833, Taiwan, ROCDepartment of Electronics, Cheng Shiu University, Kaohsiung 833, Taiwan, ROCDepartment of Chemical and Materials Engineering, National Kaohsiung University of Applied Sciences, Kaohsiung 807, Taiwan, ROC

r t i c l e i n f o

rticle history:eceived 4 May 2011eceived in revised form 22 October 2011ccepted 11 November 2011

eywords:omposite materialshotoluminescence spectroscopy

a b s t r a c t

CdSe–poly(N-vinylcarbazole) (CdSe–PVK) nanocomposite was synthesized and utilized as the electronacceptor in the active layer of polymer solar cells. The photovoltaic properties of the polymer solar cellsbased on poly(3-hexylthiophene) (P3HT):CdSe–PVK as the active layer were investigated in detail. Theeffects of annealing temperature (100–200 ◦C) and time (5–60 min) on the device performance werestudied. At annealing temperature of 150 ◦C for 30 min, the device demonstrated an optimal efficiencyof 0.235% under AM 1.5 (100 mW cm−2) solar simulated light irradiation. The improved efficiency underthe optimal conditions was confirmed by the highest light harvest in UV–vis spectra due to the increased

tomic force microscopynnealing

crystallinity of P3HT after thermal annealing. Photoluminescence of these devices also exhibited that thequench effect increases with the increasing of annealing temperature, indicating that the charge sepa-ration between electron-donating (P3HT) and electron-accepting (CdSe–PVK) molecules was increasedafter heat treatment. Atomic force microscopy (AFM) images showed that the phase segregation and3D interpenetrating networks of P3HT:CdSe–PVK were responsible for the enhancement of the device

efficiency.

. Introduction

Over the past decades, polymer solar cells (PSCs) basedn conjugated polymers have attracted considerable attentionecause of their potential use for future cheap and renewablenergy production [1–3]. Efficient polymer-based solar cells uti-ize donor–electron acceptor (D–A) bulk heterojunction (BHJ) filmss active layers [1,2]. The donor is typically a kind of conjugatedolymer, while the acceptor is generally a type of organic or inor-anic molecule. The D–A BHJ structure enables availability of thecceptor molecules in close proximity to the electron donor poly-ers, and thereby facilitates charge transfer from excited polymer

hains to the electron acceptor molecules. Recently, many bulketerojunction solar cells based on blends of conjugated poly-ers and inorganic nanocrystals that offer high electron mobility

r improved spectral coverage have been investigated [4–9]. Fur-

hermore, it has been stated that semiconductor nanocrystalscolloidally synthesized quantum dots, QDs) have the poten-ial to increase the efficiency of conversion of solar photons to

∗ Corresponding author. Tel.: +886 7 5919278; fax: +886 7 5919277.E-mail address: tlwang@nuk.edu.tw (T.-L. Wang).

254-0584/$ – see front matter © 2011 Elsevier B.V. All rights reserved.oi:10.1016/j.matchemphys.2011.11.008

© 2011 Elsevier B.V. All rights reserved.

electricity up to about 66%, and can overcome the efficiency limitcaused by carrier thermalization in the conventional solar cells [10].In particular, spherical semiconductor nanoparticles, such as CdSequantum dots, have been the subject of extensive studies over thepast decade because of their unique optical and electronic prop-erties [11–13]. Since CdSe has high electron mobility, we haveemployed this specific property for improving device efficiency ofhybrid CdSe/conjugated polymer photovoltaic (PV) system in thiswork.

Although CdSe/conjugated polymer solar cells have been stud-ied in several groups [5,14–18], the efficiency of photovoltaicdevices is majorly limited by the low compatibility betweeninorganic CdSe nanoparticles and the conjugated polymers. Con-sequently, a good dispersion of the CdSe nanocrystals in thepolymer donor phase is required for the CdSe/conjugated polymernanocomposite in solar cells to create a larger interfacial surfacearea, enhancing the charge transfer between polymer matrix andnanocrystals.

In our previous report, we have described the synthesis of

CdSe–PVK nanocomposite via ATRP approach [19]. In compari-son with CdSe or PVK, this new nanocomposite possesses uniqueoptical properties. Further effort has been made to use this com-posite as the electronic acceptor in the active layer of PSCs [20].

1 istry and Physics 132 (2012) 131– 137

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normally means increased packing of the P3HT domains. The filmheat-treated at 125 ◦C shows a similar behavior. The maximumabsorption is observed for the film annealed at 150 ◦C, indicatingan enhanced conjugation length and the more ordered structure of

300 400 500 600 70 0 800

Abs

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100 C

125 C

150 C

175 C

200 C

32 T.-L. Wang et al. / Materials Chem

ore recently, organic solar cells made from blends of poly(3-exylthiophene) (P3HT) and [6,6]-phenyl-C61-butyric acid methylster (PCBM) exhibited that both the external quantum efficiencyEQE) and power conversion efficiency (PCE) were remarkablymproved by thermal annealing and/or electrical aging on theevices [21,22]. Thermal annealing condition was the key variableor improving short circuit current, fill factor, and therefore thefficiency of the device. Herein, a comprehensive study of ther-al annealing conditions is presented on the power conversion

fficiencies of PSCs based on the active layer of P3HT:CdSe–PVK.n addition, the relationship between device performance and the

orphology of active layers is also investigated by UV–vis, photo-uminescence, and AFM.

. Experimental

The polymer photovoltaic cell in this study consists of a layerf P3HT/CdSe–PVK blend thin film sandwiched between transpar-nt anode indium tin oxide (ITO) and metal cathode. The devicetructure is ITO/PEDOT:PSS/P3HT:CdSe–PVK/Al. P3HT (FEM Tech.,n = 16, 900) acts as the p-type donor polymer and CdSe–PVK

Mn = 950 for PVK) as the n-type acceptor in the active layer.efore device fabrication, the glass substrates coated with indiumin oxide (ITO) were first cleaned by ultrasonic treatment incetone, detergent, de-ionized water, methanol and isopropyl alco-ol sequentially. The ITO surface was further coated with ca.0 nm layer of poly(3,4-ethylene dioxythiophene):poly(styrene)PEDOT:PSS) by spin coating. The substrate was dried for 10 mint 140 ◦C in air, and then moved into the nitrogen-filled glove-ox for spin coating the active layer. The P3HT:CdSe–PVK blendas prepared with 1:1 weight ratio (10 mg mL−1 P3HT) in 1,2-ichlorobenzene (DCB) as the active layer. This solution blend waspin-coated onto the PEDOT:PSS layer at 800 rpm for 30 s. Thebtained thickness for the blend film of P3HT:CdSe–PVK was ca.00 nm. The device was completed by depositing a thin Al layer asn electrode with an area of 6 mm2 as defined by a mask.

The films of active layers were annealed directly on top of a hotlate in the glove box, and the temperature was monitored by using

thermocouple touching the top of the substrates. After removalrom the hotplate, the substrates were immediately put onto a

etal plate at the room temperature. Ultraviolet-visible (UV–vis)pectroscopic analysis was conducted on a Perkin-Elmer Lambda5 UV-Vis spectrophotometer. Room temperature photolumines-ence (PL) spectrum was recorded on a Hitachi F-7000 fluorescencepectrophotometer. The film topography images of active layersere recorded with a Digital Instruments Dimension 3100 atomic

orce microscope (AFM) in tapping mode under ambient condi-ions. The J–V curves were measured using a Keithley 2400 source

eter, under illumination from a solar simulator. The intensity ofolar simulator was set with a primary reference cell and a spec-ral correction factor to give the performance under the AM 1.5100 mW cm−2) global reference spectrum (IEC 60904-9).

. Results and discussion

.1. Effects of annealing temperature

Fig. 1 shows the chemical structures of the materials used in thistudy. The effect of annealing temperature on the UV–vis absorp-ion spectra for the thin films of P3HT:CdSe–PVK (1:1 weight ratio)pun cast on quartz substrates is shown in Fig. 2. These films were

nnealed under nitrogen atmosphere inside the glove box at atmo-pheric pressure. The annealing time was kept 30 min for all of thennealing temperatures. The absorption spectra show a consider-ble change after thermal annealing of the films. For the untreated

Fig. 1. (a) Chemical structures of CdSe–PVK used in the active layer. (b) Schematicrepresentation of a typical device structure of polymer bulk-heterojunction photo-voltaic device fabricated in this study.

film, the peak absorption wavelength (�max) is 516 nm with twoshoulders ranging from 540 to 553 nm and 590 to 605 nm, respec-tively. The first two bands is attributed to the �–�* transition,whereas the last shoulder is due to the inter-chain interactions[23,24]. The absorption peaks at about 344 and 331 nm are ascribedto the �–�* transition of carbazole groups that are pendent to thePVK backbone of CdSe–PVK [19].

At annealing temperature of 100 ◦C, the intensities of threebands increase without change in the position of the three vibronicpeaks. An increase in the absorption strength after heat treatment

Wavelength (nm)

Fig. 2. UV–vis absorption spectra of P3HT:CdSe–PVK blend films after annealing atdifferent temperatures for 30 min.

T.-L. Wang et al. / Materials Chemistry and Physics 132 (2012) 131– 137 133

550 600 650

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Wavelength (nm)

RT100 C125 C150 C175 C200 C

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Fig. 4. AFM topography images (5 �m × 5 �m) of P3HT:CdSe–PVK blend films after

ig. 3. Photoluminescence spectra of P3HT:CdSe–PVK blend films after annealingt different temperatures for 30 min with excitation at 525 nm.

3HT. Since the thickness of all the films is similar (ca. 100 nm), thencrease in the peak absorption intensity and more distinguishablehoulders during thermal annealing may be attributed to the low-ring of the band gap between � and �*, the increase of the optical–�* transition, and the increased inter-chain interaction among

he P3HT chain [25]. After thermal annealing, the P3HT moleculesfford higher energy and can move more easily. Consequently, theolymer chains become mobile and self-organization can occur toorm ordering. Therefore, the peak intensity increases and bothhoulders become more distinguishable in the more ordered films.

Further increasing the annealing temperature to 175 ◦C, how-ver, results in a decrease in the intensities of the three bands. At00 ◦C, the film shows even weaker absorption bands than thatf annealing at 100 ◦C. On the other hand, two peaks at 344 and31 nm, corresponding to CdSe–PVK, increase significantly afterhermal annealing at 100 and 125 ◦C, indicating a more orderedtructure of CdSe–PVK. However, a slightly decrease of the absorp-ion is observed for the film annealed at 150 ◦C and the absorptionsf films further decrease at annealing temperatures of 175 and00 ◦C. From the above results, the optimum annealing temper-tures for P3HT and CdSe–PVK are 150 and 125 ◦C, respectively.ince the radiation flux of visible light is larger than that of ultravi-let light in the solar radiation spectrum, the optimum annealingemperature is conceivable to be 150 ◦C.

Fig. 3 shows the PL intensity for blend films annealed at dif-erent temperatures. The PL is due to photogenerated excitons in3HT that do not take part in charge separation. The phenomenonf PL quenching can be attributed to the interfacial charge transfer.ormally, PL quenching increases with the increase of interfacialrea between donor and acceptor materials in the active layer. PLuenching provides direct evidence for exciton dissociation, andhus efficient PL quenching is necessary to obtain efficient organicolar cells. However, Yang et al. indicated that a large quenchffect is no absolute guarantee of high power conversion efficiency26], as the result shown by the unannealed film. Compared to thentreated film, the blend film annealed at 100 ◦C shows an intenseeak in the PL spectrum, which may be due to the reduction ofhe interfacial area between the P3HT and CdSe–PVK and resultsn a decrease of exciton dissociation. The reduction of the interfa-ial area between the D–A materials is also confirmed by the AFM

mages, as shown in Fig. 4. This may be attributed to the enhance-

ent of phase segregation between P3HT and CdSe–PVK uponnnealing, which is consistent with the improved polymer order-ng where more pure P3HT and CdSe–PVK domains are formed. In

annealing at different temperatures for 30 min. 2D height image for the blend film(a) unannealed, (b) annealed at 100 ◦C, (c) annealed at 125 ◦C, (d) annealed at 150 ◦C,(e) annealed at 175 ◦C, and (f) annealed at 200 ◦C. Phase image for the blend film (g)unannealed, (h) annealed at 100 ◦C, (i) annealed at 125 ◦C, (j) annealed at 150 ◦C, (k)annealed at 175 ◦C, and (l) annealed at 200 ◦C.

134 T.-L. Wang et al. / Materials Chemistry

Table 1Surface roughness of blend films obtained from AFM after annealing at differenttemperatures for 30 min.

Annealing temp.

RT 100 ◦C 125 ◦C 150 ◦C 175 ◦C 200 ◦C

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Average roughness (nm) 1.02 1.46 4.40 4.26 4.69 3.40Root mean square (nm) 1.42 2.04 5.60 5.78 6.12 4.33

ddition, it can be seen that the PL intensity decreases with thencrease of annealing temperature. The PL intensity shows a min-mum at thermal annealing of 175 ◦C. This significant reduction inhe PL intensity is attributed to efficient photoinduced charge sep-ration between electron-donating (P3HT) and electron-acceptingCdSe–PVK) molecules. This can be attributed to the higher inter-acial area between D–A molecules compared with those of othernnealing temperatures, as seen in Fig. 4. However, this does notecessarily mean that the stronger the PL quenching, the betterhe performance of the solar cells. Although the PL intensity of thelend film annealed at 150 ◦C is a little higher in comparison withhat of the film annealed at 175 ◦C, the highest power conversionfficiency has been achieved by this blend film. As we will discuss inhe following part, although the phase separation reduces excitonissociation, the increase of optical absorption in the visible lightegion and the improved charge carrier transport in both donor andcceptor phases after thermal annealing offsets the former effectnd results in an overall improvement in device performance.

The BHJ structure basically consists of 3D interpenetrating net-orks of donor and acceptor materials. The morphology of theolymer blend layer plays an important role in determining deviceerformance. The typical height and phase images of blend filmsbtained from AFM before and after annealing at different temper-tures are shown in Fig. 4. Two kinds of surface roughness from theeight images are listed in Table 1. Two features have been used forhe comparison of solar cell performance. For the height images, itas been indicated that the higher roughness of the film may giveigher efficiency device [25], while a highly ordered texture fromhe phase images represents a higher optical absorption in the solaradiation spectrum. As to the effect of roughness, it might be thencreased contact area between the polymer film and the metalathode for the films with higher surface roughness, resulting in aore efficient charge collection at the interface. It is known that

he diffusion length of excitons in organic materials is often only few nanometers [27], and only those carriers generated near thelectrodes will be collected and give rise to a photocurrent. There-ore the increased contact area will definitely have an effect on thefficiency of charge collection at the metal polymer interface. Thencreased surface roughness might also enhance internal reflectionnd improve light collection, which would also increase device effi-iency. But the surface area in the roughest film was found to benly 1% more than that of a completely flat film. This implies thathe above mentioned mechanisms only play a minor role in thenhancement of device efficiency.

Similar to Yang’s work [28], the bright regions in the heightmages are attributed as P3HT domains, while the dark regions arescribed to CdSe–PVK clusters. Surface roughness from the heightmages represents the contact area between polymer film and the

etal cathode. For the as-cast film without thermal treatment theurface is very smooth with a root mean square (rms) roughnessf 1.42 nm (Table 1). This may be due to a homogeneous and ran-om distribution of P3HT chains and CdSe–PVK molecules. Afterhermal annealing, phase segregation occurs and two-phase mor-

hology is formed. The phase segregation induces the realignmentf P3HT chains and results in a more order structure for higherptical absorption. Through a sequential thermal treatment at 100,25, 150, 175, and 200 ◦C in a period of 30 min for each sample,

and Physics 132 (2012) 131– 137

the rms roughness becomes larger than that of untreated film. Thesurface roughness increases as the annealing temperature isincreased to 175 ◦C, giving rise to a more efficient exciton disso-ciation and a greater PL quenching, but the surface becomes lessrough on further increasing the temperature to 200 ◦C comparedto the film treated at 175 ◦C.

Fig. 4 also shows the change in the texture of the film afterannealing at different temperatures. Without heat treatment,the interpenetrating networks are not well developed, and thedonor–acceptor (D–A) domains are difficult to distinguish. Afterthermal annealing for 30 min at 150 ◦C, the morphology of inter-penetrating D–A networks becomes clearer and easily visible. Theheight image of the film shows a coarser as well as more homo-geneous D–A two-phase texture compared to the other films. Thetwo-phase structure results from the phase segregation of P3HTdonor and CdSe–PVK acceptor. The phase image of the heat-treatedfilm also shows a two-phase interpenetrating network to facilitatecharge separation and transport. In the case of the films treatedat 125 and 175 ◦C, both height images and phase images exhibitgreater and large-scale phase segregation and inhomogeneous dis-tribution of D–A materials, indicating smaller interfacial areas ofD–A materials in comparison with that of the film annealed at150 ◦C. The more order structure of P3HT for the film treated at150 ◦C enhances the higher optical absorption and higher powerconversion efficiency. Although the height and phase images of thefilm annealed at 200 ◦C show a homogeneous two-phase structureas well, the length scale of each phase is larger than that of thefilm treated at 150 ◦C. Therefore, the total interfacial area is stillsmaller for charge generation between the P3HT and CdSe–PVK. Asshown in a later part, the best device performance and the highestefficiency are obtained for the devices with the active layer heat-treated at 150 ◦C. It is obvious that the CdSe–PVK molecules are ableto diffuse to form a large domain upon thermal annealing and thatthe phase separation of P3HT and CdSe–PVK would provide a morecontinuous pathway for holes and electrons. An obvious co-relationbetween the device performance and the surface morphology isthat the continuous percolation pathway and larger interfacial areawould give higher efficiency device. The annealing treatment alsoimproves the crystallinity of P3HT within the phase-separated net-works and thereby enhances the optical absorption and facilitatesthe charge transport to the electrodes.

Combining the above results, the improved nanoscale morphol-ogy from AFM examination indicates that the phase segregationand 3D interpenetrating networks of P3HT:CdSe–PVK are respon-sible for the enhancement of the device efficiency. The annealingtreatment enhances the formation of ordered structure as indi-cated by vibronic peaks in the absorption spectra. Consequently,higher absorption, 3D interpenetrating networks and the increasein the carrier mobility are the most likely reasons for the efficiencyenhancement.

Fig. 5 shows the J–V characteristics under white light illumina-tion (100 mW cm−2) for photovoltaic devices subjected to thermalannealing after spin coating the active layer but before cathodedeposition (pretreatment). For pretreatment conditions, the deviceannealed at 150 ◦C shows the best efficiency of 0.235%, followedby the one that was treated at 175 ◦C with efficiency ca. 0.158%.Interestingly, the higher performance is consistent with the higherhomogeneity of D–A two-phase morphology as shown by AFMimaging (refer to Fig. 5(d)(j) and (f)(l)). The photovoltaic param-eters of these cells, such as the short-circuit current density (Jsc),the open circuit voltage (Voc), the fill factor (FF), and the powerconversion efficiency (PCE), are summarized in Table 2. The device

without heat treatment had a rather poor PCE of 0.051%. How-ever, as the heat-treatment temperature was increased to 150 ◦C,the PV parameters were remarkably improved. As a result, thePV cell prepared with heat treatment at 150 ◦C showed the best

T.-L. Wang et al. / Materials Chemistry and Physics 132 (2012) 131– 137 135

Table 2Photovoltaic characteristics of devices under different annealing temperatures for30 min.

RT 100 ◦C 125 ◦C 150 ◦C 175 ◦C 200 ◦C

Voc (V) 0.289 0.412 0.409 0.387 0.378 0.369Jsc (mA cm−2) 0.664 0.499 1.134 1.993 1.353 0.536

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Table 3Surface roughness of blend films obtained from AFM after annealing at 150 ◦C fordifferent time periods.

Annealing time

0 min 5 min 15 min 30 min 45 min 60 min

FF (%) 26.48 26.46 30.74 30.45 30.84 26.52� (%) 0.051 0.054 0.143 0.235 0.158 0.052

evice performance. Thus, it is clear that the variation of theptoelectrical/physical properties of the active layer with anneal-ng treatment affects the device performance significantly and thathere exists optimum range of the annealing temperature.

.2. Effects of annealing time

Fig. 6 shows the effect of annealing time on the UV–vis absorp-

ion spectra for the active layer (P3HT:CdSe–PVK). The annealingemperature for all of the thin films was 150 ◦C. As seen in the fig-re, the absorption peak increases with the increase of annealing

-0.1 0. 0 0.1 0. 2 0.3 0. 4 0.5 0.6

-2.0

-1.5

-1.0

-0.5

0.0

0.5ITO/PEDOT:PSS/P3HT:CdSe-PVK/Al

RT 100OC 125OC 150OC 175OC 200OC

J(m

A/c

m2 )

V(V)

ig. 5. J–V characteristics of devices under AM 1.5 simulated solar illumination atn intensity of 100 mW cm−2 after annealing at different temperatures for 30 min.

300 400 500 600 700 800

Abs

orba

nce

Wavelength (nm)

0 mi n 5 mi n 15 min 30 min 45 min 60 min

ig. 6. UV–vis absorption spectra of P3HT:CdSe–PVK blend films after annealing at50 ◦C for different time periods.

Average roughness (nm) 1.02 1.11 5.04 4.26 4.33 2.96Root mean square (nm) 1.42 1.40 6.52 5.78 5.42 3.86

time and the maximum absorption is observed at 30 min. Fur-ther increasing the annealing time, absorption peaks graduallydecrease and an obvious deterioration of absorption is observedat 60 min. Moreover, the last shoulder corresponding to the inter-chain interactions increases to a maximum intensity at annealingof 30 min. The increased inter-chain interaction among the P3HTchains results in more ordered structure via thermal annealingas mentioned above. As we will discuss later, the higher opticalabsorption seems a guarantee for higher power conversion effi-ciency, as is evident from the film annealed for 30 min.

Fig. 7 provides the PL quenching effect for the blend filmtreated at 150 ◦C for different time periods. Compared to the unan-nealed film, the blend film annealed for 5 min shows an intensepeak in the PL spectrum, resulting from the reduction of interfa-cial area between P3HT and CdSe–PVK. This is attributed to theenhancement of phase separation between the donor and accep-tor molecules upon annealing as indicated above. In contrast, alonger annealing time has a larger quench effect and a maximumPL quenching is reached at annealing time of 30 min. However,further increasing the annealing time to 45 and 60 min, the pho-toluminescence peaks increase and an obvious deterioration effectis observed. As above-mentioned, the phase separation reducesexciton dissociation, but the charge carrier transport is improvedin both donor and acceptor phases after thermal annealing. Withregard to the film annealed for 30 min, the later effect plus thehigher optical absorption offsets the former effect and results inan overall improvement in device performance. The maximum PLquenching at annealing of 30 min is in agreement with the maxi-mum device performance as shown in Table 3.

The nanoscale morphology of two-phase segregation andinterpenetrating networks of P3HT and CdSe–PVK after thermal

annealing with different time periods are shown in Fig. 8. From thedata of rms roughness (Table 3), the surface roughness increasesas the annealing time is increased to 15 min, but further increasing

Fig. 7. Photoluminescence spectra of P3HT:CdSe–PVK blend films after annealingat 150 ◦C for different time periods with excitation at 525 nm.

136 T.-L. Wang et al. / Materials Chemistry and Physics 132 (2012) 131– 137

Fig. 8. AFM topography images (5 �m × 5 �m) of P3HT:CdSe–PVK blend films afterannealing at 150 ◦C for different time periods. 2D height image for the blend film (a)annealed for 5 min, (b) annealed for 15 min, (c) annealed for 30 min, (d) annealedfor 45 min, and (e) annealed for 60 min. Phase image for the blend film (f) annealedfor 5 min, (g) annealed for 15 min, (h) annealed for 30 min, (i) annealed for 45 min,and (j) annealed for 60 min.

0.0 0.2 0. 4 0.6

-2.0

-1.5

-1.0

-0.5

0.0

0.5ITO/PEDOT:PSS/P3HT:CdSe-PVK/Al

J(m

A/c

m2 )

V(V)

0 min5 min15 min30 min45 min60 min

Fig. 9. J–V characteristics of devices under AM 1.5 simulated solar illumination atan intensity of 100 mW cm−2 after annealing at 150 ◦C for different time periods.

Table 4Photovoltaic characteristics of devices at 150 ◦C for different annealing times.

0 min 5 min 15 min 30 min 45 min 60 min

Voc (V) 0.289 0.414 0.441 0.387 0.414 0.483Jsc (mA cm−2) 0.664 1.043 1.226 1.993 1.495 1.342

FF (%) 26.48 29.97 29.98 30.45 30.07 31.27� (%) 0.051 0.129 0.162 0.235 0.186 0.203

the time to 30 min, the surface becomes less rough compared to thefilm treated for 15 min.

As stated above, the higher roughness of the film will givehigher efficiency device. In fact, the roughness only plays a minorrole in the device efficiency enhancement. On the other hand, theheight and phase images of the film annealed for 30 min exhibitmore distinguishable two-phase texture compared to those of thefilm treated for 15 min. Since the well-developed interpenetrat-ing networks facilitate charge separation and transport, the powerconversion efficiency is thus higher for the blend film annealed for30 min. Furthermore, although the film annealed for 60 min has ahomogeneous D–A two-phase structure as well, the rms roughnessis lower and the length scale of each phase seems higher as com-pared to those of the film treated for 30 min. Consequently, the filmtreated for 30 min has the highest PCE.

The J–V characteristics under illumination (100 mW cm−2) forphotovoltaic devices subjected to different annealing times at150 ◦C are shown in Fig. 9. For pretreatment conditions, the deviceannealed for 30 min shows the best efficiency of 0.235%, followedby the one that are treated for 60 min with efficiency ca. 0.203%.The performance is consistent with the optical absorption, the PLquenching, as well as the degree of homogeneity of D–A two-phasemorphology as shown by AFM images. The photovoltaic parametersof these cells are summarized in Table 4, revealing that the anneal-ing time significantly affects the device performance and that thereexists some optimum range of the annealing time.

4. Conclusions

The CdSe–PVK nanocomposite has been synthesizedand employed as the electron acceptor in the active layerof BHJ-type polymer solar cells. The BHJ PV devices have

been assembled by using a typical sandwich structure ofITO/PEDOT:PSS/P3HT:CdSe–PVK/Al. As a result of annealingtreatment at an optimum condition (150 ◦C/30 min), the PV cell

istry

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T.-L. Wang et al. / Materials Chem

erformance was dramatically improved and the power conver-ion efficiency of device reached to 0.235% under white lightllumination (100 mW cm−2). The thermal annealing contributedo the enhanced PV cell performance by optimizing both theonor/acceptor morphology in the BHJ active layer. In conclusion,nnealing treatment on the PV devices enhanced 3D interpenetrat-ng networks in the active layer, light absorption, and the carrier

obility, leading to the improvement of the device performance.

cknowledgements

We gratefully acknowledge the support of the National Scienceouncil of Republic of China with Grant NSC 97-2221-E-390-005-Y2.

eferences

[1] O. Inganäs, M. Svensson, F. Zhang, A. Gadisa, N.K. Persson, X. Wang, M.R. Ander-sson, Appl. Phys. A 79 (2004) 31.

[2] F. Zhang, W. Mammo, L.M. Andersson, S. Admassie, M.R. Andersson, O. Inganäs,Adv. Mater. 18 (2006) 2169.

[3] A. Gadisa, W. Mammo, L.M. Andersson, S. Admassie, F. Zhang, M.R. Andersson,O. Inganäs, Adv. Funct. Mater. 17 (2007) 3836.

[4] D.J. Milliron, I. Gur, A.P. Alivisatos, MRS Bull. 30 (2005) 41.[5] N.C. Greenham, X. Peng, A.P. Alivisatos, Phys. Rev. B 54 (1996) 17628.[6] W.J.E. Beek, M.M. Wienk, R.A.J. Janssen, Adv. Mater. 16 (2004) 1009.

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and Physics 132 (2012) 131– 137 137

[7] D. Qi, M. Fischbein, M. Drndic, S. Selmic, Appl. Phys. Lett. 86 (2005) 093103.[8] J. Boucle, S. Chyla, M.S.P. Shaffer, J.R. Durrant, D.D.C. Bradley, J. Nelson, Adv.

Funct. Mater. 18 (2008) 622.[9] S.D. Oosterhout, M.M. Wienk, S.S. van Bravel, R. Thiedmann, L.J.A. Koster, J. Gilot,

J. Loos, V. Schmidt, R.A.J. Janssen, Nat. Mater. 8 (2009) 818.10] A.J. Nozik, Physica E 14 (2002) 115.11] A.P. Alivisatos, Science 271 (1996) 933.12] B.L. Cushing, V.L. Kolesnichenko, C.J. O’Connor, Chem. Rev. 104 (2004) 3893.13] Y. Yin, A.P. Alivisatos, Nature 437 (2005) 664.14] W.U. Huynh, X. Peng, A.P. Alivisatos, Adv. Mater. 11 (1999) 923.15] J. Liu, T. Tanaka, K. Sivula, A.P. Alivisatos, J.M.J. Fréchet, J. Am. Chem. Soc. 126

(2004) 6550.16] W.U. Huynh, J.J. Dittmer, A.P. Alivisatos, Science 295 (2002) 2425.17] B. Sun, E. Marx, N.C. Greenham, Nano. Lett. 3 (2003) 961.18] M.D. Heinemann, K. von Maydell, F. Zutz, J. Kolny-Olesiak, H. Borchert, I. Riedel,

J. Parisi, Adv. Funct. Mater. 19 (2009) 3788.19] T.L. Wang, C.H. Yang, Y.T. Shieh, A.C. Yeh, Macromol. Rapid Commun. 30 (2009)

1679.20] T.L. Wang, C.H. Yang, Y.T. Shieh, A.C. Yeh, Euro. Polym. J. 46 (2010) 634.21] F. Padinger, R.S. Ritterberger, N.S. Sariciftci, Adv. Funct. Mater. 13 (2003) 85.22] V. Dyakonov, Appl. Phys. A: Mater. Sci. Process. 79 (2004) 21.23] X. Jiang, R. Osterbacka, O. Korovyanko, C.P. An, B. Horowitz, R.A.J. Janssen, Z.V.

Vardeny, Adv. Funct. Mater. 12 (2002) 587.24] P.J. Brown, D.S. Thomas, A. Kohler, J.S. Wilson, J.S. Kim, C.M. Ramsdale, H. Sir-

ringhaus, R.H. Friend, Phys. Rev. B 67 (2003) 064203.25] G. Li, V. Shrotriya, Y. Yao, Y. Yang, J. Appl. Phys. 98 (2005) 043704.

1636.27] P. Peumans, S.R. Forrest, Appl. Phys. Lett. 79 (2001) 126.28] X.N. Yang, J. Loos, S.C. Veenstra, W.J.H. Verhees, M.M. Wienk, J.M. Kroon, M.A.J.

Michels, R.A.J. Janssen, Nano Lett. 5 (2005) 579.