Interaction between P3HT and CdS surface - Annealing effect on Photovoltaic Performance for

Chapter 7 Annealing effect on Photovoltaic Performance for

7.5 Interaction between P3HT and CdS surface

Figure 7.5 shows the surface morphology change as monitored by AFM. The morphology can be further investigated for the origin of the long-term stability of P3HT/CdS devices with different annealing temperatures. For the as-deposited

(a) (b)

(unannealed) film with AR=4 and 16, the surface is very smooth: the root mean square (rms) roughnesses are 1.067 nm and 1.184 nm, respectively. However, after thermal treatment at 160 °C, the rms roughnesses obtained from Figure 6(c) and 6(d) for AR=4 and 16 become 7.259 nm and 4.821 nm, respectively, because the CdS nanocrystals aggregate during thermal annealing. Moreover, it can be found that the scale of the aggregations for AR=4 is larger than that for AR=16, indicating that the CdS nanocrystals of AR=4 aggregate more easily than those of AR=16, which is in good agreement with the results of Figure 3. Compared with the film with AR=4, the film with AR=16 reveals rod-like texture due to a higher AR and smaller aggregations which consist of fewer CdS nanocrystals. As mentioned above, the best device performance and the highest PCE are obtained when the devices are annealed at 160 °C. Therefore, CdS nanocrystals aggregate excessively comparable to the exciting diffusion length, which leads to the reduced charge segregation and device efficiencies [199, 200].

Figure AFM height images of P3HT/CdS films with (a) AR=4 and (b) AR=16. (c) AR=4 and (d) AR=16 annealed at 160 °C for 60 min.

To understand the interaction and molecular structure between P3HT and CdS nanocrystals during in-situ growth of P3HT/CdS, ¹H NMR measurements were performed on the pristine P3HT and P3HT/CdS nanocrystals composites with an AR of 4 and 16. As shown in Figure 7(b) and 7(c), as compared to pure P3HT in Figure 7(a), it was found that the broader proton peaks at a chemical shift of ~6.98 (thiophene ring (a)) and ~2.79 ppm (hexyl chain (b)) were clearly observed for P3HT/CdS in Figure 7(b) and 6(c), which confirms that some interaction occurs between the polymer and the CdS. Additionally, the proton peaks from the hexyl chain (e) (~0.91 ppm) was only slightly broadened compared with the proton peaks

(a) (b)

(c) (d)

from the thiophene ring (a) and hexyl chain (b). A similar phenomenon also has been reported for polymer-CNT composites; the interactions of polymers with CNTs cause broadening and reduced intensity of ¹H NMR peaks [201, 202]. The closer the protons are to the surface of CNTs, the broader and weaker the peak will be. In addition, the degree of broadening in the peak corresponding to the thiophene ring is due to the corresponding protons coming very close to the nanoparticles [203]. This indicates that the thiophene ring is much closer to the CdS surface than the hexyl chain. This conclusion was further supported by the change of the relative intensity of proton peaks (a), (b), and (e). In the pristine P3HT solution, the ratio of protons (a), (b), and (e) is almost 1:2:3, whereas in P3HT/CdS, this ratio is ~0.94:1.89:3.0 and 0.85:1.73:3.0 for AR=4 and 16, respectively. As a result, although the polymer backbone interacts with the CdS surface, the hexyl chain is relatively free in solution.

This is shown schematically in Figure 8(a), where the thiophene ring is in close spatial proximity to the CdS surface. Furthermore, a stronger adsorption between CdS nanoparticles and P3HT takes place at AR=16, indicating that the P3HT chain prefers to adsorb to a longer flat surface due to the planar P3HT conformation. This explains the observations made earlier on the basis of absorption and DSC spectra.

Figure 7.6 ¹H NMR spectra for (a) P3HT and P3HT/CdS composites with the AR of (b) 4 and (c) 16.

In previous reports [204], it has demonstrated that only a relatively small number of polymer segments within a chain are directly bound to the surface because of conformational limitations introduced by the particles, in addition to other restrictions on chain conformation as illustrate in Figures 7.7(b) and 7.7(c). For P3HT/CdS with AR=16, the polymer chain segments display a stronger interaction with the surface of the CdS, resulting in a denser layer and short loops close to the surface [205], as shown in Figure 7.7(b). On the contrary, the loops extend farther into the polymer matrix and form a region of lower density of polymer when AR is 4, as shown in Figure 7.7(c) [205]. It is noted that denser polymer chains can shield the surface from

(a)

(b)

a

b c

d e

(c)

other chains, likely resulting in a denser layer formed from fewer chains with fewer entanglements with other chains [204, 205]. Therefore, the strength of the interaction of a polymer molecule with the surface of the nanocrystals controls both the polymer molecular conformation at the surface and the entanglement distribution in a larger region surrounding the nanocrystal. Hence, a higher degree of entanglement will result in a larger number of polymer chains that are associated with a given nanocrystal, of which only a fraction is actually anchored to the surface. As a result, the strength of the interaction directly affects the aggregation rate of the CdS nanocrystals in the P3HT matrix during the annealing treatments, which causes the change in the morphology and the PCE (Figure 7.4 and 7.5).

Figure 7.7 (a) Schematic illustration of the molecular structure near the interface between CdS and P3HT, indicating the adsorption characteristics of polymer chains onto the surface of the nanoparticles. (b) A strong adsorption polymer adheres to the surface and most of the segments reside on the surface and (c) a weak adsorption polymer adheres to the surface and most of the segments reside in loops.

(a)

(b) (c)

CdS

7.5 Summary

We prepared P3HT/CdS nanocrystal nanocomposites with different ARs by the in-situ method. Upon annealing Photoluminescence measurements confirmed the formation of polymer crystallites and also showed that the photo-induced electron transfer becomes effective as the AR increases. The ¹H NMR results demonstrated that the thiophene ring is much closer to the CdS surface than the hexyl chain, and that the interaction strength between P3HT and CdS increases with the AR of the CdS nanocrystal. As a result, a lower AR drives the extensive aggregation of CdS in the polymer matrix upon annealing treatment, as evidenced by AFM. The annealing-condition-dependent PCE study revealed that this interaction shows not only a dramatic effect on the aggregation of nanocrystals during the annealing process but also device performance upon annealing. Therefore, a higher PCE (or Photovoltaic performance) can be obtained for the in-situ-growth P3HT/CdS with AR=16 upon annealing treatment at 160 °C for 60 min.

Chapter 8 Improvement in Photovoltaic Performance for Hybrid P3HT/elongated CdS nanocrystals solar cells with F-doped

SnO

Arrays

8.1 Introduction

Organic photovoltaic devices are drawing attention because of their potential for the production of flexible and large-area solar cells at very low cost. Furthermore, polymer-inorganic hybrid solar cells are of particular interest because they combine the solution process ability of polymers with the high electron mobility of inorganic semiconductors [102]. Recently, hole-conducting polymers have been combined with a wide range of inorganic nanomaterials, including CdSe quantum dots, rods, tetrapods, hyper-branched colloids [50, 85, 87, 171] PbS, PbSe, CuInS2 and CuInSe2

nanoparticles [6,14, 53, 153, 206]. In general, the optimal device thickness of a bulk heterojunction solar cell is typically 100-200 nm, depending on the combination of materials. The optimal thickness is determined by the equilibrium between absorption of the films and charge carrier transport in the device. The amount of exciton generation increases with thickness because of a greater total absorption of light.

However, the extent of recombination may also increase for very thick films due to the increased distance required for the charges to reach the electrodes, resulting in a decrease of overall power conversion efficiency. Therefore, a continuous ordered inorganic nanostructure that may help charge carrier collection and transport can potentially address this issue [107, 138, 156, 207]. For example, nanorod arrays not

only solve the above-mentioned problem, but also provide far greater surface area (over 100 times more) than thin films, which thus offer great advantages in applications where surface area plays a critical role in improving charge collection. As a result, it might be possible to achieve both large absorption and an efficient charge transport with thicker film devices. Therefore, the possibility of utilizing one-dimensional nanorods in both dye-sensitized solar cells and organic solar cells may help to boost device performance. This concept has already been attempted with ZnO and TiO2 nanorods and nanotubes [8, 119, 140, 208, 209]. For instance, Olson et al. reported a P3HT/ZnO nanorod cell exhibiting a short-circuit current density (JSC) of 2.2 mA/cm², an open-circuit voltage (VOC) of 0.44 V, a fill factor (FF) of 0.56 and an efficiency of 0.53% using nanorods grown in basic conditions [8]. They also reported that the introduction of phenyl C61-butyric acid methyl ester (PCBM) into the hybrid devices could significantly improve the efficiency of the devices by up to 2.03%. Nelson and Peiro´ et al. also reported hybrid photovoltaic devices using polymer/ZnO nanorod combinations [119, 140]. These reports revealed the importance of using nanorod array electrodes to improve the photovoltaic performance. However, up to now, very little work has been published regarding the fabrication of transparent conductive nanorods in solar cells [210, 211]. It is expected that increasing the interface area between the electrode and metal oxide will improve device performance by increasing charge capture. At present, indium tin oxide (ITO) has been the most commonly used plane anode in thin-film solar cells. Recently, Wang et al. produced ITO nanowires by electrophoretic deposition and demonstrated that nanowire electrodes could enhance the device performance of dye-sensitized solar cells [211]. We believe in the possibility of creating a nanostructured electrode to improve hybrid cell performance, especially if it can be achieved via simple and low-cost processes for the fabrication of oxide nanostructures.

Tin oxide, SnO2 is an insulator and an important, colorless, low-cost, large-bandgap (n-type) semiconductor material. When doped with Sb or F ions, F-doped SnO2 (FTO) is an ideal candidate for applications requiring a transparent conductive oxide due to its ability to adhere strongly to glass, resistance to physical abrasion, chemical stability, high optical visible transparency and electrical conductivity, such that FTO is widely used as a transparent conducting oxide substrate [212] as well as electrode material for energy conversion [213]. Recently, FTO nanorod electrodes developed by template-filling methods have been also demonstrated by Russo et al [214]. However, this method needs complex processes to remove the anodic alumina template. Therefore, in this study, a spray pyrolysis deposition process (SPD) was developed to obtain a good distribution of FTO film on ZnO nanorods grown using the aqueous solution method. The method not only makes it possible to coat ZnO nanorods with a thin, nominally uniform layer of FTO, but also controls the FTO thickness. In this study, it has been demonstrated that the efficiency of the P3HT/CdS/ZnO nanorod solar cell can be improved by nearly ~2 fold by coating the nanorod arrays in an FTOlayer using SPD. Both ZnO nanorod length and FTO layer thickness had a crucial effect on the cell output characteristics.

Moreover, we investigated effect of elongated CdS nanocrystals on the performance of FTO-coated ZnO nanorod polymer solar cells with a thick hybrid layer of CdS/P3HT.

8.2 Microstructure of FTO-coated ZnO nanorods

Figure 8.1(a) shows an FE-SEM image of the as-synthesized ZnO nanorods. The ZnO nanorods had a diameter that ranged from 80 to 90 nm and lengths of up to 380 nm depending on the growth conditions. The length of the ZnO nanorod increased with

the growth time, where the average length of the nanorods was estimated from the cross-sectional FE-SEM images (Figure 8.1(b)). The average length increased from ca.

140 to 380 nm by prolonging the growth reaction time from 90 to 180 min. Figure 8.1(c) show the image of ZnO nanorods coated with a 10 nm FTO layer by the spray pyrolysis deposition process (SPD). After the SPD process, the nanorods, as shown in the inset of Figure 8.1(c), exhibited increased diameters and rough surfaces as compared to those in Figure 8.1(a). This result indicates that FTO was successfully deposited on the ZnO nanorods. Figure 8.1(d) shows the X-ray diffraction (XRD) patterns of the FTO-coated ZnO nanorods and template ZnO nanorods. The distinct peaks corresponding to ZnO and SnO2 were also observed, and the major diffraction peaks of the coating specimen were consistent with the known tetragonal SnO2

structure with lattice constants given in the literature of a = 4.755 Å and c = 3.199 Å (JCPDS 41- 1445). Furthermore, it was noted that the peak of tetragonal SnO2 (101) almost overlapped with the main peaks of ZnO (002). It was concluded that the coated FTO layers retained their perfect crystalline phase and physical structure, which further confirmed that an FTO layer rather than an alloy was formed. The broad peaks of SnO2 in the XRD spectrum indicated that the nanoparticles were small. The average size of the nanoparticles as calculated by the Scherrer equation was about 9 nm. In addition, no obvious peaks corresponding to SnCl2, Sn or other tin oxides were observed in the pattern.

Figure 8.2(a) shows a typical low-resolution TEM image of the FTO-coated ZnO nanorods. It was observed that almost all ZnO nanorods had been fully coated with thin and uniform layers. The thickness of the coating was estimated to be about 10 nm.

The corresponding high-magnification TEM image in Figure 8.2(b) clearly shows that the FTO nanoparticles were crystalline and well-distributed on the nanorods. The inset of Figure 8.2(b) is the corresponding fast Fourier transform (FFT) diffraction pattern,

which can be indexed to tetragonal SnO2. The two clear polycrystalline rings correspond to crystal faces of (110) and (101) of tetragonal SnO2.

In order to optimize the FTO-coated ZnO nanorods for organic solar cells, ZnO nanorod arrays were grown for different periods of time by controlling the growth time. As demonstrated in Figure 8.1(b), the average length increased from 140 to 380 nm and the diameter of the nanorods was in the range of 80-90 nm for all cases, and almost independent of the reaction time. Figure 8.3(a) shows the TEM image of CdS/P3HT samples prepared at cadmium acetate concentrations of 0.83 mg/mL at 120 ^oC with a DCB-to-DMSO ratio of 8:4 that demonstrated the size of CdS nanocrystals to be 2-3 nm. The structural morphologies of the CdS nanorods synthesized from the DCB-to-DMSO volume ratios of 7:5, 8:4 and 9:3 at a lower cadmium acetate concentration of 8.3 mg/mL for 30 min are shown in Figure 8.3(b), 3(c), and 3(d), respectively. It was found that a variation of the DCB-to-DMSO volume ratio from 7:5 to 8:4 to 9:3 at a lower cadmium acetate concentrations of 8.3 mg/mL gave CdSe nanocrystals with aspect ratios increasing from ~4 to ~8 to ~16 for the samples with the same concentration under similar reaction conditions, which is corresponding to the length of CdS about 9 nm, 18 nm, and 39 nm, respectively. The role of the structural direction of the P3HT template was demonstrated in our previous work [193]. An elongated P3HT chain can be developed by increasing the DCB-to-DMSO ratio, resulting in CdS nanocrystals with a higher aspect ratio.

Figure 8.1 FE-SEM images of (a) ZnO nanorods micrograph, (b) cross-sectional ZnO nanorods with various lengths, (c) ZnO nanorods coated with a 10 nm FTO layer. (d) XRD patterns of ZnO nanorods and ZnO nanorods coated with a 10 nm FTO layer.

The scale bar in Figure 8.1(b): 300 nm.

Figure 8.2 (a) TEM image and (b) HRTEM micrograph of ZnO nanorods coated with a 10 nm FTO layer. The inset is the corresponding FFT diffraction pattern.

100 nm

Figure 8.3 (a) TEM image of CdS nanocrystals/P3HT composites synthesized at 120 ^oC. TEM image of CdS nanocrystals synthesized in P3HT with a volume ratio of DCB-to-DMSO of (b) 7:3, (c) 8:4 and (d) 9:3 at 180 ^oC.

8.3 Effect of ZnO nanorod length and FTO thickness

To investigate the dependence of the device performance on the length of the FTO-coated ZnO nanorods, power conversion efficiency (PCE), fill factor (FF), short circuit current density (JSC) and open-circuit voltage (VOC), a series of devices with 25 nm thickness FTO were prepared where the P3HT/CdSe had an aspect ratio of 1.

Figures 8.4(a) and 8.4(b) illustrate the device performance as a function of the length of the FTO-coated ZnO nanorods. We observed that the device without nanorods showed a JSC of 3.8 mA cm^-2, a VOC of 0.56 V and a FF of 34%, resulting in a PCE of

(c) (d)

10 nm

(a) (b)

10 nm 10 nm

0.70%. With the array with 320 nm FTO-coated ZnO nanorods, the PCE was improved to 1.8% under the same conditions as a result of the increased values for JSC

of 6.8 mA cm^-2, VOC of 0.62 V and FF of 44%. The JSC increased strongly with the lengths of the nanorods. In addition, the FF was improved from 34% to 44% by introducing the FTO-coated ZnO arrays, although the improvement in VOC was small.

The improvement in both JSC and FF mainly contributed to the increase of PCE from 0.70% to 1.8% as a result of the increase in the FTO-coated ZnO nanorod length. In our devices, the absorption layer was very thick (up to 450 nm), which improved the light absorption. However, when the nanorod was short, the probability of exciton recombination was higher with a thicker active layer due to the small carrier mobility.

As the nanorod length increased, the average distance from the generation point of the charge carriers to the FTO electrode decreased, resulting in an increase in the number of charge carriers collected by the FTO-coated ZnO nanorods, which leads to an improvement of FF and JSC. However, an increase in the nanorod length from 320 to 380 nm resulted in a decreased JSC, which may be due to the annealing treatment being insufficient to drive P3HT to completely fill the volume between the longer nanorods [215]. Therefore, the ability to improve performance by efficient electron collection using longer nanorods may eventually be limited by issues with P3HT infiltration.

To further clarify the role of the CdS nanocrystals, a system without CdS in the P3HT layer was compared. As shown the inset of Figure 8.4(a), it was found that the incorporation of CdS into P3HT resulted in a large increase in JSC as compared with the device without CdS. This result demonstrates that the interface between the P3HT and FTO-coated ZnO nanorods makes a small contribution to the charge separation, and the charge separation mainly occurs at the P3HT/CdS interface. The reason for the more efficient charge separation at the P3HT/CdS interface could be attributed to

100 150 200 250 300 350 400 0.6

much larger interface area and large energy difference between P3HT and FTO. From the results above, we can conclude that these excitons are mainly dissociated to electrons and holes at the interface between P3HT and CdS, and that the FTO-coated ZnO nanorod array structure permits vertical charge transport from the active layer.

Therefore, it is believed that the role of FTO-coated ZnO nanorod arrays in this study can be considered as an electron collector that collects the electrons more efficiently than a planar cathode structure.

Figure 8.4 Hybrid solar cell characteristics. (a) VOC and JSC (c) fill factor and PCE plotted as a function of nanorod length. The inset is JSC plotted as a function of nanorod length (without CdS nanocrystals).

(a)

(b)

After charge separation at the CdS/P3HT interface, the electrons generated at the conduction band of CdS (-4.0 eV) [216] can be transferred to the conduction band of SnO2 (-4.4 eV) [216], which has a high electron mobility that helps charge-carrier collection and transport. Moreover, the conduction band of ZnO is slightly higher in energy than that of SnO2, so the FTO layer dominates electron transport. Furthermore, because the SnO2 has high electron mobility compared to most organic semiconducting materials, the electrons can be quickly collected by the FTO layer.

Consequently, the FTO-coated ZnO nanorods work as electron collectors by shortening the average electron diffusion distance in the CdS network of the bulk heterojunction, resulting in a reduction in charge recombination.

To better understand the function of the FTO layer, we tested several devices with different thicknesses of the FTO layer. A nanorod length of 320 nm was used in all devices. Because the FTO layer dominates electron transport, the effect of the FTO

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