Highly Symmetrical CdS Tetrahedral Nanocrystals Prepared by Low

Highly Symmetrical CdS Tetrahedral Nanocrystals Prepared by Low-Temperature Chemical Vapor Deposition Using Polysulfide as the Sulfur Source Jie Zheng,† Xubo Song,† Nan Chen,† and Xingguo Li*,†,‡

CRYSTAL GROWTH & DESIGN 2008 VOL. 8, NO. 5 1760–1765

Beijing National Laboratory for Molecular Sciences (BNLMS) (The State Key Laboratory of Rare Earth Materials Chemistry and Applications), College of Chemistry and Molecular Engineering, Peking UniVersity, Beijing, P. R. China, 100871, and College of Engineering, Peking UniVersity, Beijing, P. R. China, 100871 ReceiVed August 24, 2007; ReVised Manuscript ReceiVed NoVember 28, 2007

ABSTRACT: In this paper, we report a new scheme for chemical vapor deposition (CVD) of CdS nanocrystals, in which polysulfide and cadmium chloride were used as the sulfur and cadmium source, respectively. During the reaction, hydrogen sulfide was generated by controlled hydrolysis of polysulfide and reacted with cadmium chloride to give CdS. Tetrahedral CdS nanocrystals with homogeneous size and nearly perfect Td point group symmetry were obtained in high yield. The source and deposition temperature were only 540 and 400 °C, respectively, which were significantly lower than those in conventional CVD methods. The low deposition temperature was found critical for the formation of the tetrahedral nanostructures. The photoluminescence spectra of the tetrahedral nanocrystals consisted of a sharp band edge emission peak and a broad band related to trapped states. Introduction Cadmium sulfide (CdS) nanocrystals are extremely attractive for their excellent luminescence properties and have found broad applications in optoelectronic nanodevices,1–3 biological labeling,4–7 and photochemistry.8,9 Meanwhile, CdS is of structural interest for the abundant morphologies and manifold crystallographic phenomena in its nanostructures.10–16 Therefore, the controlled synthesis of CdS nanocrystals has attracted intense research attention, as indicated by the extensive studies on the preparation methods in both solutions16–22 and vapor phase.23–27 Vapor phase transportation and deposition is a popular approach to prepare CdS nanocrystals and is of particular importance for CdS-based optoelectronic nanodevices. So far, most CdS nanomaterials for this purpose were fabricated by the vapor phase transportation and deposition approaches because high-quality nanocrystals with clean surfaces can be obtained.1–3 Despite its success, this method is unsatisfactory in several aspects. At present, most vapor phase transportation and deposition methods for CdS preparation were either direct thermal evaporation of raw CdS powder or chemical vapor deposition by decomposition of metallorganic precursors. Therefore, either high operation temperature (typically over 800 °C for thermal evaporation) or metalloragnic precursors that were difficult to access and manipulate were required. In addition, the morphologies of the thermal evaporation and CVD products are quite limited, mostly nanowires and nanobelts. These drawbacks primarily originated from the limited precursors and chemical routes in thermal evaporation and CVD methods. Solution chemistry methods, on the other hand, are capable of forming products with abundant morphologies and are usually operating at mild conditions due to the large freedom in choosing precursors and reaction routes.11–19,22,28 Therefore, we believe that improvement of these aspects may open opportunities for new CVD methods with mild operation conditions and better controllability of product morphologies. * Corresponding author. Phone: 86-10-6276-5930. Fax: 86-10-6276-5930. E-mail: [email protected]. † BNLMS. ‡ College of Engineering.

Ge et al. have demonstrated that transition metal sulfide nanostructures could be prepared by CVD methods using the corresponding chlorides and elemental sulfur as the metal and sulfur source, respectively.24,25,29 Inspired by their work, we designed a scheme for CVD preparation of CdS nanocrystals using the reaction between CdCl2 and H2S. From a chemical point of view, H2S is a more effective sulfur source compared to elemental sulfur. However, its application was limited due to its high toxicity and difficult manipulation. Therefore, we used polysulfide as an alternative sulfur source to release H2S by controlled hydrolysis, so that the reactivity of H2S could be utilized without using a H2S tank. In this method, tetrahedral CdS nanocrystals around 400 nm in size with nearly perfect Td point group symmetry were obtained at a low operation temperature of only 540 °C. Their structures, annealing behaviors, and photoluminescence properties were characterized. At present, the tetrahedral CdS nanocrystals were almost exclusively prepared by solution chemistry methods,14,16,22 with only very few reports on the fabrication of this structure using CVD.30,31 The results of our work demonstrated the effectiveness of proper precursors and chemical routes in reducing the reaction temperature and broadening the product morphologies in CVD methods. Experimental Section Dehydration of Chemicals. Anhydrous cadmium chloride (CdCl2) was prepared by dehydration of 6 g of CdCl2 · 2.5H2O (Beijing Chemical Factory, 99%) at 350 °C for 2 h in a N2 atmosphere. Anhydrous sodium thiosulfate (Na2S2O3) was prepared by dehydration of 30 g of Na2S2O3 · 5H2O (Beijing Chemical Factory, 99%) at 85 °C for 4 h under vacuum. Although a certain degree of hydrolysis took place during the dehydration process, the byproducts, which were primarily cadmium oxide (CdO) and sodium sulfate (Na2SO4), were inert in the CVD process. Therefore, the dehydrated chemicals were used without further treatment. Preparation of the Polysulfide Precursor. The polysulfide precursor was prepared by heating anhydrous sodium thiosulfate at 400 °C for 1 h in a N2 environment. After heating, the loosely packed white powder turned into a concrete orange solid. When exposed to moisture, a H2S odor was immediately released and the orange color gradually faded. XRD study suggested that the orange solid was composed of sodium sulfate and sodium polysulfide, primarily Na2S5.32 Since Na2SO4 was

10.1021/cg700804t CCC: $40.75  2008 American Chemical Society Published on Web 04/10/2008

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Figure 1. Schematic illustration of the experimental setup. inert to hydrolysis, the orange mixture was used as the sulfur source without further treatment. For a typical run, 1.3 g of anhydrous sodium thiosulfate was used. After the reaction, the orange color faded to light yellow. Preparation of CdS Nanocrystals. The experimental setup consisted of a conventional horizontal CVD system and a water bubble system upstream, as schematically illustrated in Figure 1. A 0.2 g sample of CdCl2 was loaded in the furnace center. The polysulfide precursor was placed at the entrance of the furnace, where the temperature was 150 °C. Si (100) wafers (0.5 by 0.5 cm2) free of any catalyst were placed 7–12 cm downstream from the CdCl2 as collecting substrates. The whole system was first purged with 200 sccm (sccm: standard cubic centimeters per minute) pure nitrogen for 30 min to remove oxygen. The furnace was then heated to 540 °C under 200 sccm N2 at 20 °C min-1. When the preset temperature was reached, 30 mL of distilled water was introduced through the filler to initiate the reaction. The water was warmed using a 70 °C water bath. The reaction lasted 20 min. After that, the water bath was removed and the system was allowed to cool to room temperature. The N2 flux was kept at 200 sccm through the whole process. A thick layer of light yellow powder was collected on the Si substrates. The product was characterized by X-ray diffraction (XRD, Rigaku D-max 200, Cu Ka), scanning electron microscopy (SEM, Hitachi S4800), and transmission electron microscopy (TEM, JEOL 200CX, and Tecani F30). Photoluminescence spectroscopy (PL) was measured using a 325 nm He-Cd laser as the excitation source at room temperature.

Figure 2. (a-c) SEM images with different magnifications of CdS tetrahedral nanocrystals. (d-f) SEM images of a single nanocrystal viewed from different angles: (d) along the growth direction of one arm, (e) from the junctions of two arms (black arrows indicate how two arms are connected by sharing their basal plane edges, arrow 1: a shared full-length edge; arrows 2: shared half-length edges; see text for details), (f) from the junctions of three arms.

Results and Discussion Figure 2 shows SEM images of the obtained CdS tetrahedral nanocrystals. The product was collected 10 to 12 cm downstream from the CdCl2, where the temperature was around 400 to 420 °C. The products were tetrahedral nanocrystals with a uniform size around 380 nm. The nanocrystals connected to their neighbors through their four arms (Figure 2b,c). Their uniform size and shape and random orientation enabled us to collect images along different viewing angels of a single tetrahedron under SEM (Figure 2d-f). Each arm of the tetrahedron was an identical hexagonal pyramid about 200 nm in length. Viewing along an arbitrary arm, the nanocrystal exhibited 3-fold symmetry (Figure 2d,f). Two-fold and reflection symmetry could be identified at the junction of two arms (Figure 2e). In the above discussion of symmetry, the arms were treated as hexagonal pyramids rather than cylinders that are rotational isotropic. Therefore, the obtained tetrahedral nanocrystals had nearly perfect Td point group symmetry even taking each facet into account. The four hexagonal pyramids were interconnected in the following way. Each pair of arms shared one-third of the periphery of their hexagonal basal plane, including a full edge (indicated by the black arrow 1 in Figure 2e) and two halves of its two neighboring edges (Figure 2e, arrows 2). The shared full-length edge was located at the position where one C2 axis passed through the tetrahedron. Three of the shared half-length

Figure 3. XRD patterns of the as-prepared and annealed CdS nanocrystals. The annealing was carried out in an Ar atmosphere at 500 °C for 2 h.

edges were symmetrically distributed around the position where one C3 axis passed through the tetrahedron. Figure 3 shows the XRD pattern of the as-obtained products. The positions of the diffraction peaks were consistent with those of the wurtzite (WZ) CdS (JCPDS 41-1409). Notably, the (0002) and (0004) peaks were extraordinarily intense compared to the standard pattern, suggesting the preferential orientation of the [0001] direction. This result was in agreement with the hexagonal pyramid-shaped arms observed in the SEM images. The amount of zinc blende (ZB) phase was beyond the detection

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Figure 4. (a) TEM image of the CdS nanocrystals. (b and c) TEM images of a single nanocrystal at the sample holder tilting angle 20° (b) and -20° (c). Insets: schematic illustration of the orientation of the tetrahedron. The solid and dashed wedges correspond to the arms on the upper and lower side of the projection plane, respectively.

limit of the equipment, as indicated by the lack of ZB (002) diffraction peak at 30.75°. Figure 4a is a typical TEM image of the tetrahedral nanocrystals, showing their homogeneous size and shape. The core part of the nanocrystal exhibited much darker contrast than the arms. By tilting the sample holder, the observed image of a single nanocrystal changed accordingly, further confirming its tetrahedral shape (Figure 4b,c). The corresponding orientations of the tetrahedron are illustrated in the insets, using a description similar to that used for the carbon tetrahedron in stereochemistry: the solid and dashed wedges correspond to the arms on the upper and lower side of the projection plane, respectively. Figure 5 shows the structure analysis of the tetrahedral nanocrystals by HRTEM. The studied nanocrystal lay in the projection plane symmetrically, showing three arms: one longer one (arm 1 in Figure 5d) and two shorter ones with identical length (arm 2 in Figure 5d). The electron diffraction spots of the long arm exhibited a typical [112j0] zone pattern (Figure 5a). The arm growth direction was [0001]. The resolved lattice interval was 0.333 nm, also in accordance with that of the (0002) planes of wurtzite CdS (Figure 5b). The growth direction was consistent with both the hexagonal pyramid-shaped arms observed in SEM and the extraordinarily intense (0002) peaks in the XRD pattern. The projection length of this arm reached its maximum in this configuration, indicating that it lay exactly in the projection plane. The symmetrical distribution of the other two shorter arms suggested that the projection plane was perpendicular to one σv of the tetrahedron (Figure 6a). The schematic illustration of the orientation of the studied nanocrystal with respect to the projection plane is shown in Figure 6. The two shorter arms observed in Figure 5a are inclined to the projection plane and are on the same side. Assuming perfect Td symmetry, they form an angle of 25.1° to the projection plane. The electron diffraction of one out-of-plane arm (arm 2 in Figure 5d) is shown in Figure 5e. The two base vectors were (01j11) and (1j103), respectively. The resolved lattice interval was 0.316

Figure 5. Single nanocrystal for the HRTEM study. (a, b) Electron diffraction patterns and HRTEM images of area 1 in (d). (c) HRTEM image of area 3 in (d). (d) TEM image of a single tetrahedral CdS nanocrystal. (e, f) Electron diffraction patterns and HRTEM images of area 2 in (d).

nm, corresponding to that of the (01j11) planes of wurtzite CdS (Figure 5f). The growth direction of the out-of-plane arm on the projection plane was close to the sum vector of the two basics, i.e., perpendicular to the (1j014) plane. The dihedral angle between the (1j014) and (0001) planes was 28.1°. This value should equal that between the out-of-plane arms and the projection plane (25.1°) since all four arms were growing along the [0001] direction (Figure 6b). The 3° discrepancy was a quite good agreement considering that the normal of the (1j014) plane was only an approximate orientation. Figure 5c is the HRTEM image taken in the junction of the two arms, where the (0002) and (1j103) lattice fringes were observed simultaneously for both the in-plane and out-of-plane arms, respectively. A model describing the structure of group II-VI tetrahedral nanocrystals suggested that the tetrahedral nanocrystal consisted of a tetrahedral core of zinc blende structure enclosed by four ZB{111} planes and four arms of wurtzite structure growing along the WZ [0001] direction. The arms epitaxially attached to the ZB{111} planes of the core using their WZ (0001) planes.33 Since the two crystal planes have identical atomic structures, there will be no lattice mismatch introduced. This model was structurally perfect and was proved by direct HRTEM observations of solution synthesized II-VI tetrahedral nanocrystals.16,29,34,35 However, we failed in resolving satisfactory HRTEM images for the core part of our sample due to the large sample thickness. The ZB phase in the nanocrystal was

Highly Symmetrical CdS Tetrahedral Nanocrystals

Figure 6. (a) Schematic illustration of the orientation of the nanocrystal in Figure 5d with respect to the projection plane. The four strongest lines represent the four arms (their exact shape was not specified for simplicity). The dark area in the projection plane is the observed shape under TEM. R is calculated using geometry assuming the tetrahedron is of perfect Td symmetry. (b) Schematic illustration of the out-ofplane indicated by the dashed ellipse in (a), showing the relationship between the (0001) and (1j014) planes. β is calculated using the lattice parameters of wurtzite CdS.

only quite a small fraction according to the XRD results and was buried deep in the core part; thus it was very difficult for HRTEM detection when the arm diameter was large. Considering their near perfect Td symmetry, our tetrahedral nanocrystals were also likely to follow the ZB-core-WZ-arms model. In addition, the low deposition temperature also favored the nucleation of the metastable ZB phase, which was expected to transform into the stable WZ phase at temperatures higher than 300 °C for bulk CdS.36 Temperature was found to be the most influential factor for the formation of tetrahedral nanocrystals. The optimum values for the CdCl2 source and the deposition temperatures were 540 and 400 °C, respectively. Increase of the deposition temperature was detrimental for the tetrahedral structure. Both the products collected closer to the CdCl2 source at 540 °C or those obtained at the same location but at higher CdCl2 temperature gradually lost the well-defined hexagonal pyramidal shapes of their arms and evolved into irregular nanoparticles (Figure 7b,c). A large number of nanocones were obtained at higher source temperature (inset, Figure 7b). These nanocones could be viewed as unconnected arms, which developed into comparable sizes to the tetrahedral nanocrystals. Unlike the arms of the tetrahedral nanocrystals, the separated nanocones did not exhibited welldefined facets, which could be attributed to the higher deposition temperature. In addition to the temperature effect, both higher source temperature and shorter distance from the source caused higher Cd/S ratio at the deposition locus. However, we believe that the temperature effect was majorly responsible for the shape transformation, as supported by the annealing behaviors of the as-grown tetrahedral nanocrystals (Figure 7d). After annealing at 500 °C for 2 h in Ar, the tetrahedral nanocrystals transformed into near-round nanoparticles, similar to products deposited at higher temperature. The XRD pattern also changed correspondingly (Figure 3b). The intensities of (0002) and (0004) diffrac-

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Figure 7. (a-c) SEM images of CdS nanocrystals obtained under different conditions: (a) CdCl2 540 °C, 10–12 cm, (b) CdCl2 570 °C, 10–12 cm, (c) CdCl2 540 °C, 7–9 cm. (d) SEM image of the tetrahedral nanocrystal in (a) after annealing in Ar at 500 °C for 2 h.

tion peaks were reduced significantly to those of the standard pattern, suggesting the loss of high shape anisotropy at elevated temperature. On the basis of the above observations, the increasing deposition temperature had two effects that were unfavorable for the tetrahedral structures. First, the hexagonal pyramidal arms were destabilized. They transformed into irregular or round nanoparticles at elevated temperature to reduce the surface energy. Second, the number of nexus to interconnect the arms was reduced. This was reasonable because according to the ZBcore-WZ-arm model, the nexus were of the metastable zinc blende structure, which was unfavorably formed at higher temperature. The second point also partly supported the correctness of the structural model for the tetrahedral nanocrystals. Therefore, the tetrahedral nanocrystals were kinetic controlled products that were formed only at low temperature (400 to 420 °C). The low deposition temperature in our method was favored by the formation chemistry CdCl2 + H2S ) CdS + 2HCl. When HCl is efficiently removed, this reaction takes place quite easily. Because it involves only component exchange without a redox process, the transition state is expected to be simple and the activation energy must be low. Therefore, a much lower deposition temperature compared to those of conventional thermal evaporation and CVD methods could be achieved. Although it is difficult to quantify the generation rates of the Cd and S species, we deduce that the tetrahedral nanocrystals were produced in a sulfur-rich environment based on two points. First, a large amount of sulfur was detected in the NaOH solution located at the end of the gas flow, indicating sulfur excess. In addition, the optimal location to collect the tetrahedral nanocrystals was 10 to 12 cm downstream to the CdCl2 source, where the CdCl2 partial pressure in the gas phase was low due to the low temperature and consumption along the gas flow. One major advantage of our method compared to the conventional ones was the freedom of adjusting the Cd/S ratios. The Cd generating rate was controlled by the CdCl2 source temperature. The S generating rate was determined by the hydrolysis rate of the polysulfide, which was adjusted by the temperature of the water thermal bath. A temperature lower than 50 °C resulted in insufficient H2S generation and thus very little deposition. When the temperature was higher than 80 °C, a large

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on the surface or even incorporated into the nanocrystal, which accounted for the 745 nm emission.23,41 The desorption of these species after annealing in Ar also caused a change of surface state. Therefore, both the change of surface geometry and chemical environment was expected to create different surface-trapped states and lead to a change in the red emission band. Conclusions

Figure 8. PL spectra of the as-prepared sample and the sample after annealing in Ar at 500 °C for 2 h. The dashed lines are the Gaussian deconvolution of the broad emission band of the annealed sample.

amount of water vapor was introduced into the reaction chamber and condensed on the wall of the quartz tube in the lowtemperature region. However, the product morphologies were not very sensitive to the H2S generating rate with respect to the temperature. Satisfactory tetrahedral nanocrystals could be obtained when the water bath temperature was 65 to 75 °C when the N2 flow was 200 sccm. The photoluminescence spectra of the as-prepared and annealed samples are shown in Figure 8. The as-prepared sample exhibited a sharp peak centered at 514 nm (green emission) and a broad band from 600 to 800 nm (red emission) centered at 745 nm with similar intensity. The sharp emission peak was located close to the band gap of bulk CdS (∼2.4 eV), which originated from the carrier recombination from the band edges or very shallow trapped states near the band edges. The broad emission band from 720 to 770 nm was quite frequently observed for CdS nanocrystals,15,23,37–40 especially those prepared by thermal evaporation.23,38–40 The origin of this emission was usually attributed to surface-trapped states23,41,42 or a trapped level formed by a Cd vacancy and a substituted Cl- at the neighboring site.37 After annealing, the position of the green emission peak did not change, while the intensity was slightly enhanced with respect to the red one. In addition, the red emission became broader and was blue-shifted compared to that of the as-prepared sample. Deconvolution of the broad band gave two Gaussian peaks centered at 724 and 646 nm, respectively (Figure 8b). The emission band around 650 nm was observed in some colloidal CdS nanocrystals, which was also believed to be surface-trapped states’ emission.43,44 The slight enhancement of the near band edge emission was reasonable because annealing was expected to improve the crystallinity of the nanocrystals and thus reduce the intensity of the trapped state emission.45 The change in the red emission band corresponded to the change of surface states after annealing, since surface-trapped states were mainly responsible for the red emission. The as-prepared tetrahedral nanocrystals with well-defined facets could be identified as enclosed by low-index crystal planes, like (0002) and (1j 100). The round surface of the post-annealing nanoparticles was expected to consist of small crystal planes with low surface energy separated by surface steps, which lowered the total surface energy. New surface states could be created due to the change of the exposed crystal planes. Another possible change during the annealing was the desorption of the surface-adsorbed species. Because the growth was carried out at relatively low temperature and a sulfur excessive environment, anions such as SH- and Cl- could be adsorbed

In conclusion, we demonstrated the effectiveness of polysulfide as a sulfur source in the chemical vapor deposition of CdS nanocrystals. Using the reaction between CdCl2 and H2S, which was generated through the hydrolysis of polysulfide, CdS tetrahedral nanocrystals with nearly perfect Td symmetry were prepared at a low deposition temperature of 400 °C. The low deposition temperature was critical for the formation of the tetrahedral structure. The photoluminescence spectra exhibited a sharp band edge green emission and a red emission band related to trapped states. Particularly, similar tetrahedral structures were almost exclusively obtained with solution chemistry methods. Our results suggested that by using proper precursors and reaction routes, mild operation conditions and abundant product morphologies could also be achieved in CVD methods. Acknowledgment. This work was supported by the National Natural Science Foundation of China (Nos. 20221101, 10335040, and 20671004), MOST of China (No. 2006AA05Z130), and MOE of China (No. 707002). The authors thank Dr. Renmin Ma from School of Physics, Peking University, for the PL measurement.

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