LETTER pubs.acs.org/NanoLett
Chemical Transformation of Au-Tipped CdS Nanorods into AuS/Cd Core/Shell Particles by Electron Beam Irradiation Marijn A. van Huis,*,†,‡,r Albert Figuerola,§,||,r Changming Fang,†,^ Armand Beche,#,‡ Henny W. Zandbergen,† and Liberato Manna*,†,§ †
Kavli Institute of Nanoscience, Delft University of Technology, Lorentzweg 1, 2628 CJ Delft, The Netherlands EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium § Istituto Italiano di Tecnologia, via Morego 30, 16163 Genova, Italy Departament de Química Inorganica, Universitat de Barcelona, Martí i Franques 1-11, 08028 Barcelona, Spain ^ Materials innovation institute (M2i), Mekelweg 2, 2628 CD Delft, The Netherlands # FEI Company, Achtseweg Noord 5, 5600 KA Eindhoven, The Netherlands
)
‡
bS Supporting Information ABSTRACT: We demonstrate that electron irradiation of colloidal CdS nanorods carrying Au domains causes their evolution into AuS/Cd core/shell nanoparticles as a result of a concurrent chemical and morphological transformation. The shrinkage of the CdS nanorods and the growth of the Cd shell around the Au tips are imaged in real time, while the displacement of S atoms from the CdS nanorod to the Au domains is evidenced by high-sensitivity energy-dispersive X-ray (EDX) spectroscopy. The various nanodomains display different susceptibility to the irradiation, which results in nanoconfigurations that are very different from those obtained after thermal annealing. Such physical manipulations of colloidal nanocrystals can be exploited as a tool to access novel nanocrystal heterostructures. KEYWORDS: Colloidal heteronanocrystals, chemical transformations, structural transformations, electron irradiation, transmission electron microscopy
W
et chemistry techniques nowadays enable the fabrication of a wide variety of heterogeneous nanostructures, consisting of two or more nanoscale materials which can be assembled in various three-dimensional configurations. Examples are core shell nanoparticles, dumbbells, nanorods with tips, and branched nanocrystals.16 The plethora of possibilities to design nanoscale particles has led to new nanomaterials with a wide range of physical properties.57 Semiconductormetal heterostructures such as the CdS/Au system are promising materials for photovoltaics, photocatalysis, and optoelectronics.813 Their interesting optical properties stem from a different spatial distribution of electron and hole charge carriers. In the case of nanoscale Au/ CdS core/shell particles it is found that the electrons are mainly confined in the Au domain, while the holes are confined in the CdS domain,14,13 which leads to long radiative lifetimes. De Paiva and Di Felice, who studied the electronic structure of the CdS/ Au interface by means of first-principles calculations, predicted orbital mixing effects between S and Au,15 which was actually confirmed by Khon et al. who found experimentally that strong electronic coupling occurs between the CdS and Au domains in CdS/Au nanorods.13 In a recent work, Chen et al. reported the synthesis of hybrid CdSAu2SAu nanostructures obtained starting from CdSAu nanocrystals via an inter-nanocrystal Cd2+Au+ cation exchange process.16 r 2011 American Chemical Society
The control over the selective location of metallic domains on the surface of semiconductor nanocrystals and the quality of the interface between them are also important issues when considering the possible use of such types of heterostructures in electronic devices. In an earlier work, we studied the ripening process of Au dots on CdSe nanorods as a result of thermal heating, and found that ripening takes place through a combined process of atomic and Au cluster diffusion, and eventually resulted in a well-defined epitaxial orientation relationship between the CdSe nanorod and the Au tips.11 Besides thermal treatments, there is an increasing number of examples in the literature of the so-called radiationinduced annealing, whereby photon or electron beams induce ripening processes leading to the selective nucleation and growth of Au domains in nanostructures.9,10,17 For instance, in the specific case of Au-tipped CdS nanorods with different seed materials (CdS, ZnTe, and CdSe), Menagen et al. studied the ripening process as a result of thermal heating and of light illumination.9,10 Recently, Carbone et al. were able to grow large gold domains exclusively on one side of CdS or CdSe/CdS quantum rods via a photoreduction process of gold ions under anaerobic conditions.18 Received: May 14, 2011 Revised: October 10, 2011 Published: October 13, 2011 4555
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Figure 1. Thermal ripening of Au-decorated CdS nanorods. (a) Initial configuration. (b) Same configuration obtained after annealing for 20 min at a temperature of 180 C. In general the Au tips at the polar ends of the nanorod grew at the expense of the side tips. Also midsectioning of the CdS rods was observed (white arrow). (c) After annealing at 240 C, all Au side tips had dissociated, and all remaining Au domains were connected to the CdS nanorods by flat interfaces. (d) The CdS/Au interfaces obey a well-defined orientation relationship between the two crystals, identical to that found for CdSe/Au interfaces.11
Furthermore, Zeng et al. showed that a chemical reaction triggered by light illumination can also be used to create new nanostructures. In that specific case they observed the formation of hollow CdCl2 nanotubes starting from CdSe nanocrystals.19 We report here a systematic study carried out “in situ” in the transmission electron microscope (TEM) on the thermal evolution and an electron-irradiation-induced transformation of CdS/Au heteronanostructures. At low electron beam intensities (below 1 106 electrons nm2 s1), only ripening of the Au dots to the tips was observed upon heating, while at intense electron beam irradiation an electrochemical transformation was observed by which CdS/Au rod-tip configurations evolved to AuS/Cd core/shell nanoparticles. New chemical compounds were thus formed (AuS and Cd) from the starting compounds (CdS, Au). This process took place already at room temperature, but at elevated temperatures it was seen to proceed at a faster rate. Other research groups investigating CdS/Au nanostructures have not reported this transformation so far, most likely because in those experiments the electron beam flux was too low and/or the irradiation times too short. Apart from the new possibility uncovered here of fabricating novel heteronanostructures by means of electron irradiation, our experiments raise a general question about the chemical stability of the various nanocrystal heterostructures that have been proposed so far by many groups as building blocks for photovoltaic devices. As all these nanostructures would be subjected to photon irradiation for long periods of time (decades), possible chemical and morphological transformations have to be taken into account. Various sets of CdS/Au heterostructures were used for the experiments: separated CdS rods dotted with small Au domains at their surfaces, separated CdS rods having larger Au domains but only at the nanorod ends, and CdS tetrapods dotted with Au domains. The synthesis procedures are described elsewhere,20,11 and additional details are given in the Supporting Information. In a first series of experiments, the influence of the electron beam was avoided by keeping the beam dispersed and the irradiation
times short. The thermal evolution of the CdS nanorods decorated with many small Au domains is shown in Figure 1. The nanostructures had been drop-cast onto thin (20 nm thick) SiN membranes, which are incorporated in a MEMS micro hot plate enabling high-resolution imaging at elevated temperature in the transmission electron microscope (TEM).21 Upon heating to temperatures of 120240 C, a ripening mechanism occurred as large Au domains grew at the expense of smaller Au domains. In the final stages of the thermal evolution, nearly all Au domains were located at the polar ends of the wurtzite CdS nanorods. Well-defined, straight CdS{0001}/Au{111} interfaces were formed, whereby the two crystal domains had a well-defined orientation relationship (OR) that was identical to the OR found by us in a previous work for CdSeAu rodtip configurations.11 Figure S1 in the Supporting Information shows the thermal evolution of Au-decorated CdS tetrapods. Also here, larger Au domains grew at the expense of smaller Au domains and preferentially developed at the tips of the pods. It is also of interest to note that occasionally, segmenting of the CdS rods took place, as is clear from Figure 1b,c and Supporting Figure S2 in the Supporting Information. In these cases a Au domain was formed within a CdS nanorod, whereby the CdS rod was segmented in two parts by the creation of two new CdS{0001}/ Au{111} interfaces. The occurrence of such Au middle segments is probably an indication that this type of interface is energetically favored. Such segmenting was reported previously by Robinson et al., where periodic arrays of Ag2S segments were formed in CdS nanorods through Cd2+ w Ag+ cation exchange.22 In the current work, though, the sectioning is induced by thermal annealing, and the new sections consist of pure Au rather than of Au2S. As mentioned above, in the thermal annealing study the influence of the electron beam was avoided by keeping the beam intensity low and/or exposing the nanostructures only for short time durations. However, when exposing the nanostructures to a high-intensity electron beam (flux 2 106 to 5 106 electrons nm2 s1) the transformation to AuS/Cd core/shell particles took place during a time interval of minutes. The transformation is shown in Figure 2ac, and a scheme is included in Figure 2d. 4556
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Figure 2. Au-tipped CdS nanorods, which transform into AuS/Cd coreshell particles upon electron irradiation, while being heated at a temperature of 200 C. (a) Initial configuration of Au-tipped CdS nanorods. (b) Formation of rectangular cubic Cd domains, incorporating one or two AuS domains. (c) Final configuration: multiply twinned particles (MTPs) of AuS with a Cd shell. (d) Schematic of the radiation-induced chemical transformation.
Figure 3. CdS nanorod with Au tip at the end facet, before (a) and after electron irradiation (b). After 330 s of electron irradiation, the Au tip had evolved into a AuS/Cd core/shell particle, while the CdS nanorod had shrunk and was breaking up. This experiment was performed at room temperature.
The Cd nanocrystals had a cubic crystal lattice and had the tendency to adopt a rectangular morphology. Each Cd nanocrystal was always enclosing a smaller AuS nanocrystal (the S had diffused into the Au domains, as will become clear below from chemical mapping). Although we will refer to these domains with the general designation AuS, the composition was actually found to vary (AuxSy). Upon continued electron irradiation and annealing, the Cd shell shrunk, probably by knockout of the Cd atoms by the electron beam or by out-diffusion of remnants of S, so that spherically shaped AuS/Cd core/shell particles were eventually formed (Figure 2c). The structures displayed in panels b and c of Figure 2 were formed by electron beam irradiation while keeping the sample at a temperature of 200 C. However, the concurrent chemical and morphological transformation also took place at room temperature, as shown in Figures 3 and 4. In Figure 3, a configuration before and after electron irradiation is shown. The CdS nanorod had shrunk and was breaking up, while a large Cd shell had formed
around the Au tip. A similar intermediate stage is shown in Supporting Figure S3 (Supporting Information). Figure 4 shows a series of stills from Supporting Movie M1 in the Supporting Information, where the growth of the Cd shell was monitored in real time. In order to assess unambiguously the chemical composition of the observed new nanodomains, high-sensitivity EDX measurements were performed using the Super-X system as implemented in the Chemi-STEM microscope located at FEI Nanoport, The Netherlands (details in the Supporting Information). The energy resolution of 0.15 keV of the system allows to distinguish well the Au, Cd, and S peaks. The result of EDX mapping is shown in Figure 5, while Figure S10 in the Supporting Information displays the Au, Cd, and S concentration profiles in a digital line scan over the structures. In both figures, the concentration of elements is displayed after quantification by fitting the X-ray spectra (in every pixel) using the CliffLorimer method. Figure 5a shows chemical mapping of the elements (Au, Cd, S) in the nanorod/tip configuration and Figure 5b the distribution of elements in the core/shell configuration. Before the transformation, the tips consist mainly of Au, while the Cd and S are present in similar concentrations in the CdS nanorod. After the transformation, the S has been displaced to the Au tips (the S maps and Au maps are clearly overlapping), while the area showing the presence of Cd is much larger. The Cd shell is surrounding the AuS domain in three dimensions, so that in projection the larger Cd shell overlaps with the smaller AuS domain. The presence of elemental Cd in the shell is a rather surprising result, since the crystal structure of the shell is clearly cubic, as evidenced by the many high-resolution transmission electron microscopy (HRTEM) images (Figures 2b, 4, and 6), while the bulk crystal structure of Cd is hexagonal (hexagonal close packed, hcp). We will show below from a comparison of experimental and calculated lattice parameters that the crystal structure is most likely face-centered cubic (fcc). From the digital EDX line scans (Figure S10, Supporting Information) it is also clear that the fraction of S in the Au dots varied, between 20 and 60 atomic %. The variation in S content can be plausibly explained by the fact that the volume of the Au 4557
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Figure 4. Stills of a movie recording (Supporting Movie M1, Supporting Information) whereby the growth of the Cd shell around the Au tip of a CdS nanorod was followed in real time, at room temperature. The stills correspond to 35 s time intervals; the total duration is 240 s.
Figure 5. Chemical mapping of elements Au, Cd, and S of (a) an Au-tipped CdS nanorod before the transformation and (b) of several AuS/Cd core/ shell particles after the transformation. From left to right: HAADF-STEM image, corresponding Au map, Cd map, S map, and combination of Au, Cd, and S maps. For the core/shell particles in (b), the Au and S maps are nearly identical, indicating that the S is confined in the AuS nanodomains. The larger Cd shells are surrounding the AuS domains (the Cd is present at the sides but also covering the front and the rear of the AuS domains), and therefore in projection the larger Cd domains are overlapping with the smaller AuS nanodomains.
dots varies considerably (Figure 1 and Figure S2 in the Supporting Information), while the dimensions of the CdS nanorods (being the “source” of S) are uniform. The value of 60 atomic % of S in the Au dot (Figure S10b, Supporting Information) is very high, though, as there are no known bulk phases with a similar composition. Possibly the S content is overestimated, as the fitting software assumes that the X-rays of S are initiated by the electron beam, while one needs to consider that the X-rays from Au are very energetic (strong Lα peak at 9.5 keV) and are able to excite secondary X-rays in S atoms (main Kα peak at 2.3 keV), which adds up to the total S signal. Figure 6 shows high-resolution TEM images of several core shell configurations, displaying clearly the orientation relationship between the AuS and Cd crystal lattices. From the (002) lattice spacings in Figure 6a, we derived lattice parameters equal to 4.9 ( 0.1 Å for Cd and 4.2 ( 0.1 Å for AuS. Bulk Au has the fcc structure with a lattice parameter of 4.08 Å. Therefore, the AuS lattice is expanded with respect to that of pure Au.
In order to estimate the lattice parameter for cubic Cd and to understand the affinity of S with Au, density functional theory (DFT) calculations were performed with the first-principles VASP code,23,24 employing plane augmented waves (PAWs) and the generalized gradient approximation (GGA) with the potentials of Perdew, Burke, and Ernzerhof (PBE).25 Details of calculation settings and calculated lattice parameters of the various phases are given in the Supporting Information. Table S1 in the Supporting Information shows the lattice parameter and formation enthalpies for elemental phases, and for a number of (defective) CdS, AuS, and CdAu compounds. Here the formation enthalpy ΔE of a composite structure AuxCdySz is defined with respect to the elemental phases (fcc Au, hcp Cd, the S2 molecule) ΔEðAux Cdy Sz Þ ¼ EðAux Cdy Sz Þ xEðAuÞ yEðCdÞ z1 =2 EðS2 Þ
ð1Þ 4558
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Figure 6. (a) HRTEM image of a AuS domain, in a cube-on-cube orientation relationship with a larger Cd domain. From the inset showing the Fourier transform of the indicated area, it is clear that their lattice parameters differ substantially. (b) Two other AuS/Cd core shell particles formed during the early stages of irradiation.
whereby all energies are expressed in eV/atom. The values for the formation enthalpies should be considered with much caution, as they are valid only for thermal equilibrium, a condition that is not satisfied during irradiation. So, e.g., the energy gain that is achieved by fusion of Cd and S into wurtzite CdS (the nanorod material) equals 1.07 eV/atom (Table S1, Supporting Information), but under intense electron irradiation the nanorod material is nonetheless decomposing. The formation enthalpies show that bonding of S with Au is energetically favorable for some compositions, e.g., Au4S (with S at a tetrahedral interstitial site in the fcc Au lattice) gives an energy gain of 0.12 eV/atom, while the Au2S phase yields an energy gain of 0.36 eV/atom. Table S1 (Supporting Information) shows that incorporation of S atoms in the fcc Au lattice (compositions Au32S, Au3S) leads to an increased lattice parameter of 4.24.3 Å, in agreement with the experimental observations. No evidence was found for the presence of the Au2S phase, which has a crystal structure that is not of fcc-type (see Supporting Information for details). Table S1 (Supporting Information) also shows the lattice parameters and energies of elemental Cd, in the hcp, fcc, and body centered cubic crystal structures. From the analysis of a large number of HRTEM images, it was found that the Cd shell has the fcc crystal structure. The particle at the right-hand side of Figure 2c, for example, displays the [011] projection of the fcc lattice. For bulk Cd, the fcc phase is less stable than the hcp phase by +0.026 eV/atom. However, it is not uncommon that nanoscale materials adopt an alternative crystal structure.2628 In bulk materials the contribution from surface energies to the total
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energy is negligible, while in nanomaterials the number of surface atoms is relatively large, and therefore the contribution from surface energies is significant. As a result, nanoscale materials can adopt structures that would be unstable as bulk material. In our experiments it was found that as the Cd shell becomes thinner during the transformation, the lattice spacings also decrease to values corresponding to an fcc lattice parameter of 4.5 Å. The experimental lattice parameter of fcc Cd found in this work (4.54.9 Å) cannot be compared with values in the literature, as this phase has never been reported experimentally. The DFT calculations (Table S1, Supporting Information) predict a lattice parameter for fcc Cd of 4.52 Å, in very good agreement with the lattice spacings found at the late stages of coreshell particles. The expanded lattice parameter found in the early stages is most likely caused by the presence of remnant S atoms in the fcc Cd lattice. Also for this purpose, we have performed additional DFT calculations and found that incorporation of S is energetically favorable (e.g., for the composition Cd4S whereby the S atom occupies a tetrahedral interstitial site) and leads to an increased lattice parameter (5.03 Å for Cd4S). Therefore, it seems likely that the shell grows as a S-rich Cd lattice, but then over a few minutes of irradiation the S is increasingly displaced to the Au domain, so that the shell eventually consists of nearly pure Cd. The structural relationship between the wurtzite CdS phase, the zinc blende CdS phase, and fcc Cd is further explained in the Supporting Information. Considering the concentration profiles as determined by EDX (Figure S10, Supporting Information), we denote the core material as AuS, and the shell material as Cd, although the latter may incorporate S, depending on the stage of the transformation. Both crystal structures are of fcc-type. Considering Figure 6a, it is clear that the two cubic lattices (AuS and Cd) are in a cube-oncube orientation relationship (i.e., {100}AuS//{100}Cd and [010]AuS//[010]Cd). The experimentally observed lattice mismatch is approximately 15% at the early stages of irradiation, and about 6% at the late stages. From the bent lattice planes in Figure 6a, it is also clear that the interface is strained. Apparently, the chemical binding energy is so strong that it overcomes the interfacial strain energy. Another feature indicating that the AuS/ Cd interface is energetically very favorable is evidenced by the observation that AuS domains were often observed to be physically drawn into the larger Cd domains, as shown in Figure S7 of the Supporting Information. The CdS tetrapods with Au tips transformed in a manner similar to the Au-tipped CdS nanorods. Figure S9 (Supporting Information) shows the case of a Au-decorated CdS tetrapod before the transformation (a) and after partial transformation (b), whereby a large cubical Cd domain has grown at the position of the pod that is directed parallel to the beam. In general, the AuS domains were not always arranged as single crystals such as shown in Figure 6a and at the right-hand side of Figure 6b; also the decahedral morphology29 was often observed, e.g., in Figure 2c. In those cases, the Cd shell also consisted of several subcrystals whose orientations were dependent on those of the AuS subcrystals in the multiply twinned particle (MTP). The question now rises why this transformation occurs at all, and why it cannot be achieved, e.g., by thermal heating (which instead gives the epitaxial CdS/Au rodtip configurations shown in Figure 1). One possibility is that the transformation is caused mainly by atomic displacements directly induced by the electron beam, whereby vacancyinterstitial pairs (Frenkel defects) for Cd and S are formed, after which the interstitial atoms diffuse to 4559
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Nano Letters the CdS/Au interface and react there to form new compounds. The current observations, though, were well reproduced at 80 keV electron beam energy, which is lower than the minimal electron beam energies of 285 and 115 keV that are required to provoke direct atomic displacements of S and Cd atoms, respectively, in bulk CdS.30 Therefore, although knockout of Cd and S atoms at higher acceleration voltages may speed up the chemical transition (this was indeed observed), it cannot serve as the main explanation for the concurrent chemical/morphological transformation. It is known from the literature that in semiconductors, atomic displacements can also be caused indirectly, through a process which is referred to as “subthreshold ionization”.31,32 Even when the energy of an incident 80100 keV electron is insufficient to displace a lattice ion, a multiple ionization process taking place in the target material leads to the expulsion of anions or cations and the formation of vacancyinterstitial point defects. When the electron beam flux is sufficiently high, the subthreshold ionization leads to decomposition, as was shown previously for irradiated CdS thin films.31 Metal nanoparticles are mostly unaffected by 100 keV electron beam irradiation as the process of subthreshold ionization is not effective within metals. On the other hand the CdS material within the present CdSAu nanostructure is very sensitive to electron beam irradiation and decomposes. (The area covered by electron beam irradiation always included both the CdS and the Au nanodomains.). The freed Cd atoms agglomerate as Cd metal, which is less susceptible to the irradiation. The S atoms instead prefer to bind with Au rather than with Cd, probably because AuxSy is more stable under these conditions than CdxSy. In conclusion, the cause of the transformation of CdSAu nanostructures observed by us should be found in the different susceptibility of the various domains of the nanostructure to irradiation. This does not answer all questions, though. When pure CdS nanorods without any Au tips were examined,33 decomposition of the CdS was not observed, also not at high beam intensities. These CdS nanorods only decomposed (sublimate) upon heating to temperatures of 400450 C, but this happened everywhere on the support (not only in areas that were examined with the beam), and therefore the sublimation is a thermal effect and not a radiation induced effect. Thus it can be concluded that the presence of the Au domains considerably affects the stability of the CdS nanorod. From recent reports in the literature,13,14 it is known that the CdS and Au domains display strong electronic coupling effects, leading to interdomain charge transfer, bending of energy bands, and modifications in the density of states near the interface. These shifts in chemical potential apparently render Au-tipped CdS nanorods much less resistant to irradiation than CdS nanorods without Au tips. Considering the thermal stability of the AuS/Cd coreshell particles, we mention that additional heating (in the absence of the electron beam) did not alter the morphology of the coreshell configuration. Whereas other core/shell systems such as PbSe/CdSe have the tendency to evolve into two domains separated by a single, flat interface,34 the current coreshell system retained its morphology upon heating. The stable coreshell morphology indicates that the fcc Cd free surface is energetically favored over the AuS free surface, so that a lower total energy of the nanostructure is obtained when the outer surface consists only of fcc Cd, while the AuS is confined to the core and therefore does not have any free surface (there is only an interface with Cd).
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In summary, here we have shown that completely new nanostructures can be formed by means of electron irradiation. The transformation is caused by the different susceptibilities of the various nanodomains to electron irradiation and takes place rather swiftly (within minutes) under an intense electron beam. This can lead to the decomposition of particular domains within the heterogeneous nanostructure, to changes in composition of other domains, and even to the growth of new types of nanodomains.
’ ASSOCIATED CONTENT
bS
Supporting Information. Details on the synthesis, chemicals, DFT calculations, and TEM measurements. Additional high-resolution TEM images in Figures S1S9, chemical mapping in Figure S10, and Supporting Movie M1 (corresponding to Figure 4) showing the chemical and structural transformations in real time. This material is available free of charge via the Internet at http://pubs.acs.org.
’ AUTHOR INFORMATION Corresponding Author
*Email:
[email protected],
[email protected]. Author Contributions r
These authors contributed equally to this work.
’ ACKNOWLEDGMENT The authors acknowledge funding from the European Union Framework 6 Program, Reference 026019 ESTEEM, and the FP7 starting ERC grant NANO-ARCH (Contract No. 240111). A.F. acknowledges financial support from the Spanish Government through a Ramon y Cajal Fellowship. The authors thank the FEI Company for access to the Chemi-STEM microscope with the Super-X EDX system and C. Mitterbauer (FEI) for help with the quantitative analysis. ’ REFERENCES (1) Carbone, L; Cozzoli, D Nano Today 2010, 5, 449–493. (2) Costi, R; Saunders, A, E; Banin, U Angew. Chem., Int. Ed. 2010, 49, 4878–4897. (3) Deka, S.; Miszta, K.; Dorfs, D.; Genovese, A.; Bertoni, G.; Manna, L. Nano Lett. 2010, 10, 3770–3776. (4) De Mello Donega, C. Chem. Soc. Rev. 2011, 40, 1512–1546. (5) Lo, S. S.; Mirkovic, T.; Chuang, C.-H.; Burda, C.; Scholes, G. D. Adv. Mater. 2011, 23, 180–197. (6) Krahne, R.; Morello, G.; Figuerola, A.; George, C.; Deka, S.; Manna, L. Phys. Rep. 2011, 501, 75–221. (7) Mocatta, D.; Cohen, G.; Schattner, J.; Millo, O.; Rabani, E.; Banin, U. Science 2011, 332, 77–81. (8) Saunders, A. E.; Popov, I.; Banin, U. J. Phys. Chem. B 2006, 110, 25421–25429. (9) Menagen, G.; Mocatta, D.; Salant, A.; Popov, I.; Dorfs, D.; Banin, U. Chem. Mater. 2008, 20, 6900–6902. (10) Menagen, G.; Macdonald, J. E.; Shemesh, Y.; Popov, I.; Banin, U. J. Am. Chem. Soc. 2009, 9, 2031–2039. (11) Figuerola, A.; van Huis, M. A.; Zanella, M.; Genovese, A.; Marras, S.; Falqui, A.; Zandbergen, H. W.; Cingolani, R.; Manna, L. Nano Lett. 2010, 10, 3028–3036. (12) Khon, E.; Hew-Kasakarage, N. N.; Nemitz, I.; Acharya, K.; Zamkov, M. Chem. Mater. 2010, 22, 5929–5936. 4560
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(13) Khon, E.; Mereshchenko, A.; Tarnovsky, A. N.; Acharya, K.; Klinkova, A.; Hew-Kasakarage, N. N.; Nemitz, I.; Zamkov, M. Nano Lett. 2011, 11, 1792–1799. (14) Yang, T.-T.; Chen, W.-T.; Hsu, Y.-J.; Wei, K.-H.; Lin, T.-Y.; Lin, T.-W. J. Phys. Chem. C 2010, 114, 11414–11420. (15) De Paiva, R.; Di Felice, R. J. Phys. Chem. C 2010, 114, 3998–4007. (16) Chen, S.; Zhang, X.; Hou, X.; Zhou, Q.; Tan, W. Mater. Lett. 2010, 64, 489–492. (17) Meyns, M.; Bastus, N. G.; Cai, Y.; Kornowski, A.; Juarez, B. H.; Weller, H.; Klinke, C. J. Mater. Chem. 2010, 20, 10602–10605. (18) Carbone, L.; Jakab, A.; Khalavka, Y.; Soennichsen, C. Nano Lett. 2009, 9, 3710–3714. (19) Zeng, J.; Liu, C.; Huang, J.; Wang, W.; Zhang, S.; Li, G.; Hou, J. Nano Lett. 2008, 8, 1318–1322. (20) Figuerola, A.; Franchini, I. R.; Fiore, A.; Mastria, R.; Falqui, A.; Bertoni, G.; Bals, S.; Van Tendeloo, G.; Kudera, S.; Cingolani, R.; Manna, L. Adv. Mater. 2009, 21, 550–554. (21) van Huis, M. A.; Young, N. P.; Pandraud, G.; Creemer, J. F.; Vanmaekelbergh, D.; Kirkland, A. I.; Zandbergen, H. W. Adv. Mater. 2009, 21, 4992–4995. (22) Robinson, R. D.; Sadtler, B.; Demchenko, D. O.; Erdonmez, C. K.; Wang, L.-W.; Alivisatos, A. P. Science 2007, 317, 355–358. (23) Kresse, G.; Furthm€uller, J. Comput. Mater. Sci. 1996, 6, 15–50. (24) Kresse, G.; Furthm€uller, J. Phys. Rev. B 1996, 54, 11169–11186. (25) Perdew, J. P.; Burke, K.; Ernzerhof, M. Phys. Rev. Lett. 1996, 77, 3865–3868. (26) Luther, J. M.; Zheng, H. M.; Sadtler, B.; Alivisatos, A. P. J. Am. Chem. Soc. 2009, 131, 16851. (27) Sahu, A.; Qi, L. J.; Kang, M. S.; Deng, D.; Norris, D. J. J. Am. Chem. Soc. 2011, 133, 6509–6512. (28) Li, H.; Zanella, M.; Genovese, A.; Povia, M.; Falqui, A.; Giannini, C.; Manna, L. Nano Lett. 2011, DOI: 10.1021/nl202927a. (29) Barnard, A. S.; Young, N. P.; Kirkland, A. I.; van Huis, M. A.; Xu, H. ACS Nano 2009, 3, 1431–1436. (30) Bryant, F. J.; Cox, A. F. J. J. Phys. C: Solid State Phys. 1968, 1, 1734–1736. (31) Loginov, Y. Y.; Brown, P. D.; Thompson, N. Phys. Status Solidi A 1991, 126, 63–83. (32) Loginov, Y. Y.; Brown, P. D.; Thompson, N. Phys. Status Solidi A 1992, 132, 323–337. (33) Baranov, D.; Fiore, A.; van Huis, M.; Giannini, C.; Falqui, A.; Bertoni, G.; Lafont, U.; Zandbergen, H. W.; Cingolani, R.; Manna, L. Nano Lett. 2010, 10, 743–749. (34) Grodzinska, D.; Pietra, F.; van Huis, M. A.; Vanmaekelbergh, D.; de Mello Donega, C. J. Mater. Chem. 2011, 21, 11556–11565.
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