Strong and Tunable Blue Luminescence from Cd1-xZnxS Alloy

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Strong and Tunable Blue Luminescence from Cd1‑xZnxS Alloy Nanocrystallites Grown in Langmuir−Blodgett Multilayers Pavan K. Narayanam,† Purvesh Soni,† R. S. Srinivasa,‡ S. S. Talwar,† and S. S. Major†,* †

Department of Physics and ‡Department of Metallurgical Engineering and Materials Science, Indian Institute of Technology Bombay, Mumbai-400076, India ABSTRACT: Cd1‑xZnxS alloy nanocrystallites were grown in organic layered matrix by H2S exposure of precursor CdA-ZnA Langmuir−Blodgett multilayers at room temperature. The multilayers containing the alloy nanocrystallites were subsequently subjected to intercalation with Cd2+ ions followed by heat treatment up to 300 °C. With variation of the Zn content of precursor multilayers, the alloy nanocrystallites exhibit changes in lattice constants as well as a shift of excitonic absorption and emission between the extreme values for CdS and ZnS. The alloy nanocrystallites also exhibit treatment process dependent characteristic changes in optical absorption and luminescence, and the optical data have been used to propose suitable energy level diagrams. A strong enhancement of blue excitonic emission and suppression of emission due to bulk and surface defects are seen after postsulphidation treatments of the multilayers containing alloy nanocrystallites. These studies have shown that the organic moieties encapsulating the Cd1‑xZnxS nanocrystallites tend to restrict their growth and aggregation, while the presence of cadmium species in the organic matrix in proximity with the nanocrystallites is responsible for the passivation of surface defects as well as the reduction of bulk defects.

1. INTRODUCTION Semiconducting chalcogenide nanocrystallites have been extensively investigated during the past two decades, owing primarily to their size dependent and tunable luminescence behavior and promising applications in optoelectronics.1,2 Tunability of their optical properties has also been attempted by changing the composition of ternary chalcogenide nanocrystallites, which results in the extension of their emission into spectral regions that are not easily accessible through the size variation of binary quantum dots.3−5 CdZnS is one such ternary compound that forms a continuous series of solid solutions and allows a controlled and systematic variation of the optical band gap from 2.42 eV (CdS) to 3.7 eV (ZnS). Owing to their high photosensitivity, CdZnS thin films have been widely used as photodetectors.6−8 CdZnS has also been used as a window layer in thin film heterojunction solar cells, primarily due to the larger band gap, reduced absorption losses, and lattice matching with various quaternary semiconductors.9−11 Resistive switching in CdZnS thin films has been studied for nonvolatile memory applications.12 In recent years, nanostructured CdZnS © XXXX American Chemical Society

has also attracted a lot of interest. CdZnS nanowires have been investigated for their field emission behavior13 and CdZnS nanoribbons have exhibited wavelength controlled lasing in UV-spectral region.14 CdZnS nanocrystallites have been reported15 to exhibit up-conversion luminescence, which has important applications in biological imaging with reduced tissue photodamage, phototoxicity and photobleaching.16,17 High quality CdZnS nanocrystallites have also been recently investigated for potential applications as photocatalysts18 and blue emitters.19 Various synthetic approaches have been used for the preparation of CdZnS nanocrystallites in liquid media, such as, coprecipitation,20 colloidal method,15 chemical reduction route,21 reverse micelles22 and solvothermal synthesis.23 There are also a few reports on the preparation of CdZnS nanocrystallites in solid state media by different Received: December 20, 2012 Revised: January 29, 2013

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dx.doi.org/10.1021/jp312546a | J. Phys. Chem. C XXXX, XXX, XXX−XXX

The Journal of Physical Chemistry C

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was spread on the subphase, which was maintained at 10 ± 1 °C. The monolayer was compressed at a constant barrier speed of 3 mm/min and was transferred at a surface pressure of 30 mN/m. LB deposition was carried out by vertical dipping and lifting method at a speed of 3 mm/min. Fused quartz, CaF2 substrates, and carbon coated copper grids were used to transfer typically 15 monolayers of CdA/ZnA/CdA-ZnA in each multilayer. The CdA-ZnA multilayers will be referred to as Cd40Zn60A and Cd20Zn80A multilayers, hereafter, based on the corresponding metal ion concentrations in the subphase. The precursor multilayers were exposed to H2S gas for 5 min at a constant flow rate at room temperature to form sulphide nanocrystallites, as described in detail elsewhere.26 These multilayers will be referred to as “as-sulphided” multilayers in the rest of this work. The as-sulphided multilayers were immersed in aqueous CdCl2 solution for 60 min and subsequently washed in deionized water for the next 60 min to remove excess cadmium ions. These multilayers will be referred to as “intercalated” multilayers in the rest of this work. Heat treatment studies of intercalated multilayers were carried out in the temperature range of 100−300 °C. The structure, chemical composition, surface morphology, and optical properties of the multilayers were investigated at different stages of processing, namely, sulphidation, intercalation, and heat treatments. The chemical composition was studied by FT−IR spectroscopy, using Perkin-Elmer Spectrum One instrument in the wavenumber range of 1300−3000 cm−1. The layered structure of the composite multilayers was studied by X-ray reflection (XR), using an X’pert Pro diffractometer with Cu Kα radiation in the 2θ range of 4−20°. XPS studies were carried out using a Thermofisher-VG Scientific (Multilab 2000) photoelectron spectrometer equipped with 150 W monochromatic Al Kα X-ray source. The phase and size of nanocrystallites were determined by HRTEM, using a JEOL model JEM-2100F HRTEM instrument operated at 200 kV. AFM measurements were carried out in tapping mode using a Nanoscope IV multimode SPM to study the changes in surface morphology. Perkin-Elmer Lambda 950 UV−vis spectrometer was used to study the absorption of the sulphide nanocrystallites in the wavelength range of 200−1100 nm. JobinYvon iHR550 monochromator and Kimmon He−Cd laser (λ = 325 nm) were used to study the room temperature photoluminescence from sulphide nanocrystallites in the wavelength range of 350−800 nm.

methods, including self-assembly18 and Langmuir−Blodgett (LB) route.24,25 Langmuir−Blodgett technique has been extensively used for the preparation of binary sulphide nanocrystallites, such as, CdS26−28 by postdeposition H2S exposure of precursor cadmium arachidate/stearate multilayers. The size, shape and distribution of CdS nanocrystallites have been reported to be significantly influenced by the presence of layered structure and molecular order in the LB multilayer.26,29−31 This method of preparation has also been extended to grow ternary sulphide nanocrystallites such as CdMnS 32 and CdZnS24,25 by sulphidation of the corresponding mixed fatty acid salt LB multilayers. It has been shown25 that the formation of alloy sulphide nanocrystallites begins in the first 5 min of H2S exposure and is completed in ∼3 h. The partial sulphidation of about 5 min results in the formation of nanocrystallites of relatively smaller size, without causing substantial damage to the layered structure. Although the optical properties of CdZnS nanocrystallites in solid state matrix18,24,25,33 have been studied to some extent, the luminescence of CdZnS nanocrystallites formed in the LB layered matrix, its dependence on the composition and microstructure of the composite multilayer and its enhancement by surface modification and postgrowth treatments have not been investigated. Recent studies34,35 on heat treatment of CdS nanocrystallites grown in LB multilayer matrix have shown that the availability of cadmium species within the organic matrix results in substantial enhancement of photoluminescence owing to the improvement in crystallinity and suppression of defect luminescence due to bulk and surface trap states. However, such studies have not been attempted on alloy sulphide nanocrystallites. In the present work, CdZnS alloy nanocrystallites have been grown by H2S exposure of mixed CdA-ZnA LB multilayers, which were prepared by varying the composition of Cd and Zn ions in the subphase. The composite multilayers containing CdZnS nanocrystallites were subjected to intercalation with aqueous CdCl2 solution followed by heat treatment up to a temperature of 300 °C in air. The effect of postsulphidation treatments on the optical absorption and luminescence of CdZnS nanocrystallites has been investigated and analyzed in the light of the concurrent changes in the composition, microstructure, and surface morphology of composite multilayers. These results have been used to develop a comprehensive understanding of the influence of the organic matrix and the presence of cadmium species within it on the optical behavior, particularly the luminescence of CdZnS alloy nanocrystallites.

3. RESULTS AND DISCUSSION The changes in chemical composition of CdA-ZnA multilayers after sulphidation, intercalation, and subsequent heat treatment have been studied by FT−IR spectroscopy, and the results are shown in Figure 1. The assignments of various bands described below are broadly based on earlier reports on CdA and ZnA LB multilayers.36,37 The spectra of both Cd40Zn60A and Cd20Zn80A multilayers show characteristic vibrational bands due to CH3 (2955 cm−1) and CH2 (2918 cm−1 and 2850 cm−1) stretching modes. The asymmetric stretching band of carboxylate (COO−) group appears as a single peak in the range of 1541−1542 cm−1, which is between the extreme values of 1546 cm−1 for CdA38 and 1538 cm−1 for ZnA.37 For Cd40Zn60A multilayer, the CH2 scissoring vibration mode appears as a doublet (1473 and 1463 cm−1), which is characteristic of the herringbone type molecular packing of CdA multilayer.39 However, in the case of the Cd20Zn80A multilayer, it appears as a dominantly single peak at ∼1464

2. EXPERIMENTAL SECTION LB multilayers of cadmium arachidate-zinc arachidate (CdAZnA) were prepared using a KSV 3000 LB instrument. The subphase was prepared using aqueous solutions of CdCl2 and/ or ZnCl2 in ultrapure deionized water (resistivity: 18.2 MΩcm) with total salt concentration of 5 × 10−4 M. The pH of subphase was maintained at 6.3 ± 0.1 with dilute HCl/ NaHCO3. The composition of subphase was varied with different cadmium and zinc ion concentrations, namely, 0 mol % (CdA), 60 mol %, 80 mol %, and 100 mol % (ZnA) of Zn in aqueous solution. The Cd and Zn concentrations as well as the deposition conditions were chosen on the basis of earlier work from our group on the preparation and characterization of CdA-ZnA mixed LB multilayers.25 150 μL solution (1 mg/mL) of arachidic acid (Aldrich, 99%) in chloroform (HPLC grade) B



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Figure 1. FT−IR spectra of (a) Cd40Zn60A and (b) Cd20Zn80A multilayers transferred on CaF2 substrates: (i) as-deposited, (ii) after sulphidation, (iii) after intercalation and after subsequent heat treatment at (iv) 100 °C, (v) 150 °C, (vi) 200 °C, (vii) 250 °C, and (viii) 300 °C. Figure 2. XR patterns of (a) Cd40Zn60A and (b) Cd20Zn80A multilayers transferred on quartz substrates: (i) as-deposited, (ii) after sulphidation, (iii) after intercalation and after subsequent heat treatment at (iv) 100 °C and (v) 150 °C. The (00l) peaks refer to a bilayer period of 55.2 Å and the (00l)′ peaks refer to a bilayer period of 51.6 Å. The (003) peak of ZnA (bilayer period, 47 Å) is also indicated.

cm−1, which has been attributed to “loose molecular packing in a hexagonal layer cell” of ZnA multilayer.38,39 These results indicate that with an increase in Zn fraction, the molecular packing of CdA-ZnA multilayers is significantly altered due to changes in metal ion coordination with COO− at the headgroup interfaces. It is also noticed from the abovedescribed FT−IR results that in both cases, separate molecular domains with characteristic spectral features of CdA and ZnA are not seen, and hence, it is inferred that molecular level mixing of CdA and ZnA molecules is prevalent in the mixed CdA-ZnA multilayers. After 5 min H2S exposure, the CdA-ZnA multilayers show a substantial decrease in intensity of the carboxylate peak and appearance of a peak at ∼1700 cm−1 due to carbonyl stretching of carboxylic acid group. These changes are attributed to formation of arachidic acid, during the partial sulphidation of arachidate molecules, as reported earlier.25 After intercalation of the partially sulphided multilayers with Cd2+ ions, the carbonyl peak nearly disappears and a significant increase in the intensity of carboxylate peak is observed. Interestingly, the carboxylate peak after intercalation appears in the range of 1544−1545 cm−1. The shift toward 1546 cm−1 is attributed to the increase in cadmium content of the multilayers and formation of CdA during the intercalation process. Subsequent heat treatment in the range of 100−300 °C results in a monotonous decrease of the intensities of all of the vibrational bands of both Cd40Zn60A and Cd20Zn80A multilayers. This is indicative of the progressive desorption of organic moieties, which is nearly complete at 300 °C. The changes in the layered structure of CdA-ZnA mixed multilayers due to sulphidation, intercalation, and subsequent heat treatment have been studied by X-ray reflection (XR). Figure 2 shows the corresponding XR patterns. The XR pattern of as-deposited Cd40Zn60A multilayer shows well-defined (00l) Bragg peaks (third order onward), which appear as single peaks corresponding to an average bilayer period of 55.2 Å. These features indicate that the layered structure of Cd40Zn60A multilayer is similar to that of a Y-type CdA LB multilayer in which the alkyl chains are nearly perpendicular to the layer plane.39 The presence of a single type of layered

structure is attributed to the molecular level mixing of CdA and ZnA molecules, as inferred above from FT−IR results. After 5 min H2S exposure (partial sulphidation), the XR pattern of Cd40Zn60A multilayer shows a decrease in the intensities of Bragg peaks along with the appearance of shoulders on the larger 2θ side. These features are indicative of reduction of the structural order of the multilayers and the formation of molecular domains with tilted alkyl chains, as reported earlier in the case of sulphided CdA multilayers.26 After intercalation with Cd2+ ions, although the intensities of Bragg peaks decrease substantially, the layered structure corresponding to the bilayer period of 55.2 Å is restored, which is attributed to the formation of CdA during the intercalation process (as indicated by FT−IR results) and its influence on the layered structure. After heat treatment of the intercalated multilayer at 100 °C, the intensities of the Bragg peaks reduce drastically and a small peak appears at 2θ = 5.6°, which corresponds to the third order Bragg peak of the reported layered structure of ZnA, in which the arachidate molecules are tilted at ∼32° with respect to the layer normal. These features indicate a substantial reduction of the structural order of the multilayer along with the formation of molecular domains of ZnA molecules after heat treatment at 100 °C.39 On further heat treatment at 150 °C, no Bragg peaks are seen, indicating a complete destruction of the layered structure. The structural changes seen after heat treatment are attributed to melting and recrystallization of cadmium and zinc arachidate molecules, which are reported to melt in the temperature range of 100−130 °C.40,41 Figure 2(b) shows the changes in the structure of Cd20Zn80A multilayer during the above processing steps. The XR pattern of as-deposited Cd20Zn80A multilayer shows a single type of layered structure corresponding to an average bilayer period of 51.6 Å, which is attributed to molecular level C



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mixing of CdA and ZnA molecules packed with alkyl chains tilted at ∼23° with respect to the layer normal, as reported earlier.38 As in the case of Cd40Zn60A multilayer, partial sulphidation results in significantly smaller and broader Bragg peaks, indicating substantial reduction of the structural order and formation of molecular domains with alkyl chains having a distribution of tilt angles about 23°. Interestingly, after intercalation with Cd2+ ions, the Bragg peaks become sharper and represent an average bilayer period of 55.2 Å, corresponding to the layered structure of Y-type CdA. Thus in both Cd40Zn60A and Cd20Zn80A multilayers, the formation of CdA during intercalation has a dominating influence on the layered structure of the composite multilayer. After subsequent heat treatments at 100 and 150 °C, the intercalated Cd20Zn80A multilayer exhibits a behavior similar to that observed in the case of Cd40Zn60A multilayer. Figure 3 shows the Cd-3d, Zn-2p, and S-2p core level XPS spectra of sulphided multilayers. The sulphided CdA multilayer shows the Cd-3d5/2 peak at 405.0 eV, which is within the range of values (405.0−405.5 eV) reported for CdS.42,43 Similarly, the sulphided ZnA multilayer shows the Zn-2p3/2 peak at 1022.4 eV, which is within the range of values (1021.9−1022.6 eV) reported for ZnS.42,44,45 For both the sulphided CdA-ZnA multilayers, the binding energies of Cd-3d5/2 and Zn-2p3/2 peaks are in the range of 404.6−405.2 eV and 1022.0−1022.6 eV, respectively. The positions of the S-2p peaks in all of the cases are in the range of 161.2−161.8 eV, which are within the range of reported values for CdS (161.1 eV) and ZnS (162.3 eV).42,44−47 These results indicate the formation of sulphide during H2S exposure of CdA-ZnA multilayers. The Zn (cationic) fractions in the sulphided multilayers were also estimated from XPS data and are listed in Table 1. The Zn fractions of 28% for Cd40Zn60A and 54% for Cd20Zn80A are significantly lower than the corresponding molecular fraction of ZnA in the subphase. This is attributed to the preferential transfer of Cd over Zn from a subphase containing both cadmium and zinc ions, as reported earlier.24 The significantly larger uptake of CdA relative to ZnA is also seen as the reason for the dominating influence of CdA molecules in determining the layered structure of the composite multilayers, as the abovedescribed XR results have shown. It may however be pointed out that the nature of molecular packing in Cd40Zn60A multilayers (as evidenced by the FT−IR results) is strongly dominated by CdA molecules, while in Cd20Zn80A multilayers, although it is dominated by ZnA molecules, marginal influence of CdA molecules is observed, as reported earlier.38 The size, distribution, and structure of sulphide nanocrystallites formed in the CdA, ZnA, and CdA-ZnA multilayers were studied by HRTEM. In all of the cases, the low resolution micrographs of sulphided multilayers (Figure 4) show the presence of nanocrystallites of size ∼3 nm, distributed uniformly within the organic matrix. Typical lattice images of nanocrystallites are also shown as inserts. The electron diffraction patterns (not shown) as well as direct measurement of d-spacings on high resolution images were used to identify the phase of sulphide nanocrystallites and the corresponding values of lattice constants (“a” and “c”) are listed in Table 1. The lattice constants of both CdS and ZnS are in good agreement with the standard values for the hexagonal wurtzite phases of CdS (a = 4.12 Å, c = 6.68 Å, JCPDS file 80−0006) and ZnS (a = 3.78 Å, c = 6.18 Å, JCPDS file 80−0007). Table 1 also shows that the lattice constants of the nanocrystallites formed in CdA-ZnA multilayers shift monotonously from CdS

Figure 3. XPS spectra showing (a) Cd-3d core levels, (b) Zn-2p core levels, and (c) S-2p core level for CdA (solid curve), Cd40Zn60A (dash dotted curve), Cd20Zn80A (dotted curve), and ZnA (dashed curve) multilayers after sulphidation.

values to ZnS values with increase in Zn content of the multilayers. These results indicate that the formation of alloy sulphide nanocrystallites with wurtzite phase takes place during H2S exposure at room temperature. The lattice constants of alloy sulphide nanocrystallites have been used to estimate the corresponding alloy compositions, assuming Vegard’s law48 and the estimated values of Zn fractions are listed in Table 1. These values are in reasonable agreement with those obtained from XPS data, considering the large uncertainties (±10%) in the measured values of lattice constants. The changes in surface morphology of Cd40Zn60A and Cd20Zn80A multilayers after sulphidation, intercalation, and subsequent heat treatments were studied by AFM. Figure 5 shows that the as-sulphided Cd40Zn60A multilayer has a uniformly porous surface. The depths of pores range from a D



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Table 1. Zn (Cationic) Fraction in Cd40Zn60A and Cd20Zn80A Multilayers after Sulphidationa multilayer CdA Cd40Zn60A Cd20Zn80A ZnA

Zn (mol. %) in subphase 60 80

Zn (cationic) fraction (XPS)

“a”(Å)

“c”(Å)

Zn (cationic) fraction (HRTEM) (±10%)

± ± ± ±

6.69 ± 0.03 6.64 ± 0.35 6.42 ± 0.24 6.12

0.08 0.54

4.15 4.09 3.93 3.82

0.28 0.54

0.03 0.09 0.21 0.02

Lattice constants, “a” and “c” obtained from HRTEM studies are listed for the mixed multilayers along with those for CdA and ZnA multilayers, for the purpose of comparison.

a

significantly different surface morphological features. Instead of the spiral- and wire-like features, the Cd20Zn80A multilayer shows large lumps or agglomerates of organic moieties on the surface, which desorb with increase in heat treatment temperature. Interestingly, after heat treatment at 300 °C, particle-like features are seen on the surface, as in the case of Cd40Zn60A multilayer, which are also attributed to the presence of alloy sulphide. Figure 7 shows the absorbance spectra of the multilayers before and after H2S exposure. All of the H2S exposed multilayers exhibit a substantial increase of absorbance in the shorter wavelength region (below 450 nm), which is attributed to sulphide formation. The spectra of the sulphided multilayers also show the presence of a small hump in the absorbing region. The corresponding derivative spectra have been used to estimate the positions of the humps, which are listed in Table 2. For the sulphided CdA and ZnA multilayers, the humps at 400 and 296 nm are respectively ascribed to the excitonic absorption of sulphide nanocrystallites.49,50 The position of the excitonic absorption hump may be considered to represent the first e−h state and thus the excitonic energy gap (separation between the ground state and the first e−h state) of CdS and ZnS nanocrystallites can be nominally taken as 3.1 and 4.2 eV, respectively. Following the same approach, the humps at 374 and 351 nm in the absorption spectra of sulphided Cd40Zn60A and Cd20Zn80A multilayers are also attributed to excitonic absorption of Cd1‑xZnxS alloy nanocrystallites with respective excitonic energy gaps of 3.3 and 3.5 eV. Since TEM studies have shown that the average size of nanocrystallites is nearly the same (∼3 nm) in all of the cases, the increase in excitonic energy gap of Cd1‑xZnxS nanocrystallites between the extreme values of 3.1 eV (CdS) and 4.2 eV (ZnS) is ascribed to the increase in their Zn content. This has been taken into account while proposing the energy level diagrams for Cd1‑xZnxS nanocrystallites, which are shown in Figure 8. The excitonic energy gaps of the alloy sulphide nanocrystallites formed in Cd40Zn60A and Cd20Zn80A multilayers have accordingly been shown as 3.3 and 3.5 eV, respectively. The other features of the energy level diagrams will be discussed later in this work. The emission behavior of Cd1‑xZnxS nanocrystallites was studied by room temperature PL spectroscopy. Figure 9 shows the PL spectra of the sulphided CdA, Cd40Zn60A, and Cd20Zn80A multilayers. It was not possible to obtain PL signal from the sulphided ZnA multilayer, due to limitations of excitation wavelength (325 nm) of He−Cd laser. All of the three spectra show a broad emission band centered at ∼500 nm and a weak emission shoulder at ∼400 nm. In order to gain deeper insight into the nature of the emission, the PL spectra have been deconvoluted by fitting multiple Gaussian peaks using a nonlinear curve-fitting program. The results of deconvolution are included in Figure 9 and the corresponding peak positions are listed in Table 2. In all of the cases, the weak

Figure 4. HRTEM images of (a) CdA, (b) Cd40Zn60A, (c) Cd20Zn80A, and (d) ZnA multilayers after sulphidation. The insert in each case shows a high resolution image of an individual nanocrystallites.

monolayer to bilayer thickness. After intercalation, rod-like features appear on the surface, which are attributed to the formation of CdA (indicated by FT−IR results) during the intercalation process. Heat treatment at 100 °C (not shown) does not result in any significant change in the rod-like surface features. After heat treatment at 150 °C, curved spiral-like features appear on the surface with heights in the range of 20− 50 nm, which are attributed to recrystallization of the organic moieties. After heat treatment at 200 °C, the spiral-like features are transformed into uniformly distributed wire-like features, which appear to partially desorb after heat treatment at 250 °C. After heat treatment at 300 °C, particle like features of size 50− 100 nm are seen, which are attributed to the presence of alloy sulphide on the surface. This inference is based on the FT−IR observation that the organic moieties practically disappear at this temperature and is further supported by UV−vis studies (presented later) which show the presence of sulphide at this temperature. Figure 6 shows that the as-sulphided Cd20Zn80A multilayer also has a porous surface, but with slightly larger average size of pores compared to those seen in the case of as-sulphided Cd40Zn60A multilayer. After intercalation and subsequent heat treatment at 100 °C (not shown), the morphological changes are similar to those observed in the case of Cd40Zn60A multilayer. However, after heat treatment in the temperature range of 150−250 °C, the Cd20Zn80A multilayer shows E



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Figure 5. AFM images of Cd40Zn60A multilayer after (a) sulphidation, (b) intercalation and after subsequent heat treatment at (c) 150 °C, (d) 200 °C, (e) 250 °C, and (f) 300 °C.

shoulder has contributions from a small component at ∼390 nm due to quartz, along with a dominant component which peaks at 415 nm for CdS, but shifts to shorter wavelengths for the Cd1‑xZnxS nanocrystallites. The blue shift of the emission peak with increase in Zn content of the nanocrystallites is consistent with the corresponding blue shift of the excitonic absorption peak. Hence this component of luminescence is attributed to excitonic emission from CdS and Cd1‑xZnxS alloy nanocrystallites. The blue shifts observed for both excitonic absorption and emission peaks reiterate that the excitonic energy gap of alloy nanocrystallites increases monotonously with an increase in Zn content. Table 2 also shows that the Stokes shifts for alloy nanocrystallites (∼300 meV) are significantly larger than that for CdS (120 meV) nanocrystallites. The larger Stokes shifts for alloy nanocrystallites may be attributed to structural and compositional inhomogeneities.51 Figure 9 also shows that in all the cases, the broad defect emission band centered at ∼500 nm has two components. The dominant component peaks at 533 nm in case of CdS nanocrystallites and shifts to lower wavelengths of 516 and 480 nm as the Zn content of the alloy nanocrystallites increases. The weak and broad component centered at 640 nm for CdS nanocrystallites also shifts to lower wavelengths of 619 and 591 nm with increasing Zn content of alloy nanocrystallites. The

emission at ∼600 nm and larger wavelengths has usually been attributed to the presence of dangling bonds and sulphide ions at the surface of CdS52,53 and ZnS47,54 nanocrystallites. Hence, this component of defect emission is attributed to radiative transitions from conduction band edge to surface trap states, as shown in Figure 8. The blue shift of this defect emission component from 640 nm (CdS) to 591 nm with increase in the Zn content of alloy nanocrystallites is attributed primarily to the increase in excitonic energy gap of the nanocrystallites. Extensive luminescence studies on bulk and thin films of CdS55−59 have shown that the prominent and broad defectrelated emission peaking at ∼550 nm, arises from the presence of native defects, such as interstitials and vacancies of cadmium and sulfur. The broad defect emission peaking at 533 nm for CdS nanocrystallites is thus attributed to bulk defects, as reported earlier.34,56,57 It has also been argued34 that during the sulphidation process of LB multilayers, there is unlimited availability of H2S gas but limited availability of Cd at the headgroup interfaces, where sulphide nanocrystallites are formed. This is due to the limited concentration of Cd present in the multilayers as well as limited possibility of its diffusion at room temperature. It can therefore be assumed that in the case of alloy sulphide nanocrystallites also, cationic (Cd or Zn) vacancies are more likely to be present, rather than sulfur vacancies. Further, considering that the nanocrystallites are F



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Figure 6. AFM images of Cd20Zn80A multilayer after (a) sulphidation, (b) intercalation, and after subsequent heat treatment at (c) 150 °C, (d) 200 °C, (e) 250 °C, and (f) 300 °C.

formed at room temperature, the presence of cationic interstitials is also not ruled out, at least in the case of assulphided multilayers. For bulk and thin films of CdS, the shallow donor states due to cadmium interstitials are reported to be located 0.12−0.26 eV below the conduction band,60,61 while the acceptor states due to cadmium vacancies have been reported to be located 0.26−1.2 eV above the valence band.62−64 The radiative transition between cadmium interstitial states to cadmium vacancy states has been reported prominently in CdS thin films.55,57,58 It may be emphasized here that being deep trap states, the defect levels due to cationic vacancies in nanocrystallites are expected to remain practically unchanged during the increase in the excitonic energy gap due to decrease of particle size.65 In contrast, the shallow donor levels are expected to shift to higher energies with the conduction band edge, owing essentially to certain degree of spatial delocalization.65 On the basis of the above inputs, in the case of Cd1‑xZnxS nanocrystallites, if the defect states due to cationic interstitials are nominally assumed to be ∼0.1 eV below the conduction band edge, then the dominant defect emissions (2.4 and 2.6 eV for the two alloy compositions) can be attributed to radiative transitions between cationic interstitial states and cationic vacancy states, as shown in Figure 8. The blue shift of the dominant defect emission from 2.4 to 2.6 eV with increase in

Zn content of alloy nanocrystallites is a consequence of the increase in the excitonic energy gap of nanocrystallites, which is similar to the behavior of emission peak due to surface trap states. The positions of the defect states due to cationic vacancies for both alloy compositions are thus estimated to be ∼0.8 eV above the valence band edge, as shown in Figure 8. This is in reasonable agreement with the reported positions of cationic vacancy states in CdS (0.26−1.2 eV)34,62−64 and ZnS (0.62−1.2 eV).66,67 The effect of intercalation and subsequent heat treatment on the absorption and emission of Cd1‑xZnxS nanocrystallites formed in the organic matrix has also been investigated. Figures 10 and 11 show the absorbance and corresponding derivative spectra of the two alloy compositions after intercalation and subsequent heat treatment in the temperature range of 150− 300 °C. The corresponding positions of the excitonic absorption peaks are listed in Table 2, which show that the alloy nanocrystallites of both compositions exhibit similar behavior. After intercalation, a marginal increase in absorption and a small red shift of the excitonic absorption peak are seen in both cases. After heat treatment at 150 °C, the excitonic absorption peaks red shift further, and also become relatively sharper. After heat treatments at 200, 250, and 300 °C, the absorbance due to the presence of sulphide remains unchanged, but the excitonic absorption peaks red shift monotonously with G



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peaks also exhibit red shift, consistent with the red shifts of the corresponding excitonic absorption peaks. This is attributed to an increase in Cd-content of the alloy nanocrystallites during intercalation, although marginal contribution due to the increase in average size of nanocrystallites (which could not be ascertained by TEM measurements) cannot be ruled out. Heat treatment of the intercalated multilayers at 150 °C [Figures 12(c) and 13(c)], results in substantial enhancement and narrowing of the excitonic emission peaks in both cases, along with small but significant blue shifts, as compared to the intercalated state (Table 2). It may be recalled from the corresponding absorption spectra [Figures 10(c) and 11(c)] that after heat treatment at 150 °C, the alloy nanocrystallites show prominent and relatively sharp excitonic absorption peaks, which are slightly red-shifted with respect to their positions after intercalation (Table 2). As a consequence, the values of Stokes shift (Table 2) for alloy nanocrystallites of both compositions decrease significantly after heat treatment at 150 °C. The Stokes shifts for both alloy compositions are also found to remain nearly unchanged after heat treatments at higher temperatures. It is inferred from these observations that heat treatment of the intercalated multilayers at 150 °C results in the formation of structurally and compositionally homogeneous alloy nanocrystallites having a relatively narrow size distribution. After heat treatment at 200 °C, the alloy nanocrystallites of both compositions show substantial enhancement of excitonic emission along with a large red shift. The red shifts of excitonic emission peaks are consistent with the red shift of corresponding excitonic absorption peaks (for both alloy compositions, as seen from Table 2). This behavior is attributed to a decrease in the band gap of the nanocrystallites due to aggregation and particle growth. It may be recalled from FT− IR and AFM studies, that the organic moieties begin to desorb substantially at about this temperature. It is thus inferred that the removal of organic moieties facilitates the aggregation and growth of nanocrystallites. Heat treatment at 250 °C results in a drastic reduction and red shift of the excitonic emission of the alloy nanocrystallites formed in the Cd40Zn60A multilayer. In the case of nanocrystallites formed in the Cd20Zn80A multilayer, after heat treatment at 250 °C, the emission is practically from quartz, as is the case for both compositions after heat treatment at 300 °C (not shown). Although the absorption spectra at this stage continue to show the presence

Figure 7. UV−vis absorption spectra along with the corresponding derivative spectra of (a) CdA, (b) CdA40Zn60A, (c) CdA20Zn80A, and (d) ZnA multilayers after sulphidation. The vertical arrows shown on the derivative spectra in all cases indicate the excitonic absorption peak positions. The dashed curves in all of the cases show the absorption spectra of the corresponding as-transferred multilayers.

an increase in heat treatment temperature. It may however be noted from Table 2 that at all stages of postsulphidation treatments, the positions of excitonic absorption peaks appear at lower wavelengths for the alloy nanocrystallites with larger Zn content. Figures 12 and 13 show the PL spectra of Cd1‑xZnxS nanocrystallites after intercalation and subsequent heat treatment in the temperature range of 150−250 °C. After intercalation, a significant enhancement of excitonic emission is observed for both alloy compositions. The excitonic emission

Table 2. Excitonic Absorption and Emission Peak Positions, Stokes Shifts and Defect Emission Peak Positions for CdA, Cd40Zn60A and Cd20Zn80A Multilayers after Different Stages of Processing sample description CdA Cd40Zn60A

Cd20Zn80A

sulphidation sulphidation intercalation 150 °C 200 °C 250 °C 300 °C sulphidation intercalation 150 °C 200 °C 250 °C 300 °C

excitonic absorption peak (nm) [photon energy] (eV) 400 374 378 384 404 418 439 351 354 370 388 415 430

excitonic emission peak (nm) [photon energy] (eV)

[3.11] [3.32] [3.29] [3.23] [3.07] [2.97] [2.83] [3.54] [3.51] [3.36] [3.20] [2.99] [2.89] H

Stokes shift (meV)

bulk defect emission peak (nm) [photon energy] (eV)

415 408 411 403 419 435

[2.99] [3.04] [3.02] [3.08] [2.96] [2.86]

120 280 270 150 110 110

533 516 490 501 531 549

[2.33] [2.41] [2.53] [2.48] [2.34] [2.26]

395 399 391 410

[3.14] [3.11] [3.18] [3.03]

400 400 180 170

480 464 476 499

[2.59] [2.68] [2.61] [2.49]



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Figure 8. Proposed energy level schemes for alloy sulphide nanocrystallites formed in (a) Cd40Zn60A and (b) Cd20Zn80A multilayers. The numbers in parentheses give the corresponding wavelengths in “nm”.

Figure 9. Photoluminescence spectra (circles) of (a) CdA, (b) CdA40ZnA, and (c) CdA20ZnA multilayers after sulphidation. The corresponding deconvoluted components are shown as “dashed curves” and the consolidated curve of the deconvoluted components is shown as a “solid line”.

of sulphide in both cases, the drastic reduction of photoluminescence is attributed to significant particle growth due to nearly complete desorption of organic moieties, as indicated by FT−IR and AFM studies. This effect is more pronounced in the alloy nanocrystallites with larger Zn content, since in this case, the recrystallized organic moieties form lumps or agglomerates that may not be effective in restraining particle growth. The PL spectra of nanocrystallites after intercalation and heat treatment show interesting changes in defect emission. Figure 12 and Table 2 show that after intercalation, the dominant defect emission of alloy nanocrystallites formed in Cd40Zn60A multilayer blue shifts to 2.5 eV with respect to its position of 2.4 eV after sulphidation. Similarly, Figure 13 and Table 2 show that the dominant defect emission of alloy nanocrystallites

Figure 10. UV−vis absorption spectra of Cd40Zn60A multilayer after (a) sulphidation, (b) intercalation, and after subsequent heat treatment at (c) 150 °C, (d) 200 °C, (e) 250 °C, and (f) 300 °C.

formed in the Cd20Zn80A multilayer blue shifts to 2.7 eV with respect to its position of 2.6 eV after sulphidation. The blue shifts of the dominant defect emission in both cases can be explained by assuming the reduction/disappearance of shallow defect states due to cationic interstitials, during the intercalation process. This implies that after intercalation, the corresponding radiative transition takes place between the conduction band edge (instead of cationic interstitial states) and cationic vacancy states, as shown in Figure 8. Interestingly, no such blue shifts are observed for the longer wavelength defect emission, which I



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Figure 12. Photoluminescence spectra (circles) of Cd40Zn60A multilayer after (a) sulphidation, (b) intercalation, and after subsequent heat treatment at (c) 150 °C, (d) 200 °C, and (e) 250 °C. The corresponding deconvoluted components are shown as “dashed curves” and the consolidated curve of deconvoluted components is shown as “solid line”.

Figure 11. UV−vis absorption spectra of Cd20Zn80A multilayer after (a) sulphidation, (b) intercalation, and after subsequent heat treatment at (c) 150 °C, (d) 200 °C, (e) 250 °C, and (f) 300 °C.

causes the passivation of surface trap states and improvement in excitonic emission after intercalation. The much stronger enhancement of excitonic emission after heat treatment in the range of 150−200 °C is attributed to effective suppression of defect emission and nonradiative transitions due to improvement in the crystalline quality of the nanocrystallites and substantial reduction of bulk defects. It may be noted that the dissociation of CdA is known to take place at ∼180 °C.70 Heat treatment in this temperature range may thus result in the improved access of cadmium to the alloy nanocrystallites, which facilitates the reduction of cadmium vacancies. The smaller enhancement of excitonic emission seen in the case of nanocrystallites formed in Cd20Zn80A concurs with this view, since in this case, the substantially different morphology of recrystallized organic moieties may not facilitate the availability of Cd-species in close proximity with the nanocrystallites.

supports the view that the corresponding radiative transition takes place between the conduction band edge and the surface trap states. After heat treatment of the intercalated multilayers, the emission peaks due to bulk defects as well as surface trap states show monotonous red shifts with increase in heat treatment temperature, for both alloy compositions. This behavior is attributed to a decrease in the excitonic energy gap of alloy nanocrystallites due to aggregation and particle growth. The enhancement of the excitonic emission of alloy nanocrystallites after intercalation and heat treatment (Figures 12 and 13) up to 200 °C can be explained by assuming that the presence of cadmium in the organic matrix plays a key role in the passivation of surface defects and reduction of bulk defects. Surface modification of nanocrystallites and polycrystalline thin films of CdS by various cadmium sources is known to substantially improve their luminescence characteristics.52,55,68,69 The FT−IR spectra recorded after intercalation (Figure 1(iii)) have shown that during this process, nearly complete conversion of AA into CdA takes place. The alloy nanocrystallites, which were encapsulated largely by AA molecules formed in their vicinity during the sulphidation process, thus get surrounded entirely by CdA molecules after the intercalation process. It is inferred that the availability and proximity of cadmium species to the surface of nanocrystallites

4. CONCLUSIONS Postdeposition H2S exposure of CdA-ZnA LB multilayers results in the formation of Cd1‑xZnxS alloy nanocrystallites of average size ∼3 nm within the organic layered matrix. The formation of alloy nanocrystallites at room temperature is facilitated by molecular level mixing of CdA and ZnA molecules in the precursor multilayers. With increase in Zn content of the precursor multilayers, the alloy nanocrystallites exhibit J



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stages of processing have been used to propose energy level diagrams for the two alloy compositions studied in this work. The optical properties of Cd1‑xZnxS alloy nanocrystallites grown in CdA-ZnA multilayers are significantly influenced by the melting and desorption behavior of the organic matrix and availability of cadmium within it. During the postsulphidation treatment processes, the organic moieties primarily restrict the growth and aggregation of alloy nanocrystallites and the presence of cadmium in the organic matrix plays a key role in passivation of surface defects and reduction of bulk defects. The present work establishes the potential of the LB approach for the growth of ternary alloy nanocrystallites, in general, and fine-tuning of the excitonic energy gap of Cd1‑xZnxS nanocrystallites over a wide range of energy values, in particular, by manipulating the growth conditions and postsulphidation treatments.

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AUTHOR INFORMATION

Corresponding Author

*Tel: +91-22-25767567; fax: +91-22-25767552; e-mail: syed@ iitb.ac.in. Notes

The authors declare no competing financial interest.

■ Figure 13. Photoluminescence spectra (circles) of Cd20Zn80A multilayer after (a) sulphidation, (b) intercalation, and after subsequent heat treatment at (c) 150 °C, (d) 200 °C, and (e) 250 °C. The corresponding deconvoluted components are shown as “dashed curves” and the consolidated curve of deconvoluted components is shown as “solid line”.

ACKNOWLEDGMENTS This work was supported by Project No. INT/RFBR/P-58 of Department of Science and Technology (DST), Government of India under the MoU with Russian Foundation for Basic Research (RFBR). Prof. K.C. Rustagi of the Department of Physics, IIT Bombay is gratefully thanked for useful discussions. The authors thankfully acknowledge FIST (PHYSICS)-IRCC Central SPM Facility of IIT Bombay for AFM measurements, Central Surface Analytical Facility of IIT Bombay for XPS measurements and Centre for Research in Nanotechnology and Science of IIT Bombay for HRTEM measurements.

monotonous shift of lattice constants as well as blue shifts of excitonic absorption and emission between the extreme values for CdS and ZnS. Analysis of room temperature PL data suggests that the alloy nanocrystallites present in the assulphided multilayers exhibit a dominant broad defect emission contributed by radiative transitions from cationic interstitial states to cationic vacancy states in the bulk and from conduction band edge to surface trap states due to dangling bonds and sulphide ions. The intercalation of the sulphided multilayers with Cd2+ ions results in enhancement of excitonic emission and changes in defect emission, which indicate passivation of surface defects and disappearance of bulk defects due to cationic interstitials, as a result of which radiative transition from conduction band edge to cationic vacancies takes over as the dominant defect emission. Subsequent heat treatment of intercalated multilayers up to 200 °C results in a strong enhancement of the blue excitonic emission from alloy nanocrystallites and suppression of emission due to bulk defect states. Heat treatment at higher temperatures results in a strong red shift of both excitonic absorption and emission as well as a drastic decrease in the excitonic emission intensity, which is attributed to aggregation and particle growth, following the desorption of organic moieties at these temperatures. The characteristic changes observed in the absorbance and photoluminescence spectra of the alloy nanocrystallites at different

(1) Alivisatos, A. P. J. Phys. Chem. 1996, 100, 13226−13239. (2) Yao, W. T.; Yu, S. H. Adv. Funct. Mater. 2008, 18, 3357−3366. (3) Bailey, R. E.; Nie, S. J. Am. Chem. Soc. 2003, 125, 7100−7106. (4) Korgel, B. A.; Monbouquette, H. G. Langmuir 2000, 16, 3588− 3594. (5) Li, Y.; Ye, M.; Yang, C.; Li, X.; Li, Y. Adv. Funct. Mater. 2005, 15, 433−441. (6) Torres, J.; Gordillo, G. Thin Solid Films 1992, 207, 231−235. (7) Sebastian, P. J.; Gomez-Dasa, O.; Nair, P. K. J. Phys. D. 1992, 25, 1848−1850. (8) Averin, S. V.; Kuznetsov, P. I.; Zhitov, V. A.; Alkeev, N. V.; Kotov, V. M.; Zakharov, L. Y.; Gladysheva, N. B. Tech. Phys. 2012, 57 (11), 1514−1518. (9) Yamaguchi, T.; Yamamoto, Y.; Tanaka, T.; Demizu, Y.; Yoshida, A. Thin Solid Films 1996, 281, 375−378. (10) Oladeji, I. O.; Chow, L. Thin Solid Films 2005, 474, 77−83. (11) Bhattacharya, R. N.; Contreras, M. A.; Egaas, B.; Noufi, R. N. Appl. Phys. Lett. 2006, 89, 253503. (12) Zheng, W.; Griffin, P. B.; McVittie, J.; Wong, S.; McIntyre, P. C.; Nishi, Y. IEEE Electron Device Lett. 2007, 28 (1), 14−16. (13) Lin, Y. F.; Hsu, Y. J.; Lu, S. Y.; Chen, K. T.; Tseng, T. Y. J. Phys. Chem. C 2007, 111, 13418−13426. (14) Liu, Y.; Zapien, J. A.; Shan, Y. Y.; Geng, C. Y.; Lee, C. S.; Lee, S. T. Adv. Mater. 2005, 17, 1372−1377. (15) Ouyang, J.; Ripmeester, J. A.; Wu, X.; Kingston, D.; Yu, K.; Joly, A. G.; Chen, W. J. Phys. Chem. C 2007, 111, 16261−16266. (16) Mohanta, D.; Narayanan, S. S.; Pal, S. K.; Raychaudhuri, A. K. Eur. Phys. J. Appl. Phys. 2008, 43, 263−268.

■

K

REFERENCES



Article

(50) Rossetti, R.; Hull, R.; Gibson, J. M.; Brus, L. E. J. Chem. Phys. 1985, 82, 552−559. (51) Efros, A. L.; Rosen, M.; Kuno, M.; Nirmal, M.; Norris, D. J.; Bawendi, M. Phys. Rev. B 1996, 54, 4843−4856. (52) Spanhel, L.; Haase, M.; Weller, H.; Henglein, A. J. Am. Chem. Soc. 1987, 109, 5649−5655. (53) Kim, D.; Miyamoto, M.; Mishima, T.; Nakayama, M. J. Appl. Phys. 2005, 98, 083514. (54) Bhattacharjee, B.; Lu, C. H. Thin Solid Films 2006, 514, 132− 137. (55) Abken, A. E.; Halliday, D. P.; Durose, K. J. Appl. Phys. 2009, 105, 064515. (56) Agata, M.; Kurase, H.; Hayashi, S.; Yamamoto, K. Solid State Commun. 1990, 76, 1061−1065. (57) Vigil, O.; Riech, I.; Rocha, M. G.; Angel, Z. J. Vac. Sci. Technol. A 1997, 15, 2282−2286. (58) Morales, L. R.; Angel, Z. O.; Delgado, T. G. Appl. Phys. A: Mater. Sci. Process. 2001, 73, 61−65. (59) Tsai, C. T.; Chuu, D. S.; Chen, G. L.; Yang, S. L. J. Appl. Phys. 1996, 79, 9105−9109. (60) Mochizuki, K.; Satosh, M.; Igaki, K. J. J. Appl. Phys. 1983, 22, 1414−1417. (61) Hershman, G. H.; Krisger, F. A. J. Solid State Chem. 1970, 2, 483−490. (62) Ferrer, I. J.; Salvador, P. J. Appl. Phys. 1989, 66, 2568−2577. (63) Bube, R. H. J. Phy. Chem. Solids 1957, 1, 234−248. (64) Susa, N.; Watanabe, H.; Wada, M. J. J. Appl. Phys. 1976, 15, 2365−2370. (65) Chestnoy, N.; Harris, T. D.; Hull, R.; Brus, L. E. J. Phys. Chem. 1986, 90, 3393−3399. (66) Samelson, H.; Lempicki, A. Phys. Rev. 1962, 125 (3), 901−909. (67) Wang, X.; Shi, J.; Feng, Z.; Li, M.; Li, C. Phys. Chem. Chem. Phys. 2011, 13, 4715−4723. (68) Wang, C. W.; Moffitt, M. G. Langmuir 2004, 20, 11784−11796. (69) Wan, L.; Bai, Z.; Hou, Z.; Wang, D.; Sun, H.; Xiong, L. Thin Solid Films 2010, 518, 6858−6865. (70) Schreck, M.; Schier, H.; Gopel, W. Langmuir 1991, 7 (10), 2287−2292.

(17) Allen, P. M.; Liu, W.; Chauhan, V. P.; Lee, J.; Ting, A. Y.; Fukumura, D.; Jain, R. K.; Bawendi, M. G. J. Am. Chem. Soc. 2010, 132, 470−471. (18) Wang, W.; Zhu, W.; Xu, H. J. Phys. Chem. C 2008, 112 (43), 16754−16758. (19) Bae, W. K.; Kwak, J.; Lim, J.; Lee, D.; Nam, M. K.; Char, K.; Lee, C.; Lee, S. Nanotechnology 2009, 20, 075202. (20) Chen, J.; Lin, S.; Yan, G.; Yang, L.; Chen, X. Catal. Commun. 2008, 9, 65−69. (21) Wang, W.; Germanenko, I.; El-Shall, M. S. Chem. Mater. 2002, 14, 3028−3033. (22) Kim, J. U.; Lee, M. H.; Yang, H. Nanotechnology 2008, 19, 465605. (23) Chen, Y.; Zhang, X.; Jia, C.; Su, Y.; Li, Q. J. Phys. Chem. C 2009, 113 (6), 2263−2266. (24) Moriguchi, I.; Matsuo, K.; Sakai, M.; Hanai, K.; Teraoka, Y.; Kagawa, S. J. Chem. Soc. Faraday Trans. 1998, 94 (15), 2199−2204. (25) Parhizkar, M.; Kumar, N.; Nayak, P. K.; Talwar, S. S.; Major, S. S.; Srinivasa, R. S. Colloids Surf., A 2005, 257, 405−410. (26) Dhanabalan, A.; Kudroli, H.; Major, S. S.; Talwar, S. S. Solid State Commun. 1996, 99, 859−862. (27) Smotkin, E. S.; Lee, C.; Bard, A. J.; Campion, A.; Fox, M. A.; Mallouk, T. E.; Webber, S. E.; White, J. M. Chem. Phys. Lett. 1988, 152, 265−268. (28) Moriguchi, I.; Tanaka, I.; Teraoka, Y.; Kagawa, S. J. Chem. Soc. Chem. Commun. 1991, 19, 1401−1402. (29) Vitta, S.; Metzger, T. H.; Major, S. S.; Dhanabalan, A.; Talwar, S. S. Langmuir 1998, 14, 1799−1803. (30) Dultsev, F. N.; Sveshnikova, L. L. Thin Solid Films 1998, 322, 303−307. (31) Kumar, N. P.; Narang, S. N.; Major, S. S.; Vitta, S.; Talwar, S. S.; Dubcek, P.; Amenitsch, H.; Bernstorff, S. Colloids Surf., A 2002, 198, 59−66. (32) Yamaki, T.; Yamada, T.; Asai, K.; Ishigure, K. Thin Solid Films 1998, 327, 586−590. (33) Wada, Y.; Niinobe, D.; Kaneko, M.; Tsukahara, Y. Chem. Lett. 2006, 35 (1), 62−63. (34) Mandal, P.; Talwar, S. S.; Srinivasa, R. S.; Major, S. S. Appl. Phys. A: Mater. Sci. Process. 2009, 94, 577−584. (35) Narayanam, P. K.; Srinivasa, R. S.; Talwar, S. S.; Major, S. S. Colloids Surf., A 2011, 380, 292−299. (36) Rabolt, J. F.; Burns, F. C.; Schlotter, N. E.; Swalen, J. D. J. Chem. Phys. 1983, 78, 946−952. (37) Parhizkar, M.; Kumar, N. P.; Shukla, B.; Srinivasa, R. S.; Vitta, S.; Kumar, N.; Talwar, S. S.; Major, S. S. Colloids Surf., A 2005, 257, 177−182. (38) Nayak, P. K.; Talwar, S. S.; Major, S. S.; Srinivasa, R. S. Colloids Surf., A 2006, 284, 187−191. (39) Kumar, N. P.; Major, S.; Vitta, S.; Talwar, S. S.; Dubcek, P.; Amenitsch, H.; Bernstorff, S.; Ganesan, V.; Gupta, A.; Dasannacharya, B. A. Colloids Surf., A 2002, 198, 75−81. (40) Saperstein, D. D. J. Phys. Chem. 1986, 90, 1408−1412. (41) Nayak, P. K.; Singh, S.; Talwar, S. S.; Srinivasa, R. S.; Major, S. S. Colloids Surf., A 2006, 284, 192−197. (42) Wagner, C. D.; Riggs, W. M.; Devis, L. E.; Moulder, J. F.; Mulenberg, G. E. Handbook of X-Ray Photoelectron Spectroscopy; Perkin-Elmer Corporation: Eden Prarie, Minnesota, 1979. (43) Briggs, D.; Seah, M. P. Practical Surface Analysis, second ed.; John Wiley & Sons: New York, 1993; Vol. 1. (44) Xu, J. F.; Ji, W.; Lin, J. Y.; Tang, S. H.; Du, Y. W. Appl. Phys. A: Mater. Sci. Process. 1998, 66, 639−641. (45) Nanda, J.; Sarma, D. D. J. Appl. Phys. 2001, 90 (5), 2504−2510. (46) Winkler, U.; Eich, D.; Chen, Z. H.; Fink, R.; Kulkarni, S. K.; Umbach, E. Phys. Stat. Solidi A 1999, 173, 253−259. (47) Zhu, Y.; Yuan, C. L.; Ong, P. P. J. Appl. Phys. 2002, 92 (11), 6828−6832. (48) Vegard, L.; Schjelderup, H. Phys. Z 1917, 18, 93. (49) Brus, L. E. J. Chem. Phys. 1984, 80, 4403−4409. L


Strong and Tunable Blue Luminescence from Cd1-xZnxS Alloy

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