J. Phys. Chem. 1996, 100, 14037-14047
14037
Electrochemical Fabrication of CdS Nanowire Arrays in Porous Anodic Aluminum Oxide Templates Dmitri Routkevitch,† Terry Bigioni,† Martin Moskovits,*,† and Jing Ming Xu‡ Department of Chemistry, Department of Electrical Engineering, and the Ontario Laser and LightwaVe Research Centre, UniVersity of Toronto, Toronto, Ontario M5S 1A1, Canada ReceiVed: September 29, 1995; In Final Form: March 3, 1996X
A technique is described for fabricating arrays of uniform CdS nanowires with lengths up to 1 µm and diameters as small as 9 nm by electrochemically depositing the semiconductor directly into the pores of anodic aluminum oxide films from an electrolyte containing Cd2+ and S in dimethyl sulfoxide. The nanowire arrays were characterized by powder X-ray diffraction (XRD) and electron microscopy. The deposited material is found to be hexagonal CdS with the crystallographic c-axis preferentially oriented along the length of the pore. The effects of annealing on the crystallinity of the deposited semiconductor were investigated by XRD and resonance Raman spectroscopy. The deposition technique is, in principle, generalizable as a means of fabricating nanowires of a wide range of semiconductors.
Introduction The large range of potential new materials applications made possible by the fabrication of uniform nanoscale structures has generated a tremendous amount of interest recently. The technological implications for the design and manufacture of new classes of optoelectronic and electronic devices, as well as for the emergence of new physics, have resulted in this field’s being one of the most active in the physical sciences today. Fundamental to its development is the evolution of novel techniques, among them electrochemical methods, for fabricating uniform structures with dimensions in the nanometer range. Traditionally, electrochemistry has been used to grow thin films on conductive surfaces. Because the growth is controllable almost exclusively in the direction normal to the substrate surface, pseudo-0- and 1-dimensional nanostructures are realizable electrochemically only if the deposition is confined within the cells of an appropriate template. Electrodeposition has been used successfully for the formation of ceramic,1a metallic,1b,c and semiconductor1d superlattices. Attempts were also made to use sequential underpotential electrodeposition to grow epitaxial compound semiconductor layers.2 Anodic aluminum oxide (AAO) films grown in strong acid electrolytes possess very regular and highly anisotropic porous structures3 with pore diameters, dp, ranging from below 10 to 200 nm, pore lengths, lp, from 1 to 50 µm, and pore densities in the range 109-1011 cm-2. The pores have been found to be uniform and nearly parallel,3 making AAO films ideal templates for the electrochemical deposition of fairly monodispersed nanometer-scale particles. Other porous films such as polymeric membranes, manufactured by etching nuclear tracks, have also been used.4 It is now feasible to produce AAO films with pore sizes on the scale of the Bohr diameters of bulk semiconductor excitons, suggesting that quantum confinement effects might be observed in the radial, although not necessarily in the axial, direction. In using templates to produce nanostructures, one must take into account the template’s chemical stability, its insulating properties, the minimal diameter and uniformity of the pores, and the pore density. Pore sizes small enough to ensure the observation of quantum size effects in the deposited structures are not readily available with commercial anodic aluminum * To whom correspondence should be addressed. † Department of Chemistry. ‡ Department of Electrical Engineering. X Abstract published in AdVance ACS Abstracts, July 15, 1996.
S0022-3654(95)02910-8 CCC: $12.00
oxide membrane filters such as those produced by Anopore. The smallest mean pore diameter in nuclear track polycarbonate membranes used for the fabrication of nanowires has been reported to be 18 nm.4 Recently, membranes with nominal 10 nm diameter pores have become available from Poretics and from other sources. These have been used to fabricate nanoelectrode assemblies.5 However, direct measurement of the diameter distribution function for metallic wires deposited into these pores determined6 that the mean pore diameters of nominal 10 nm and the nominal 30 membranes were respectively 36 ( 3 and 57 ( 3 nm. The nominal size specified for these membranes corresponds, in fact, to the value of the smallest pores in the distribution, rather than to the mean. Additionally, the nuclear track pores are not parallel, and the pore density (6 × 108 cm-2) is substantially lower than in AAO films. The pore diameters and densities of anodic aluminum oxide films such as those described in this paper can be easily varied by changing anodization parameters such as the electrolyte used, its concentration, and the anodizing voltage Ua. The main constraint in using porous alumina films directly after anodization is the insulating, dense oxide barrier layer separating the Al substrate and the porous portion of the aluminum oxide. The thickness of the barrier layer is a function of the anodizing voltage (10-14 Å/V). This imposes a limitation on the use of dc electrodeposition to fill the pores. However, the inherent rectifying properties of the barrier layer allow the pores to be filled uniformly by ac electrolysis without simultaneously depositing material on the surface or into the macroscopic defects of the film. The possibility of electrochemical (EC) deposition of certain metals into these pores resulting in the formation of the highly anisotropic, aligned particle arrays with unusual optical7 and magnetic8 properties has been known for some time. These particles faithfully reproduce the shape of the pores.9 It was also demonstrated recently that porous templates can be filled electrochemically with semiconductor. Cadmium sulfide/selenide wires were deposited into the 0.2 µm pores of Anopore membranes,10 and selenium nanotubules were synthesized in the 2.5 µm pores of nuclear track membranes.11 In the present paper we report an electrochemical technique for fabricating CdS nanowire arrays based on single step ac electrolysis into the pores of AAO films. Although different in execution, the technique is similar in scope to those used by Chakavarti and Vetter11 and by Sailor’s group.10 However, the membranes used in refs 10 and 11 contained pores of consider© 1996 American Chemical Society
14038 J. Phys. Chem., Vol. 100, No. 33, 1996
Routkevitch et al.
TABLE 1: Anodization Conditions and Resultant Film Parameters anodizing conditions electrolyte
U a, V
ta, °C
τa, min
1.2 M H2SO4 1.2 M H2SO4 1.2 M H2SO4 0.23 M H2C2O4 0.87 M H3PO4
15 15 15 15 20
1-2 6-8 12-14 20 20
90 80 40 60 60
film parameters AAO film pore diameter thickness, µm dp, nm 2.8 2.9 3.0 1.5-2 1.3-1.5
9 ( 0.9 10 ( 1 11 ( 1.1 16 ( 1.5 35 ( 3.5
ably larger dimensions than those used in the present work. Specifically, the dimensions were significantly larger than the Bohr diameter of the 1s exciton in CdS, which is estimated to be approximately 4 nm.12 For the current study we produced CdS nanowires with diameters in the range 9-35 nm, the lower end of which is sufficiently small so as to observe the onset of quantum confinement, the manifestation of which will be reported elsewhere.13 In this paper we concentrate on the deposition technique and on the effects of the confinement on the crystalline structure of CdS nanowires. Experimental Section Materials and Sample Preparation. Anodic oxide films were grown on 0.05 mm thick 99.8% aluminum foil (Aldrich) or on 0.1 mm thick 97% aluminum foil (Alcan). Samples were degreased ultrasonically in trichloroethylene for 1 h, electropolished to a mirror finish in a 5:1 v/v solution of EtOH (95%)/ HClO4 (70%) at 18 V and 8-10 °C for 2-3 min, etched in 0.24 M Na2CO3 at 80-85 °C for 1 min to remove the native oxide layer, and then anodized in an acid electrolyte with graphite counter electrodes. The anodization conditions and the parameters of the ensuing AAO films are given in Table 1. When the anodization was complete, the anodizing voltage was decreased stepwise in 5 V intervals over a period of 5 min in order to reduce the thickness of the barrier layer. This thinning of the barrier layer was found to improve the subsequent electrodeposition of the semiconductor into the AAO pores. The diameters of the pores could be increased by the partial dissolution of the oxide in an acidic solution. Widening in a solution of 0.1 M H3PO4 at 40 °C8b,c for varying times, τw, can produce pores ranging from 9 to 25 nm in AAO films initially anodized in sulfuric acid. Films with larger pores (diameter 35 nm) were produced by anodization in phosphoric acid (Table 1). Pore diameters and film thicknesses were determined by SEM and TEM in a manner similar to what was reported previously, that is, by assuming that nanowires removed from the pores were faithful replicas of their matrix.9 The diameters of several wires (at least 10) were measured for each average value reported. Additionally, at least five measurements were made along the length of each wire. The mean pore diameter of some samples was also determined from N2 adsorptiondesorption measurements, giving similar results. The standard deviation of pore/wire diameter was never higher than 10% of the mean diameter. Deposition was carried out in a solution contained 0.055 M CdCl2 and 0.19 M elemental sulfur dissolved in dimethyl sulfoxide (DMSO)14 at 100-160 °C and 30-50 V ac (60-500 Hz) applied between the Al\AAO working electrode and a graphite counter electrodes for 5-60 min. Films of CdS were also deposited on a Pt polycrystalline electrode by dc electrolysis (j ) 1-0.5 mA cm-1) in order to produce samples for comparison and control measurements. The rectifying properties of the oxide barrier layer that separates the bottom of the pores from the underlying aluminum substrate make it possible to use
ac deposition. Although the applied voltage is sinusoidal and symmetrical, the current is greater during the cathodic halfcycles, making deposition possible without subsequent stripping during the anodic half-cycles. Of course, the oxide barrier layer can be removed and a metal electrode deposited on the exposed lower orifices of the pores and ordinary dc electrodeposition used to fill the pores as was done in refs 10 and 11. However, for samples with very narrow pore diameters this strategy only works when the anodic films are free from cracks and defects; otherwise, the electrochemistry becomes dominated by processes in the much more accessible cracks. Fabricating templates of such structural perfection is very difficult. The pore diameters of the templates used in refs 10 and 11 were sufficiently large to allow the electrodeposition in the pores to compete successfully with that occurring at cracks and defects. By retaining the barrier layer and using ac electrolysis, this problem is avoided even with templates of very small pore diameter since the rectification only occurs inside the pores while at defects the electrochemical processes occurring during the cathodic halfcycle are reversed, more or less, during the anodic half-cycle. Following this line of reasoning it is clear that ac electrodeposition is not ideal for conducting metal substrates such as the Pt surface used to produce the control samples. Electrodeposition was carried out in a glass cell fitted with adapters for the counter electrodes, the working electrode, a platinum resistance temperature detector, and a condenser with moisture trap. The temperature of the electrolyte was maintained to within 0.5 °C with a heater controlled by a CN76000 temperature controller and solid state relay (Omega Technologies). The smallest pore diameters were achieved with lowtemperature anodization which in turn affected the choice of electrodeposition parameters. After deposition, the electrodes were washed with warm DMSO, methanol, and distilled water. Some samples were annealed under oxygen-free N2 for 1 h. In order to minimize oxidation, samples were removed from the annealing furnace approximately 1 h after it had cooled under continuous nitrogen flow. The electrodeposited semiconductor particles were examined by electron microscopy by first liberating them from their AAO matrix by dissolving the anodic oxide in 0.1 M NaOH at 40 °C, which does not affect CdS. The particles were freed from the AAO completely after 5 min of dissolution or longer. They could then be easily detached from the film surface and suspended in water to form fairly stable colloids. Instrumentation and Characterization. Scanning electron microscopy (SEM) was carried out on Hitachi S-570 or S-4500 (Field Emission) instruments operated at 10-20 kV and equipped with energy dispersive X-ray (EDX) fluorescence microanalysis. SEM and transmission electron microscopy (TEM, Philips-430, operated at 100 kV) were used to obtain morphological information as well as the size and the atomic composition of the semiconductor particles. For SEM, Al\AAO, CdS samples with partially removed AAO were mounted on Al stubs with conductive silver paint. The conductivity of the CdS was sufficient enough so that high-resolution SEM micrographs could be obtained without applying an additional conductive coating. For TEM imaging, the CdS nanowire suspension was centrifuged and washed several times, then a small drop of the colloid was placed on the carbon/formvar films supported by Cu grids, and the excess water was blotted off after 2-5 min of partial sedimentation. Cross sections of the Al\AAO,CdS foils were prepared by embedding small pieces of the foil in epoxy resin, curing, and microtoming. X-ray powder diffraction (XRD) was carried out on a D5000 SIEMENS diffractometer with solid state (KEVEX) detection
Electrochemical Fabrication of CdS Nanowire Arrays
J. Phys. Chem., Vol. 100, No. 33, 1996 14039
Figure 1. SEM micrographs of CdS nanowires exposed by dissolution of the oxide film in 0.1 M NaOH at 40 °C. Anodizing was carried out in 1.2 M H2SO4 at 8 °C followed by the 15 min of pore widening (A, B) and in 0.87 M H3PO4 at 23 °C (C, D). Deposition conditions: Ud ) 41 V (A, B) and 35 V (C, D); 60 Hz, tdep ) 120 °C, time (min) ) 15 (A, B) and 6 (C, D). Oxide dissolution time ) 3 min. Numbers on the micrographs are discussed in the text.
using filtered Cu KR radiation. For routine measurements the counting time per step and the step size were respectively 1.5 s and 0.02°. For more precise measurements those parameters were changed to 5 s and 0.01°. The diffraction peaks of the aluminum substrate were used as an internal standard to ensure proper 2Θ calibration. Peak positions, areas, and fwhm were determined by fitting the diffraction profiles to Lorentzian, Voigt, pseudo-Voigt, or Pearson functions,15,16 as required. Unit cell refinement was performed using a modified AppelmanEvans routine.15 Resonance Raman spectra were excited with an argon ion laser and collected using a computer-driven double-pass monochromator (SPEX 1401) equipped with photon counting. Results and Discussion Direct electrochemical deposition of CdS films on planar electrodes is a well-known procedure.17 The deposition of CdS into the pores of anodic oxide films is complicated by side reactions involving the anodic oxide such as reanodization, dissolution, and Al corrosion which can even result in the destruction of the porous film. In fact, attempts to use the conventional aqueous acid electrolytes containing Cd2+ and sources of colloidal sulfur such as S2O32- or (NH2)2CS to deposit CdS on an Al\AAO electrode resulted entirely in pitting corrosion without any CdS deposition.18 Nonaqueous solutions
and, in particular, the solution developed by Baranski14 which contains Cd2+ and elemental S dissolved in DMSO, however, gave good results. When a voltage greater than the minimum ) 25-35 V rms was applied across deposition voltage Umin d the cell, the AAO electrode developed a uniform color starting as a light lemon-yellow which ultimately became a deep yellow film presumably as a result of the deposition of CdS into the were found to AAO pores. The deposition rate and Umin d depend on the pore diameter, barrier layer thickness, ac frequency, and electrolyte temperature. SEM micrographs of nanowires exposed from the template (Figure 1) show that the deposited semiconductor, indeed, fills the pores uniformly and that the nanowires are apparently continuous. The measured diameter of the CdS wires corresponded closely to the pore diameter, vindicating the previous assumption. Transmission electron micrographs of CdS particles liberated from the template are shown in Figure 2. Similar images were recorded for many of the samples used in this study. The electron micrographs suggest a uniform and uninterrupted wire structure for the CdS deposit, and the TEM images of the AAO\CdS cross sections verified the film structure (Figure 3). The growth of CdS nanowires inside the pores begins at the bottom of the pores and grows outward. Small gaps between the end of some wires and the pore bottom are probably artifacts
14040 J. Phys. Chem., Vol. 100, No. 33, 1996
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Figure 2. TEM micrographs of the CdS nanowires deposited in a sulfuric acid AAO film shown after the complete dissolution of anodic oxide supporting matrix in 0.1 M NaOH at 50 °C. Widening time (min) ) 20 (A) and 0 (B).
Figure 4. X-ray diffractograms of an as-deposited (120 °C) and annealed (500 °C, 1 h, N2) AAO\CdS nanowire array (mean diameter 35 nm, length 0.8 µm), Pt\CdS film (1 µm thickness), and polycrystalline hexagonal CdS with mean crystal domain size ∼500 nm. The vertical lines at the bottom of the figure indicate the standard position and relative intensities of hexagonal CdS (ASTM 6-314, greenockite).
Figure 3. TEM micrograph of a microtomed cross section of an Al\AAO\CdS film. Anodization: 0.87 M H3PO4, 23 °C, 20 V. CdS deposition conditions: 35 V, 60 Hz, 125 °C, 6 min. Micrographs A and B are different areas of the same sample.
due to mechanical damage during microtoming. If ac electrodeposition is continued beyond the point where the pores are filled, the deposition continues on the surface of the anodic film, and the wire tips coalesce to form hemispherical mushroomlike surface structures, which are clearly seen in Figure 1B (marked as 3). In spite of the large number of nanowires that are damaged during the sample preparation, the EM micrographs show that while locked inside the template, the nanowires are continuous (Figure 1A,B). The wire lengths are also found to be quite uniform (Figure 3) with a standard deviation approximately 10-11% of the mean length. Figure 1 shows that the wires can be microns in length, resulting in aspect ratios (Figure 1A,B) as large as 120 (length ∼2400 nm and diameter
∼20 nm), limited mainly by the template thickness (Table 1). The wire length can be controlled by varying the deposition time. The particles appear to be more segmented in the TEM images (Figure 2). This may be due to the fact that TEM is more sensitive to density variations than SEM, or it may reflect a greater level of particle damage resulting from the more severe sample preparation conditions required for this form of microscopy. Cross sectioning also results in sample damage due to the fact that the alumina matrix is a hard material which does not cut easily without fracturing. The electron micrographs also show that some of the wires, especially those grown in nonwidened templates, have morphological imperfections such as branching, nonuniform diameters along a single wire, and differences in mean diameter from wire to wire. We believe that these reflect irregularities in the pore structure of the template in which the wires were made or perhaps nonuniform filling of the pores rather than changes due to samples handling. These might result from temperature fluctuations during the anodization. Nevertheless, these imperfections result in a diameter standard deviation which is never greater than 10% of the mean diameter. Precisely controlled anodizing conditions and uniform current distribution across the electrode surface are probably the key factors in establishing the pore uniformity. Quantitative electron microprobe analysis from a 10 × 10 µm area indicates a composition of 53.6 at. % Cd and 44.7 at. % S, which is close enough to a 1:1 stoichiometry for measurements carried out without an internal standard. (An appropriate standard, reflecting the mesoscopic structure of the sample, moreover, is difficult to devise.) The phase composition of the deposited material was further verified by XRD. All diffraction peaks could be assigned to CdS, Al, and AAO, without any trace of elemental Cd or S. Additionally, the resonance Raman spectra and the optical band gap13 correspond well to those of CdS. All of this together with the available data on the electrodeposition of CdS from DMSO14,17,19,20 makes us confident that the deposited material is n-CdS, probably doped with Cd. Although the presence of a dopant may affect
Electrochemical Fabrication of CdS Nanowire Arrays
J. Phys. Chem., Vol. 100, No. 33, 1996 14041
TABLE 2: Peak Positions, dhkl, and Peak Areas, Ihkl, for the (100), (002), and (101) Planes of CdS and Unit Cell Parameters (100) diffraction plane sample
(002) diffraction plane
(101) diffraction plane d, Å
I, %
a, Å
c, Å
Vcell, Å3
34 22
AAO\CdS (pore diameter 35 nm) 3.331 100 3.142 3.353 100 3.167
71 58
4.087(9) 4.135(7)
6.67(2) 6.70(1)
96.6(2) 99.3(3)
3.530 3.578 3.579
30 8 63
3.307 3.351 3.355
3.126 3.156 3.158
100 28 100
4.069(1) 4.134(1) 4.134(5)
6.61(3) 6.709(2) 6.714(1)
94.8(7) 99.3(4) 99.4(2)
3.58
75
3.36
3.16
100
4.14
6.71
99.60
d, Å
I, %
unannealed annealed
3.526 3.576
unannealed annealed polycrystalline CdS ASTM file (6-314)
d, Å
I, %
Pt\CdS 32 100 47
the crystal structure of the deposited CdS slightly, we do not expect this effect to depend of the nanowire diameter. The XRD data (Figure 4) shows that, even before annealing, the nanowires have diffraction patterns corresponding to the hexagonal greenockite phase of CdS (ASTM standard 6-314). The interplanar diffraction spacing (dhkl) and the unit cell parameters of CdS, deposited into porous AAO with mean pore diameter 35 nm, differ slightly from those reported for polycrystalline CdS (Figure 4, Table 2). This was also found to be the case for a control sample consisting of CdS deposited on a polycrystalline Pt electrode. In both cases the unit cell of the as-deposited CdS was found to be compressed with respect to that of polycrystalline CdS. This compression was even more pronounced for the CdS film on Pt. Annealing at 500 °C for 1 h successfully relieved this distortion, except, as will be shown below, for wires with small diameters. The diffraction data also suggest that electrodeposition orients the CdS crystallites with respect to the substrate plane. For AAO\CdS deposited by ac electrolysis, the relative intensity of the 002 diffraction peak which corresponds to interplane distances d ) 3.33-3.35 Å is greater than that of the corresponding peak for the Pt\CdS film deposited by dc or for polycrystalline CdS powder (Figure 4). This is unequivocally due to the fact that the samples are textured, that is, their crystallites are not oriented randomly; rather, the c-axis of hexagonal crystals is preferentially aligned along the direction normal to the substrate. For the AAO\CdS sample this means that the crystalline c-axis lies preferentially along the length of the pore. The intensity of 002 reflection is increased with annealing, slightly for the AAO\CdS samples and very greatly for the Pt\CdS films. The contributions of the other main diffraction planes ((100), (101), (110)) are suppressed significantly in both sets of samples in comparison with the powder standard. This orientational effect for Pt\CdS films has been reported previously; however, our data differ from that of Baranski et al.,19 who observed only one intense diffraction peak (d002 ) 3.34 Å) for a 5 µm Pt\CdS film electrodeposited from the same electrolyte as was used in this study. Fatas et al.20 report that increasing the deposition temperature in the range 70-160 °C caused the grain size of the CdS deposit to increase and the degree of c-axis orientation to decrease as evidenced by the appearance of new diffraction peaks belonging to the hexagonal phase. In contrast, the relative intensities of the diffraction peaks of our AAO\CdS samples are not observed to change significantly from 100 to 160 °C (Figure 5, mean pore diameter 9 nm), implying that the preferred c-axis orientation is preserved. However, after the pores become filled, the further growth of CdS on the film surface produces changes in the intensities of the X-ray diffraction lines as shown in Figure 6 which illustrates the behavior of AAO with 35 nm mean pore diameter. The 002 diffraction peak appears immediately after deposition begins. As deposition proceeds, this peak increases slightly, its width remains almost constant, and
60
Figure 5. X-ray diffractograms of AAO\CdS (mean wire diameter 9 nm), deposited at 35 V, 60 Hz for 20 min at the deposition temperatures shown.
Figure 6. X-ray diffractograms of as-deposited AAO\CdS with 35 nm mean wire diameter, deposited at 120 °C, 35 V, 60 Hz for the lengths of time shown.
two new diffraction peaks (100 and 101) appear. With 4-5 min deposition which is enough to just fill the pores, practically no other CdS diffraction features were observed. At this point SEM shows no trace of CdS on the AAO surface. When the deposition time is long enough for the wires to form hemispherical structures on the AAO surface, additional peaks belonging to polycrystalline CdS appear in the diffraction profile (Figure 6, 21 min deposition time), which grew stronger as more material was deposited outside the pores. An illustration of the
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early stages of the growth of these mushroom cap-like structures is shown in Figure 1A,B. (For the XRD study of the influence of pore diameter on the crystallite size and d-spacing, the deposition time was chosen so as to produce samples in which the wires were of comparable length (0.8 ( 0.1 µm), and less than the membrane thickness, in order to avoid deposition on the surface of the film.) Using Scherrer’s formula (eq 1), the dimensions of the crystallites of which the CdS nanowires and the CdS films on Pt are comprised were estimated from the widths of the major diffraction peaks observed. These were the 002, 100, and 101 for the AAO\CdS sample with 35 nm diameter wires and for the Pt\CdS control sample. For AAO\CdS with smaller pore diameters only the 002 line was seen; hence, only the crystallite dimension along the pore axis could be determined.
Dhkl ) kλ/(∆hkl cos Θ)
(1)
where λ is X-ray wavelength in angstroms, ∆hkl is the corrected fwhh of the peak in radians, Θ is the Bragg angle, k is a crystal shape constant (0.94) which assumes that the sample peak is Gaussian and that all of the crystallites are of uniform size, and Dhkl is the linear dimension of the coherent diffracting domain along a direction normal to the diffraction plane (hkl).13,14 When the term “crystallite size” is used, we will be referring to the dimensions of the coherent diffracting domains. The crystallite dimension can, clearly, be less than the dimension of the nanowire in both the axial and the radial direction, although because the aspect ratio of the wire is large the number of crystallites in the axial direction will, in general, be greater than that in the radial direction. Very narrow wires can, in fact, be comprised of a single crystallite in the radial direction but several crystallites in the axial direction. The method may be used for crystalline domains as small as 20 Å, providing small enough angular increments and long enough count times are used.15 Instrumental broadening effects were determined using diffraction from crystalline CdS powder with a mean particle size of approximately 0.5 µm as the standard. The corrected peak width was calculated as the square root of the difference between the squares of the sample’s and reference’s fwhm. The KR2 intensity was removed from the reference’s and samples’ profiles. Equation 1, which is applicable to samples where lattice strain is absent, assumes that the reference peak which expresses the instrumental broadening is Lorentzian and that the sample peak has a Gaussian profile. Since actual profiles are never purely Gaussian but are better represented by a linear combination of Gaussian and Lorentzian (Voigt profiles), this approach for estimating the particle dimensions is of necessity approximate.16 The peak shape limitation did not appear to be serious in our case since similar values for the peak widths and, more importantly, the same trends as a function of the template’s average pore size were obtained by fitting the diffraction profile to a number of different peak profile functions (see Figure 9). More serious problems might arise from the fact that the unannealed samples possess significant strains, as indicated by the shift of the 002 peak position, which could also be contributing to the width of the peak, thereby affecting the estimates of the crystallite size of the as-deposited CdS. The effects of size and distortion broadening may be separable if at least two, and preferably several, orders of the 00l reflections are available,16 which is not the case with AAO\CdS for which only the few first diffraction peaks are detectable. Therefore, for the unannealed samples the size of the crystalline domains determined from the XRD peak widths will be used only as a comparative measure among samples. Estimates for the annealed samples are more quantitative.
Figure 7. X-ray diffractograms of as-deposited AAO\CdS as a function of AAO pore diameter.
TABLE 3: Mean Crystallite Sizes Dhkl (nm) along the [100], [002], and [101] Zone Axes AAO\CdS (dp ) 35 nm) D100 D002 D101
Pt\CdS film
unannealed
annealed
unannealed
annealed
6.6 17 7.5
10 21 10
7 11 9
19 76 51
The dimensions of the CdS crystallites, extracted from the widths of the observed XRD reflections for AAO\CdS (dp ) 35 nm) and Pt\CdS, are listed in Table 3. The results indicate that CdS deposited in AAO and that deposited on Pt behave significantly differently upon annealing. In the former the crystallite dimensions and c-axis orientation change very little with annealing; in the latter these parameters change a great deal. This is almost certainly due to the confinement of the CdS crystallites in the pores of the template. As expected, the crystalline domains of the dp ) 35 nm AAO\CdS samples were found to be anisotropic with an approximate aspect ratio of 2:1, calculated from the dimensions determined along the [002] and [100] axes. Because the deposition of AAO\CdS and Pt\CdS were carried out under very similar conditionssthe same temperature, electrolyte, and comparable deposition rates (0.07 µm s-1)swe are tempted to conclude that the main factor determining the special structural features of AAO\CdS, such as its crystalline orientation, is the spatial confinement of the deposition within the pores. However, the possible influence of other parameters such as the differing substrate material and the differing electrochemical influence of ac versus dc deposition cannot be ruled out. The 002 diffraction peak obtained for AAO\CdS samples with mean pore diameters between 9 and 35 nm (Figure 7) range from very broad and weak peak for 9 nm to quite narrow and well-defined for the 35 nm sample. This results from the combined effect of changing the crystal size and the total quantity of the CdS in the AAO\CdS sample with changing pore diameter. We attempted to estimate the behavior of this XRD reflection as a function of the pore and cell diameters of the templates by considering only the changes in the crystallite dimension and in the quantity of the CdS contained in the AAO film but omitting the influence of lattice strain. (The cell diameter is the average linear dimension of the portion of the AAO surface that can be associated with the a single pore. It is approximately
Electrochemical Fabrication of CdS Nanowire Arrays
Figure 8. Calculated XRD profile for AAO\CdS as a function of the pore diameter: (A) cell diameter changes along with pore diameter and dc ≈ 2dp mimicking the use of varying the anodization conditions to change the pore diameter; (B) the cell diameter is constant at 40 nm, mimicking the situation in which postanodization etching is used to widen the pores.
equal to the average distance between pores.) The details of the calculation are summarized in the Appendix. The influence of the pore diameter on the XRD profile of the material filling the pores is illustrated in Figure 8 for both constant and varying cell size. The two parts of Figure 8 mimic the two strategies used in varying the pore size. The average pore diameter can be varied by varying the anodization parameters including the electrolyte temperature, in which case the cell dimension increases with increasing pore diameter. Alternatively, the pore diameter can be increased by etching after anodization is complete. In that case the cell dimension remains constant. The broad and low-intensity peaks predicted for the smallest CdS nanowires (Figure 8) agree well with experiment, confirming our suggestion that our particle sizes are approaching the limit of XRD analysis. The crystallite size (D002) and interplane distance or d-spacing (d002) obtained from the 002 reflections show an interesting dependence on the mean pore diameter (Figure 9). The crystallite size at first appears to increase with increasing pore diameter up to dp ) 15 nm, beyond which the particle dimension appears to be approximately constant or even to decrease somewhat with increasing pore diameter. The XRD data imply that the mean linear dimension of the crystallites along the length of the nanowires (where the growth is presumably much less restricted) is larger than their mean diameter, dp, in the range 9 < dp < 12 nm, and it becomes almost constant for dp > 12 nm. For the as-deposited CdS nanowires the d-spacing increases with increasing template pore diameter, reaching a maximum at dp ) 12-15 nm, and then drops slightly, perhaps indicating an increase in lattice distortion. For annealed samples there is a steady increase in d002 with increasing pore diameter, approaching the value for bulk hexagonal CdS. In the asdeposited Pt\CdS samples, where the growth is the least restrictive, the d-spacing (3.307 Å) is even lower than that for the AAO\CdS samples (Table 2, Figure 9), which suggests that
J. Phys. Chem., Vol. 100, No. 33, 1996 14043
Figure 9. (A) Mean crystalline linear dimension in the [002] direction determined from the 002 diffraction peak width as a function of the average pore diameter of the AAO template. The result of using different line-shape functions to fit the observed diffraction lines is shown. The data refer to annealed samples. (B) The (002) interplane distances obtained for CdS nanowires as a function of the average pore diameter of the AAO template. Data for both the as-deposited and the annealed samples are shown.
the growth of CdS is more ordered and less strained when spatial confinement is imposed by the AAO template. Seemingly contradictory, this observed template effect is reconcilable when the experimental conditions and the deposition mechanism are considered. Electrodeposition quite often produces deposits with high strain, primarily because of the mismatch between the crystallographic parameters of the substrate and the deposit and because the new phase formation is a nonequilibrium process due to the low temperatures used. Additionally, the mechanism of CdS deposition consists of a heterogeneous electrochemi2+ 0 0 + 2e- f Cdads ) and a chemical step (Cdads + cal step (Cdsln 0 Ssln f CdSads) which may result in the formation of a large number of nucleation sites and therefore in a low degree of longrange order and in high strain in both the template-deposited and the Pt-deposited samples. The relatively high deposition rate may also contribute to the growth of disordered dendritic particles with a great deal of lattice distortion. When the growth is confined to narrow parallel pores, several key parameters of affecting the buildup of the deposit are altered. The diffusion front becomes more uniform, and fewer nuclei encounter other growing nuclei before encountering the static pore wall. The influence of the template is expected to be less pronounced inside very large pores where the growth mechanism may more closely resemble that at a planar interface, with many nucleation sites contributing to the growth of any given wire. On the basis of the crystallite size for the CdS film and wires in the 35 nm pores, we conclude that with decreasing pore diameter the condition is ultimately reached when very few or even one nucleation site is responsible for the growth of any given wire. We estimate at what crystallite size this condition is reached by noting that comparable crystallite sizes were obtained for the AAO\CdS (pore diameter 35 nm) and Pt\CdS samples (Table 3): D002 ) 14 nm, D100 ) 7 nm, and D101 ) 8
14044 J. Phys. Chem., Vol. 100, No. 33, 1996
Routkevitch et al.
Figure 10. Resonance Raman spectra of as-deposited and annealed AAO\CdS nanowire arrays with mean diameter (A) 10, (B) 12, and (C) 22 nm. (D) RRS of a Pt\CdS film. Samples were excited with s-polarized, 488 nm, 75 mW Ar+ laser light. RRS were collected through an s-polarized analyzer.
nm. When the pore diameter approaches the crystallite dimension in the [h00] direction, the pore wall may be assumed to impose a limit on the crystallite size, leading to another kind of lattice distortion which manifests itself as a decrease in d002 for nanowires produced in pores with diameters dp < 12 nm (Figure 9). Annealing relieves the bulk strain, almost eliminating the variations in d002 for AAO\CdS with different pore diameters (Figure 9, Table 2), and partially reduces the strain due to the constraint imposed by the pore walls, resulting in only a slight decrease in d002 for samples produced in pores with dp < 12 nm (Figure 9). The most pronounced changes upon annealing occur in the AAO\CdS samples with dp ) 24 and 35 nm as well as in the CdS film on Pt, which have more strain of the first kind. By contrast, much smaller changes are observed for samples with dp < 24 nm. Interestingly, when the pore diameter is reduced below approximately 12 nm, the apparent crystallite size is reduced not only in the lateral direction, which for oriented crystallites corresponds to the [h00] direction, but there also seems to be a reduction in the particle size in the axial [002] direction. This may be a consequence of several factors, one of which is the orientational influence of the template, presumably leading to a small increase in D002 with decreasing pore diameter when dp < 12 nm. A decreasing number of nuclei may be another reason for slightly more coherent growth. When the pore diameter becomes small enough, the pore walls become an obstacle to the propagation of coherent growth along the length of the pore, and the influence of the CdS-AAO interface at the edges of a growth front becomes more pronounced. Therefore, the length over which coherent growth occurs is likely to be reduced in the axial as well as in the lateral directions. This may also explain the presence of the small maximum at approximately dp ) 12 nm in the graph of the mean crystallite size in the [002] direction as a function of mean pore diameter (Figure 9). At that value of the pore diameter the increase in the mean particle length due to the restriction in the number of nucleation sites has not yet been totally offset by the reduction in the
average axial coherence length induced by the larger number of grain boundaries or other defects resulting from the increase in the surface-to-bulk ratio. Polarized resonance Raman spectroscopy (RRS) was also used to characterize the AAO\CdS samples before and after annealing. Annealing causes dramatic changes in the intensities of the resonance Raman spectra of both the AAO\CdS and the Pt\CdS samples (Figure 10). The changes are much more pronounced than those in XRD. Annealing at 300 °C produced a 2-3-fold increase in the signal-to-noise ratio (S/N) of the RRS of the AAO\CdS samples, while annealing at 500 °C increased their S/N ratio by a factor of approximately 20. The overall intensity of the Raman scattering and the number of overtones observable also increased with annealing. The quality of the spectra was also particle diameter dependent. After annealing to 500 °C, the S/N of the sample with the largest diameter CdS particles was approximately 1.5 times higher than that for the sample with the smallest. The Raman spectrum is dominated by overtones of the longitudinal optical phonon (LO). The 12 nm particles produced a RRS spectrum with three clearly discernible LO modes, while the RRS spectrum of the 22 nm nanowires contains a fourth LO mode. Samples annealed at 500 °C proved to have reasonably strong Raman spectra which could be analyzed to yield values of the band gap and the homogeneous line width of the excited state as described in ref 13. The RRS spectra all possess a background signal which increases with increasing Raman frequency. The background observed in the spectra of the 10 nm AAO\CdS samples (Figure 10a) is due to luminescence from the anodic aluminum oxide. It increases with annealing temperature. The same sort of luminescence is observed from annealed but unfilled aluminum oxide films. Luminescence originating from the CdS is, therefore, hard to discern from that of the oxide film for the small diameter samples. However, the aluminum oxide luminescence decreases sharply for pore-widened samples in which the anodic alumina
Electrochemical Fabrication of CdS Nanowire Arrays walls are partially dissolved. The decrease is greater than that expected entirely on the basis of the reduction in the oxide volume. This leads us to believe that the luminescent material lies primarily near the surface of the pore walls of the anodic alumina so that it is disproportionately greatly removed during pore widening. The CdS luminescence band maximizing at approximately 19 800 cm-1 is observed with progressively increasing intensity as the nanowire diameter is increased. The luminescence background is much more intense for the Pt\CdS samples. Their spectrum consists of weak Raman peaks superimposed on a broad, intense luminescence background (Figure 10d). The low-frequency Raman band in the Pt\CdS sample and its very high luminescence may be caused by impurities incorporated during dc deposition which are either not taken up or taken up to a lesser extent by the AAO\CdS samples during ac deposition. The apparently contrasting behavior of the RRS and XRD intensities of the AAO\CdS samples with annealing is not difficult to reconcile. The intensity and line widths of the XRD reflections are determined by the particle sizes and the quantity of crystalline CdS probed (Figures 8 and 9). Neither of these changes much with annealing. The degree of crystallinity contributes more significantly to the XRD peak position which is influenced dramatically by annealing (Figure 9). Contrariwise, the RRS intensity is greatly influenced by the degree of crystallinity and by the annealing of strains and bulk and surface defects. By exciting CdS nanowires with s- (the electric field vector lies across the wires) and p- (the electric field vector has components both along and across the wires) polarized laser radiation and analyzing the s- and p-polarized Raman scattered light, one might, in principle, observe different physical properties along the axial and radial directions of the CdS nanowires, as a result of unequal confinement conditions following excitation along and across the nanowire. Additionally, the anisotropy may result from the crystallographic orientational propensity of the hexagonal CdS along the wire. The frequency of LO phonon is known to differ slightly depending on propagation direction.21 The Bohr diameter of the 1S exciton of CdS is 4 nm,12 approximately half the smallest pore diameter we can make. Consequently, incipient weak quantum confinement effects is expected to be observable.22 The intensities of the overtones in the Raman spectra of semiconductor nanoparticles depend strongly on their size.23,24 The intensity and number of overtones depend both on the value of the band gap and also on the displacement of the minimum of the ground state potential with respect to that of the excited state along the normal mode of the crystal corresponding to the LO phonon. In our case the ratio of the fundamental mode intensity to that of the first overtone increases with decreasing particle diameter (Figure 10), suggesting that the quantum size effect regime may already have been approached for the AAO\CdS samples described here. Quantitative analysis of these phenomena is presented in a separate publication,13 where the value of the exciton energy of the CdS nanowires was determined as a function of the wire diameter from the excitation wavelength dependence of their resonance Raman spectra in the vicinity of the CdS adsorption edge. AC electrodeposition may be generalized so as to fabricate nanowires of other semiconductors including other metal chalcogenides, GaAs and InSb. Materials with smaller electronic effective masses and larger exciton diameter such as InSb (with an electronic mass of 0.014m0 and exciton diameter well over 10 nm) or ZnTe (with an exciton diameter of 14 nm) should show strong quantum confinement if deposited into pores with
J. Phys. Chem., Vol. 100, No. 33, 1996 14045 diameters of 10 nm or less and would, therefore, display a great deal of tunability as a function of pore diameter. InSb has an absorption threshold in the near-infrared which could be pushed toward the visible through size reduction. Semiconductor lasers manufactured from these materials would have potentially variable output frequencies over the near-infrared to red wavelength range. Moreover, the oxide film might act as a waveguide containing the active laser medium. The possibility also exists of reducing the pore diameters further by proper the choice of electrolyte concentration, current density, and temperature. We have already succeeded in producing samples with a large fraction of its pore diameter distribution below 5 nm. Finally, it is possible to expose the tips of the semiconductor nanowires for the purposes of making electrical contact or for other purposes by selectively etching the aluminum oxide or the underlying aluminum without harming the semiconductor. Conclusions We have described a simple electrochemical method for depositing CdS into the pores of anodic aluminum oxide films, thereby producing arrays of semiconductor nanowires with diameter as small as 9 nm and with high aspect ratios. The approach is applicable to several other semiconductors. Among its advantages are its simplicity and the structural stability of the samples generated by it, resulting from the refractory nature of the oxide matrix. The optical transparency of the oxide matrix suggests the possibility of producing optoelectronic devices based on these relatively monodispersed semiconductor nanostructures. The controlled variability of the particle diameters and composition might also be used to tune in desired optical properties. Cadmium sulfide nanowires, fabricated using this method, were characterized by electron microscopy, X-ray diffraction, and Raman spectroscopy. The confinement of the semiconductor particles to the interior of the oxide matrix also allowed us to improve their crystallinity by annealing without changing the particle dimensions (at least in the direction perpendicular to the length of the pore). The strong resonance Raman spectra recorded for CdS nanowires with mean diameters from 9 to 35 nm showed some particle size-dependent features. Acknowledgment. We thank the Natural Sciences and Engineering Research Council of Canada for financial support. D.R. thanks NSERC for a NATO Science Fellowship. The authors also acknowledge the help of Dr. T. Haslett and Dr. C. Douketis with the Raman measurements, Dr. S. Petrov for the XRD measurements, and Dr. N. Coombs for electron microscopy. Appendix The intensity of a diffraction line obtained from an array of nanowires embedded in an AAO film, Iw, will be lower than that characterizing a homogeneous film of the material of the same thickness, I0,15 by the ratio of the mass absorption of the material to that of the entire film:
() µ
Iw ) IO
F
0
no. of elements
∑ i)1
()
Xw
µ F
i
or in our case
14046 J. Phys. Chem., Vol. 100, No. 33, 1996
Routkevitch et al.
TABLE 4: X-ray Adsorption Coefficients X-ray cross section for Cu KR radiation (µ/F), cm calculated absorption coefficient (µ/F), cm2 g-1
ICdS ) I0
2
g-1
(ref 25)
(µF)
XCdS
CdS
S
Al
O
H
229
92.5
50.2
11.0
0.39
CdS
γ-Al2O3
Al2O3‚XH2O
198.7
31.8
28.4
(A1)
µ µ XCdS + XAAO F CdS F AAO
()
Cd
()
where (µ/F) is mass absorption coefficient and Xi is the mass fraction of component i in the film. The mass absorption coefficients of the compound can be expressed in terms of the known values for the individual elements,15 presented in Table 4. Anodic aluminum oxide is a mixture of γ-Al2O3 and amorphous hydrated aluminum oxide, Al2O3‚H2O. The difference in adsorption coefficients between the two is negligible due to the low element numbers for AAO. The mass absorption coefficients of the materials contained in our CdS loaded AAO films are as follows:
(µF)
CdS
(µF)
) XCd
(µF)
γ-Al2O3
(µF)
Al2O3‚H2O
Cd
(µF)
) XAl
) XAl
(µF)
Al
+ XS
Al
(µF)
S
(µF)
+ XO
(µF)
+ XO
O
O
(µF)
+ XH
H
(A2)
The mass fractions of the components of the film (XCdS and XAAO) and that of the elements (XCd, XS, XAl, XO, and XH) are given by
XCdS )
mCdS RFCdS ) mCdS + mAAO aFCdS + (1 - a)FAAO XAAO )
(1 - a)FAAO
XCd ) ACd/MCdS, XS ) AS/MCdS XAl ) AAl/MAl2O3 (or MAl2O3‚H2O), etc.
volume fraction of the pores as follows:
ICdS ) I0
(A3)
aFCdS + (1 - a)FAAO
Figure 11. Calculated X-ray cross section (A) and intensity of the XRD profile (B) for AAO\CdS as a function of the pore/cell diameter ratio.
(A4)
where a is the volume fraction of the pores which can be derived from the known geometrical parameters of the template:
a ) Vw/Vtot ) spNph/Stoth ) spNp/Stot ) (1/4)πdp np, np ) Np/Stot (A5) 2
where Vw is the total volume of the pores and there are Np pores in total. This assumes that the pores are fully loaded with the deposited material. Vtot is the volume of the entire film containing Np pores, sp the average cross-sectional area of a pore, np the pore area density, dp the pore diameter, h the average film thickness, and Stot the total area of a film containing Np pores. The X-ray cross sections of the elements present in our samples and the calculated X-ray cross sections for the CdS, γ-Al2O3, and Al2O3‚H2O are listed in Table 4. Using these values and eqs A1 and A3, we can determine the expected intensity for the diffraction lines of CdS. Equation A1 can be written in a more convenient form in terms of the
(µF)
aFCdS
CdS
µ µ aFCdS + (1 - a)FAAO F CdS F AAO
()
()
(A6)
The mass adsorption coefficient for AAO\CdS can be determined as a function of pore and cell diameters, by first determining the volume fraction a in terms of these parameters. First, let us consider situation where the average pore diameter is increased by etching. The maximum pore diameter obtainable in this way is the so-called cell diameter, dc. When dp ) dc, the pores merge and the template (ideally) becomes an array of AAO columns with near-triangular cross sections. Using eq A5 and assuming hexagonal close packing of the cells, we obtain following expression for Stot and a:
1 2 πd N πdp2 4 p p π x3 2 ) d N wa) Stot ) dc Np sin ) 3 2 c p x3 2 2x3dc2 dc Np 2 (A7) 2
amax 1 (at dp ) dc) )
π 2x3
(A8)
Using (A7), one can determine the mass absorption coefficient of CdS in AAO, (µ/F)AAO\CdS, in terms of dp and dc (Figure 11a), which then allows one to calculate the XRD line intensity for CdS in AAO in terms of the same two variables using (A6) (Figure 11b). (The curves were corrected for pore overlap in calculating the portion of the curves shown in Figure 11 corresponding to dp/dc > 1. This situation is not achievable experimentally and is included in order to illustrate the changes
Electrochemical Fabrication of CdS Nanowire Arrays
J. Phys. Chem., Vol. 100, No. 33, 1996 14047
in the X-ray adsorption coefficient over the full range of semiconductor loading including the limit of pure CdS.) The line shape chosen to represent the X-ray diffraction peaks is the pseudo-Voigt function:
I0 ) ηL + (1 - η)G
(A9)
where L and G refer to Lorentz and Gauss functions, respectively. The line width was obtained from the Scherrer equation (1), assuming, for the purposes of producing the profiles shown in Figure 8 that the domain dimension equals the pore diameter, which is not, experimentally, the case (Figure 9). Nevertheless, this assumption helps illustrate the XRD profile behavior when approaching smallest wire diameter, where the diffraction peaks are very broad and of low intensity. References and Notes (1) (a) Switzer, J. A.; Shane, M. J.; Phillips, R. J. Science 1990, 247, 444. (b) Lashmore, D. S.; Dariel, M. P. J. Electrochem. Soc. 1988, 135, 1218. (c) Yahalom, J.; et al. J. Mater. Res. 1989, 4, 755. (d) Streltsov, E. N.; Osipovitch, N. P.; Routkevitch, D.; Sviridov, V. V. Manuscript in preparation. (2) Gregory, B. W.; Stickney, J. L. J. Electroanal. Chem. 1991, 300, 543. (3) (a) Diggle, J. W.; Downie, T. C.; Goulding, C. W. Chem. ReV. 1969, 69, 365. (b) O’Sullivan, J. P.; Wood, G. C. Proc. R. Soc. London, A 1970, 317, 511. (c) Wernick, S.; Pinner, R.; Sheasby, P. G. The Surface Treatment and Finishing of Aluminum and its Alloys; Finishing Publishing: Teddington, 1987; Vol. 1. (4) (a) Martin, C. R. Science 1994, 266, 1961. (b) Whitney, T. M.; Jiang, J. S.; Searson, P. C.; Chien, C. L. Science 1993, 261, 1316. (5) Menon, V. P.; Martin, C. R. Anal. Chem. 1995, 67, 1920. (6) Chlebny, I.; Doudin, B.; Ansermet, J.-Ph. Nanocryst. Mater. 1993, 2, 637. (7) (a) Foulke, D. G.; Stoddard, W. B. In Modern Electroplating; Lowenhiem, F. A., Ed.; Wiley: New York, 1963; p 632. (b) Preston, C. K.; Moskovits, M. J. Phys. Chem. 1988, 92, 2957. (c) Preston, C. K.; Moskovits, M. J. Phys. Chem. 1993, 97, 8495. (d) Saito, M.; Kirihara, M.; Taniguchi, T.; Miyagi, M. Appl. Phys. Lett. 1989, 55, 607.
(8) (a) Kawai, S. In Proceedings of the Symposium on Electrochemical Techniques in Electronics; Romankiw, L. T., Osaka, T., Eds.; Electrochem. Society: Pennington, NJ, 1987; PV 88-23, p 389. (b) AlMawlawi, D.; Coombs, N.; Moskovits, M. J. Appl. Phys. 1991, 70, 4421. (c) Dunlop, D. J.; Xu, S.; Ordemir, O.; AlMawlawi, D.; Moskovits, M. Phys. Earth Planet. Inter. 1993, 76, 113. (9) Pontifex, G. H.; Zhang, P.; Wang, Z.; Haslett, T. L.; AlMawlawi, D.; Moskovits, M. J. Phys. Chem. 1991, 95, 9989. (10) Klein, J. D.; Herrick, R. D., II; Palmer, D.; Sailor, M. J.; Brumlik, C. J.; Martin, C. R. Chem. Mater. 1993, 5, 902. (11) Chakarvarti, S. K.; Vetter J. Micromech. Microeng. 1993, 3, 57. (12) Alivisatos, A. P.; Harris, T. D.; Carrol, P. J.; Steigerwald, M. L.; Brus, L. E. J. Chem. Phys. 1989, 90, 3463. (13) Routkevitch, D.; Haslett, T. L.; Ryan, L.; Bigioni, T.; Douketis, C.; Moskovits, M. Special issue of Chem. Phys. on Confined Excitation in Molecular and Semiconductor Nanostructures, in press. (14) Baranski, A. S.; Fawcett, W. B. J. Electrochem. Soc. 1980, 127, 766. (15) Modern Powder Diffraction, ReViews in Mineralogy; Bish, D. L., Post, J. E., Eds.; Mineral. Society of America: Washington, DC, 1989; Vol. 20. (16) Klug, H. D.; Alexander, L. E. X-Ray Diffraction Procedures for polycrystalline and amorphous materials, 2nd ed.; J. Wiley: New York, 1974. (17) (a) See, for example, table of references in ref 2. (b) Electrochemical Deposition of Semiconductors. DeMattei, R. C., Feigelson, R. S. In Electrochemistry of Semiconductors and Electronics; McHardy, J., Ludwig, F., Eds.; Noyes Publishing: Parc Ridge, 1992. (18) AlMawlawi, D.; Routkevitch, D.; Moskovits, M. Unpublished results. (19) Baranski, A. S.; Fawcett, W. R.; McDonald, A. C.; de Nobriga, R. M. J. Electrochem. Soc. 1981, 128, 963. (20) Fatas, E.; Herrasti, P.; Arjona, F.; Camarero, E. C.; Medina, J. A. Electrochim. Acta 1987, 32, 139. (21) Scott, J. F.; Leite, R. C. C.; Damen, T. C. Phys. ReV. 1969, 188, 1285. (22) Kayanuma, Y. Phys. ReV. B 1988, 38, 9797. (23) Shiang, J. J.; Goldstein, A. N.; Alivisatos, A. P. J. Chem. Phys. 1990, 92, 3232. (24) Hayashi, S.; Sanda, H.; Agata, M.; Yamamoto, K. Phys. ReV. B 1989, 40, 5544. (25) CRC Handbook of Chemistry and Physics, 71st ed.; CRC Press: Boca Raton, FL, 1992.
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