How and Why Alumina Matrix Architecture Influence the Shape and Composition of Nanowires Grown by AC Deposition?

According to description, nanowires are one-dimensional materials with size ≤100 nm in two out of directions from which various architectures can be organized for recent devices offering new and sometimes unique oppurtunities. Among numerous methods can be ap‐ plied to date for densely packed nanowire (nw) arrays formation the template-assisted elec‐ trochemical deposition is attributed to the most widespread allowing simple control on the geometrical, morphological and crystallographic properties of various nanowire arrays in an independent manner. Note that in case of formation extremely thin and densely packed nanowires, demonstrating a significant improvement of their thermoelectric, photovoltaic, catalytic and optical properties, due to huge surface size and quantum-size effects, this pos‐ sibility becomes crucial (Bejenari et al.2011). Besides uniformity in wire diameter, morpholo‐ gy and composition, the crystallinity and crystallographic orientation also strongly influence the properties of metallic and semiconductor nanowires and their arrays (Lan et al.2009; Yan et al. 2010). However, most works to date have reported the growth of metallic and semicon‐ ductor nanowire arrays inside the alumina and polycarbonate (PC) templates with pores larger than 40 nm, especially in case of polymeric templates, and only few address the im‐ pact of pore diameter and deposition regime on the peculiarities of extremely thin nano‐ wires growth. For example, variables of the morphology, surface roughness and crystalline orientation of the Bi2Te3 nanowires with the PC membrane channel diameter decrease from 150 nm down to ~10 nm have been right now reported (Pitcht et al. 2012) demonstrating the possibility of obtaining the nanowire arrays with preferential growth of either {205}, {015}, or {110} planes perpendicular to the nanowire axis from a same composition of deposition sol‐ ution enabling us to tune their figure of merit and to improve the device performance.


Introduction
According to description, nanowires are one-dimensional materials with size ≤100 nm in two out of directions from which various architectures can be organized for recent devices offering new and sometimes unique oppurtunities. Among numerous methods can be applied to date for densely packed nanowire (nw) arrays formation the template-assisted electrochemical deposition is attributed to the most widespread allowing simple control on the geometrical, morphological and crystallographic properties of various nanowire arrays in an independent manner. Note that in case of formation extremely thin and densely packed nanowires, demonstrating a significant improvement of their thermoelectric, photovoltaic, catalytic and optical properties, due to huge surface size and quantum-size effects, this possibility becomes crucial (Bejenari et al.2011). Besides uniformity in wire diameter, morphology and composition, the crystallinity and crystallographic orientation also strongly influence the properties of metallic and semiconductor nanowires and their arrays (Lan et al.2009;Yan et al. 2010). However, most works to date have reported the growth of metallic and semiconductor nanowire arrays inside the alumina and polycarbonate (PC) templates with pores larger than 40 nm, especially in case of polymeric templates, and only few address the impact of pore diameter and deposition regime on the peculiarities of extremely thin nanowires growth. For example, variables of the morphology, surface roughness and crystalline orientation of the Bi 2 Te 3 nanowires with the PC membrane channel diameter decrease from 150 nm down to ~10 nm have been right now reported (Pitcht et al. 2012) demonstrating the possibility of obtaining the nanowire arrays with preferential growth of either {205}, {015}, or {110} planes perpendicular to the nanowire axis from a same composition of deposition solution enabling us to tune their figure of merit and to improve the device performance.
In this chapter, the influence of alumina template barrier layer thickness on the electrochemical growth of copper and cobalt nanowires is demonstrated. Our results obtained investigating the peculiarities of bismuth selenide electrodeposition by alternating current (AC) treatment in alumina templates varied in pore diameter within 10 to 100 nm range are presented in subsection 5.1 showing, for the first time, a strong dependency of formed nws composition, morphology and their optical properties on the diameter of pores (Ø pore ).

Filling of porous alumina templates
Porous oxide films (alumina) prepared via anodization of high pure and smooth aluminum surface in an aqueous solution of sulfuric, oxalic, and o-phosphoric acid at proper concentration, temperature, and voltage are typical templates for fabrication arrays of various nanowires in quite uniform diameter and spacing with well-defined product dimensions at packing density of 10 9 -10 11 species/cm 2 Jassensky et al. 1998; Li et al. 1998;Nielsch et al. 2002). To date high-ordered structure alumina with broad range of pore diameters as large as 300 nm (Quin et al. 2005) and as small as 5 nm (AlMawlawi et al. 1991) can be fabricated in unlimited size area. The pore diameter, cell size, and barrier-layer thickness positioned at the bottom of alumina pores ( Fig. 1) are all linearly dependent on the anodizing voltage (O'Sullivan & Wood 1970), while the depth of pores can simple be controlled by anodizing time (Metzger et al. 2000). According to the final applications, the thickness of alumina templates is usually limited to 20-30 μm, but thicker templates (Römer & Steinem 2004), as well as very thin (Kokonou et al. 2007), are sometimes required for uniform filling with various nw arrays. Till now, direct current (DC), alternating current (AC) and pulsed current depositions have been developed for filling of alumina pores by various materials. For DC depositions, enabling substantial control over composition and crystallinity of nws, the detachment of alumina film from the substrate, opening the pore bottoms via voltage decrease and chemical etching and conductive layer formation are usually required. To remove the barrier-layer only at the bottom of alumina nanochannels or perforate it, various etching solutions (Zheng et al. 2002) and different methods (Mardilovich et al. 1995) have been proposed during last decade. Also, the sputtering of gold (Yoo & Lee 2004), silver ) and platinum (Panet al. 2005) layer onto the back-side of perforated template as a conducting and well-adhesive layer have been applied. On the other hand, there always exists the possibility of filling alumina pores by AC modes, leaving intact the barrierlayer at the bottom of pores because anodic aluminum oxides conduct preferentially in cathodic direction. According to some opinions (Goad & Moskovits 1978;Clebny et al. 1993), AC electrolysis is an ideal method for deposition of metals and semiconductors, starting from the bottom of pores. Moreover, AC depositions through the rectifying barrier-layer require fewer processing steps and are more amenable to scall-up but currently provide far less control over the structure and the material deposited (Gerein & Haler 2005). As a result, different groups using this method (Preston & Moskovits 1993; Sheasby & Cook 1974) have observed interrupted growth of various polycrystalline materials and just partial depositions, namely only in a fraction of pores. Note that for alumina templates formed at higher anodizing voltages and consequently having a thicker barrier-layer, some degree of this layer thinning is essential to enable deposition even under AC treatment. Furthermore, the chemical composition of alumina films differs depending on the composition of anodizing solution as well as anodizing conditions due to incorporation of acid anions and water molecules into the outer part of alumina cells. For example, the sulfuric acid alumina films contained high amount of anion species (12-14 wt.% sulfate), while the phosphate and oxalate contents in corresponding alumina films are respectively 6-8 and 2-4 wt.% (Thompson 1997). Note that incorporated anion species produce a negative surface charge of the pore walls influencing the pore filling process by one or another material. Therefore, the hydrophobic/ hydrophilic pore wall properties could play a significant role for nanowires growth in the precursor solution.

The influence of porous alumina barrier-layer
Potentially nanowires in the alumina template pores can be synthesized by adsorbing and decomposing precursor species, high-pressure injection of a melt, electro less and electro deposition routes. However, the dominant synthesis technique in this area remainsAC deposition of metals and semiconductors copying exactly the pores configuration. However, the optimal AC electrolysis conditions differ for various solutions and various templates whereas under others the following phenomena as the alumina barrier-layer spalling (Sheasby & Cooke 1974), pitting corrosion (Routkevich et al. 1996) and the template peeling off from the substrate (Doughty et al. 1975) take place. Besides, the filling of alumina templates grown at higher voltages in the phosphoric or oxalic acid baths is more problematic, since at higher AC voltages, required in this case for metal ions discharge, the alumina barrier-layer breakdown is inevitable. A few investigations devoted to uniform growth of metallic nanowires by means of pulse or AC electrolysis deal in this case with requirement to decrease the thickness of the alumina barrier-layer (d), (Forrer et al. 2000;Xu et al. 2002;Sauer et al. 2002). However, there is much uncertainty about the optimal d for uniform filling of alumina pores with metal nanowires. It is still far from clear which barrier-layer thickness and AC voltage value will be optimal for the uniform filling of one or another alumina template for given material. Reported in Paulus et al. study (2001) d values for deposition of Fe, Ni, and Co nanowires range from 10 to 20 nm. According to recent study (Sausa et al. 2011), the optimal barrier layer thickness of alumina templates for homogenous and complete filling of all pores with Ni by AC treatment approximates 10 nm. It is likely, however, that highly uniform deposition of copper into oxalic acid grown pores by AC treatment is significantly more challenging than deposition into sulfuric acid grown template pores because of the different chemistry and structure of their barrier-layers (Gerein & Haber 2005). Therefore, the influence of dis still an open issue since the composition and properties of various alumina barrier-layers are complex and ill defined, especially after thinning through voltage decrease and chemical etching.
To shed light on this problem, in this study we focused on the use of the dependency the plots of the quantity of copper deposited within the template pores on the alumina growth and subsequent AC deposition conditions expressed as AC electrolysis and the template forming voltages ratio.
The amount of copper was determined after the complete dissolution of Cu 0 nanowires from a 4.5 cm 2 template surface in HNO 3 : H 2 O (1:2) solution (2 cm 3 ) for about 3 min. All solutions used for nanowires dissolution, sample rinsing and further double soaking in triply distilled water for 3 min were mixed together, diluted to constant volume and analyzed quantitatively using a Perkin Elmer Lambda 35 UV/Vis spectrometer. To increase absorbance detection sensitivity of copper analyte at 450 nm, 1 % sodium diethyldithiocarbamate (2.0 cm 3 ) was used as a complexing agent. Reproducibility of the analysis was checked by 3 repeated experiments. Standard solutions were made from 99.999 % grade copper.
The electrochemical impedance spectroscopy (EIS) was utilized to characterize the alumina barrier-layer properties upon the anodizing bath voltage decrease. The EIS spectra were recorded using a VoltaLab 80 (Radiometer Analytical, France) electrochemical system for frequencies between 1 and 10 5 Hz with ten measurements and are presented in the complex capacitance form, that is, the plots of ImY/ω versus Re Y/ω, where Y is the admittance and ω is the angular frequency [ω = 2πf, and f is the frequency in hertz (Hz)]. These plots allow simple models of the interface to be readily inferred when the electrochemical response exhibits capacitive behaviour. The amplitude of the applied AC signal was set to 10 mV. The spectral data were analyzed/fitted with Z View software (Scribner Associates, South Pines, NC, USA). High-frequency stray capacitance cross-talk was eliminated from the EIS spectra as described earlier (Vanderah et al. 2003).

Peculiarities of the alumina barrier-layer thinning
It is well-documented that the thickness of the barrier-layer of alumina templates is proportional to the voltage applied in the anodizing process (Diggle et al. 1969). Going through the voltage drop, at the end of anodizing, one could expect the decrease in the barrier-layer thickness due to the field-assisted ejection of Al +3 from the oxide lattice to the solution (Harkness & Young 1966). However, this process can proceed only in some pores if a large voltage decrement is applied in a single step (Furneaux et al. 1989) causing the increased heterogeneity in the physical properties of the films. Therefore, one might expect the concomitant changes in the EIS spectra reflecting these phenomena. Figure 2 displays a set of the EIS spectra obtained for the alumina films of different anodization end-voltage values. As seen, all the spectra possess a typical semicircular shape, which suggest simple capacitive behavior (Vanderah et al. 2003). This indicates the dielectric film that can be modeled as an equivalent circuit of a series of R s C elements, where R s is the solution resistance and C is the capacitance of the alumina barrier-layer. Though being similar in shape the spectra in Fig.2 show several trends and differences. First, as the final anodizing voltage decreases the diameter of the semi-circles increases, signaling about the increase in capacitance of the barrier layer. Second, the plots of ImY/ω versus Re Y/ω of the films obtained using high end-voltage values (10 -40 V) tend to approach the limit:limReY / ω → 0 ω→0 at the low frequency edge of the EIS spectra indicating nearly perfect capacitive behavior. While, the films obtained using low final anodizing voltages do not. This observation suggests that the low end-voltage films exhibit significant micro-heterogeneity of the physical properties, while the opposite is true for the high end-voltage films. Taking into account these observations, for fitting the experimental EIS spectra data to the R s C model, we replaced the capacitance with the constant phase element (CPE) as Macdonald (1987) suggested. The CPE reflects the deviation of alumina impedance from the ideal behavior. In the presence of a CPE, the film impedance exhibits a law frequency dependence: The results of the fitting to this model are summarized in Table 1.As seen from Table 1, the numerical values of C' and α are consistent with the qualitative features of the Re Y/ω vs. ImY/ω spectra, i.e., the gradual lowering of the anodizing end-voltage yields to the increase in the CPE's coefficient and concomitant decrease in the CPE's exponent value.  It is well documented, that aluminum anodizing voltage and the thickness of alumina barrier-layer are related through the equation δ b = κU a,fin where κ is the "anodizing ratio" coefficient close to 1 nm V -1 (Diggle et al. 1969). The estimated thickness values are tabulated in Table 1, column 1(numbers in the brackets). On the other hand, the capacitance and thickness of the barrier-layer are related through the following equation: where: ε 0 is the vacuum permitivitty, 8.85 10 -14 F/cm, ε is the dielectric constant of alumina, hereinafter, the value 9.8 is accepted (Harkness & Young 1966), A is the surface area of the electrode, 0.33 cm 2 , d -is the thickness of the barrier-layer and β' is the factor that accounts for the fraction of the surface occupied by the alumina pores (0 <β' < 1), and β" is the roughness factor of the aluminum surface (β" > 1) (Saif et al. 2002); β =β'β". Using Eq.(1) and assuming the approximate equality C≈C' it is possible to verify whether or not the experimental capacitance (constant phase element) values follow the expected barrier-layer thickness sequence. Figure 2A shows the experimental dependence of the C vs.d -1 , which is linear in the range 40 to about 5 nm. The slope of the line is 2.81 10 -13 F cm, which yields a quite realistic value of β= 0.98. However, below 5 nm there is a clear deviation from linearity (not shown). These results imply that during re-anodizing the barrier-layer thickness linearly decreases with (U a,fin ) -1 only down to about 5.0 nm. The further decrease in U a,fin results in slower reduction of d because the thickness of Al native oxide film exceeds 5 nm at room temperature (Saif et al. 2002). In other words, at U a,fin < 5.0 V, the chemical interaction between aluminum and the ambient changes the relationship between the U a,fin and d. In addition, it is likely that the constant phase element decrease from 0.98 to about 0.94 reflects the enhanced fluctuations of d from pore to pore, yielding more heterogeneous distribution of the physical and chemical properties of the barrier-layer.

The m Cu versus U v /U a,fin plots for copper nanowires growth
Our experimental data suggest that a broad range of AC voltages can be used for deposition of copper within the alumina template pores if a suitable composition of the solution is chosen.
Unfortunately, this is applicable only for alumina grown in sulfuric acid bath where AC voltages U v 6.0 to 18.0 V can be successfully applied (Fig. 3, curve 2). All attempts to fill more completely alumina templates grown in oxalic and phosphoric acid solutions varying AC voltage and the deposition time have failed since only limited voltages can be used in these cases. At higher AC voltages or somewhat longer electrolysis durations the barrier-layer breakdown of these templates was found to be inevitable (curves 1 in Figs. 4 and 5). As seen from curve 1 in Fig From these plots one can see that the behavior of alumina templates in acidic Cu(II)solution manifests itself through parabolic m Cu vs.U v /U a,fin dependencies. Moreover, we have found that a range of AC voltages at which deposition of copper nanowires proceeds within the alumina pores strongly depends on the U a,fin , decreasing with the alumina barrier-layer thickness. It should be also noted, that at AC voltages higher than m Cu vs.U v /U a,fin curve peak, (U v /U a,fin ) peak , the alumina spalling and peeling off from the substrate becomes critical, especially when U a,fin > 20 V.Therefore, attention was focused on the range of AC voltages suitable for Cu 0 nanowires growth without the barrier-layer breakdown. As clearly seen from the experimental results depicted for various templates in Figs. 3 to 5, the range of voltages suitable for copper deposition shifts to the higher ratio of U v /U a,fin , the lower U a,fin . On the other hand, despite the fact that a much wider range of AC voltages can be used for copper deposition within the pores of alumina with a quite thin barrier-layer, it seems impossible to completely fill such a matrix even at the U v /U a,fin ratio 3.0 when d< 5 nm (U a,fin < 5 V). This is due to a sharp decay in the current strength even during first 10-20 s of AC electrolysis up to a few mA cm -2 ; characteristic feature for films with d lower than 5.0 nm. In these cases only imperceptible quantity of copper can be deposited within the alumina pores at the electrode edges since the gas bubbles evolved at the Al│oxide interface push away the alumina film from the substrate. The appearance of the gas bubbles at the Al│oxide interface was clearly observed by the naked eye. It is also common for alumina templates having thicker barrier-layers if prolonged AC electrolysis and higher AC voltages are used. Consequently, it seems difficult to estimate one optimal U v /U a,fin for Cu 0 nanowires deposition within the pores of different templates. Evidently, the optimal AC voltage is lower (U v / U a,fin ) peak and the final choice is up to uniformity and completeness of the filling of the alumina pores. On the other hand, it has been found that the range of AC voltages suitable for copper nanowires fabrication depends also on the composition of Cu(II) solution. The m Cu vs.U v /U a,fin plots presented in Fig. 6 show that an increase in the solution pH widens the range of AC voltages suitable for copper deposition. Although the reasons of such behavior are unclear yet, we believed that this is most probably accomplished by a sharp decrease in the gas quantities evolved simultaneously with copper deposition from neutral and alkaline Cu(II) solutions at the Al│oxide boundary (Jagminas et al. 2002).

Concluding remarks
The above results show that in order to grow copper nanowires within the pores of alumina template obtained by Al anodizing at higher voltages the alumina barrier-layer thickness should be lowered. Using the acidic anodizing baths, a linear dependence of d on the final anodizing voltage, U a,fin , is observed down to 5 V. The linearity was verified by impedance spectroscopy data, so, this technique might be used to monitor the parameters of the alumina template formation. However, further U a,fin decrease below 5 V results in significant barrier-layer thickness fluctuations, which are possibly due to uneven native oxide formation at the bottom of the pores. Spectrometric analysis of deposited copper content has shown that the range of AC voltage suitable for copper nanowires growth within the alumina pores sharply increases with decrease in U a,fin and hence d. The most favourable U a,fin range for uniform copper nanowires growth is 15 to 7.0 V.

Alumina template-dependant growth of cobalt nanowire arrays by AC deposition 4.1. Depositions into as-grown templates
In this study, different electrochemical regimes and porous alumina fabricated by aluminum anodizing in either sulphuric or oxalic acid solutions were applied for template synthesis of cobalt nanowire arrays, revealing several peculiar cases. By this way, we found that the growth of cobalt nanowires depends much stronger on the conditions for fabricating the alumina template than other metals like copper, silver or tin. For example, only sulfuric acid alumina templates can be successfully filled in the optimized deposition solutions by Co nws using AC, while the use of the same solution for uniform growth of Co nws in oxalic or phosphoric acid alumina pores is problematic. Therefore, in this study we focus on the investigation the peculiarities of the Co nws electrochemical growth using oxalic and sulfuric acid alumina templates at different regimes.
In this study, the composition of solution for Co nws electrosynthesis within the alumina pores was organized using CoSO 4 , as a precursor for cobalt ions, and H 3 BO 3 , as a buffering ingredient, to prevent any pH variation within the alumina pores and to inhibit the formation of cobalt hydroxide species, as a result of hydrogen evolution ( Figure 7 demonstrates typical variation of the amount of Co assembled inside the alumina pores (m Co ) with the deposition time and AC current frequency (f) used for depositions. As seen, the amount of deposited Co increases linearly during the first 10-20 min of AC treatment at constant voltage. However, with further processing the rate of deposition inside the alumina pores progressively decreases. Moreover, it was observed that this solution allows the growth of cobalt nws within a wide range of AC frequencies, ca. from 10 to 200 Hz, coloring the template uniformly from bronze to deep black. A further increase in frequency, up to 1000 Hz, results in a smaller amount of deposited cobalt and therefore in a lighter template color intensity. As seen from the inset of Fig. 7, the maximum amount of cobalt can be deposited using 100 Hz frequency. The bath temperature within 10 to 40°C range was found to have negligible effect on the amount of deposited cobalt as well as on the uniformity of depositions.  In order to visualize the uniformity of Co nws growth by AC deposition, cross-sections of the alumina templates were investigated using field emission scanning electron microscopy (FESEM). Figures 8A-C show the arrangement of Co nws deposited inside the pores of sulfuric acid alumina templates at the same AC frequency (50 Hz) and peak-to-peak voltage (U p-p ) during 5, 15 and 60 min, respectively. The obtained data show quite uniform growth of Co nws from the bottom of almost all pores only at the onset of the process. The pores filling rate depends on AC voltage and of the pores diameter, Ø pore . In case of sulfuric acid alumina templates formed at 25 V (Ø pore 25-30 nm after pore widening), the uniform growth corresponds to a filling rate v Co~ 6 μm/h for the first 10 min of deposition at a constant U p-p of 32 V. Processing further, v Co decreases due to an increase of the template resistance, causing a reduced AC current. A smaller Ø pore results in a faster growth of cobalt nws under the same AC treatment conditions. For example, v Co~ 9.3 μm/h was detected at U p-p of32 V for alumina templates with average Ø pore of 15 nm. It can be observed from Fig. 8B that some cobalt nws grew faster than nws front. For prolonged AC treatment, this progressively leads to the formation of more and more uneven lengths of Co nws emerging onto the template surface in cobalt caps ( Fig. 8C). The nws height uniformity was found to be independent on the f. Typical morphology of cobalt nw arrays assembled inside the sulfuric acid alumina pores by short-term AC deposition after template etching is presented in Figure 9. As seen, in case of 1-2 μm length of Co nws they are densely packed and quite uniform.

Depositions through a reconstructed barrier-layer
The key feature of AC deposition process of the densely packed Co nws is that only sulfuric acid alumina templates can be successfully applied. Fabrication of Co nws in the nanochannels of alumina template formed in the oxalic or phosphoric acid anodizing baths, however, is problematic byAC deposition way. One possible explanation for this effect is the adsorption and incorporation of acid anions at some depth of the alumina barrier-layer, changing the state of alumina/solution interface at the bottom of pores (surface charge, free energy, etc.) and preventing the discharge of Co 2+ ions. Besides, highly ordered oxalic and phosphoric acid alumina templates are usually formed at higher voltages Masuda et al. 1997;Li et al. 2000) and, therefore, present much thicker barrier-layers at the metal│oxide interface. To use these templates for AC deposition of various materials, the step-wise voltage decreasing at the end of anodizing process has been proposed (Furneaux et al. 1989) and successfully used in several works. We found herein that this is helpful also for the Co case, however, only for short-time processing. The Co nw array produced by longAC treatment, i.e. longer than 15 min, viewed nonuniform from pore to pore with some mushroomed Co fragments (Fig. 10) outgrowing from the breakdown sites of the alumina barrier-layer. The modification of alumina barrier-layers through (i) the decreasing of anodizing voltage (U a ) at the end of oxalic acid alumina growth down within 13 to 5 V, (ii) the cathodic treatment in the same anodizing bath for 3 min at various potentials and (iii) the chemical etching in the solution of sulfuric acids inhibit the uniformity of the depositions (see Fig. 11). The most uniform alumina color was obtained after decreasing U a and chemical etching in the solution of sulfuric acid. Nevertheless, in this case the SEM crosssectional observations of templates revealed the formation of Co nws tufts in random areas of template (Fig. 12A). It is worth to note that these tufts were found to arise from cobalt balls (Figs 11B and 12 C) formed at the metal/template interface.  Variations in the conditions of cathodic treatment and chemical etching of the oxalic acid template as well as variations of the deposition potential were found to be ineffective for rod-like Co formation by AC deposition through the remained barrier-layer. Moreover, in case of DC deposition, the detachment of alumina template from the substrate even after several minutes of treatment took place.

Galvanostatic DC deposition
An alternative approach for Co nws deposition was further examined for oxalic acid alumina templates by a DC constant current density deposition, after removing or perforating the barrier-layer on the pores bottoms. In this setup, we used either an electrochemical/chemical method for the barrier layer perforation, or we detached the alumina from the substrate, removed the barrier layer and made a subsequent evaporation of Cr/Au layer, acting later as a conducting pad. The resulting Co nws released from the as-filled templates are shown in images A and B of Fig. 13. We found here that by applying a low current density during the entire deposition process, ca. ≤ 0.2 mA cm -2 , well-ordered, densely packed, continuous and highly aligned Co nw arrays, faithfully reproducing the shape of the pores and with height up to several tens of micrometers, can be synthesized by this way. Furthermore, after alumina dissolution these nws seems not to collapse and stuck together, as in the case of AC synthesis, implying an easier their application in future nanoelectronics and novel efficient sensors. A typical XRD profile of a template filled with Co nws via DC deposition at a constant current density of 0.12 mA cm -2 for 5 hours is shown in Figure 14. Only a single peak is observable at 2Θ = 41.59. According to the XRD library patterns for bulk Co (PDF 89-4308), this peak corresponds to the (100) reflection of the hexagonal closely packed Co lattice. Further, some additional weak signal situated at 2Θ = 75.89, ascribed to hexagonal Co phase in (110) direction, can be observed. This weak feature probably inferred that Co nws are not single crystals but consist of oriented polycrystals with a preferred (100) growth direction perpendicular to the substrate. We note that the preferential growth of hexagonal Co phase in (100) direction is not a trivial case and differs from the growth of Co nws inside the alumina pores via AC and potentiostatic depositions reported by Kartopu et al. (2008) where the formation of Co nw arrays with a preferred (110) orientation has been demonstrated.   In this article, we report the results of investigation on peculiarities of bismuth selenide electrode position byAC treatment in different alumina templates varied in pore diameter, Ø pore , within 10 to 100 nm range demonstrating, for the first time, a strong dependency of formed nws composition, morphology and their optical properties on the Ø pore .

Compositional, structural and optical properties of bismuth selenide nws synthesized by template approach
Porous alumina templates, 5.0 to 10.0 μm thick, were grown via two-step direct current (dc) anodizing of specimens for 0.5 to 20 hours in thermostated and vigorously stirred solutions under conditions indicated in Table 2. First anodization lasted two hours following the alumina film stripping in 0.2 M CrO 3 + 0.7 M H 3 PO 4 at 60 ºC for 5 hours, while the second oneas indicated in Table 2. The barrier layer of alumina films after the second anodizing was set to about 15 nm by reduction of the anodizing voltage step-wise (2-1 V per 30-60 s) as in the study of Furneaux et al. (1989). All depositions of bismuth selenide nws within the alumina template pores were performed at room temperature in a glass cell where two graphite rods were used as the auxiliary and Al/alumina as a working electrode. The solution containing 0.02Bi 2 (SO 4 ) 3   To increase crystallinity of the final-products, samples were annealed for 3 hours in vacuum.
The optimal annealing temperatures (T ann ) for various alumina templates were found in this study experimentally and approximated to: 250 o C for 100 and 50 nm, 200 o C for 28 and 13 nm and 170 o C for 10 nm Ø pore templates. The different values of T ann were chosen due to the well-known dependency of melting and crystallization temperatures of nanomaterials on their size (Noh et al. 2007).
The alumina templates intended for XRD and UV-vis-IR investigations were separated from the electrode surface by one-side sequential etching of the electrode window in a solution of 1.5 mol/LNaOHand then in 10 wt% HCl and 0.1 mol/L CuCl 2 followed by thorough rinsing and drying in a nitrogen stream. X-ray diffraction studies were performed with a diffractometer D8 (Bruker AXS, Germany) equipped with a Göbel mirror (primary beam monochromator) for CuK α radiation. To prepare TEM samples, deposited species were liberated by dissolving alumina template in 0.1 mol/L sodium hydroxide at 40 o C. The released products were then rinsed many times and finally dispersed in ethanol. At each stage, solvent exchange was carried out by centrifuging, extracting the supernatant and adding fresh solvent. Finally, free-standing nanospecies were re-dispersed in ethyl alcohol. For TEM observations, specimens were prepared by placing a drop of suspension on a Lacey carbon grid and left overnight at room temperature to evaporate the solvent. Nanostructured products were examined with a TEM microscope (model MORGAGNI 268) operating at 80 kV.
Optical properties of bismuth selenide arrays fabricated inside the alumina template pores were studied by recording the transmittance spectra within the 190 to 3150 nm wavelength range with respect to pure alumina template using a Shimadzu UV-3101PC spectrophotometer. The transmission data were manipulated for calculating the absorption coefficient dependency on the photon energy.

Results
Uniform filling of alumina pores by densely packed Bi 2 Se 3 nws with an average pore diameter Ø pore = 40-50 nm in the solution and conditions adopted in this study were demonstrated by us in (Jagminas et al. 2008). Figure15 presents typical FESEM images of the final products encapsulated within the alumina template pores with average diameter 10, 13, 25 and 50 nm showing that the diameters of nws grown inside the alumina pores by ac deposition are in agreement with the nominal pore diameter of templates while the height of deposited products depends on the current density, j ac , deposition time, τ dep , and Ø pore . Under the same deposition conditions, an increase in the Ø pore resulted in filling of pores of a lower height. Furthermore, variations in Ø pore do not noticeably altered the completeness of the pore fill-ing; for all cases most of the pores, especially at the metal/oxide interface with heights up to 1 μm, seem nicely filled. On the other hand, the dispersity of nws lengths increases with τ dep , especially at higher j ac , and AC voltages. Surprisingly, the pitting of alumina films, frequently observed in other solutions as a result of the alumina barrier-layer breakdown and crystallization of salts (Jagminas 2002), was not observed for the solution of this study under a wide range of deposition conditions: j ac up to 0.5 A dm -2 , τ dep up to 45 min, and Ø pore up to 100 nm. The influence of the solution temperature on the composition of products was also investigated here. In all cases increase in the solution temperature higher than 40 °C led to non-uniform depositions within the alumina pores of products in the lower quantity most likely due to alumina pore sealing, as could be expected. Figure 16 shows XRD patterns of porous alumina templates with average Ø pore 100, 50, 25, 13 and 10 nm filled with Bi x Se y Nw arrays by AC electrolysis in the same optimized solution at the same constant j ac , ca. 0.3 A/dm 2 , for 25 min. From the patterns, it has been found that under the same electrolysis conditions Bi 2 Se 3 , Bi 3 Se 2 or both phases of selenides can be deposited. As seen from Fig. 16a, a quite pure, Se-rich phase, Bi 2 Se 3 , grows when templates with Ø pore ≥ 50 nm are used. However, Bi-rich phase, Bi 3 Se 2 , appears to form more readily when fine structure templates with Ø pore ≤ 13 nm are employed (Fig. 16b). Moreover, the Nws array grown under the same conditions within extremely tiny pores, with Ø pore = 10 nm, was found to be composed of Bi 3 Se 2 and some Bi 0 inclusions while in the case of Ø pore = 28 nm the nws are composed of Bi 2 Se 3 with some amount of Bi 3 Se 2 . In the case of tiny pores (Ø pore = 13 and 10 nm), increase in the molar ratio of the selenium and bismuth precursors (α Se/Bi ) from 1.0 to 2.0 results in the formation of BiSe (α Se/Bi = 1.5) and finely grained Bi 3 Se 2 (α Se/Bi = 2.0) without Bi 0 inclusions. Also, with α Se/Bi increase a somewhat slower bath voltage growth has been determined during the deposition process at a constant AC current density.
The influence of the solution temperature on the composition of products was also investigated here. In all cases increase in the solution temperature higher than 40 °C led to nonuniform depositions within the alumina pores of products in the lower quantity most likely due to alumina pore sealing, as could be expected. For the same Ø pore , however, no changes in the phase composition of deposited products have been observed within 5 to 35 °C. The compositional variations of deposited nws with the size of alumina template pores can be explained as follows: It is known that Se-rich thin films of bismuth selenide, namely Bi 2 Se 3 , may be electrodeposited only when enough Se is present in the reaction zone. However, as it has been shown earlier by us (Jagminas et al. 2005), the discharge of SeO 3 -2 ions from aqueous solutions of selenious acid at the bottom of tinny pores under ac bias is hampered. As a result, the size of a-Se species that can be deposited drastically decrease with Ø pore shrinking. Surprisingly, the increase in the ac current density and electrolysis time influenced the content of deposited a-Se in these pores only negligible. In contrast, the content of selenium deposited in wider pores, ca. 40-50 nm, under the same conditions increased many folds. Thus, the formation of Se-rich bismuth selenide nanowires in the oxalic acid alumina pores (Ø pore ≥ 40 nm) can simply be released. Again, only Bi-rich phases can be deposited in the sulfuric acid alumina templates with Ø pore ≤ 13 nm.   In this study, the optical properties of bismuth selenide Nw arrays fabricated within the pores of various alumina templates were investigated by recording the transmittance UVvis-NIR spectra using the same pure alumina templates as reference. In this set-up, the thickness of alumina templates varied between 5.0 and 7.5 ± 0.5 μm and both as-grown and annealed in vacuum at 170, 200 and 250 o C templates with various Ø pore were studied. To achieve more precise results, the working and reference samples were anodized, post treated and annealed together. Again, the optical band-gap, E g , for as-grown and annealed arrays was calculated using a well-known Tauc's relation: where α is the absorption coefficient, A is a constant, hν is the photon energy, E g is the band gap, and n depends on the nature of transition, being equal to 1 or 3 for direct-allowed or direct-forbidden and 4 or 6 for indirect-allowed or indirect-forbidden transitions, respectively. The absorption coefficient was calculated from the transmittance spectra using a simple relation: The average height of bismuth selenide nws, h BiSe , was determined from the FESEM observations of cross-sectioned templates after the optical measurements. By this way, the predominant direct band-to-band transition across the gap of different wavelengths was verified for all bismuth selenide nw arrays fabricated in this study. Typical plots of a 2 versus the corresponding values of photon energy, hν, for Bi x Se y nws deposited inside the alumina pores with different Ø pore are given in Fig. 17. As seen, variation of α vs. hν demonstrates a wide light absorption region from NIR to UV. In case of alumina templates with average Ø pore of 100 nm (curve 4) extrapolating the straight line part of the curve α 2 vs. hν to the energy-axis, the value of E g,dir equal to 0.4 eV was obtained for as-formed Bi 2 Se 3 nw arrays that is close to E g = 0.35 eV of bulk bismuth selenide. Note that with decrease in the diameter of Bi 2 Se 3 nws, the absorption of higher energy light increases. In case of Ø pore = 25 nm, the shape of α 2 vs. hν plot implied two absorption edges, perhaps due to the deposition within such alumina pores of species composed of a Bi 2 Se 3 and Bi 3 Se 2 mixture, as it has been shown above by XRD investigations. For Bi 3 Se 2 nanoscaled products encased within the alumina template pores with Ø pore 10 and 13 nm, the α 2 vs. hν plots demonstrate similar shapes (see curves 1 and 2). However, we found that the straight parts both in α vs. hν and α 2 vs. hν plots are not clear and thus numerous tangents can be extrapolated to the energy-axis from these plots indicating, for example, that E g value for 13 nm nws could be between 0.9 and 1.7 eV, while for 10 nm nws E g approximated to from ~2.2 to 2.7 eV. Consequently, the effective band gaps of these arrays cannot be precisely distinguished from the absorption spectra. Notice that these results are in line with the results presented in the recent publication (Sun et al. 2008) where the same problem raised analyzing the absorption spectra of CdTe quantum wires. Nevertheless, an obvious blue shift of α 2 vs. hν plot is observed for 10 nm nanowired products, e.g. when the diameter of alumina template pores approach to the Bohr radius of bismuth selenides (see Inset in Fig. 17).  To understand absorption variables of bismuth selenide arrays fabricated herein, we further studied the morphology of products, deposited within the alumina templates with Ø pore 50 nm and 13 nm using modern high-resolution FESEM and TEM techniques. Shown in Figure  18 are the high resolution panoramic (A) and top-side (B) FESEM images of alumina films encased with Bi 2 Se 3 species before (A) and following the template etching with a drop of 0.5 mol/LNaOH, while (C) depicts the TEM image of the same product disengaged from the template through the template dissolution and collection of the remained species by centrifugation and several washings. As seen from images B and C, Bi 2 Se 3 deposited within the 50 nm pores has a granular shape. This granular structure can be also visualized from the high magnification FESEM observation of nanowired product (see Inset). Through the TEM observation the size of granules varied within the 17 to 40 nm range implying that 50 nm nws of Bi 2 Se 3 are composed of weakly connected nanocrystals. A similar morphology was also observed for disengaged species of bismuth selenides deposited within 25-28 nm pores. In the case of tiny pores, however, the structure of the deposited material after dissolution of alumina matrix was found to differ significantly from the 50 nm products in that it consists of short and tinny nw fragments in length of from 0.2 to 1.0 μm (see Fig. 19) even after the template etching and liberation from the matrix procedures. It is some confusing knowing that in the case of Ø pore = 13 nm the deposited material is not a phase pure material but according to XRD results consists of a Bi 3 Se 2 and Bi 2 Se 3 mixture.

Discussion
Bi 2 Se 3 is reported as a direct band gap semiconductor material. Of special note is that the band gap (E g ) values of the Bi 2 Se 3 crystals determined from the optical measurements by different scientists disagree strongly both for bulk materials and nanomaterials. For example, for bulk Bi 2 Se 3 Novoselova(1978) reported E g = 0.35 eV, while Lide(1991) only E g = 0.16 eV. In the case of Bi 2 Se 3 films fabricated by electroless deposition, the presence of two edges corresponding to E g = 0.354 eV and E g = 1.03 eV has been reported by Bhattacharya and Pramanik(1980). Moreover, for thin Bi 2 Se 3 films the band gaps as high as 2.3 eV has been reported by Pejova and Grozdanov (2002) linking such high values with a nanocrystalline film nature. The great E g variations have been also reported for nm-scaled Bi 2 Se 3 including 1.59 eV (Jiang et al. 2006) and 2.25 eV (Ota et al. 2006). Noteworthy that the reported variations in the band gap values of bismuth selenide thickfilms and nanomaterials frequently are ascribed more to the morphology and purity of this semiconductor than to the size quantization effects. Consequently our findings of high E g of bismuth selenide species deposited within alumina pores under conditions of this study cannot be considered as surprising. Figure 19. A top-side FESEM view of alumina surface chemically etched by a drop of 0.5 mol/LNaOH for a case when alumina template with average Ø pore = 13 nm is encased with bismuth selenide nws by AC deposition as in Fig. 15. In the Inset, 13 nm in diameter a product fragment acquired at 1 kV using a concentric BackScatter detector is shown.

Concluding remarks
We initiated this study to show for the first time that the composition and morphology of bismuth selenide deposited within the alumina template pores by means of alternating current deposition depend on the diameter of pores. Under the same electrolysis conditions nw arrays from Se-rich (Bi 2 Se 3 ), Bi-rich (Bi 3 Se 2 ) or both phases can be successfully fabricated if the templates with average Ø pore ≥ 50 nm, Ø pore = 28 nm, and Ø pore ≤ 13 nm, respectively, are used. Also, the optical properties of bismuth selenide nws differing in size and phase composition, differ significantly. The band-edge absorption at 0.4 eV, characteristic to bulk E g of Bi 2 Se 3 ,was clearly evidenced only for bismuth selenides deposited within the alumina pores approximated to 100 nm. With decrease in Ø pore and diameter of bismuth selenide nws the blue shift of absorption edge is obvious although in the case of very tiny pores, when Ø pore approaches the Bohr radius, the determination of effective band gaps for deposited Bi x Se y nw arrays was found to be somewhat problematic using only the experimental transmission spectra perhaps due to at least dual composition of nws.