NANO LETTERS
Tailoring the Optical Properties of Silicon Nanowire Arrays through Strain
2002 Vol. 2, No. 8 811-816
Daniel M. Lyons, Kevin M. Ryan, Michael A. Morris, and Justin D. Holmes* Materials Section & Supercritical Fluid Centre, Department of Chemistry, UniVersity College Cork, Cork, Ireland Received May 14, 2002; Revised Manuscript Received July 2, 2002
ABSTRACT A unique supercritical fluid inclusion-phase technique has been developed to embed silicon nanowires, with size monodispersed diameters, within the pores of mesoporous silica hosts. These nanocomposite materials displayed quite intense room temperature ultraviolet and visible photoluminescence (PL), and the emission wavelength maximum was found to be dependent on the diameter of the encased nanowires. This previously unobserved wavelength dependence of the ultraviolet PL with decreasing nanowire size is explained using a continuum strain model resulting from confinement of the wires within the host lattice.
Currently there is intense interest in nanoscale structures with low dimensionality such as quantum wells (2-D), nanowires (1-D), and quantum dots (0-D).1 Such materials exhibit optical and electronic properties that offer a potential means of investigating quantum phenomena as a function of nanostructure size. The ability to manipulate quantum effects is considered to be critical for technologies that will form the basis for the next generation of computing, optical, and electronic devices.2 Nanowires have particular relevance as “building blocks” for nanoscale structures in the microelectronics industry, as they can function as both devices and electrical contacts. The reduction in semiconductor feature size to allow increased speed and computing power may potentially be achieved by using silicon3 nanowires to give greater device densities. A number of techniques for preparing pseudo 1-D wires have been reported, including laser ablation of silicon and germanium targets,4 liquid crystal templating methods, and vapor-liquid-solid growth mechanisms.5 However, these routes typically yield disordered entanglements of nanowires that are difficult to manipulate into useful architectures. Considerable effort has been expended by Lieber et al.6 to manipulate these nanowire meshes into orientations that can be exploited as nanoelectronic devices. One promising technique for the integration of nanowires into well-defined architectures is the deposition of nanowires within the pores of mesoporous silica. Methods such as electrochemical deposition and chemical vapor deposition7 have previously been used to fill porous matrices with precursor material to form nanowires. Our work has focused on using supercritical fluid (SCF) methods to * Corresponding author. E-mail:
[email protected]. Tel: +353 (0)21 4903608. Fax: +353 (0)21 4274097. 10.1021/nl0256098 CCC: $22.00 Published on Web 07/23/2002
© 2002 American Chemical Society
reproducibly form a range of tailor-made high quality silicananowire composites.8,9 The discovery of unusual quantum-induced electronic properties, including photoluminescence (PL), from nanocrystalline silicon (and that is not seen for larger silicon structures) has aroused huge scientific interest.10 The origin of the PL in these silicon systems is hotly debated, although much of the present controversy arises due to difficulty in identifying the role that both quantum confinement11 and surface states12 have on the band structure in these materials. Lack of control over crystallite diameters and the inability to create monodispersed nanocrystal samples often leads to broadening of PL emission peaks, making investigations into the origin of specific PL transitions in nanocrystalline silicon difficult.13 In this work hexagonal mesoporous silica, synthesized by a hydrothermal nonionic surfactant templating route, was used as a template to form 2-D arrays of embedded crystalline silicon nanowires of different diameters and with narrow size distributions ((10%). In a typical synthesis, 0.5 g of a poly(ethylene oxide) (PEO)-polypropylene oxide (PPO)-poly(ethylene oxide) (PEO) triblock copolymer surfactant (e.g., Pluronic P85, PEO26PPO39PEO26, BASF) was dissolved in 1.8 g tetramethoxysilane to which was added 1 g hydrochloric acid (0.5 M). Methanol generated as a reaction byproduct during silica framework condensation, and crosslinking was removed using a rotary evaporator. The resulting gel was allowed to condense for 1 week at 40 °C. Calcination of the organic-inorganic composite in air at 450 °C was carried out for 24 h. Any residual surfactant in the as-calcined product was removed by placing the products in a flowing ozone stream for 30 min. Mesoporous silicas were also prepared using P65 (PEO20PPO30PEO20) and P123 (PEO20-
Figure 2. High angle PXRD data collected for Si50 displaying crystalline silicon reflections after inclusion of silicon within the pores. Figure 1. Low angle PXRD data collected for a series of hexagonal mesoporous silicas prepared using P65, P85, and P123 triblock copolymer surfactants. Inset shows a transmission electron micrograph of HMS-P85. Scale bar represents 60 nm.
PPO69PEO20) surfactants, yielding materials with different pore diameters.14 The resulting calcined and ozone-treated white products, with pore diameters of 45, 50, and 73 Å ((10%), will be referred to hereafter as hexagonal mesoporous silica (HMS-P65, HMS-P85, HMS-P123 respectively). Silicon nanowires were grown within the silica mesopores using a supercritical fluid inclusion-phase technique. This technique involves the degradation of the silicon nanowire precursor at elevated temperatures and pressures as reported in previous publications.9 In a typical preparation, 0.5 g HMS-P85 was degassed at 200 °C in nitrogen for 6 h and then placed in a high-pressure titanium reaction cell under an oxygen-free nitrogen atmosphere. Diphenylsilane (0.022 mol, 4.05 mL) was placed in the reaction vessel. The reaction cell was connected, via a three-way valve, to a piston and cylinder assembly that was filled with anhydrous hexane. The cylinder was connected to a high-pressure ISCO pump. Carbon dioxide was used to drive the piston and force pressurized anhydrous hexane into the reaction cell. The cell was placed in a furnace and held at 500 °C and at 375 bar for 30 min. The assembly was then allowed to cool at ambient temperature, and the product washed out of the cell with anhydrous hexane. The isolated product was further washed with anhydrous hexane and ethanol, filtered, and dried. The products had a reddish-brown appearance. The silica remained white in the absence of diphenylsilane. The 2-D arrays of silicon nanowires encased in the mesoporous silica template will be referred to as Si45, Si50, and Si73, representing the encased silicon nanowire diameters of 45, 50, and 73 Å ((10%), respectively. Low angle powder X-ray diffraction (PXRD) data for the as-calcined hexagonal mesoporous silica materials is shown in Figure 1. For all of the mesoporous silica templates synthesized, intense features were observed and are indexed in the figure as the 〈100〉, 〈110〉, and 〈200〉 reflections. These features are consistent with highly ordered long-range mesoporous arrays with a repeat distance of 72, 76, and 98 812
Å for HMS-P65, HMS-P85, and HMS-P123, respectively, based on the position of the 〈100〉 reflection, and corresponding to pore center to pore center distances of 83, 88, and 113 Å. Transmission electron micrographs allow direct observation of high quality mesoporous materials for all samples prior to embedding nanowires, as shown in the inset of Figure 1. After filling the mesoporous framework with silicon, extinction of these PXRD reflections was evidenced due to the reduced scattering contrast between the mesoporous silica host and pore included silicon and not due to increased disorder in the hexagonal arrangement of the pores.9 In all cases the hexagonal mesoporous silica displayed a single broad featureless high angle peak, centered at about 22° 2θ, and can be assigned to amorphous silica. After embedding silicon nanowires within the mesoporous hosts using the supercritical fluid phase inclusion method, high angle PXRD diffraction data show the broad amorphous silica peak and peaks that are referenced as the 〈111〉, 〈022〉, and 〈113〉 reflections of crystalline silicon as shown in Figure 2 for Si50 templated from HMS-P85. Surface analyses by nitrogen sorption measurements were carried out for all mesoporous silica and mesoporous silicasilicon nanowire composites. Prior to nanowire inclusion, the hexagonal mesoporous silicas displayed type IV adsorption and desorption isotherms characteristic of mesoporous materials and had surface areas of 521-789 m2 g-1 calculated by the BET method. After silicon inclusion in the pores, the surface area of all samples decreased by between 92 and 98% of the original and pore volume decreased by 96-99%, as summarized in Table 1. Pore size distribution data were calculated by the BJH method and are shown for the mesoporous silicas in Figure 3. The pore diameters at peak maxima were found to be 45, 50, and 73 Å for HMS-P65, HMS-P85, and HMS-P123, respectively. After nanowire inclusion in the silica matrix extinction of these pore size distribution peaks was observed. It is clear from PXRD data that hexagonal mesoporous silicas with long-range well-ordered arrays have been prepared, with pore sizes in the 45-73 Å range as determined by BJH analysis. Inclusion of guest nanowires within the mesoporous host matrix is indicated by the extinction of low Nano Lett., Vol. 2, No. 8, 2002
Table 1: Surface Analysis by Nitrogen Sorption Data for Hexagonal Mesoporous Silicas and for Silica-Nanowire Compositesa mesoporous silica (HMS)
HMS-nanowire composite
SABET (m2 g-1) V (cm3 g-1) SABET (m2 g-1) V (cm3 g-1) P65 P85 P123
617 789 521
0.834 0.949 0.591
49 15 16
0.029 0.011 0.017
a Data was collected at 77.4 K on a Micromeritics Gemini 2375 instrument after degas under flowing nitrogen for 6 hours at 473 K. Surface area (SABET) and pore volume (V) were calculated by BET and BJH methods, respectively.
Figure 4. Room-temperature optical absorbance spectra for nanowires included in a mesoporous host and dispersed in anhydrous hexane. Nanowires with diameters of 45, 50, and 73 Å are referenced in the figure as Si45, Si50, and Si73, respectively.
Figure 3. Pore size distribution data for a series of hexagonal mesoporous silicas prepared from P65, P85, and P123 triblock surfactants.
angle PXRD features and pore size distribution peaks and with a consequent loss of pore volume and surface area. High angle PXRD data indicates the presence of crystalline silicon in all samples. 29 Si nuclear magnetic resonance (NMR) data of the mesoporous silica-silicon nanowire composites also clearly indicate that the encapsulated silicon nanowires are anchored at the pore walls and intimately bonded to the mesoporous silica matrix, consistent with previously reported results.9 Therefore, by nanoengineering of pore dimensions14 we have prepared silicon nanowires with diameters of 45, 50, and 73 Å ((10%), and with lengths of >1 µm. The embedded nanowires displayed the same high monodispersivity exhibited by the host lattice, allowing PL as a function of controlled crystallite diameter to be investigated. Figure 4 shows the UV-vis absorbance spectra for the nanowire-host composites, displaying absorbance in the 4-5 eV range with an absorption edge at about 4.2 eV. This optical absorption edge for these mesoporous silica-silicon nanowire materials is observed to be strongly blue shifted with respect to the direct-gap absorption edge of bulk silicon (3.5 eV), although the energy at which the trailing edge appears corresponds to the bulk silicon direct-gap absorption edge.13,15 Each material prepared showed PL at wavelengths in both the visible red and ultraviolet regions of the electromagnetic Nano Lett., Vol. 2, No. 8, 2002
Figure 5. (A) Visible PL emission spectra for silicon nanowires embedded in a mesoporous silica matrix and dispersed in anhydrous hexane. (B) Ultraviolet PL emission spectra (excitation energy of 4.92 eV) collected for the same materials.
spectrum. In our studies each nanowire-silica composite was interrogated at an excitation energy of 4.92 eV. Figure 5A shows that visible PL was observed for all of the nanowire samples at energies between 1.5 and 2.1 eV. PL emission maxima were observed at 1.79 and 2.01 eV for nanowires 813
with 50 Å diameters (Si50) and emission at 1.98 and 2.05 eV for 45 Å diameter (Si45) nanowires. These two emission peaks are characteristic of indirect-gap nanowires and are the slow band-edge emission band and the higher energy fast emission band, respectively.16 For the silica material containing 73 Å diameter silicon nanowires (Si73), we observed only very weak emission in the same range with a measured quantum yield (ΦPL) value of 0.024.17 This weak emission for Si73 is ascribed to phonon-assisted relaxation and not to quantum confinement.18 As the diameter of the nanowire decreases from 50 to 45 Å there is an increase in intensity with a slight shift toward higher energy. This result is consistent with the expectation that below the Bohr free exciton radius of silicon (50 Å) quantum confinement effects become more pronounced and smaller nanowires would be expected to show not only an increase in PL intensity but also a shift toward higher energy. The measured ΦPL for Si50 and Si45 increased to 0.09 and 0.13 ((0.03), respectively.17 However, these quantum yields are a factor of 2 lower than those previously reported for silicon nanocrystals.19 Figure 5B reveals the high-energy PL emission peaks observed for all of the encapsulated silicon nanowire samples. Intense PL in the ultraviolet region of the electromagnetic spectrum was observed for all materials with a blue shift in the wavelength of the peaks as a function of decreasing nanowire diameter. Luminescence emission maxima were observed at 2.93, 3.3, and 3.49 eV for Si73, Si50, and Si45, respectively. Hence, a linear decrease in emission wavelength with decreasing nanowire diameter was found. Smaller PL emissions were observed at higher and lower energies than the PL emission peak maxima and which mirror discrete features that appear in the absorbance spectra shown in Figure 4. Photoluminescence excitation spectra displayed two discrete peaks at 4.25 and 4.8 eV for all of the mesoporous silica-silicon nanowire samples. The origin of visible PL (1.5-2.1 eV) in silicon nanowires is uncertain, although proposals have been put forward ascribing it to surface states.20 At the surfaces of nanocrystals reconstruction occurs and any such changes in atomic positions will give rise to localized energy states in the band gap of the bulk material. Surface states thus provide a ready trap for electrons, and so any enhanced electronic properties in the nanocrystalline bulk will be degraded due to these states. We have sought to investigate the visible luminescence as a function of the surface properties by terminating the nanowire surface with another material that bonds directly to the nanowire and prevents the surface from reconstructing.21 The abrupt difference in electronic potential between the silicon nanowire and bulk silica matrix surrounding the nanowire essentially confines electrons within the nanowire to energy levels that do not lie in the forbidden band gap region.15 This would be expected to enhance size-related quantum effects for nanowires with dimensions smaller than the free-exciton Bohr radius of silicon (3 nm diameter) investigated. PL in the ultraviolet region of the electromagnetic spectrum has been previously reported for both oxidized silicon nanowires22 and completely amorphous silica nanowires.23 The photophysics properties of these systems have been ascribed to an electron-hole recombination model, where radiative recombination occurs at bulk silica defect centers arising from oxygen deficient sites such as 2-fold coordinated silicon lone pair centers and neutral oxygen vacancies. In the case of silicon-silica systems excitation of the electron-hole pairs is thought to occur in the nanocrystalline silicon and subsequent tunneling of these pairs into the silica bulk allows radiative recombination at defect sites.24 It was found by Qin et al. that by varying the diameter of a crystalline silicon core in a silica sheath the intensity of the PL could be enhanced.24 However, compared to the results presented here, this increase in intensity was not accompanied by any shift in the PL emission energy. It is possible that this lack of control over the tuneability of PL is related to the inability of other synthetic techniques to produce relatively monodispersed samples. Importantly, in the work reported herein, we have created size-monodispersed nanocrystalline silicon with uniform geometry and have determined an ultraviolet PL energy dependence on the nanowire diameter as shown in Figure 5. The tunability of the ultraviolet PL emission energy therefore suggests changes in the band structure for these silicon nanowire arrays as the critical dimensions get increasingly smaller. Unlike other silicon-silica systems, the ultraviolet PL cannot be simply due to radiative recombination in silica. The silicon nanowires reported here do not have an infinitely small diameter, and have a diameter considerably greater than that of polysilane chains. Therefore, the band structure of silicon nanowires would be expected to deviate substantially from the idealized model of 1-D structures such as polysilane chains.25 In such cases mechanisms related to the coupling of crystal vibrations and mixing of wave vectors (k-states) have been suggested as being responsible for relaxing the k selection rule and modifying the band structure. Although phonon-assisted electronic transitions in nanostructures such as nanowires may become possible by enhancing the oscillator strength, such a process is not sufficiently strong to give the intense high-energy transitions in the ultraviolet region seen in this work.26 Therefore, if phonon-assisted transitions are only minor contributors to the PL, other radiative mechanisms must be responsible for the ultraviolet luminescence seen here. In addition, Fourier transform infrared spectroscopy did not evidence peaks due to Si-H stretch modes characteristically found at about 2300-2100 cm-1 and suggests that PL does not arise due to hydride presence.27 It has been previously reported by Nishida28 that relatively flat silicon surfaces favor the formation of dimers for minimizing surface energy contributions. The interaction of neighboring pairs of silicon atoms causes a buckling of the surface layers, resulting in increased strain fields and Nano Lett., Vol. 2, No. 8, 2002
anisotropic stress in the structure.29-31 Theoretical calculations of the electronic effects of such dimer formation in nanocrystals have indicated the appearance of quantized surface state energy levels either well below the valence band edge or higher than the conduction band edge.28 The calculations clearly indicate that band gaps much greater (2.5-3.5 eV) than the band gap of bulk silicon (1.1 eV) are achievable in silicon nanowires with diameters larger than expected from theoretical calculations.32 Such effects have been predicted to become significant even for silicon nanocrystals of 100 Å diameter.33 The generation of such surface states was found to be closely related to strain within the nanocrystals and, in particular, strain induced in the backbonds from surface to internal silicon atoms. These variations in surface state derived band gap energies subsequently result in higher energy electronic transitions in strained nanocrystals than for bulk silicon. Other authors such as Liu et al.34 have reported that stresses introduced to nanocrystalline silicon cores due to oxide shell growth may also alter the band structure, and hence the electronic properties, of nanocrystalline silicon. It is apparent that if strain could be introduced in a controlled way into nanoscale systems, the optical properties of these materials could be manipulated and a clearer understanding of the PL role of surface states in these materials may be provided. There is evidence for both strain-related effects and manipulation of this strain to achieve tunable ultraviolet PL in the work reported here. Crystalline silicon nanowires of varying diameter, with relatively narrow size distributions, were anchored within a mesoporous matrix. As the degree of curvature of the nanowire is inversely related to radius, decreasing the nanowire diameter resulted in increased surface curvature and hence strain.35 The nanowires have a uniform cylindrical shape and an aspect ratio of 100:1 or greater, hence the curvature may be considered insignificant longitudinally along the nanowire and ignored for the purposes of strain-induced quantum confinement effects. Therefore, the diameter proves to be the critical dimension. The silicon nanowire-silica samples exhibited ultraviolet PL transitions with a linear dependence on nanowire diameter, and hence curvature and strain (Figure 5). In this study, direct control over nanocrystal size has allowed sharper PL transitions to be observed, compared to the broad featureless PL emission peaks previously reported for silicon and silica nanowires.22,23 The transitions clearly arise from discrete energy levels that have energy separation greater than the theoretical valence-conduction band gap calculated for bulk silicon. Previously calculated energies in silicon nanocrystals28 for transitions between surface strain derived energy levels were 2.5-3.5 eV, compared to peak maxima PL emission energies of 2.93-3.49 eV observed for the silica-nanowire composites reported here. Fine structure is resolved in the ultraviolet PL spectra for the smallest diameter and most highly strained nanowires. This fine structure is consistent with the expectation that energy levels become more separated and resolved as the silicon nanowire diameter decreases. These sidebands may be ascribed to characteristic phonon-assisted absorption and luminescence processes in Nano Lett., Vol. 2, No. 8, 2002
nanocrystalline silicon13 and provide the most compelling evidence that the PL data presented here arise from within the nanowire and not at defect centers. Previous kinetics experiments suggest that such phonon-related events are confined within the nanocrystalline core16 and that these bands may be related to coupling between silicon nanowire core and surface states.36 More work will be required probe the exact origins and to assign these bands definitively. Furthermore, inhomogeneous strain fields at the siliconsilica boundary will allow relaxation of the forbidden transition rule,37 in agreement with the upper limit lifetime decay measurements of ∼20 ns obtained in this work.17 Quantum yields for the composites prepared here were typically of the order of ∼0.14 for the ultraviolet PL, and this enhanced ΦPL may be attributed to a reduction in radiationless processes as opposed to an increased radiative rate.13 Continuum strain offers a more general explanation for the origin of high-energy light emission from siliconsilica composites that is not based on simple band-edge or silica defect radiative recombination but on transitions between strain-induced surface states. As strain is an intrinsic property of all nanostructured materials this explanation will have general applicability in any future discussions of ultraviolet PL from nanocrystalline silicon. In conclusion, nanowires have been engineered to yield atomically abrupt surfaces intimately bonded to a mesoporous silica host that inhibits surface reconstruction. Quantum effects due to the creation of strain-induced surface states offer an efficient pathway for radiative recombination, and this work indicates for the first time evidence for a theoretical “continuum strain” model, describing a phenomenon that gives intense ultraviolet light that is tunable over a wide energy range for these composite nanoscale materials. It should be clearly stated, however, that while photoluminescence from strain-induced surface states is the most reasonable explanation for the data reported here, further work is required before a definitive assignment can be made. The work reported herein demonstrates the ability to manipulate quantum confinement properties in 1-D semiconductor nanowires and offers experimental insight into a mechanism of light emission from nanocrystalline silicon that has not been widely investigated to date. The fabrication of such nanoscale composites offers new and exciting possibilities for a range of future optoelectronic applications and emerging technologies, from light emitting diodes to full color intelligent image displays. Acknowledgment. The authors acknowledge financial support from the European Union under the Future and Emerging Technologies Program (Project No. IST-200025469) and Intel (Ireland) Ltd. The authors thank Andy Ruth for measuring PL lifetimes and David Nikogosyan for the use of his laser equipment. We also acknowledge Eoin O’Reilly and Kirk Ziegler for useful discussions. References (1) Handbook of Nanophase Materials; Goldstein, A. N., Ed.; Dekker: New York, 1997 (see also references therein). (2) Green, M. A.; Zhao, J.; Wang A.; Reece, P. J.; Gal, M. Nature 2001, 412, 805. 815
(3) Holmes, J. D.; Johnston, K. P.; Doty, R. C.; Korgel, B. A. Science 2000, 287, 1471. (4) Tang, Y. H.; Zhang, Y. F.; Peng, H. Y.; Wang, N.; Lee, C. S.; Lee, S. T. Chem. Phys. Lett. 1999, 314, 16. (5) Wu, Y.; Yang, P. J. Am. Chem. Soc. 2001, 123, 3165. (6) Huang Y.; Duan, X.; Cui, Y.; Lauhon, L. J.; Kim, K.-H.; Lieber, C. M. Science 2001, 294, 1313. (7) Leon, R.; Margolese, D.; Stucky, G.; Petroff, P. M. Phys. ReV. B 1995, 52, 2285. (8) Coleman, N. R. B.; Ryan, K. M.; Spalding, T. R.; Holmes, J. D.; Morris, M. A. Chem. Phys. Lett. 2001, 343, 1. (9) Coleman, N. R. B.; Morris, M. A.; Spalding, T. R.; Holmes, J. D. J. Am. Chem. Soc. 2001, 123, 187. (10) Canham, L. T. Appl. Phys. Lett. 1990, 57, 1046. (11) Glinka, Y. D.; Lin, S. H.; Huang, L. P.; Chen, Y. T.; Tolk, N. H. Phys. ReV. B 2001, 64, 5421. (12) Ledoux, G.; Guillois, O.; Porterat, D.; Reynaud, C.; Huisken, F.; Kohn, B.; Paillard, V. Phys. ReV. B 2000, 62, 15942. (13) Brus, L. E.; Szajowski, P. F.; Wilson, W. L.; Harris, T. D.; Schuppler, S.; Citrin, P. H. J. Am. Chem. Soc. 1995, 117, 2915. (14) Ryan, K. M.; Coleman, N. R. B.; Lyons, D. M.; Hanrahan, J. P.; Spalding, T. R.; Morris, M. A.; Steytler, D. C.; Heenan, R. K.; Holmes, J. D. Langmuir 2002, 18, 4996. (15) Littau, K. A.; Szajowski, P. F.; Muller, A. J.; Kortan, A. R.; Brus, L. E. J. Phys. Chem. 1993, 97, 1224. (16) Calcott, P. D. J.; Nash, K. J.; Canham, L. T.; Kane, M. J.; Brumhead, D. J. Phys. Condens. Mater. 1993, 5, L91. (17) Colloidal suspensions of silica-silicon samples in anhydrous hexane were excited using the fourth harmonic wave of a Nd:glass laser (Twinkle, Light Conversion Ltd., Vilnius, Lithuania) at 264 nm with a pulse duration of ∼200 fs and energy of ∼5-30 µJ per pulse. PL decay lifetimes were measured at detection wavelengths of 400 ( 18 nm and 640 ( 60 nm and a photomultiplier tube operating at 800 V. PL quantum yields were measured on a Perkin-Elmer LS50B spectrofluorimeter with respect to β-naphthol and rhodamine B that have known quantum yields of 0.15 and 0.7, respectively. (18) Sanders, G. D.; Stanton, C. J.; Chang, Y. C. Phys. ReV. B 1993, 48, 11067. (19) Holmes, J. D.; Ziegler, K. J.; Doty, R. C.; Pell, L. E.; Johnston, K. P.; Korgel, B. A. J. Am. Chem. Soc. 2001, 123, 3743. 20. (20) Koch, F.; Petrova-Koch, V.; Muschik, T.; Nikolov, A.; Gavrilenko, V. Mater. Res. Soc. Symp. Proc. 1993, 283, 197. (21) PXRD and surface analysis confirms the ability to control the diameter of the pores and hence the diameter of the nanowire guests. Since these wires are embedded within the pores of a silica matrix (and all evidence suggests that intimate contact exists between the wall atoms and the silicon atoms), these nanocomposite materials should minimize contributions to the PL originating from surface reconstruc-
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(36) (37)
tion. This does not suggest that reconstruction does not occur but will be determined by the precise difference in the band structure at the interfacial sites. The abrupt difference in electronic potential between the guest nanowire and host matrix should act to confine electrons within the nanowire to energy levels that do not lie in the forbidden band-gap between valance and conduction bands. All three silicon-silica composite materials were synthesized under identical conditions in the same apparatus. This was used to prevent contamination of the product materials due to introduction of impurities. In addition, it should be pointed out that if all materials were prepared in the same way and an impurity was present it would not lead to the changes in the PL emission wavelengths observed in this work. Qin, G. G.; Lin, J.; Duan, J. Q.; Yao, G. Q. Appl. Phys. Lett. 1996, 69, 1689. Yu, D. P.; Hang, H. L.; Ding, Y.; Zhang, H. Z.; Bai, Z. G.; Wang, J. J.; Zou, Y. H.; Qian, W.; Xiong, G. C.; Feng, S. Q. Appl. Phys. Lett. 1998, 73, 3076. Lin, J.; Yao, G. Q.; Duan, J. Q.; Qin, G. G. Solid State Commun. 1993, 86, 559. Brus, L. J. Phys. Chem. 1994, 98, 3575. Sanders, G. D.; Chang, Y. C. Phys. ReV. B 1992, 45, 9202. Coleman, N. R. B.; O’Sullivan, N.; Ryan, K. M.; Crowley, T. A.; Morris, M. A.; Spalding, T. R.; Steytler, D. C.; Holmes, J. D. J. Am. Chem. Soc. 2001, 123, 7010. Nishida, M. Solid State Commun. 2000, 116, 655. Osanai, M.; Yasunaga, H.; Natori, A. Surf. Sci. 2001, 493, 319. Wu, F.; Lagally, M. G. Phys. ReV. Lett. 1995, 75, 2534. Chang, C. S.; Tsong, T. T. Prog. Surf. Sci. 1997, 54, 387. Iyer, S. S.; Xie, Y.-H. Science 1993, 260, 40. Thean, A.; Leburton, J. P. Appl. Phys. Lett. 2001, 79, 1030. Liu, H. I.; Biegelsen, D. K.; Ponce, F. A.; Johnson, N. M.; Pease, R. F. W. Appl. Phys. Lett. 1994, 64, 1383. Curved instead of flat surfaces, and the enhanced strain in both silicon dimers at such surfaces and in back-bonds from surface atoms, would be expected to modulate band gap energies. The term “curved” does not imply that the surface consists of a series of “bent” surface planes. Clearly even at these small dimensions these crystalline particles will adapt to the pore confinement by terminating in a series of different surface planes. However, the atoms will be highly strained due to surface tension and confinement related factors due to their unusual structure. English, D. S.; Pell, L. E.; Yu, Z.; Barbara, P. F.; Korgel, B. A. Nano Lett. 2002, 2, 681. Zacharias, M.; Fauchet, P. M. J. Non-Cryst. Solids 1998, 227, 1058.
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