Novel Method for Site-Controlled Surface Nanodot Fabrication by Ion

Feb 26, 2005 - ... below the surface, while at higher doses Ga nanodots form on the surface as metallic Ga droplets. Possible applications include def...
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NANO LETTERS

Novel Method for Site-Controlled Surface Nanodot Fabrication by Ion Beam Synthesis

2005 Vol. 5, No. 4 771-776

Ryan Buckmaster,*,† Takashi Hanada,† Yoshiyuki Kawazoe,† Meoung-whan Cho,†,‡ Takafumi Yao,†,‡ Nobuaki Urushihara,§ and Akira Yamamoto§ Institute for Materials Research, Tohoku UniVersity, Katahira 2-1-1, Aoba-ku, Sendai 980-8577, Japan, Center for Interdisciplinary Research, Tohoku UniVersity, Aramaki, Aoba-ku, Sendai 980-8578, Japan, and ULVAC-PHI, Analytical Laboratory, 370 Enzo, Chigasaki, Kanagawa, 253-0084, Japan Received November 25, 2004; Revised Manuscript Received February 18, 2005

ABSTRACT By using a Ga FIB system to spatially control the implantation of Ga into SiO2 followed by vacuum annealing, we have fabricated selfassembled surface Ga nanodots with a high degree of control of nucleation location. The morphology of the Ga nanodots is closely related to Ga dose, showing a critical dose needed for nucleation that results in Ga nanodot formation just below the surface, while at higher doses Ga nanodots form on the surface as metallic Ga droplets. Possible applications include defining nucleation sites for subsequent growth, use as Ga source for GaN or GaAs quantum dots, or as catalyst for nanowire growth.

Ion implantation technology is a mainstay of the electronics industry for doping semiconductors, but has also been used in ion-beam synthesis (IBS) to implant ions that can then be used as a growth source for nanostructures. Examples of IBS fabricated nanostructures include self-assembled quantum dots formed in a bulk matrix1-4 and surface precipitates5 that form after annealing ion-implanted materials. Focused ion beam systems (FIB), typically Ga+ ion sources for engineering reasons, allow the movement of the ion beam to be controlled by a pattern generator. Thus, FIB systems are most commonly used for maskless lithography and micromachining, where the FIB is used to sputter away surface layers, such as for transmission electron microscopy (TEM) sample preparation.6 Also, FIB patterning has been demonstrated as a method to control the nucleation sites of quantum dots7-9 and other nanostructures.10 In addition, an FIB system implants ions at the same time as sputtering the surface, and thus can be used to locally implant micro- or nanometer scale areas. By combining IBS with the spatial control of a Ga+ FIB we have locally implanted areas of SiO2 with high doses of Ga followed by annealing to fabricate self-assembled Ga nanodots at arbitrary locations. The Ga nanodots have potential applications, including defining nucleation sites for the growth of other materials, as a Ga source for Ga * Corresponding author. E-mail: [email protected]. † Institute for Materials Research, Tohoku University. ‡ Center for Interdisciplinary Research, Tohoku University. § ULVAC-PHI, Analytical Laboratory. 10.1021/nl048044j CCC: $30.25 Published on Web 02/26/2005

© 2005 American Chemical Society

compound semiconductor quantum dot fabrication, and as a catalyst for nanowire growth. In addition, the IBS nanofabrication techniques shown here, which we will refer to as implant source growth (ISG), can easily extended to other material systems besides Ga/SiO2. For all experiments, 60 nm of oxide was thermally grown on Si(111) wafers. Patterning and Ga implantation was carried out using a liquid Ga source Hitachi 2100 focused ion beam system at an acceleration voltage of 40 keV with ion doses ranging from 1 × 1015 to 2 × 1017atoms/cm2. Typical patterns had depths from 0.5 nm to 15 nm, and widths and lengths from 150 nm to 2 µm. After removal from the FIB system, the samples were ultrasonically cleaned in acetone and ethanol before being loading into an ultrahigh vacuum (UHV) chamber with a base pressure of 1 × 10-7 Torr and annealed typically at 600 °C for 1 h. After removal from the vacuum chamber, the samples were observed using an atomic force microscope (AFM) and a field effect scanning electron microscope (SEM). Elemental analysis of the nanodots was performed using a Phi-700 scanning auger nanoprobe (AES) capable of 6 nm spatial resolution. We initially FIB patterned a series of 150-300 nm wide trenches from 2.5 nm to 11 nm deep and confirmed with AFM that the trenches had sharp geometric shapes. However, after vacuum annealing we observed that self-assembled nanodots 50-150 nm in diameter had formed in the 5 nm and 7.7 nm deep trenches. No nanodots were observed for

Figure 1. SEM image of 200 nm by 2 µm trenches patterned by systematically increasing the Ga doses for each trench show a critical Ga dose for nucleation and also a linear decrease in the average number of nanodots (average of three identical sets of patterns) per trench with increasing dose.

annealing temperatures 500 °C and below. Suspecting that the nanodots were related to the implanted Ga, we next fabricated and annealed the pattern in Figure 1, consisting of 200 nm by 2 µm patterns with depths and Ga ion doses systematically increasing from 0.4 nm at 4.0 × 1015/cm2 to 13.6 nm at 5.3 × 1016/cm2, respectfully. Three features are most apparent as the Ga ion dose increases, one is the critical Ga ion dose required for nanodot formation. Second, after the critical Ga ion dose is reached, the number of selfassembled nanodots per trench (from an average of three identical sets of patterns) linearly decreases with increasing Ga ion dose. And third, the characteristics of the nanodots also change with increasing dose as the diameter of the nanodots increases from 40 nm wide and 6-7 nm high to over 100 nm wide and 80-100 nm high. Furthermore, the nanodots change from gray with fuzzy contrast in SEM images to white with sharp contrast as the ion dose increases. To verify the composition of the nanodots we used AES, which is extremely surface sensitive, to both perform point probes and also map the relative Ga, Si, and O concentrations on the surface of a smaller “gray” nanodot formed in a 5 nm deep trench by a Ga dose of 2.0 × 1015/cm2 and a larger “white” nanodot formed in a 7.7 nm deep trench by a Ga dose of 3.2 × 1016/cm2, Figure 2. Point probes of the “gray” nanodot indicated that it was mostly composed of Ga, although surprisingly a point probe at the bottom of the trench next to the nanodot caused a smaller nanodot to form. After the point probes, both the original and smaller nanodot formed by the point probe were bright in SEM images, with a sharp contrast similar to the larger “white” nanodots formed at higher Ga doses. Subsequent AES Ga elemental mapping showed high Ga concentrations and low O and Si concentrations at the locations of the point probed nanodots. However, mapping of a Ga nanodot that had not been first exposed to a point probe, Figure 2a, showed no Ga signal, Figure 2b, although the surface of the dot was slightly O poor compared to the surrounding SiO2. In contrast, a larger “white” Ga nanodot that formed in a 7.7 nm deep trench, Figure 2c, did 772

Figure 2. SEM, left, and corresponding AES Ga concentration map, right, of self-assembled Ga nanodots. A Ga nanodot formed in a 5 nm deep trench (a, top) shows no significant Ga concentration on the surface (b), while the Ga nanodot formed in a 7.7 nm deep trench (c, bottom) shows a strong surface Ga signal, (d). Thus the Ga nanodot in (a) actually formed just below the surface, while the nanodot in (c) is basically a droplet of metallic Ga on the surface.

show a high Ga concentration, Figure 2d, during mapping without being first exposed to a point probe. From these results we can conclude that the “gray” nanodots actually nucleate slightly below the surface and are covered by a thin layer of SiO2 which is removed by the auger point probe and exposes the Ga below. In contrast, the larger “white” nanodots are basically Ga droplets on the surface. The structure of the SiO2-capped nanodot is particularly interesting as it self-assembles with its own thin layer of insulation. The formation mechanism of the Ga nanodots is directly related to the diffusion and solubility of Ga in SiO2. From TRIM calculations,11 we estimate the peak and straggle (lateral standard deviation) of the implanted Ga distribution to be 32 nm and 9.2 nm, respectively, so we conclude that most of Ga was implanted within the SiO2 film. Furthermore, for doses sputtering less than 15 nm of SiO2, the Ga concentration immediately at the surface should be low before annealing. Also, we anticipate that the implanted Ga is chemically stable and remains mobile in SiO2, as Ga will not reduce Si to form Ga2O3,12 but a small amount of Ga will almost certainly be incorporated into the SiO2 matrix as the formation of Ga oxides is energetically favorable and some free oxygen atoms are created as a result of implant damage.13 Due to interest in using Ga as a dopant in electronic device technology and related issues, Ga diffusion in SiO2 has been studied using vapor14,15 and ion-implanted Ga sources.13,16,17 The most applicable to our case and the only report on Ga diffusion at supersaturation is by van Ommen,16 who found that there are several different diffusion coefficients and Nano Lett., Vol. 5, No. 4, 2005

mechanisms of Ga diffusion depending on the implantation and annealing conditions. The most important result of van Ommen’s is that for high Ga doses around 1 × 1016/cm2, the Ga solubility limit is exceeded and the Ga diffuses to the surface and segregates when annealing in a N2 ambient. As N2 is generally inert, this result is applicable and in excellent agreement with our own findings of the critical Ga dosage required for nanodot formation. Thus diffusion and segregation of supersaturated Ga is the main phenomenon responsible for Ga nanodot formation, although several other factors also influence the morphology of the nanodots. A study of implantation profiles of Mn, Ni, Cr, and Xe in Al2O3 single crystals5 by Ohkubo et al. provides many interesting insights into the diffusion of supersaturated implants. One of the most critical factors that determined whether the implants significantly diffused to the surface or not was if the top layer of the Al2O3 substrate was rendered amorphous. The recrystallization of the amorphous layer and supersaturation provide the driving force for diffusion to the surface. While our own SiO2 films were already amorphous before implantation, the recovery of local order in the SiO2 from implant damage may play a role in the surface diffusion of Ga. Gas-phase studies of Ga diffusion in quartz and much more porous bulk-fused silica indicate that microchannels enhance the diffusion of Ga as well.18 Damage from the Ga implantation process almost certainly results in the SiO2 structure becoming more porous and open, possibly offering low-resistance, fast-diffusion pathways which may affect the rate and directional preference of Ga diffusion. Also of interest is that Ohkubo et al. observed MnO micrometersized surface precipitates crystalgraphically aligned with the Al2O3 substrate after annealing Mn implanted Al2O3 in oxygen. Thus lattice matched quantum dots and nanostructure growth should be possible with ISG. The new nanodot that formed in the unoccupied portion of the trench near an existing nanodot after exposure to the AES electron probe beam can be explained by the decrease of Ga solubility in SiO2 at lower temperatures. First the GaSiO2 system comes into equilibrium at 600 °C by Ga diffusion to the surface and segregating to form Ga nanodots. After annealing, the temperature drops and the Ga again becomes supersaturated, but is immobile at low temperatures and unable to further segregate, leaving a supersaturated solid solution. The AES electron beam then causes local heating and provides kinetic energy, allowing the room-temperature supersaturated Ga in the SiO2 to become mobile and come out of solution and form a new nanodot. This phenomenon could potentially be used as a fabrication method by making room temperature supersaturated thin films and using e-beam lithography to selectively mobilize the solute atoms at a particular location. The reason for the disappearance of the Ga dots at doses above 4.1 × 1016/cm2 is not as clear. Based on the linear trend of the decreasing number of nanodots, perhaps the diffusion distances simply became long enough to allow all of the Ga to diffuse away, which would suggest that the surface energy of the SiO2 trench significantly changes with sputtering. Another possibility is that at sputter depths beyond Nano Lett., Vol. 5, No. 4, 2005

11 nm a much greater concentration of Ga is present on the surface that is oxidized after being removed from the FIB system and forms a Ga-silicate-glass layer19,16 that blocks Ga diffusion to the surface during annealing. It is also possible that changes in surface energy caused by sputtering increase the SiO2 surface energy enough to make the Ga wet the substrate and form a thin film. This issue will be further explored in a future work. Perhaps the most interesting feature of the Ga nanodots is the transition from nucleating just under the surface to nucleating on the surface as the Ga ion dose increases. Typically, heterogeneous nucleation at an interface is favored over homogeneous nucleation, so the formation of nanodots under the surface is somewhat unexpected. Two possible general sets of factors for the transition from capped to surface nanodots are kinetic factors and stability factors. One kinetic factor is that, for the low dose case, 5 nm less SiO2 is sputtered away than in the case for surface dots, increasing the time for the Ga to diffuse to the surface. This case can be easily dismissed, as even in the case where samples are ramped up to 600 °C and immediately cooled, the results are similar to Figure 1, although the nanodot sizes tend to be more varied and less ordered. Essentially there is enough time for the Ga to diffuse to the surface if that is where it is most stable. Another possible kinetic factor is that at low doses the topmost SiO2 layer is very dense and prevents diffusion to the surface, forcing the Ga to segregate inside of the SiO2. While this effect is seen for deep implants into crystalline substrates where the topmost layers are not rendered amorphous,5 our implants are relatively shallow and our substrate is already amorphous. In addition, as the SiO2 capping layer was easily removed by the concentrated AES electron beam, it is probably extremely thin. While we cannot rule out the effects of ion damage to the SiO2 on diffusion and nucleation, we do not believe this effect is responsible for the nanodot nucleation just below the surface. In the case where stability factors dominate, the configuration with the lowest free energy will be favored. For a new phase nucleating, the total change of free energy, ∆Gtotal, is typically expressed as sum of the reduction of free energy per volume, ∆Gv and the increase of surface energy, γ, due to the creation of a new surface. The greater the decrease of free energy of a process, the more favorable it is. While we will not try to quantitatively evaluate the free energies for the Ga nanodots, free energy and classical nucleation theory provide a useful framework to examine some of the factors that could shift the balance between surface or buried nanodots. The change of free energy of a Ga nanodot buried relatively deeply in the SiO2, i.e., no thin capping layer, would be simply ∆Gburied ) 4/3πr3∆Gvb + 4πr2γGa-SiO2 where r is the radius of a nanodot. For a surface Ga nanodot, ∆Gsurface ) Vd∆Gv + SdγGa + SbγGa-SiO2 - SbγSiO2 where Vd ) 1/3πr3(cos θ + 2)(cos θ - 1)2 is the nanodot 773

volume, Sd ) 2πr2(1 - cos θ) is the portion of the surface not in contact with the substrate, Sb ) πr2sin2θ is the area of the nanodot in contact with the SiO2, and θ is the contact angle between the nanodot and substrate. Finally, for a capped nanodot with a SiO2 capping layer, ∆Gcapped ) 4/3πr3∆Gvc + 4πr2γGa-SiO2 + ∆SSiO2‚γSiO2 where ∆SSiO2 is the increase of area of the SiO2 surface due to the capping layer. Assuming the respective surface energies are the same for each case, that ∆Gvb is equal to ∆Gv, and that θ > 45°, then for the same nanodot volumes ∆Gsurface will always be greater than ∆Gburied proportional to r2. Thus we should not expect to observe any surface nanodots. However, if ∆Gv is less than ∆Gvb, then initially ∆Gburied will be less than ∆Gsurface, but as the radius continues to increase then ∆Gsurface will be less than ∆Gburied and surface nanodots should be favored. In addition, if capped nanodots are observed, then ∆Gvc must be less than ∆Gvb. While we do not know the precise surface energies of GaSiO2 or of the SiO2 surface of the FIB patterned trenches, several experimental values have been tabulated which are useful for comparing the relative magnitude of the surface energies involved. The surface energy of Ga in a vacuum is well-known and is approximately 700 erg/cm2.20 While the surface energy of SiO2 is sensitive to the exact surface structure and termination, a value of 200 erg/cm2 is reasonable,21 and the Ga-SiO2 interfacial energy has been reported to be 590 erg/cm2 for slide glass and 520 erg/cm2 for quartz.22 From these tabulated values we can make two assertions, one is that Ga will tend to ball up on the surface of SiO2 (θ > 45°), which seems to be the case for the surface nanodots and other observations of Ga on SiO2.22,23 Second, unless the ∆Gv values vary for the different types of nanodots, then the buried nanodots will always be the most stable phase, particularly for larger dots. The increasing Ga ion dose has several simultaneous effects on the SiO2 which include sputtering, ion damage, and most importantly, more Ga implanted into the SiO2 film. The changes that higher Ga dose would have on the driving force of nucleation, ∆Gv, are difficult to quantitatively estimate, although generally higher Ga concentrations, i.e., higher degree of supersaturation, should increase the magnitude of ∆Gv. Considering only the effects of supersaturation, a Ga atom going from the SiO2 bulk at a particular Ga concentration to either a surface nanodot or a buried nanodot would experience the same decrease in chemical potential, so the Ga concentration alone is probably not a major factor in determining if a buried or surface nanodot is ultimately more stable, although the Ga concentration will affect the diffusion rate and number of initial nuclei. A higher Ga dose will result in more Ga diffusing, which should lead to more and bigger nanodots. In the case of nucleation in a solid, the ∆Gv term is also affected by strain resulting from misfit of the precipitate and surrounding matrix. In our case there is no strain resulting from mismatch as Ga is a liquid at 600 °C. But, for a buried Ga nanodot to expand, work must be done on the SiO2 to cause plastic deformation. This work 774

required acts as an energy barrier to further expansion of the nanodot. As more Ga atoms join a growing Ga cluster, the SiO2 would resist expansion and exert pressure and compress the nanodot, effectively decreasing the magnitude of ∆Gv and increasing the free energy of larger nanodots buried in the SiO2. Thus, by forming so close to the surface, the capped nanodots can have lower surface energy by being embedded in the SiO2 while avoiding the high cost of deforming the SiO2 by only deforming a thin surface layer of SiO2 to relieve compressive stress; the best of both worlds. In this case, after the critical dose is reached, capped nanodots form as explained above; however, as the Ga dose increases, the dots become larger and larger, stretching the SiO2 layer more and more. A growing nanodot will experience an effective net force pushing it toward the surface. At some point the extra surface area of the increasing large capping layer, stress on the capping layer, and compression of the nanodot make surface nucleation more favorable, or the capping layer may simply break from being stretched too thin. Another possible transition mechanism is that after a capped nanodot becomes too large, a surface nanodot will nucleate on top of it. By either route, once the surface nanodots achieve a certain size, any remaining capped nanodots will be cannibalized by the more stable surface nanodots, probably under Ostwald ripening conditions. However, a better understanding of the formation and transition mechanisms of the nanodots will require more information on the microstructure a well as the nucleation dynamics of the nanodots, perhaps from TEM and in-situ observations. The increase in distance between the nanodots with increasing dose can be understood as a result of faster diffusion resulting in longer diffusion lengths that increase the average spacing between dots during ripening. From AFM measurements, the average heights of the nanodots steadily increase with Ga dose; thus so does the nanodot area above the surface, while the SiO2 capping layer thins. Both of these factors should result in diffusion distances between nanodots to be greatly enhanced as the diffusion mechanism goes from being dominated by bulk diffusion in the SiO2 to being dominated by much faster surface diffusion. By taking advantage of the critical Ga ion dose needed before Ga nanodots are formed, we have been able to control the nucleation sites of individual nanodots. An example of this can be seen in Figure 3, where we made a set of intersecting trenches patterned to a depth of 2.5 nm which corresponds to just half of the critical Ga dose needed for Ga nanodot formation. After annealing, a single capped Ga nanodot approximately 40 nm wide and 5-6 nm high forms at each trench intersection. By increasing the dose, surface Ga nanodots can be formed in the same way. By using this method, the nucleation site can be almost completely arbitrarily controlled, while the size and morphology of the Ga nanodots can be further controlled by the exact dose and geometry of the implanted areas as well as annealing conditions. An interesting aspect of the Ga nanodots is the similarity they have to MOCVD-grown ZnO nanodots on SiO2 Nano Lett., Vol. 5, No. 4, 2005

Figure 3. AFM and corresponding SEM images of an array of self-assembled capped Ga nanodots formed at the intersection of 2.5 nm deep FIB patterned trenches. Only at the intersection is the critical Ga dose needed for nanodot formation exceeded. Thus, the nanodot nucleation sites can be predetermined by choosing the appropriate geometry of the implanted area.

fabricated on very similar patterns made using a 30 keV Ga+ FIB reported by Kim, et al.9,24,25 These authors suggested that Ga atoms diffused along the FIB patterned lines and had a surfactant effect that created nucleation sites for the ZnO nanodots. Based on our findings, Ga is probably the mechanism for the high degree of order in the ZnO nanodots, although not directly as a surfactant. We propose that the mechanism is that the FIB implanted Ga diffuses to the near surface, forming small SiO2-capped Ga nanodots, which then act as preferred nucleation sites for the ZnO dots in a fashion similar to the vertical alignment observed in the formation of many stacked self-assembled quantum dot systems. Such an interaction may account for the similarity of spacings of Kim et al.’s ZnO and our Ga nanodots. From our own experiments, not shown, at the same implantation energy and sputter depth as Kim et al. surface nanodots will form. However, short growth times and the nitrous oxide used in the growth process probably supplied oxygen that immobilized a portion of the Ga, thus reducing the size of the Nano Lett., Vol. 5, No. 4, 2005

Ga nanodots compared to vacuum annealing.16 It is likely that this technique could be adapted for a wide variety of other materials systems as well, although the diffusion of Ga may impose a strict thermal budget on subsequent processing steps in many device applications. There are several potential applications for both ISG in general and the Ga nanodots presented in this paper specifically. Ordered self-assembled quantum dots of magnetic materials fabricated by ISG are one especially interesting possible application. The Ga nanodots can be potentially applied in several different ways. One is to define the nucleation site for the growth of other materials by forming capped dots that would act as preferential sites for the deposition of other materials, as the tops of the capped Ga nanodots should offer higher energy sites than virgin SiO2 or other oxides, in a fashion similar to SiO2 window methods or stacked quantum dots where the window or buried dot is a more favorable nucleation site.21,26 Second, the Ga can be directly used as a source for fabricating Ga-based semiconductor quantum dots such as GaN and GaAs by droplet epitaxy27-29 by simply annealing in the appropriate ambient. Given the ability of ISG to control nucleation sites, this approach may have significant optoelectronic applications. Ga droplets have also been successfully used as a catalyst for the growth of nanowires.30-32 The surface Ga nanodroplets fabricated by ISG should fulfill this purpose quite well as the size and nucleation sites of the nanodots can be controlled. Finally, Ga nanodots may be useful in all-optical devices due to nonlinear properties of liquefying Ga.33 Based on the results of Ohkubo et al.5 and van Ommen16 and on our own results, we propose three guidelines for successful application of ISG to surface nanostructure fabrication. First is that the implanted material must be implanted at levels above the solubility limit of the host material, thus it is advantageous to choose the implant and host so that the solubility is low. Second, the chemical reactivity of the host material and implant is also important. While chemical reactivity, such as the implant reducing the host, could be useful, generally implants that are chemically stable in the host material are desirable. Finally, the top layer of the host material must be rendered amorphous to allow the implanted material to move to the surface and segregate. Other factors such as relative surface energies of the segregated material and host surface, annealing ambient, implant depth, and size of implanted areas will additionally affect the characteristics of surface precipitates. While most FIB systems are Ga based, virtually any atomic species can be implanted with spatial control by using masking combined with low-energy ion implantation. While the maskingimplanter approach is more complicated, it is much better suited for use in mass production. In conclusion, by spatially controlling Ga implantation into SiO2 using a FIB followed by vacuum annealing, we have successfully fabricated Ga nanodots. Furthermore, we have identified a critical nucleation dosage and also found that the Ga nanodots initially nucleate just under the surface, while at higher doses the Ga dots begin to nucleate on the surface. Also, the nucleation site of the Ga nanodots can be 775

arbitrarily controlled by the choice of geometry of the implanted areas. There is no reason not to believe that using ISG to form ordered nanoscale surface precipitates cannot be applied to many other material systems. Finally, other potential applications of the Ga nanodots include defining the nucleation sites for the subsequent growth of other materials, a material source for Ga-based semiconductor quantum dots, and also as a catalyst for nanowire growth. Acknowledgment. We thank the Japanese Ministry of Education, Culture, Sports, Science and Technology for support of this research and the 21st Century COE Program Tohoku University International Center of Research and Education of Materials for access to the FIB system. References (1) Magrader, R. H., III.; Haglung, R. F., Jr; Yang, L.; Wittig, J. E.; Zuhr, R. A. J. Appl. Phys. 1994, 76, 708-715. (2) White, C. W.; Budai, J. D.; Zhu, J. G.; Withrow, S. P.; Zuhr, R. A.; Hembree, D. M., Jr.; Henderson, D. O.; Ueda, A.; Tung, Y. S.; Mu, R.; Magruder, R. H. J. Appl. Phys. 1996, 79, 1876-1880. (3) Ahn, C. G.; Jang, T. S.; Kim, K. H.; Kwon, Y. K.; Kang, B. Jpn. J. Appl. Phys. 2003, 42, 2382-2386. (4) Mazzoldi, P.; Arnold, G. W.; Battaglin, G.; Bertoncello, R.; Gonella, F. Nucl. Instrum. Methods Phys. Res., B 1994, 91, 478-492. (5) Ohkubo, M.; Hioki, T.; Kawamoto, J. J. Appl. Phys. 1986, 60, 13251335. (6) Langford, M.; Petford-Long, A. K.; Rommeswinkle, M.; Egelkamp, S. Mater. Sci. Technol. 2002, 18, 743-748. (7) Kammler, M.; Hull, R.; Reuter, M. C.; Ross, F. M. Appl. Phys. Lett. 2003, 82, 1093-1095. (8) Weller, R.; Ryle, W.; Newton, A.; McMahon, M.; Miller, T.; Magruder, R., III. IEEE Trans. Nanotechnol. 2003, 2, 154-157. (9) Kim, S. W.; Ueda, M.; Kotani, T.; Fujita, S.; Fujita, S. Jpn. J. Appl. Phys. 2003, 42, L568-L571. (10) Sun, Y. T.; Rodriguez Messmer, E.; Lourdudoss, S.; Ahopelto, J.; Rennon, S.; Reithmaier, J. P.; Forchel, A. Appl. Phys. Lett. 2001, 79, 1885-1887. (11) Ziegler, J. F.; Biesack, J. P.; Littmark, U. The Stopping and Range of Ions in Solids; Pergamon Press: New York, 1985.

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Nano Lett., Vol. 5, No. 4, 2005