In Situ TEM Near-Field Optical Probing of

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In Situ TEM Near-Field Optical Probing of Nanoscale Silicon Crystallization Bin Xiang,†,‡,# David J. Hwang,§,# Jung Bin In,∥,# Sang-Gil Ryu,∥ Jae-Hyuck Yoo,∥ Oscar Dubon,†,⊥ Andrew M. Minor,*,†,‡ and Costas P. Grigoropoulos*,∥ †

Department of Materials Science & Engineering, University of California, Berkeley, Berkeley, California 94720-7118, United States National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States § Department of Mechanical Engineering, Stony Brook University, Stony Brook, New York 11794, United States ∥ Department of Mechanical Engineering, University of California, Berkeley, Berkeley, California 94720-7118, United States ⊥ Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States ‡

S Supporting Information *

ABSTRACT: Laser-based processing enables a wide variety of device configurations comprising thin films and nanostructures on sensitive, flexible substrates that are not possible with more traditional thermal annealing schemes.1 In near-field optical probing, only small regions of a sample are illuminated by the laser beam at any given time.2 Here we report a new technique that couples the optical near-field of the laser illumination into a transmission electron microscope (TEM) for real-time observations of the laser−materials interactions. We apply this technique to observe the transformation of an amorphous confined Si volume to a single crystal of Si using laser melting. By confinement of the material volume to nanometric dimensions, the entire amorphous precursor is within the laser spot size and transformed into a single crystal. This observation provides a path for laser processing of single-crystal seeds from amorphous precursors, a potentially transformative technique for the fabrication of solar cells and other nanoelectronic devices.3−5 KEYWORDS: In situ, TEM, laser, near-field probing, nanoscale confinement, amorphous Si, crystallization he crystallization of amorphous thin films is a critical fabrication step for enhancing the performance of thinfilm transistors (TFTs)3,4 and thin-film solar cell devices.5 The use of bulk single-crystal silicon (sc-Si) offers better efficiency yet at high material costs; therefore, it would be beneficial to find an alternative paradigm for scalable single-crystal growth that can be generalized for amorphous thin-film precursors. Typical thin-film materials offer cost-effective device fabrication routes but intrinsically suffer from low degree of crystallinity, leading to necessary improvements by subsequent thermal annealing. Using a traditional furnace to increase crystallinity not only needs a large thermal budget but also has limitations in adopting inexpensive substrates such as Pyrex, soda-lime glass, or polymer substrates.6 Annealing by pulsed lasers can significantly mitigate these issues by taking advantage of precisely localized heating.7 Some groups presented in situ synthesis of 1-D nanostructures by far-field optics.8,9 Recently, it has been demonstrated that laser-induced nanoscale melting of amorphous silicon (a-Si) can effectively lead to the formation of single crystalline nanodomains after resolidification on single crystal Si substrates.10,11 Until now, however, direct observation of the thickness-wise nanoscale transformation of amorphous Si

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© 2012 American Chemical Society

to single crystal Si on an epitaxially nonparticipating substrate has not been reported. In this paper, we have implemented a novel in situ transmission electron microscopy (TEM) monitoring technique to observe the crystallization of a-Si during laser irradiation by directly coupling a laser beam into a transmission electron microscope through a fiber-optic probe. By realizing a near-field scanning optical microscopy (NSOM) fiber probe scheme,12 this in situ technique opens up a wide variety of new characterization possibilities, considering the atomic-scale spatial and analytical resolution of modern transmission electron microscopes. As opposed to prior in situ TEM observations of laser-induced phase transformations,13,14 our approach uses a near-field technique that only illuminates a nanoscale region of the sample, allowing for multiple experiments on one sample through nanomanipulation of the optical fiber probe from one spot to the next while precisely maintaining the probe−sample distance (i.e., optical near-field) under in situ TEM inspection. This configuration allows for Received: February 22, 2012 Revised: March 27, 2012 Published: April 3, 2012 2524

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laser pulses and energy levels (except for a first pulse shot that was too high and led to slight ablation at the beginning). Figure 3 shows an a-Si pillar (Figure 3a) that was transformed back and forth between polycrystalline (Figure 3b,d) and single crystal (Figure 3c,e) caps at the top of the pillar. The pillar structure can be alternatively switched between polycrystalline and single crystal phases by the increasing of the number of laser pulse shots. Table 1S demonstrates the effect of pillar size and the effect of film thickness. Based on the experiment data shown in Table S1, 580 nm is a critical size of pillar would be concluded, above which the a-Si cannot be completely transformed into single crystal instead of polycrystalline structure after a number of laser shots. In principle, the grain size resulting from the laser-induced crystallization process can be estimated by considering the nucleation rate and growth speed. For instance, slow nucleation (or low undercooling) and fast growth rate facilitate the formation of larger grains. Thus, in the case of a constant temperature process, we can readily calculate an average grain size. As the nucleation rate is much more sensitive to the amount of undercooling than growth speed is, a lower undercooling (or smaller temperature drop) is more favorable to achieve single crystalline silicon (sc-Si). Because of the localized heating, however, pulsed laser-induced crystallization is accompanied by a highly transient heat transfer such that a constant temperature condition cannot be established during crystallization. Therefore, the temperature evolution during crystallization has a critical influence on the resulting microstructure. In addition, the stochastic nature of spontaneous nucleation is a crucial factor especially for confined material15 as the randomness of our sc-Si formation event implies. While the classical nucleation theory provides a fundamental tool for control of microstructure, it only estimates a mean value of nucleation events in a deterministic way, which is more adequate in a larger domain. Because of the spatial confinement, however, our crystallization system highlights the stochastic behavior via formation of sc-Si. Thus, a stochastic nucleation model should be combined with thermal analysis in order to allow more quantitative analysis on the phenomenon. For instance, Leonard et al.16,17 proposed that Poisson statistics can reasonably describe the stochastic nature of laser-induced crystallization in thin silicon films, based on Skripov’s formula.18 They observed that the temperature of the first nucleation event varies by approximately ±40 K after complete melting of the silicon layer. Considering that a nucleation rate is very sensitive to the nucleation temperature, this highly stochastic nature of nucleation would result in a remarkably different kinetic behavior of crystal growth in our system. In order to gain a comprehensive insight into the crystallization of a single crystal Si cap, we carried out a 3D transient heat transfer simulation combining the classical homogeneous and heterogeneous nucleation mechanisms in an additive manner, based on the finite difference method (FDM) that was proposed by Leonard et al.16,17 and similarly by Kisdarjono et al.19 (also see Supporting Information) This method provides a fast simulation tool by coupling a stochastic nucleation model with the heat equation; for the domain size of practical laser-induced crystallization is beyond the computational capability that MD simulations generally accommodate. In our simulation, both (homo and hetero) terms were used simultaneously, but heterogeneous nucleation prevailed over the homogeneous mechanism at relatively higher temperature,

precise control of the laser energy absorption in the irradiated Si dot without affecting the supporting structure and inducing unwanted thermomechanical effects. With this novel integrated approach, we have been able to conduct a fundamental investigation into the nanoscale crystallization of a-Si under different laser energy levels. Our results show that the intrinsic trend of multiple nucleation sites forming within the molten volume can be effectively suppressed, and melting and resolidification of the entire volume enabled the seedless growth of a Si single crystal. Figure 1 shows a schematic of the NSOM probe coupled in a transmission electron microscope. Our system is capable of true

Figure 1. Schematic of the in situ TEM optical near-field setup. The fiber-optic probe can be manipulated in three dimensions with a piezo motion control system to approach the sample surface in a direction orthogonal to the electron beam.

near-field conditions below the diffraction limit, demonstrated in the examples shown in Figure S1. However, in the Si crystallization work we intentionally located a-Si pillars at an increased distance from the NSOM probe end (∼150 nm gap distance) for the purpose of in situ monitoring the initiation of a well-controlled crystalline structure throughout the depth of the a-Si film. The fiber probe tip is precisely positioned to illuminate amorphous a-Si/SiO2/c-Si structures prepared by focused ion-beam (FIB) processing. This sample configuration was chosen to simulate a patterned amorphous Si film on a SiO2 substrate. Upon alignment of the fiber probe to the sample, the probe delivers 532 nm laser irradiation with an effective illumination spot size dependent on the distance from the tip to the sample. The effective laser spot diameter on top of the Si pillar is estimated to be ∼800−1000 nm depending on the relative direction with respect to the laser polarization. Finite difference time domain (FDTD) simulation supports this estimation. The calculated intensity profiles are displayed in Figure S2. It is shown that while tight laser confinement is achieved at the near-field, the effective beam spot diameter increases up to ∼800−1000 nm, and the spatial beam profile is noticeably flat at a sample−probe gap distance of ∼150 nm. Thus, this relatively large spot size provides uniform illumination to our samples. For the a-Si experiments described in this paper, the tip to sample distance was adjusted to result in an effective spot size on the order of ∼800−1000 nm in diameter. Figure 2 shows the results from in situ laser-induced crystallization of an a-Si pillar structure with a nanosecond (NS) laser where the entire sample width is within the spot size of the laser. Progressing from amorphous Si (Figure 2a) to polycrystalline (Figure 2b) to a final cap of single crystal Si (Figure 2c,d) was realized through increasing the number of 2525

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Figure 2. (a) As-prepared nanoscale a-Si pillar fabricated by FIB. (b) TEM image of polycrystalline cap structure (pc-Si) observed after irradiation by several nanosecond (NS) laser pulse shots. (c) Bright field TEM image of single crystalline nanoscale Si cap (sc-Si) achieved after the polycrystalline Si cap shown in (b) irradiated by a single pulse shot of NS laser beam; in the inset, electron diffraction pattern of the nanoscale single crystalline Si as shown in (c), zone axis is [−112]. (d) The corresponding dark field TEM image of the nanoscale single crystalline Si grain as shown in (c), using the [220] reflection to only illuminate the lone Si grain.

Each curve in Figure 4a represents a thermal history of formation of sc-Si and pc-Si. With only a thin film of silicon dioxide underneath the a-Si cap, cooling by radiation is negligible, and most of the heat is transferred down the pillar through the silicon oxide to the silicon substrate below it by conduction. Because of the rapid cooling (∼7 × 109 K/s), the temperature decreases significantly lower than the melting point with no occurrence of crystal nucleation. Once nucleation of crystalline silicon is initiated, the temperature reaches local minima with delay of several nanoseconds, at around 1240− 1330 K. These minima are marked by ii and ii′ in the curves depicted in Figure 4a where thermal history differentiates depending on the nucleation events. Subsequently, the growth of crystalline silicon generates latent heat (1320 J/g) that results in the temperature rising almost to the melting temperature again (Figure 4d), via the so-called recalescence. As mentioned previously, a homogeneous nucleation does not occur since heterogeneous nucleations prevail over homogeneous mechanism. During the deep undercooling (Figure 4c) the amorphous phase may nucleate, although it was not considered in the simulation. However, we believe that the

not allowing more cooldown. In other words, homogeneous nucleation did not occur at all in the simulation. To simplify the geometry bearing no effect on the predictions, we assumed a rectangular Si−SiOx−Si pillar with the tip volume of 400 × 400 × 200 nm3, which would describe the crystallization process of Figure 3. As we applied laser repeatedly on as-produced polysilicon as well as on the initial amorphous silicon, the laser power of the simulation was adjusted to be high enough to heat the target above the melting temperature (1685 K) of crystalline silicon and thereby result in a molten silicon tip volume. We also assumed that the quasi-steady-state (QSS) value of the nucleation rate was reasonably applicable for our case. More elaborate nucleation kinetics such as athermal20,21 and non-quasi-steady-state (non-QSS)21 nucleation may emerge during the relatively fast cooling rate. However, several theoretical papers confirm that the calculated cooling rate (∼7 × 109 K/s) is in the quasi-steady-state regime21,22 (or moderately close to QSS20). The simulation result of the heat profile from our experiments is shown in Figure 4a, with selected times shown in Figure 4b−f, starting with cooling from the melting point. 2526

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Figure 3. Demonstration of an alternative process of laser-induced transformation between polycrystalline and single crystal Si, under laser pulse energy of 24 nJ. The as-prepared a-Si pillar (a) (∼420 nm in diameter; ∼230 nm thick a-Si) is transformed into polycrystalline Si (b) and, eventually, into single crystal Si pillar (c) by after ∼50 laser pulse shots. Upon continuing the irradiation, the single crystal Si pillar (c) can be transformed back into a polycrystalline Si pillar (d). The cap can then be recrystallized into a single crystal Si pillar (e) again. Insets in (c) and (e) are the corresponding electron diffraction pattern of the single crystal Si caps.

small fraction of the amorphous silicon possibly disappears during the recalescence process,23 as previously observed experimentally.21 This is also consistent with our experimental results where we did not observe any amorphous phase in the

Si cap after laser melting. Upon recalescence, the probability of additional nucleation decreases quickly and the latent heat generated by the crystal growth is balanced with the loss of heat by conduction. The result is an almost constant temperature 2527

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Figure 4. Representative transient temperature profiles of the silicon pillar where the nucleation event started at the interface during solidification. (a) The temperature evolution plot with time corresponds to the top center of the silicon melt (the black solid line traces sc-Si formation whereas the dashed gray indicates pc-Si formation.) (b−f) 3D presentations of temperature and phase data near the molten tip. Transparency was adopted to distinguish liquid (semitransparent) and solid (opaque) phase. (b) Complete liquid tip starts to be cooled, which we set to zero time. (c) The sample has cooled to ∼1320 K, at which the small nucleus starts to grow. (d, e) The temperature of the Si rises due to the latent heat of crystallization and stays at almost constant temperature until the Si cap has fully crystallized and the temperature drops again (f).

(Figure 4d,e) until the crystallization is completed, and finally the solid silicon is quenched quickly (Figure 4f). The variation in the nucleation temperatures presents the stochastic behavior we have mentioned. Indeed, higher nucleation temperature facilitates fewer numbers of grains; the solid line in Figure 4a results in a sc-Si while the dashed one ends up with a pc-Si of 25 grains. In this respect, we would like to comment that our confined crystallization system has an additional potential as an excellent tool in scrutinizing stochastic aspects of nucleation mechanism. For example, when we tried the simulation multiple times to roughly get probability information for sc-Si, we obtained 2 of sc-Si events out of 100 trials (Figure S3), which is a good agreement with the experiment. We would like to emphasize that the specific nucleation mechanism (homo- or heterogeneous nucleation) does not undermine the stochastic nature of spontaneous nucleation we focus on in this study. Indeed, when we disable the heterogeneous term, the nucleation temperature downshifts by ∼110 K, but we observe very similar stochastic behavior that allows formation of sc-Si. We would like to highlight that one of the advantages of our in situ technique is that the effective beam size adjustment via sample−probe gap distance adjustment could be monitored using the TEM imaging. It should be noted that the active gap distance adjustment is an extremely difficult task in the traditional scanning probe microscope (SPM) environment and typically relies on the dynamic response of oscillating probes with significant uncertainties. With our current in situ configuration we have a unique and significant advance in the near-field optical probing field in that it provides an actual means to monitor and control the gap distance. The configuration implemented in the present paper ensures accurate coupling of the laser energy with the Si dot experiencing a phase change without affecting the supporting

structure. This would have been practically impossible with farfield optics that even under extremely tight focusing would have resulted in substantial damage to the underlying materials since the depth of focus would have been on the order of 500 nm to 1 μm. Our in situ TEM observations lead to a possibility of a new path toward realizing monocrystalline Si seeds for optoelectronic device fabrication. Recently, efforts in utilizing nanocrystalline Si (nc-Si) for photovoltaic technologies have been made for improved silicon solar cell device efficiency.24,25 A greater degree of crystallinity leads to higher electron mobility, better absorption in solar spectrum, and a more stable structure for long-term exposure to the solar light. To date, laser-assisted nc-Si formation has been mainly explored using excimer lasers.26,27 As imagined, the device-processing scheme enabled by spatial confinement and laser crystallization could involve a variety of massive and cost-effective lithography techniques (possibly nanoimprint lithography or e-beam lithiography) to fabricate the base structures and subsequent laser crystallization steps to complete the processes. The subsequent nanoscale crystalline dot array could then serve as controlled seeds for polycrystalline or even epitaxial film growth by subsequent chemical vapor deposition (CVD) steps.28 Our templating approach would then produce rationally patterned polycrystalline film where the location of grain boundaries would be specified by controlling growth through this seeded process. It is also significant to appreciate that the herein presented method is not limited to homoepitaxial growth as the process of generating single crystalline seeds from amorphous precursors is truly independent of the kind or nature of the nonparticipating substrate material. Presumably, the base nanoscale structures need not be limited to amorphous or Si materials, and the spatial confinement plus the laser-processing scheme 2528

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NSEC. The in situ experiments were performed at the National Center for Electron Microscopy at the Lawrence Berkeley National Laboratory, which is supported by the Office of Science, Office of Basic Energy Sciences, Scientific User Facilities Division, of the U.S. Department of Energy under Contract DE-AC02-05CH11231.

could be adapted to different materials and laser spot sizes to generate arbitrary fabrication schemes. Lastly, this work demonstrates that in situ TEM observation of laser−materials interaction at the nanoscale can be a powerful tool for achieving rapid feedback for the optimization of laser-processing parameters. Methods. Amorphous silicon nanopillars were prepared from a bulk 100 nm SiO2/single-crystal silicon wafer sputtercoated with 2 μm thick amorphous silicon on the top of SiO2 layer. The nanopillars were fabricated by using a FEI 235 dual beam focus ion beam (FIB). A 20 000 pA ion beam current was applied for the large area milling at the beginning of two opposing trenches on the substrate. A thin window on the substrate was left for the nanopillar fabrication. Pillars with an aspect ratio ranging from 5:1 to 10:1 were fabricated on the window platform. With the decrease of ion beam current, a 10 pA ion beam current during the final nanopillar milling stage was employed. The amount of ion beam damage to the sample is as minimal as can be effectively achieved. Fibers (SM980-5.8-125, Thorlabs, Inc.) with 1/2 wavelength end size pulled by using P-2000 puller (Sutter Instrument Co.) can be installed into our in situ TEM holder, which is equipped with a piezo-driven control system for the fiber probe fine positioning and a manual three-axis translation stage for the fiber probe coarse positioning. The spacing between the probe tip and the sample was ∼150 nm. During laser-induced crystallization process, a nanosecond Nd:YAG laser beam of 532 nm wavelength with 5 ns pulse duration was delivered onto the surface of the sample through the fiber. For precise measurement of the pulse energy emitted from the fiber probes, two energy meters were used. Since in the actual process the output laser energy emerging from the probe apex cannot be measured, only the input energy at the split side is measured and the output energy is estimated using the precalibrated curve. The slope of the curve was defined as the fiber probe transmission efficiency, and the transmission efficiency of around ∼0.8% was achieved under the experimental conditions used. This relationship was checked before and after the experiment to ensure that there was no probe damage.

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(1) Park, I.; Ko, S. H.; Pan, H.; Grigoropoulos, C. P.; Pisano, A. P.; Fréchet, J. M. J.; Lee, E. S.; Jeong, J.-H. Adv. Mater. 2008, 20, 489. (2) Pan, H.; Hwang, D. J.; Ko, S. H.; Clem, T. S.; Fréchet, J. M. J.; Bäuerle, D.; Grigoropoulos, C. P. Small 2010, 6, 1812. (3) Im, J. S.; Kim, H. J. Appl. Phys. Lett. 1994, 64, 2303. (4) Brotherton, S. D. Semicond. Sci. Technol. 1995, 10, 721. (5) AndraÈ, G.; Bergmann, J.; Falk, F.; Ose, E. Thin Solid Films 1999, 337, 98. (6) Sposili, R. S.; Im, J. S. Appl. Phys. Lett. 1996, 69, 2864. (7) Hatano, M.; Moon, S.; Lee, M.; Suzuki, K.; Grigoropoulos, C. P. J. Appl. Phys. 2000, 87, 36. (8) Taheri, M. L.; Reed, B. W.; Lagrange, T. B.; Brownin, N. D. Small 2008, 4, 2187. (9) Taheri, M. L.; Lagrange, T.; Reed, B. W.; Armstrong, M. R.; Campbell, G. H.; Dehope, W. J.; Kim, J. S.; King, W. E.; Masiel, D. J.; Brownin, N. D. Microsc. Res. Tech. 2009, 72, 122. (10) Chimmalgi, A.; Hwang, D. J.; Grigoropoulos, C. P. Nano Lett. 2005, 5, 1924. (11) Arora, H.; Du, P.; Tan, K. W.; Hyun, J. K.; Grazul, J.; Xin, H. L.; Muller, D. A.; Thompson, M. O.; Wiesner, U. Science 2010, 330, 214. (12) Hwang, D. J.; Grigoropoulos, C. P.; Yoo, J.; Russo, R. E. Appl. Phys. Lett. 2006, 89, 254101. (13) Taheri1, M. L.; McGowan, S.; Nikolova, L.; Evans, J. E.; Teslich, N.; Lu, J. P.; LaGrange, T.; Rosei, F.; Siwick, B. J.; Browning, N. D. Appl. Phys. Lett. 2010, 97, 032102. (14) Nikolova, L.; LaGrange, T.; Reed, B. W.; Stern, M. J.; Browning, N. D.; Campbell, G. H.; Kieffer, J. C.; Siwick, B. J.; Rosei, F. Appl. Phys. Lett. 2010, 97, 203102. (15) Evans, P. V.; Devaud, G.; Kelly, T. F.; Kim, Y. K. Acta Metall. Mater. 1990, 38, 719. (16) Leonard, J. P.; Im, J. S. Appl. Phys. Lett. 2001, 78 (22), 3454. (17) Leonard, J. P.; Im, J. S. Mater. Res. Soc. Symp. Proc. 2000, 580, 233. (18) Skripov, V. P. Cryst. Growth Mater. 1997, 10, 328. (19) Kisdarjono, H.; Voutsas, A. T.; Solanki, R. J. Appl. Phys. 2003, 94 (7), 4374. (20) Im, J. S.; Gupta, V. V.; Crowder, M. A. Appl. Phys. Lett. 1998, 72 (6), 662. (21) Shneidman, V. A. J. Appl. Phys. 1999, 85 (3), 1981. (22) Shneidman, V. A. J. Appl. Phys. 1996, 80 (2), 803−811. (23) Stiffler, S. R.; Thompson, M. O.; Peercy, P. S. Phys. Rev. Lett. 1998, 60 (24), 2519. (24) Svrcek, V.; Slaoui, A.; Muller, J. C. Thin Solid Films 2004, 451− 452, 384. (25) Stupca, M.; Alsalhi, M.; Saud, T.; Almuhanna, A.; Nayfeh, M. H. Appl. Phys. Lett. 2007, 91, 063107. (26) Ehbrecht, M.; Ferkel, H.; Smirnov, V. V.; Stelmakh, O. M.; Zhang, W.; Huisken, F. Rev. Sci. Instrum. 1995, 66, 3833. (27) Huisken, F.; Hofmeister, H.; Kohn, B.; Laguna, M. A.; Paillard, V. Appl. Surf. Sci. 2000, 154−155, 305. (28) Vanamu, G.; Datye1, A. K.; Zaidi, S. H. Appl. Phys. Lett. 2006, 88, 204104.

ASSOCIATED CONTENT

S Supporting Information *

Figures S1−S3, Table S1. This material is available free of charge via the Internet at http://pubs.acs.org.

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REFERENCES

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (A.M.M.); cgrigoro@me. berkeley.edu (C.P.G.). Author Contributions #

The authors contributed equally to this publication.

Notes

The authors declare the following competing financial interest(s): CPG has a certain financial interest in Appliflex LLC.

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ACKNOWLEDGMENTS This work was supported by the DARPA/MTO under the TBN grant N66001-08-1-2041. The authors acknowledge that the research was supported in part by a US Department of Energy SBIR grant (DE-FG02-07ER84813) awarded to Appliflex, LLC. J.B.I. was supported by the NSF SINAM 2529

dx.doi.org/10.1021/nl3007352 | Nano Lett. 2012, 12, 2524−2529