Search query: 2009 Nano . 9 1636

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abstract: Based on a comparative study between the piezoelectric outputs of n type nanowires (NWs) and n core/p shell NWs along with the previous study (Lu et al 2009 Nano. Lett. 9 1223) we demonstrate a one step technique for identifying the conductivity type of individual ZnO nanowires (NWs) based on the output of a piezoelectric nanogenerator without destroying the sample. A negative piezoelectric output voltage indicates an NW is n type and it appears after the tip scans across the center of the NW, while a positive output voltage reveals p type conductivity and it appears before the tip scans across the central line of the NW. This atomic force microscopy based technique is reliable for statistically mapping the majority carrier type in ZnO NWs arrays. The technique may also be applied to other wurtzite semiconductors, such as GaN, CdS and ZnS.
author_names: S S Lin

J H Song

Y F Lu

Zhong Lin Wang
title: Identifying individual n and p type ZnO nanowires by the output voltage sign of piezoelectric nanogenerator.
venue: Nanotechnology
year: 2009

abstract: We give a concise introduction into the radiative heat transfer at the nanoscale discussing the contribution of propagating, frustrated and coupled surface modes [1] Especially, the latter contribution results in a heat flux, which can exceed the heat flux between two black bodies by several orders of magnitude for distances in the nanometer regime [1] The prediction of such an enormous heat flux enhancement is usually based on Rytov's fluctuational electrodynamics [2] and has been verified in some very recent experiments [3,4,5] Our aim is to show how the theoretical expression describing the nanoscale heat flux can be interpreted in terms of transmission coefficients and the universal quantum of thermal conductance by means of concepts of mesoscopic physics [6] Such a formulation allows for studying the fundamental limits of radiative heat transfer [7,8] emphasizing the trade off between the number of contributing modes and their transmission coefficient. [1] K. Joulain, J. P. Mulet, F. Marquier, R. Carminati, and J. J. Greffet, Surface Science Reports 57, 59 (2005) [2] S. M. Rytov, Y. A. Kravtsov, and V. I. Tatarskii, Principles of Statistical Radiophyics (Springer, New York) Vol. 3. (1989) [3] S. Shen, A. Narayanaswamy, and G. Chen, Nano Lett. 9, 2909 (2009) [4] E. Rousseau, A. Siria, G. Jourdan, S. Volz, F. Comin, J. Chevrier, and J. J. Greffet, Nature Photonics 3, 514 (2009) [5] R. Ottens, V. Quetschke, S. Wise, A. Alemi, R. Lundock, G. Mueller, D. H. Reitze, D. B. Tanner, B. F. Whiting, Phys. Rev. Lett. 107, 014301 (2011) [6] S. A. Biehs, E. Rousseau, and J. J. Greffet, Phys. Rev. Lett. 105, 234301 (2010) [7] P. Ben Abdallah and K. Joulain, Phys. Rev. B 82, 121419 (R) (2010) [8] S. Basu and Z. M. Zhang, J. Appl. Phys. 105, 093535 (2009)
author_names: Emmanuel Rousseau

Alessandro Siria

Guillaume Jourdan

Sebastian Volz

Fabio Comin

Joel Chevrier

Jean-Jacques Greffet
title: Radiative heat transfer at the nanoscale
year: 2009

abstract: Financial support from the Spanish MICINN "MAGPLAS" Grants No. MAT2008 06765 C02 01/NAN and No.NAN2004 09195 C04 02, Funcoat Consolider Ingenio Grant No. 2010 CSD2008 00023) Comunidad de Madrid (References No. S 0505/MAT/0194 and No. S 0505/TIC/0191) CSIC (E. Ferreiro Vila acknowledges JAE) and (Reference No. 2005MX0040) and European Comission (Grant No.NMP3 SL 2008 214107 Nanomagma) is acknowledged. Authors are grateful to D. Ramirez y G. Ramirez for technical assistance. This work was partially supported by CONACYT Mexico under Grants No. 48300 S 3907 and No. J110.474/2006 S 3904. 1 J. R. Sambles, G. W. Bradbery, and F. Yang, Contemp. Phys. 32, 173 1991 2H. Raether, Surface Plasmons on Smooth and Rough Surfaces and on Gratings, Springer Tracts in Modern Physics, Vol. 111 Springer Verlag, Berlin, 1988 3 S. A. Maier, Plasmonics: Fundamentals and Applications Springer, Berlin, 2007 4 J. Homola, Chem. Rev. 108, 462 2008 5 J. N. Anker, W. P. Hall, O. Lyandres, N. C. Shah, J. Zhao, and R. P. Van Duyne, Nature Mater. 7, 442 2008 6B. Sepulveda, A. Calle, L. M. Lechuga, and G. Armelles, Opt. Lett. 31, 1085 2006 7 J. C. Weeber, Y. Lacroute, and A. Dereux, Phys. Rev. B 68, 115401 2003 8B. Steinberger, A. Hohenau, H. Ditlbacher, A. L. Stepanov, A. Drezet, F. R. Aussenegg, A. Leitner, and J. R. Krenn, Appl. Phys. Lett. 88, 094104 2006 9 S. I. Bozhevolnyi, V. S. Volkov, E. Devaux, J. Y. Laluet, and T. W. Ebbesen, Nature London 440, 508 2006 10 I. I. Smolyaninov, J. Elliott, A. V. Zayats, and C. C. Davis, Phys. Rev. Lett. 94, 057401 2005 11 Z. Liu, S. Durant, H. Lee, Y. Pikus, N. Fang, Y. Xiong, C. Sun, and X. Zhang, Nano Lett. 7, 403 2007 12T. Nikolajsen, K. Leosson, and S. I. Bozhevolnyi, Appl. Phys. Lett. 85, 5833 2004 13 J. Gomez Rivas, J. A. Sanchez Gil, M. Kuttge, P. H. Bolivar, and H. Kurz, Phys. Rev. B 74, 245324 2006 14D. Pacifici, H. J. Lezec, and H. A. Atwater, Nat. Photonics 1, 402 2007 15B. Sepulveda, L. M. Lechuga, and G. Armelles, J. Lightwave Technol. 24, 945 2006 16 J. B. Gonzalez Diaz, A. Garcia Martin, G. Armelles, J. M. Garcia Martin, C. Clavero, A. Cebollada, R. A. Lukaszew, J. R. Skuza, D. P. Kumah, and R. Clarke, Phys. Rev. B 76, 153402 FERREIRO VILA et al. PHYSICAL REVIEW B 80, 125132 2009 125132 8
author_names: Elias Ferreiro-Vila

Juan B Gonzalez-Diaz

R Fermento

Maria Ujue Gonzalez

Antonio Garcia-Martin

Jose Miguel Garcia-Martin

Alfonso Cebollada

Gaspar Armelles

David Meneses-Rodriguez

Emilio Munoz Sandoval
title: Intertwined magneto optical and plasmonic effects in Ag/Co/Ag layered structures
year: 2009

abstract: Coaxial manganese oxide/carbon nanotube (CNT) arrays deposited inside porous alumina templates were used as cathodes in a lithium battery. Excellent cyclic stability and capacity of MnO2/CNT coaxial nanotube electrodes resulted from the hybrid nature of the electrodes with improved electronic conductivity and dual mechanism of lithium storage. The reversible capacity of the battery was increased by an order compared to template grown MnO2 nanotubes, making them suitable electrodes for advanced Li ion batteries. A nanoscale approach to electrochemical energy storage applications such as in lithium batteries and supercapacitors has been of great interest because of their unique properties leading to improved performances.1 5 Innovative materials chemistry has been the key to various advancements in lithium rechargeable batteries.6,7 Various transition metal oxides have been widely studied as electrode materials for rechargeable lithium ion batteries because of their high theoretical capacity, safety, environmental benignity, and low cost.8 15 However, poor electronic conductivity of metal oxides limits them from use in high performance lithium ion batteries. One of the other challenging issues is to tackle their capacity decay with cycling, leading from large volume expansion during the lithium uptake/release process. Extensive research efforts are presently devoted to overcome these problems by using electronically conductive additives and carbon coatings.16 19 Among transition metal oxides, manganese oxide (MnO2) has been studied as an electrode for lithium batteries with a high storage capacity, in addition to its low cost, environmental friendliness, and natural abundance.20 23 However its potential applications in practical Li ion batteries are limited due to its poor electrical conductivity and large volume expansion during repeated lithium cycling processes.21 One dimensional (1D) nanostructured morphologies of these electrodes with controlled size, crystallinity, and chemical composition have been designed to overcome some of these challenges.24,25 However, conductivity of these 1D nanowires remains an existing issue. Coaxial 1D nanowires of tin oxide core and indium oxide shell nanostructures have been shown to be promising Liion battery electrodes.26 Coaxial nanowires/nanotubes will lead to multiple functionalities by combining the physical properties of different materials. In order to build Li batteries with improved capacity and power capabilities, coaxial nanowires/nanotubes of multiple materials with specific electrochemical and physicochemical properties need to be fabricated. Here, we demonstrate the fabrication of hybrid coaxial nanotubes of MnO2 and carbon nanotubes (CNTs) as high performance electrodes of lithium batteries. CNTs due to their outstanding electrical properties apart from their high chemical stability, high aspect ratio, strong mechanical strength, and high activated surface area, are attractive electrode materials in energy storage devices, such as electrochemical capacitors, fuel cells, and lithium batteries.27 33 Careful engineering of hybrid coaxial electrode material having both high storage capacity metal oxide and highly conducting CNTs will lead to enhanced Li storage properties. The tubular morphology offers a unique combination of high porosity and low internal resistance. The work is aimed to increase the electronic conductivity and lithium storage capacity of MnO2 nanotubes by using the coaxial nanotubes with MnO2 nanotubes in the outer shell and highly conducting CNT in the inner core. The coaxial electrode configuration will have (i) improved electronic conductivity due to the presence of CNT, (ii) homogeneous electrochemical accessibility and high ionic conductivity by avoiding agglomerative binder and other conductive additives, (iii) To whom correspondence should be addressed, ajayan@rice.edu. Department of Mechanical Engineering Materials Science. Department of Chemical Biomolecular Engineering. NANO LETTERS 2009 Vol. 9, No. 3 1002 1006 10.1021/nl803081j CCC: $40.75 2009 American Chemical Society Published on Web 02/02/2009 well directed 1D conductive paths due to perfect coaxial alignment, and (iv) a dual lithium storage mechanism (insertion/deinsertion in the case of CNTs and formation and decomposition of Li2O in the case of MnO2 nanotubes) The coaxial nanotube structures have been prepared by a combination of simple vacuum infiltration and chemical vapor deposition techniques through a template approach (Figure 1) MnO2 nanotubes are first fabricated by vacuum infiltration inside the channels of commercially available AAO templates (nanopore diameter of ~200 nm and length of ~50 mm) which is followed by the growth of CNTs using chemical vapor deposition (see Supporting Information for method and synthesis, section 1S) The sample was plasma etched for 30 min to remove the amorphous carbon layer that forms during chemical vapor deposition. A layer of Au film ~100 nm) was sputtered onto one side of the template which serves as the current collector for the electrode. The MnO2/CNT hybrid coaxial structures are then released from the alumina templates by dissolving the templates in 3 M NaOH solution for 1 h. The presence of Au film holds the coaxial hybrid structures from collapsing, after the removal of templates. The morphology of the prepared MnO2/CNT Figure 1. Schematic diagram showing the fabrication of MnO2/CNT hybrid coaxial nanotube arrays inside AAO template using a combination of simple vacuum infiltration and chemical vapor deposition techniques. A thin layer of gold ~100 nm) was sputter coated to act as current collector for the electrodes. Figure 2. (a) X ray diffraction patterns of MnO2/CNT coaxial nanotubes, MnO2 nanotubes, and CNTs. XPS spectra of MnO2/CNT coaxial nanotubes (b) before and after (c) first lithium discharge process. Formation of manganese metal and Li2O during the initial lithiation process has been confirmed using XPS. Nano Lett. Vol. 9, No. 3, 2009 1003 coaxial nanotubes was characterized by scanning electron microscopy (FEI Quanta 400 ESEM FEG) and energydispersive X ray spectroscopy (EDX) The powder X ray diffraction patterns were obtained by a Rigaku D/Max Ultima II using Cu KR radiation. For the X ray photoelectron spectroscopy (XPS) analysis, PHI Quantera XPS was used. Flexible films of MnO2/CNT coaxial nanotubes obtained after the removal of templates were directly used as electrode in the lithium battery. Electrochemical measurements were performed in a Swagelok type cell using AUTOLAB PGSTAT 302N potentiostat/galvanostat (see Supporting Information for details, section 2S) An electrochemical test cell was assembled in argon filled glovebox using the active material (MnO2/CNT coaxial nanotubes) as working electrode, lithium metal foil as the counter/reference electrode, and 1 M solution of LiPF6 in 1:1 (v/v) mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC) as electrolyte. A glass microfiber filter was used as separator. The cells were charged and discharged galvanostatically at a rate of 50 mA/g between 3.2 and 0.2 V vs Li/Li+ The crystallinity of the coaxial nanotubes was confirmed by X ray diffraction (XRD) XRD patterns of the MnO2/ CNT coaxial nanotubes show three well resolved peaks that can be indexed as (100) (110) and (211) reflections associated with tetragonal symmetry (Figure 2a) An additional peak at 26.4deg attributed to the (002) plane of hexagonal graphite structure, indicates the presence of CNTs. XRD patterns taken from MnO2 nanotubes and CNTs are in good agreement with the reported literature (PDF#01 0714824(RDB) Elemental analysis by energy dispersive X ray spectroscopy (EDX) revealed the presence of Mn and O (see Supporting Information, Figure 1S) The additional reflections of Al and Au in the EDX spectra are due to AAO template and gold back coating, respectively. MnO2/CNT coaxial nanotubes before and after Li storage were characterized by X ray photoelectron spectroscopy (XPS) XPS profiles of MnO2/CNT hybrid coaxial nanotubes before Li storage (Figure 2b) and after Li storage (Figure 2c) are analyzed by focusing on the regions where the signals of Mn 2p3/2, Mn 2p1/2, C 1s, and Li 1s are expected. The peaks of Mn 2p3/2 and 2p1/2 which are centered at 642 and 653.8 eV, respectively, with a spin energy separation of 11.8 eV are in good agreement with reported data of Mn 2p3/2 and Mn 2p1/2 in MnO2 (Figure 2b).34 Scanning electron microscopy (SEM) images of MnO2/CNT hybrid nanostructures clearly shows uniform coaxial nanostructure (Figure 3) Electrodes are highly porous as revealed from SEM images of individual MnO2 nanotubes and CNTs (see Supporting Information, Figures 1S and 2S) MnO2 nanotube walls were found to be ~10 nm thick with good crystallinity and smooth surface. The structure of coaxial nanotubes such as MnO2 shell thickness and nanotube length could be easily controlled by varying the infiltration time, which will enable us to tune the electrochemical properties of the coaxial nanotubes. CNTs protrude on both ends of the nanotubes, enabling an effective contact with the gold thin film current collector. Recently, coaxial nanowires of MnO2/PEDOT have been synthesized by the coelectrodeposition method.35 However, in the present study, presence of a CNT core has superior advantages in terms of electrical conductivity, mechanical stability, and electrochemical performance. Also, hightemperature annealing results in high crystallinity of the coaxial nanostructure. The growth process of the coaxial nanotubes follows a typical template synthesis, wherein the manganese precursor initially combines with the template by impregnation and results in the nucleation and growth following annealing treatment. This forms the MnO2 shell. The chemical vapor deposition process allows CNTs to grow in the inner cores left by the MnO2 shells. Upon removal of templates, coa
author_names: Leela Mohana Reddy

Manikoth M Shaijumon

Sanketh R Gowda

Pulickel M Ajayan
title: Coaxial MnO 2 Carbon Nanotube Array Electrodes for High Performance Lithium Batteries Arava
year: 2009

abstract: Biologically functional cationic phospholipid gold nanoplasmonic carriers have been designed to simultaneously exhibit carrier capabilities, demonstrate improved colloidal stability, and show no cytotoxicity under physiological conditions. These carriers are able to retain their unique nanoscale optical properties under physiological conditions, making them particularly useful in a wide range of imaging, therapeutic, and gene delivery applications that utilize selective nanoplasmonic properties. Introduction Gold nanoparticles (GNPs) in the near infrared (NIR) spectral region, due to their size and core material, display unique optical properties that make them attractive candidates for drug delivery, gene delivery, biomedical and molecular imaging, and therapeutics. Because of their large surface area, GNPs are ideal carriers of biomolecules for these applications. While attached to carriers, biomolecules are in an inactive state. When GNP carriers are specifically used to convert light into heat, otherwise known as photothermal conversion, released biomolecules enter an active state to freely interact with the environment. Such optically activated GNP carriers are referred to as nanoplasmonic carriers. In particular, the NIR wavelength regime is well suited for biomedical applications because tissues and cells are essentially transparent at 800 1300 nm. It is possible to obtain very efficient photothermal conversion of energy when the NIR light is matched to the plasmon resonance wavelength of the GNP. Additionally, heat transfer from the surface of GNPs to the surrounding cellular environment is highly localized, decaying exponentially within a few nanometers, and therefore is thought to have minimal adverse effects on cells. Among the GNPs, rod shaped GNPs, known as nanorods, are of particular interest due to their large absorption cross section, a narrow spectral width of the longitudinal plasmon resonance band, and tunability of the longitudinal plasmon resonance wavelength based on aspect ratio. The unique optical properties of gold nanorods arise due to their nanoscale asymmetric geometry and gold core material. University of California Berkeley. Sandia National Laboratories. (1) Sershen, S. R. Westcott, S. L. Halas, N. J. West, J. L. J. Biomed. Mater. Res. 2000, 51, 293 298. (2) Ren, L. Chow, G. M. Mater. Sci. Eng. 2003, C23, 113 116. (3) Skirtach, A. G. Javier, A. M. Kreft, O. Kohler, K. Alberola, A. P. Mohwald, H. Parak, W. J. Sukhorukov, G. B. Angew. Chem. Int. Ed. 2006, 45, 4612 4617. (4) Wijaya, A. Schaffer, S. B. Pallares, I. G. Hamad Schifferli, K. ACS Nano 2009, 3, 80 86. (5) Chen, C. Lin, Y. Wang, C. Tzeng, H. Wu, C. Chen, Y. Chen, C. Chen, L. Wu, Y. J. Am. Chem. Soc. 2006, 128, 3709 3715. (6) Horiguchi, Y. Niidome, T. Yamada, S. Nakashima, N. Niidome, Y. Chem. Lett. 2007, 36, 952 953. (7) Lee, S. E. Liu, G. L. Kim, F. Lee, L. P. Nano Lett. 2009, 9, 562 570. (8) Copland, J. A. Eghtedari, M. Popov, V. L. Kotov, N. Mamedova, N. Motamedi, M. Oraevsky, A. A. Mol. Imaging Biol. 2004, 6, 341 349. (9) Javier, D. J. Nitin, N. Roblyer, D. M. Richards Kortum, R. J. Nanophotonics 2008, 2, 1 12. (10) Choi, Y. Kang, T. Lee, L. P. Nano Lett. 2009, 9, 85 90. (11) Nallathamby, P. D. Lee, K. J. Xu, X. N. ACS Nano 2008, 2, 1371 1380. (12) Huang, X. El Sayed, I. H. Qian, W. El Sayed, M. A. J. Am. Chem. Soc. 2006, 128, 2115 2120. (13) Huang, Y. F. Sefah, K. Bamrungsap, S. Chang, H. Tan, W. Langmuir, in press. (14) Pissuwan, D. Valenzuela, S. M. Killingsworth, M. C. Xu, X. J. Nanopart. Res. 2007, 9, 1109 1124. (15) Choi, M. Stanton Maxey, K. J. Stanley, J. K. Levin, C. S. Bardhan, R. Akin, D. Badve, S. Sturgis, J. Robinson, J. R. Bashir, R. Halas, N. J. Clare, S. E. Nano Lett. 2007, 7, 3759 3765. (16) Loo, C. Hirsch, L. Lee, M. Chang, E. West, J. Halas, N. Drezek, R. Opt. Lett. 2005, 30, 1012 1014. (17) Skabalak, S. E. Chen, J. Au, L. Lu, X. Li, X. Xia, Y. AdV. Mater. 2007, 19, 3177 3184. (18) Gobin, A. M. O'Neal, D. P. Watkins, D. M. Halas, N. J. Drezek, R. A. West, J. L. Laser Surg. Med. 2005, 37, 123 129. (19) Hauck, T. S. Jennings, T. L. Yatsenko, T. Kumaradas, J. C. Chan, W. C. AdV. Mater. 2008, 1 7. (20) Norman, R. S. Stone, J. W. Gole, A. Murphy, C. J. Sabo Attwood, T. L. Nano Lett. 2008, 8, 302 306. (21) Cortie, M. Xu, X. Chowdhury, H. Zareie, H. Smith, G. Proc. SPIE 2005, 5649, 565 573. (22) Khlebtsov, B. Zharov, V. Melnikov, A. Tuchin, V. Khlebtsov, N. Nanotechnology 2006, 17, 5167 5179. (23) Liu, G. L. Kim, J. Lu, Y. Lee, L. P. Nat. Methods 2006, 5, 27 32. (24) Svoboda, K. Block, S. M. Annu. ReV. Biomol. Struct. 1994, 23, 247 285. (25) Hirsch, L. Stafford, R. Bankson, J. Sershen, S. Price, R. Hazle, J. Halas, N. West, J. Proc. Second Joint EMBS BMES Conf. 2002, 1, 530. (26) Skirtach, A. Dejugnat, C. Braun, D. Susha, A. Rogach, A. Parak, W. Mohwald, H. Sukhorukov, G. Nano Lett. 2005, 5, 1371 1377. (27) Lu, Y. Liu, G. L. Kim, J. Mejia, Y. X. Lee, L. P. Nano Lett. 2005, 5, 119 124. ohio2/yja yja/yja yja/yja99907/yja1740d07z xppws 23:ver.3 8/27/09 21:20 Msc: ja 2009 04326j TEID: dmadmin BATID: 00000 10.1021/ja904326j CCC: $40.75 XXXX American Chemical Society J. AM. CHEM. SOC. XXXX, xxx, 000 9 A
author_names: Somin Eunice Lee

Darryl Y Sasaki

Thomas D Perroud

Daniel Yoo

Kamlesh D Patel

Luke P Lee
title: Biologically Functional Cationic Phospholipid Gold Nanoplasmonic Carriers
year: 2009

abstract: Semiconductor nanocrystals have received significant attention for their potential as low cost materials in electronic devices. Nanocrystals may be processed with many of the inexpensive techniques that make organic materials attractive, yet they are expected to have superior mobilities and resistance to degradation. In addition, new opportunities arise at the nanoscale such as tunable absorption and emission via quantum confinement, and multiple exciton generation that may allow for entirely new device architectures. However, in order to be integrated into devices, films of nanocrystals must be developed and the films would ideally have the same properties as their constituent crystals. Significant headway has been made with II VI and VI IV solutionsynthesized nanocrystals such as CdSe and PbSe, but progress with the popular group IV materials Si and Ge has lagged behind because of a lack of a satisfactory synthesis route. We present on plasma synthesized Si and Ge nanocrystals, thin films of these crystals, and electronic devices based on the films. Nanocrystals are synthesized using a nonthermal plasma approach in which precursor gases are dissociated in a radiofrequency discharge and atomic clustering leads to nanoparticle nucleation [1] The plasma environment is uniquely suited to synthesizing nanocrystals since a unipolar negative charge is dispensed on the particles, suppressing agglomeration, and the particles are selectively heated, allowing for crystallization of high melting point materials. Silicon and Ge nanocrystals are produced with respectable mass yields, yet retain a narrow size distribution as is necessary to avoid diluting the tunability of their optical properties. Additionally, the nanocrystals are synthesized as a powder with hydrogen terminated surfaces, which makes this process particularly clean and allows for flexibility in subsequent processing. Several methods have been developed for depositing thin films of nanocrystals. In one technique, Si or Ge nanocrystalline powder is transferred into solution and films are spun onto substrates. The nanocrystals are not solubilized by most solvents, leading to flocculation and poor film morphology. However, solvents with similar dielectric constants as Si and Ge suppress the van der Waals attraction between nanocrystals, resulting in a stable colloid. Films cast from these solutions are particularly smooth and have been used for thin film field effect transistors. Germanium nanocrystal transistors exhibit n type behavior with electron mobilities as large as 7x10 cm/Vs and on off ratios of 10. We have also investigated inorganic organic hybrid devices based on Si nanocrystals and poly3(hexylthiophene) (P3HT) P3HT was added to Si nanocrystal solutions and blend films were spun as the active layer in photovoltaic devices. These devices, in which the Si nanocrystals act as the electron conductor and the P3HT acts as the hole conductor, have achieved open circuit voltages of 0.8 V, short circuit currents of 3.9 mA/cm, and power conversion efficiencies of 1.5 under solar irradiation [2] References [1] L. Mangolini, E. Thimsen, and U. Kortshagen, Nano Lett. 5, 655 (2005) [2] C.Y. Liu, Z.C. Holman, and U.R. Kortshagen, Nano Lett. 9, 449 (2009)
author_names: Chin-Yi Liu

Uwe R Kortshagen
title: Silicon and Germanium Nanocrystal Electronic Devices
year: 2009

abstract: We report solar cells based on highly confined nanocrystals of the ternary compound PbSxSe1 x. Crystalline, monodisperse alloyed nanocrystals are obtained using a one pot, hot injection reaction. Rutherford back scattering and energy filtered transmission electron microscopy suggest that the S and Se anions are uniformly distributed in the alloy nanoparticles. Photovoltaic devices made using ternary nanoparticles are more efficient than either pure PbS or pure PbSe based nanocrystal devices. Colloidal semiconductor nanocrystals display a wealth of size dependent physical and chemical properties, including quantum confinement effects, shape dependent electronic structure,1,2 and control over assembly through modification of surface functionalization.3,4 Photovoltaic devices are an easily recognized potential application for nanocrystals due, in part, to their high photoactivity, solution processability and low cost of production. Several schemes for using nanocrystals in solar cells are under active consideration, including nanocrystal polymer composites,5 nanoparticle array solar cells,6 films of partially sintered nanoparticles,7 and nanocrystal analogues to dye sensitized solar cells.8 A persistent challenge for any nanoparticle based solar cell is to take advantage of quantum confinement effects to improve the optical absorption process without overly hindering the subsequent transport of charge to the electrodes. Various binary semiconductor nanoparticles, like CdSe, CdTe, Cu2S, InP, and InAs, have been explored for photovoltaic devices but the reported efficiencies remain low, mostly limited by poor charge transport between the nanocrystals.5,7 12 With so many parameters to adjust in terms of size and shape, little work has focused on ternary or quaternary compositions of nanoparticles for solar cells. Yet it is well known from thin film solar cell studies that such compositional tuning can sometimes yield significant improvements in performance. The Pb chalcogenide family of nanocrystals is actively investigated for nanocrystal solar cell applications because they have such large exciton Bohr radii (PbS 18 nm, PbSe 47 nm, and PbTe 150 nm) In the limit where the nanocrystals are only a tenth or so of the bulk exciton diameter, electrons and holes can tunnel through a thin organic surface coating, and therefore strong electronic coupling between particles facilitates transport of charge between nanocrystals. So far, Schottky solar cells based on binary compositions of PbSe and PbS nanocrystals have been investigated. PbSe nanocrystal solar cells generate larger short circuit photocurrents while PbS nanocrystal devices with similar bandgap have shown a larger VOC. This opens the possibility to better engineer particles, by creating ternary PbSxSe1 x to simultaneously optimize both carrier transport and voltage. Moreover, the properties of PbS and PbSe lead to an ideal substitutional alloy; the atomic anion radii are within 15% of each other, the lattice mismatch factor is only 2% between PbS and PbSe (see Supporting Information for the similarity of the XRD patterns) and of course the anions are isovalent. Although the synthesis of ternary PbSxSe1 x nanocrystals13,14 remains underdeveloped compared to the widely studied cadmium chalcogenides alloys,15 17 we successfully obtained monodisperse, highly crystalline nanocrystals using a one pot, hot injection synthesis. Lead oxide (PbO, 99.999% selenium (99.99% oleic acid (OA, technical grade, 90% diphenylphosphine (DPP, 98% 1,3 benzenedithiol (BDT, >98% bis(trimethylsilyl) sulfide (TMS2S, purum) 1 octadecene (ODE, 90% anhydrous solvents and aluminum shot (99.999% were purchased from Aldrich and used as received. Trioctylphosphine (TOP, >97% was acquired from Strem. Nanocrystal synthesis was performed under argon atmosphere using standard air free Schlenk line techniques. The synthesis scheme is as follows: A solution of 446 mg PbO (2 mmol) 1.4 g oleic acid (5 mmol) and 10 g ODE was degassed and heated to 150 degC in a 50 mL three neck flask for one hour. A mixture of the proper amount 1 M TOP/Se solution, TMS2S, DPP (40 mg) and ODE was then rapidly injected into this hot solution. The Se and S precursor ratio was tuned to specific values, but the total amount was kept at 1 mmol. ODE was added to dilute the precursor Towhomcorrespondenceshouldbeaddressed.E mail:alivis@berkeley.edu. NANO LETTERS 2009 Vol. 9, No. 4 1699 1703 10.1021/nl900388a CCC: $40.75 2009 American Chemical Society Published on Web 03/12/2009 solution to 2 mL total. The nanocrystals were grown at 150 degC for 90 s, and the reaction was rapidly quenched by placing the flask in a room temperature water bath and injecting 5 mL of anhydrous hexane. The nanocrystals were purified by precipitation twice in hexane/ethanol and once in hexane/ acetone and stored in a glovebox. Achieving and characterizing a uniformly alloyed nanocrystal remains difficult.16 Complications arise from the difference in precursor solubility and reactivity at a given temperature, and in the difference between nucleation and growth of nanocrystals. Since TMS2S is more reactive than TOP/Se, the stoichiometric ratio of S to Se in the resulting nanocrystal sample was greater than the injected precursor ratio. The composition of the resulting nanocrystals was characterized using energy filtered transmission electron microscopy (EF TEM)18 to determine whether the nanocrystals resulted in separately nucleated PbS and PbSe, core shell architectures, or alloyed composites. Figure 1A C show zero loss and EF TEM images of a sample of ~7 nm PbS0.7Se0.3 taken at the same position on a TEM grid. To achieve strong elemental signals, for EF TEM we found it necessary to use large nanocrystals and exchange the oleate ligands on the nanocrystal surface by adding a small amount of butylamine and washing the nanocrystals the following day such that the nanocrystals will not self assemble into close packed arrays. All nanocrystals in Figure 1A (zero loss) appear in both the S mapping of Figure 1B and at a corresponding location in the Se mapping of Figure 1C. The selected areas in Figure 1A C make the comparison easier and are shown in greater detail in Figure 1D with the S and Se maps overlaid. The TEM results indicate that to some extent both Se and S are distributed inside each nanocrystal without apparent phase separation. Note that S is more prevalent than Se in the sample (i.e. PbS0.7Se0.3) Figure 1A inset shows a high resolution TEM image of a single nanocrystal. Uniform lattice structure with no obvious stack faults or core shell structure is observed. See Supporting Information for additional TEM images of well packed oleate capped ternary PbSSe nanocrystals. Rutherford backscattering spectroscopy (RBS) was then used to investigate the actual anion ratio. Figure 2 shows results from RBS data taken on a series of samples where the relative amount of S in the precursor ratio (S/(S+Se) was systematically varied from 0 to 1. The graph shows a clear nonlinearity in the percent incorporation of anions in the nanocrystals compared to fraction present in the original precursor solution. For example, only 30% S in the precursor is needed to make nanocrystals with 70% S composition. Presumably, this nonlinearity results from the different reactivity of the chalcogen precursors. We also find that for longer reaction times, more Se is incorporated (see Figure S5 in Supporting Information) indicating a possible radial gradient in composition; however, to be consistent, all nanocrystals used in devices were only allowed to grow for 90 s, thus suppressing such a gradient. RBS data shows that all samples display Pb rich composition regardless of whether or not Pb was in excess during synthesis. For optical characterization, the alloyed nanocrystals were suspended in tetrachloroethylene. Absorbance spectra for nanocrystals with different compositions are displayed in Figure 1. (A) Bright field TEM image showing 7 nm PbS0.7Se0.3 nanocrystals. Scale bar represents 10 nm. Inset shows the high degree of crystallinity of a single ternary nanocrystal without an obvious core shell configuration. (B) Energy filtered TEM image at the same location as in panel A, showing in red regions containing sulfur. (C) EF TEM showing selenium map in blue. (D) The outlined region in panels A C is enlarged and overlaid to show sulfur and selenium in each nanocrystal. 1700 Nano Lett. Vol. 9, No. 4, 2009 Figure 3A. Arising from the smaller bandgap of PbSe relative to PbS for a given size,19 we notice the red shift of the first excitation peak with reduced S composition. This trend can be observed more clearly in the inset of Figure 3B, which shows a linear relationship between the nanocrystal bandgap energy and the composition ratios. Vegard's Law predicts the structure and function of many alloyed materials: Ealloy EA (1 )EB, where is the mole fraction, EA, EB, and Ealloy are the band gap energy (or other properties) of pure composition A, pure composition B, and the alloyed material, respectively. However, this linear relationship does not apply to several classes of semiconductor alloys. For example, both bulk and nanocrystal CdSexTe1 x alloys display pronounced nonlinear "optical bowing" effects.16,17 Zunger and co workers explain this type of observation by identifying three structural and electronic factors leading to nonlinearity of ternary compounds: different atomic size, electronegativity values of ions, and different lattice constants of the binary structures.20,21 A substantial lattice mismatch (11% also exists between the binary semiconductors CdS and CdTe that leads to enhanced nonlinear effects there also. However, in the case of PbSxSe1 x there is only a 2% lattice mismatch between PbS and PbSe, so it is reasonable to observe less nonlinearity with composition in this alloy system, considering also that the difference in atomic size
author_names: Joseph M Luther

Haimei Zheng

Yue Wu

A Paul Alivisatos
title: Letter Photovoltaic Devices Employing Ternary PbS
year: 2009

abstract: We report on experimental results of non resonant two photon absorption induced photoluminescence in ZnO nanostructures, which may act as a possible route to excite ZnO nanostructure based lasers. Epitaxial ZnO nanorod like nanostructure was grown on pre seeded Si (100) substrates by chemical vapor deposition (CVD) method with a mixed ZnO/C solid source. Crystalline ZnO seeds were prepared and controlled by the rapid thermal annealing (RTA) treatment of e beam deposited amorphous ZnO thin films with various thicknesses. INTRODUCTION ZnO is very interesting and promising optical material because it has a room temperature direct band gap of 3.37 eV [1] and excitonic binding energy of 60 meV [2] However, it remains very challenging to make it as a truly useful optoelectronic material. One of the challenges is to epitaxially grow ZnO on a suitable substrate although thin film deposition on various substrates such as silicon or sapphire wafers using different methods has been successful [3] Another challenge is to understand and control the high level of non intentionally doped residual electron density. This not yet controllable residual electron density [4, 5] is the reason behind the inability of making ZnO a good p type material necessary optoelectronic devices. It also limits the full potential of large exciton binding energy because of carrier screening effect. Despite this drawback, there are increasing amount of works [6 8] on further exploring its optical properties of this material with hope that the material issue will be soon solved in joining the success of GaN related materials. Most of reports, however, are based on single photon excitation such as single photon pumped ZnO UV lasers. On the other hand, the none centralsymmetry of the wurtzite structural ZnO should have good nonlinear multi photon effects which have potential application in multi photon absorption fluorescence microscopy, threedimensional optical data storage, frequency up conversion luminescence and lasing, and optical limiting owing to high spatial resolution and large penetration depth of excitation light. In this work, we report on the experimental results of non resonant two photon absorption induced photoluminescence in a ZnO nanostructures from at room temperature. The ZnO nanorod like nanostructure was epitaxially grown on pre seeded Si (100) substrates by Chemical Vapor Deposition (CVD) method with a mixed ZnO powder/C powder solid source. Crystalline ZnO seeds were prepared and controlled by the rapid thermal annealing (RTA) treatment of e beam deposited amorphous ZnO thin films with various thicknesses. EXPERIMENTAL RESULTS AND DISCUSSION Amorphous films with thickness from 200 to 600 nm were deposited on Si (100) substrates by e beam evaporator and then rapid thermal annealed under N2 with various annealing temperature and time period. These films were then studied by scanning electron microscopy and photoluminescence. Fig. 1a and 1b display the seed film used to epitaxially grow nanorods reported in this work. Fig. 1a is a relatively large scale image indicating clearly domain size on order of 100 mm; while Fig. 1b is 200 nm scale SEM image showing most of crystallites have a well defined orientation of (0001) direction. Fig. 1a SEM image of rapid thermally annealed ZnO amorphous film. Fig. 1b Same as Fig. 1a but with large magnification. Rapid thermal annealed films were then characterized by photoluminescence (PL) PL spectra were measured by Horiba NanoLog system coupled with an optical cryostat (4 to 350 K) from Advanced Research System. The 9 K photoluminescence spectrum is displayed in Fig. 2 for the ZnO seeded film, RTA 18, on which nanorods were grown. From the PL intensity variation with temperature, three main PL peaks are identified as A band free exciton, XA (3.365 eV) donor bound exciton, XD (3.328 eV) and phonon replica of XA (3.255 eV) The so called green is very weak indicating very low optically active defect density in these films. A commercially built vapor deposition (CVD) system (First Nano, ET2000) was used to epitaxially grow (we call it nano epi) nanostructures on these pre seeded substrates. This system was equipped with a separated solid source heater, a three zone (load, center, and end zone) furnace, a gas injector, a vacuum pump, and a quartz tube. The reaction tube was controlled by a three zone furnace to obtain a uniform temperature across the substrate at the collecting area over 3 inches by 2 inches. The source material was mixed ZnO powder (Alfa Aesar, 99.99% and graphite powder (Alfa Aesar, 99.995% with mass ratio of 1:4. The solid source was placed at the load zone in the quartz tube and heated up to high temperature by an additional solid source heater to generate Zn vapor which was then carried into the center zone by the Ar carrying gas. The reacting gas (O2) was introduced into the system by gas injector. Fig. 2 Photoluminescence spectrum from the ZnO seeded film at 9 K. Fig. 3a (with a 1 mm scale) and 3b (200 nm scale) display the SEM images of irregular nanorods epitaxially grow on the ZnO seeds. Room temperature PL spectrum was taken from these epitaxially grown nanorods. Its peak intensity is at least 3 times stronger than that from bare seed film as shown in Fig. 4a. The normalized 9 K PL spectrum from ZnO nanorods is also displayed in Fig. 4b as compared with that from the bare ZnO seed layer. It appears that the PL originated from donor bound exciton is either much weaker or completely disappeared suggesting these ZnO nanorods epitaxially grown on the ZnO seeds might be residual donor free material. Two photon absorption induced PL was excited by 532 nm laser line for an Nd: YVO laser (8 ps pulse width, 76 MHz repetition rate) A log intensity plot of excitation dependence of PL spectra taken at room temperature is displayed in Fig. 5a. It should be noticed that PL peak, 378 nm, is at the same position as the PL spectra excited by one photon excitation. This behavior is distinctly different than the two photon absorption induced PL from thick films, in which PL peak is red shifted due to self absorption. A linear plot of the two photon induced PL spectrum at the excitation power of 488 mW is shown in Fig. 5b with negligible green emission at 550 nm, consistent to our suggestion of the high optical quality of these nanorods. Fig. 3a SEM image of ZnO nanorods epitaxially grown on ZnO seeded layer. Fig. 3b Same as Fig. 4a except with high magnification. Fig. 4a Room temperature PL spectrum of ZnO nanorods as compared to that from the bare seed layer (green curve) Fig. 4b Normalized 9 K PL spectrum of ZnO nanorods as compared to that from the bare see layer (green curve) Fig. 5a Two photon absorption induced PL spectra at different excitation power. Fig. 5b Two photon induced PL spectrum at the excitation power of 488 mW. Finally, the integrated PL intensity excited by 532 nm laser is plotted as a function excitation power in Fig. 6. It clearly shows that its variation with the excitation power has an exact slope of 2 indicating unambiguously two photon absorption process is the excitation mechanism for the measured PL excited below their band gap in these ZnO nanorods. 10 10 10 10 10 10 10 10 10 Slope 2 Exc. Wavelength: 532 nm Temperature: 300 K In te gr at ed P L (c ps n m Excitation Power (mW) Fig. 6 Integrated 2 photon absorption induced PL intensity as a function of excitation power at room temperature. CONCLUSIONS Our preliminary experimental data on two photon absorption induced photoluminescence are reported from ZnO nanorods epitaxially grown on a ZnO pre seeded film. This system might provide a platform to study nonlinear optical processes in various ZnO nanostructures. ACKNOWLEDGMENTS The authors would like to acknowledge the partial financial support from NYSTAR grant for this work through the CUNY center of advanced technology on photonic applications. REFERENCES 1. E.Mollwo, in Semiconductors: Physics of II VI and I VII Compounds, Semimagnetic Semiconductors, edited by O.Madelung, M. Schulz, and H. Weiss (Springer, Berlin, 1982) vol. 17 of Landolt B ornstein New Series, p. 35. 2. E. O'Kane, Phys. Rev. B 18, 6849 (1978) 3. X. W. Sun and H. S. Kwok, J. Appl. Phys. 86, 408 (1999) 4. U. Ozgur et al. J. Appl. Phys. 98, 043101 (2005) 5. D. L. Look, J. Electronic Materials 35, 1299 (2006) 6. C. F. Zhang et al. Appl. Phys. Lett. 87, 051920 (2005) 7. C. F. Zhang et al. Appl. Phys. Lett. 89, 042117 (2006) 8. G. P. Zhu et al. Appl. Phys. Lett. 94, 051106 (2009)
author_names: Zhuo Chen

Tom Salagaj

C Jensen

Karlheinz Strobl

Vitchanetra Hongpinyo

Boon Siew Ooi

Mim Lal Nakarmi

Kai Shum
title: Two photon absorption induced photoluminescence in a ZnO nanostructure
year: 2009

abstract: The wetting of a solid surface by a liquid depends on the surface free energy, but surface free energy does not depend on the chemistry of the surface alone. Topography can enhance hydrophobicity so that an almost complete roll up of droplets to create superhydrophobicity occurs or it can enhance hydrophilicity so that water spreads into a film; the same is true for other liquids. 1 2 In this talk I will first present examples of superhydrophobic surfaces from our laboratory and then describe how surface free energy considerations can explain their wetting behaviour. I will then describe the counter intuitive (and controversial) implication that providing the energy associated with the bending of a solid can be neglected, all solids, including Teflon, are in fact hydrophilic in an absolute sense. 3 Experiments, originating from basic ideas of adhesion, supporting this contention will be described. 4 6 I will then describe four examples of how a combination of topography and surface chemistry can be used to promote "slippy" and drag reducing properties of surfaces. In the first case, the notion that hydrophobic solids can adhere to the water air interface will be developed as a method to convert a water solid contact into a solid solid contact thereby creating freely rolling liquid marbles. 8 in the second case, I will describe how protein adsorption and bio fouling on the walls of microfluidic channels can be reduced and flow enhanced detachment can be enhanced by using hydrophobic nano scale topographic features. 9 In the third case, I will describe experiments that suggest the drag associated with the flow of simple liquids through milli metric diameter copper tubes can be reduced by creating a hydrophobic copper nano ribbon internal surface coating. 1 deg In the fourth case, I will describe a recreation of a classic experiment in which a spherical object falls through a column of water under the action of gravity. However, in this case, I will present data suggesting that when the objects possess a superhydrophobic surface and retain a layer of air when immersed, they fall faster rather than slower and so show a significant reduction in drag. 11 Acknowledgements The author acknowledges the financial support of the UK EPSRC and MODIDstI and the support of coworkers. References 1. Roach, P. Shirtcliffe, N.J. Newton, M.I. "Progress in superhydrophobic surface development" Soft Matter 2008, 4, 224 240. 2. McHale, G. Shirtcliffe, N.J. Agil, S. Perry, C.C. Newton, MA. "Topography driven spreading" Phys. Rev. Lett. 2004, 93, art. 036102. 3. McHale, G. "All solids, including Teflon, are hydrophilic, but some have roughness induced hydrophobic tendencies" Langmuir 2009, in press. 4. Gao, L. McCarthy, T.J. "Teflon is hydrophilic. Comments on definitions of hydrophobic, shear versus tensile hydrophobicity, and wettability characterization" Langmuir 2008, 24, 9183 9188. 5. Py, C. Reverdy, P. Doppler, L. Bico, J. Roman, B. Baroud, C.N. "Capillary Origami: Spontaneous wrapping of a droplet with an elastic sheet" Phys. Rev. Lett. 2007, 98, art. 156103. 6. Py, C. Reverdy, P. Doppler, L. Bico, J. Roman, B. Baroud, C.N. "Capillarity induced folding of elastic sheets" Fur. Phys. J. Special Topics 2009, 166, 67 71. 7. Aussillous, P. Quere, D. "Liquid Marbles" Nature 2001, 411, 924 927. 8. McHale, G. Herbertson, D.L. Elliott, S.J. Shirtcliffe, N.J. Newton, M.I. "Electrowetting of non wetting liquids and liquid marbles" Langmuir 2007, 23, 918 924. 9. Shirtcliffe, N.J. Koc, Y. de Mello, A.J. McHale, G. Newton, M.I. Roach, P. "Nano scale superhydrophobicity: Suppression of protein adsorption and promotion of flow induced detachment" Lab on a Chip 2008, 8, 582 586. 10. Shirtcliffe, N.J. McHale, G. Newton, M.I. Zhang, Y. "Superhydrophobic copper tubes with possible flow enhancement and drag reduction" Appl. Mater. lnterf. 2009, submitted. 11. McHale, G. Shirtcliffe, N.J. Evans, C.R. Newton, M.I. "Terminal velocity and drag reduction measurements on superhydrophobic spheres" Appl. Phys, Lett. 2009, 94, art. 064104. Development of a novel fibronectin based biomaterial: A slick solution to a sticky problem Mudera V, Branford 0, Grobbelaar A, McGrouther DA and Brown R. Tissue Repair and Engineering Centre (TREC) University College London Institute of Orthopaedics Royal National Orthopaedic Hospital Brockley Hill Stanmore, Middlesex, HA7 4LP Phone: 44 (0)208 909 5634, Fax: 44 (0)208 954 8560 email: v.muderan. ucl.ac.uk
author_names: Stuart Anson Brewer
title: Easy release and "non sticking" An overview of fundamental principles
year: 2009

abstract: In this study, we demonstrate that periods of metallic gratings on elastomeric substrates can be tuned with external strain and hence are found to control the resonance condition of surface plasmon polaritons [1] The periods of the gratings are increased up to 25% by the use of applied mechanical strain. The tunability of the elastomeric substrate provides the opportunity to use such gratings as efficient surface enhanced Raman spectroscopy substrates. It's been demonstrated that the Raman signal can be maximized by tuning the period of the elastomeric grating. The surface plasmon resonance (SPR) phenomena observed on metal surfaces or nanoparticles has been a great interest in several fields of research such as nanoscale photonics and biological sensing. Continuous metallic films possessing a periodic perturbation exhibit strong extinction and scattering spectra when excited at the SPR condition. The challenge of designing effective structures to manipulate plasmonic fields and utilize them in functional devices still remains. In particular, the use of SPR in surface enhanced Raman spectroscopy (SERS) and biological sensing require an intelligent design in order to maximize the plasmonic enhancement. In this regard, the tunability of the SPR wavelength provides flexibility in many plasmonic sensing applications. Flexible designs utilizing electronic [2] ferroelectric [3] or thermal [4] tuning mechanisms are also reported in the literature. Those methods are reversible and can be applied after the plasmonic structure is fabricated. Such a repeatable process can find wide applications in the field of Raman spectroscopy and plasmonic sensing. It was reported that by controlling the geometry of the nanoshells films, the SERS enhancements can be optimized [5] A repeatable thermal tuning mechanism using silver nanoparticles for achieving a tunable SERS substrate was reported by Lu et al. [6] In this study we use an elastomeric grating structure in order to excite surface plasmon polaritons (SPP) on its metallic surface. We report a way of tuning the SPR by applying mechanical strain on the elastomeric grating structure. The elongation of the elastomer effectively changes the period of the metallic grating. It can be seen that the SPR wavelength also shifts as the external strain changes the period of the elastomeric grating coated with a thin metallic layer. We fabricated two silicone elastomers with gratings on top using two different methods. The first elastomeric grating was generated using holographic lithography with 665 nm period. The elastomeric grating is then obtained using the replication procedure. Note that the thickness of the elastomer is kept around 5 mm. To generate SPP, the PDMS grating is coated with a 55 nm of silver using thermal evaporation. For the second elastomer we used a commercially available ruled grating with 530 nm period as the master grating. The fabrication of the elastomeric grating is the same as in the first case. For demonstration of the tunable periodicity of the elastomeric grating we measured the angle of the diffracted beam when the it is excited by a 632 nm He Ne laser for the case of 665 nm grating (A) and by a 514.5 nm Ar laser for the case of 530 nm grating (B) It's been found that the periods of the gratings increase linearly up to 25% of mechanical strain. To demonstrate the tunability of the SPR condition we used both gratings, A and B. The optical normal incidence reflection spectrum of the PDMS gratings is measured using an ellipsometer (JA Woolam VASE) As seen in Fig.1, the SPR wavelengths on silver coated gratings A and B are approximately 560 nm and 670 nm, respectively, in the absence of applied strain. As the elastomeric grating is stretched, SPR wavelength red shifts due to the increased grating periodicity. The elastomeric grating with 530 nm period is used as a SERS substrate to measure SERS signal of Rhodamine 6G (R6G) molecule. The precision mechanical strain setup is used under the objective of the spectrometer. Raman signal is maximized when the strain is 20.8\ which corresponds to a 633 nm grating period. We have demonstrated the use of elastomeric gratings with tunable surface plasmon resonance condition. We have tuned the surface plasmon resonance wavelength by applying a mechanical strain on elastomeric gratings coated with a thin layer of metal. We have shown the use of the elastomeric grating with a SERS experiment of R6G molecule. An enhancement factor of more than 10 is achieved when the grating period is optimized. Note that the presented method is compatible with Raman and Micro Raman Spectroscopy methods which utilize a fixed incident angle. We believe that the method can be used not only in SERS experiments, but also in biosensing and plasmonic enhancement applications. S.O. acknowledges the support of TUBITAK and ASELSAN A.S. for their Ph.D. Scholarship Programs. [1] S. Olcum, A. Kocabas, G. Ertas, A. Atalar, A. Aydinli, Optics Express, 2009, 17(10) 8542 8547. [2] W. Dickson, G. A.Wurtz, P. R. Evans, R. J. Pollard, and A. V. Zayats, Nano Lett. 2008 8, 281 286. [3] H. L. Chen, K. C. Hsieh, C. H. Lin, and S. H. Chen, Nanotech. 2008 19(435304) [4] G. Xu, Y. Chen, M. Tazawa, and P. Jin, J. Phys. Chem. B 2006, 110, 2051 2056 [5] J. B. Jackson and N. J. Halas, Proc. Nat. Acad. Sci. 2004, 101, 17,930 17,935. [6] Y. Lu, G. L. Liu, and L. P. Lee, Nano Lett. 2005, 5, 5 9. *Correspondance should be sent to: selim@ee.bilkent.edu.tr
author_names: Selim Olcum

Askin Kocabas

Gulay Ertas

Abdullah Atalar

Atilla Aydinli
title: Tunable Surface Plasmon Resonance on an Elastomeric Grating
year: 2009