EMC 2008 14th European Microscopy Congress 1–5 September 2008, Aachen, Germany
Sivia Richter · Alexander Schwedt Editors
EMC 2008 14th European Microscopy Congress 1–5 September 2008, Aachen, Germany Volume 2: Materials Science
13
Dr. Silvia Richter Dr. Alexander Schwedt RWTH Aachen Central Facility for Electron Microscopy Ahornstr. 55 52074 Aachen Germany
[email protected] [email protected] ISBN 978-3-540-85225-4
e-ISBN 978-3-540-85226-1
DOI 10.1007/978-3-540-85226-1 © 2008 Springer-Verlag Berlin Heidelberg This work is subject to copyright. All rights are reserved, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilm or in any other way, and storage in data banks. Duplication of this publication or parts thereof is permitted only under the provisions of the German Copyright Law of September 9, 1965, in its current version, and permission for use must always be obtained from Springer. Violations are liable to prosecution under the German Copyright Law. The use of general descriptive names, registered names, trademarks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. Typesetting: digital data supplied by the authors Production: le-tex publishing services oHG, Leipzig, Germany Cover design: eStudioCalamar S.L., F. Steinen-Broo, Girona, Spain Printed on acid-free paper 987654321 springer.com
Preface Volume 2: Materials Science With the 14th European Microscopy Congress the European Microscopy Society (EMS), the German Society for Electron Microscopy (DGE) and the local organizers continue this successful series of conferences. Since the first congress in Stockholm in 1956 (still named European Congress on Electron Miscroscopy - EUREM), this conference series has become a well-accepted platform for academic and industrial scientists not only from Europe, but from all over the world, to discuss and exchange their latest results in the field of electron and other microscopies. The congress is subdivided into the three main areas, “Instrumentation and Methods”, “Materials Science” and “Life Sciences”. This 2nd volume of the conference proceedings collects the contributions related to the application of electron microscopy to the large field of Materials Science, again structured in five symposia covering all kinds of various materials: “Materials for Information Technology”, “Nanomaterials and Catalysts”, “Structural and Functional Materials”, “Soft Matters and Polymers”, and finally “Materials in Mineralogy, Geology and Archaeology”. All in all, this volume contains more than 400 contributions to this wide field of applications. Therefore, at this point we would like to express our deepest thanks to all, who contributed to making this a successful conference: the invited speakers and chairpersons, as well as all authors of contributed papers, may they be presented as oral communication or as poster. We are sure, that by the contributions of all of them the congress will reach an excellent level of scientific quality. Last, but not least, we want to thank all, who assisted in the organization of this conference, i.e. Tobias Caumanns, Achim Herwartz, Helga Maintz, Evi Münstermann, Daesung Park, Thomas Queck, Stefanie Stadler, Sarah Wentz, and especially the staffs of the Eurogress Conference Centre and of the Aachen Congress service. We wish all of you an exciting EMC2008! And after all the days of hard work, don’t forget to enjoy the marvellous city of Aachen. Silvia Richter and Alexander Schwedt Editors, Volume 2 of the EMC 2008 proceedings
Content M
Materials Science
Direct observation of atomic defects in carbon nanotubes and fullerenes ................ 1 K. Suenaga
Atomic studies on ferroelectric oxides by aberration corrected transmission electron microscopy........................................................................................................ 3 K. Urban and C.L. Jia
M1
Materials for Information Technology
M1.1 Si-based semiconductors Dark-field electron holography for the measurement of strain in nanostructures and devices ....................................................................................... 5 M.J. Hÿtch, F. Houdellier, F. Hüe, and E. Snoeck
Some device challenges towards the 22nm CMOS technology................................... 7 F. Andrieu, T. Ernst, O. Faynot, V. Delaye, D. Lafond, and S. Deleonibus
Off-axis electron holography for the analysis of nm-scale semiconductor devices. .............................................................................. 9 D. Cooper, R. Truche, L. Clement, S. Pokrant, and A. Chabli.
Influence of the oxide thickness................................................................................... 11 P. Donnadieu, V. Chamard, M. Maret, J.P. Simon, and P. Mur
Challenges to TEM in high performance microprocessor manufacturing.............. 13 H.J. Engelmann, H. Geisler, R. Huebner, P. Potapov, D. Utess, and E. Zschech
Strain study in transistors with SiC and SiGe source and drain by STEM nano beam diffraction................................................................................. 15 P. Favia, D. Klenov, G. Eneman, P. Verheyen, M. Bauer, D. Weeks, S.G. Thomas, and H. Bender
Cluster growth and luminescence in ion-implanted silica ........................................ 17 H.-J. Fitting, R. Salh, L. Kourkoutis, and B. Schmidt
Coherence Measurements of Bulk and Surface Plasmons in Semiconductors by Diffracted Beam Holography ................................................................................. 19 R.A. Herring
Comparison of 3D potential structures at different pn-junctions in FIB-prepared silicon and germanium samples measured by electron-holographic tomography ......................................................................... 21 A. Lenk, D. Wolf, H. Lichte, and U. Mühle
II
Content
EELS analyses of metal-inserted high-k dielectric stacks......................................... 23 M. MacKenzie, A.J. Craven, D.W. McComb, C.M. McGilvery, S. McFadzean, and S. De Gendt
Low voltage SEM observations of the dopant contrast in semiconductors ............. 25 K. Masenelli-Varlot, S. Luca, G. Thollet, P.H. Jouneau, and D. Mariolle
NiSi2/Si interface chemistry and epitaxial growth mode........................................... 27 S.B. Mi, C.L. Jia, K. Urban, Q.T. Zhao, and S. Mantl
Detailed investigation of a tunnel oxide defect in a flash memory cell using TEM-tomography............................................................................................... 29 U. Muehle, M. Krause, F. Goetze, D. Wolf, and U. Gaebler
Overgrowth of the Mn4Si7 phase on/around the hexagonal SiC and cubic MnSi impurity phases in the Mn4Si7/Si films............................................ 31 A. Orekhov, T. Kamilov, and E.I. Suvorova
HR-STEM EELS analysis of silicon 32 nm technology using a TITAN with a probe Cs corrector ............................................................................................ 33 R. Pantel, J.L Rouvière, E. Gautier, S. Denorme, C. Fenouiller-Beranger, F. Boeuf, G. Bidal, and M. Cheynet.
Tomographic analysis of a FinFET structure............................................................ 35 O. Richard, P. Van Marcke, and H. Bender
STEM EELS/EDX dopant analysis of nm-scale Si devices....................................... 37 G. Servanton, R. Pantel, M. Juhel, and F. Bertin
M1.2 Compound semiconductors Electron beam induced damage: An atom-by-atom investigation with TEAM0.5 .............................................................................................................. 39 C. Kisielowski, R. Erni, and J. Meyer
The atomic structure of GaN-based quantum wells and interfaces ......................... 41 C.J. Humphreys, M.J. Galtrey, R.A. Oliver, M.J. Kappers, D. Zhu, C. McAleese, N.K. van der Laak, D.M. Graham, P. Dawson, A Cerezo, and P.H. Clifton.
Using TEM to investigate antiphase disorder in GaP films grown on Si(001)........ 43 T.B. Adams, I. Nemeth, G. Lukin, B. Kunert, W. Stolz, and K. Volz
TEM characterization of InAs/GaAs quantum dots capped by a GaSb/GaAs layer.................................................................................................. 45 A.M. Beltrán, T. Ben, A.M. Sánchez, D.L. Sales, M.F. Chisholm, M. Varela, S.J. Pennycook, P.L. Galindo, J.M. Ripalda, and S.I. Molina
The microstructure of (0001)GaN films grown by molecular beam epitaxy from a nanocolumn precursor layer ........................................................................... 47 D. Cherns, L. Meshi, I. Griffiths, S. Khongphetsak, S.V. Novikov, N. Farley, R.P. Campion, and C.T. Foxon
Content
III
Compositional and Morphological Variation in GaN/AlN/AlGaN Heterostructures......................................................................... 49 P.D. Cherns, C. McAleese, M.J. Kappers, and C.J. Humphreys
TEM/STEM/EFTEM imaging and Valence Electrons Spectroscopy analysis of Ultra Low-K dielectrics ........................................................................................... 51 M. Cheynet, S. Pokrant, M. Aimadeddine, V. Arnal, and F. Volpi
Epitaxial Orientations of GaN Grown on R-plane Sapphire.................................... 53 J. Smalc-Koziorowska, G.P. Dimitrakopulos, Ph. Komninou, Th. Kehagias, S.-L. Sahonta, G. Tsiakatouras, and A. Georgakilas
Microstructure and growth model of MBE-grown InAlN thin films....................... 55 S.-L. Sahonta, A. Adikimenakis, G.P. Dimitrakopulos, Ph. Komninou, H. Kirmse, E. Pavlidou, E. Iliopoulos, A. Georgakilas, W. Neumann, and Th. Karakostas
Silicon carbide modulated structures as a result of the introduction of 8H bands in a 4H matrix ......................................................................................... 57 N. Frangis, M. Marinova, I. Tsiaoussis, E.K. Polychroniadis, T. Robert, S. Juillaguet, and J. Camassel
HRTEM study of AlN/3C-SiC heterointerfaces grown on Si(001) and Si(211) substrates .................................................................................................. 59 T. Isshiki, K. Nishio, Y. Abe, J. Komiyama, S. Suzuki, and H. Nakanishi
Electron microscopy of GaAs/AlGaAs quantum cascade laser................................ 61 A. Łaszcz, J. Ratajczak, A. Czerwinski, K. Kosiel, J. Kubacka-Traczyk, J. Muszalski, M. Bugajski, and J. Kątcki
Solving the crystal structure of highly disordered Sn3P4 by HRTEM ..................... 63 O.I. Lebedev, A.V. Olenev, and G. Van Tendeloo
Determination of In-distribution in InGaAs quantum dots...................................... 65 D. Litvinov, H. Blank, R. Schneider, D. Gerthsen, and M. Hetterich
Nanoanalytical investigation of the dielectric gate stack for the realisation of III-V MOSFET devices............................................................................................ 67 P. Longo, A.J. Craven, M.C. Holland, and I.G. Thayne
Phase mapping of uncapped InN quantum dots........................................................ 69 J.G. Lozano, M. Herrera, R. García, N.D. Browning, S. Ruffenach, O. Briot, and D. González
STEM investigations of (In,Ga)N/GaN quantum wells............................................. 71 P. Manolaki, I. Häusler, H. Kirmse, W. Neumann, P. Vennéguès, P. De Mierry, and P. Demolon
Defects in m-plane GaN layers grown on (100) γ-LiAlO2 ......................................... 73 A. Mogilatenko, W. Neumann, T. Wernicke, E. Richter, M. Weyers, B. Velickov, and R. Uecker
IV
Content
Improvements on InP epilayers by the use of monoatomic hydrogen during epitaxial growth and successive annealing ................................................................. 75 F.M. Morales, A. Aouni, R. García, P.A. Postigo, C.G. Fonstad, and S.I. Molina
Study of microstructure and strain relaxation on thin InXGa1-xN epilayers with medium and high In contents.............................................................................. 77 F.M. Morales, J.G. Lozano, R. García, V. Lebedev, S. Hauguth-Frank, V. Cimalla, O. Ambacher, and D. González
Convergence of microscopy techniques for nanoscale structural characterization: an illustration with the study of AlInN......................................... 79 A. Mouti, S. Hasanovic, M. Cantoni, E. Feltin, N. Grandjean, and P. Stadelmann
TEM analyses of microstructure and composition of AlxGa1-xN/GaN distributed Bragg reflectors ........................................................................................ 81 A. Pretorius, A. Rosenauer, T. Aschenbrenner, H. Dartsch, S. Figge, and D. Hommel
TEM study of quaternary InGaAsN/GaAs quantum well structures grown by molecular beam epitaxy.......................................................................................... 83 T. Remmele, M. Albrecht, I. Häusler, L. Geelhaar, H. Riechert, H. Abu-Farsakh, and J. Neugebauer
Lattice polarity and capping of GaN dots studied by Z-contrast imaging .............. 85 J.L. Rouviere, C. Bougerol, J. Coraux, B. Amstatt, E. Bellet-Amalric, and B. Daudin
Investigation of swift ions damage in wide band gap wurtzite semiconductors...... 87 S. Mansouri, I. Monnet, H. Lebius, G. Nouet, and P. Ruterana
A TEM analysis of the damage formation in thin GaN and AlN layers during rare earth ion implantation at medium range energy .................................. 89 F. Gloux, M.P Chauvat, and P. Ruterana
Characterization and modelling of semiconductor quantum nanostructures grown by droplet epitaxy ............................................................................................. 91 D.L. Sales, J.C. Hernandez, P.A. Midgley, A.M. Beltran, A.M. Sanchez, T. Ben, P. Alonso-González, Y. Gonzalez, L. Gonzalez, and S.I. Molina
Transmission Electron Microscopy Investigation of Self-Organized InN Nano-columns......................................................................... 93 H. Schuhmann, C. Denker, T. Niermann, J. Malindretos, A. Rizzi, and M. Seibt
Investigations on a dilute magnetic semicondutor (Ga1-xMnxAs) by conventional TEM and EELS ................................................................................ 95 M. Soda, U. Wurstbauer, M. Hirmer, W. Wegscheider, and J. Zweck
About the determination of optical properties using fast electrons ......................... 97 M. Stöger-Pollach
Content
V
M1.3 Data storage/ non-volatile memories Mapping uncompensated spins in exchange-biased systems by high resolution and quantitative magnetic force microscopy.............................. 99 H. J. Hug, M. Marioni, S. Romer, I. Schmid, and S. Romer
Ferroelectric materials and structures suitable for data storage: The role of microscopies in establishing preparation-microstructure-property relations ...................................................... 101 D. Hesse, M. Alexe, K. Boldyreva, H. Han, W. Lee, A. Lotnyk, B.J. Rodriguez, S. Senz, I. Vrejoiu, and N.D. Zakharov
Electronic structures at Magnetic Tunnel Junction interfaces: EELS experiments and FEFF calculations .............................................................. 103 K. March, D. Imhoff, G. Krill, and C. Colliex
Stability and reaction of magnetic sensor materials studied by atom probe tomography ....................................................................................... 105 G. Schmitz, C. Ene, H. Galinski, and V. Vovk
Transmission Electron Microscopy Analysis of Tunnel Magneto Resistance Elements with Amorphous CoFeB Electrodes and MgO Barrier .......................... 107 Michael Seibt, Gerrit Eilers, Marvin Walter, Kai Ubben, Karsten Thiel, Volker Drewello, Andy Thomas, Günter Reiss, and Markus Münzenberg
Study of the intermixing of Fe–Pt multilayers by analytical and high-resolution transmission electron microscopy........................................... 109 W. Sigle, T. Kaiser, D. Goll, N.H. Goo, V. Srot, P.A. van Aken, E. Detemple, and W. Jäger
Exploring structural dependence of magnetic properties in FePt nanoparticle by Cs-corrected HRTEM ....................................................... 111 Z.L. Zhang, J. Biskupek, U. Kaiser, U. Wiedwald, L. Han, and P. Ziemann
M1.4 Nanotubes, nanowires and molecular devices Understanding the Chemistry of Molecules in Nanotubes by Transmission Electron Microscopy .................................................................................................. 113 A.N. Khlobystov, M.W. Fay, and P.D. Brown
Electrical and mechanical property studies on individual low-dimensional inorganic nanostructures in HRTEM....................................................................... 115 D. Golberg, P.M.F.J. Costa, M. Mitome, Y. Bando, and X.D. Bai
Atomic structure of SW-CNTs: correlation with their growth mechanism and other electron diffraction studies....................................................................... 117 R. Arenal, M.F. Fiawoo, R. Fleurier, M. Picher, V. Jourdain, A.M. Bonnot, and A. Loiseau
VI
Content
TEM investigation ofSe nanostructures in/on Acetobacter xylinum cellulose gel-film ......................................................................................................... 119 N. Arkharova, V.V. Klechkovskaya, and E. Suvorova
In-situ electron irradiation studies of metal-carbon nanostructures ..................... 121 L. Sun, Y. Gan, J.A. Rodriguez-Manzo, M. Terrones, A.V. Krasheninnikov, and F. Banhart
Application of 80kV Cs-corrected TEM for nanocarbon materials ...................... 123 A. Chuvilin, U. Kaiser, D. Obergfell, A. Khlobystov, and S. Roth
Control of gold surface diffusion on Si nanowires................................................... 125 M.I. den Hertog, J.-L. Rouviere, F. Dhalluin, P.J. Desré, P. Gentile, P. Ferret, F. Oehler, and T. Baron
Nanowires of Semiconducting Metal-oxides and their Functional Properties...... 127 M. Ferroni, C. Baratto, E. Comini, G. Faglia, L. Ortolani, V. Morandi, S. Todros, A. Vomiero, and G. Sberveglieri
Phase relations in the Fe–Bi–O system under hydrothermal conditions............... 129 A. Gajović, S. Šturm, B. Jančar, and M. Čeh
Dose dependent crystallographic structure of InAs nanowires.............................. 131 F. Gramm, E. Müller, I. Shorubalko, R. Leturcq, A. Pfund, R. Wepf, and K. Ensslin
HRTEM simulations of planar defects in ZnTe nanowires .................................... 133 I. Häusler, H. Kirmse, W. Neumann, S. Kret, P. Dłużewski, E. Janik, G. Kraczewski, and T. Wojtowicz
A universal method for determination of helicities present in unidirectional groupings of graphitic or graphitic-like tubular structures ................................... 135 H. Jiang, D.P. Brown, A.G. Nasibulin, and E.I. Kauppinen
Microstructure of (112) GaAs nanorods grown by MBE ....................................... 137 E. Johnson, S.A. Jensen, L.P. Hansen, C.B. Sørensen, and J. Nygård
Structural characterization of ZnO nanorods grown on sapphire substrate by MOCVD ................................................................................................................. 139 P.-H. Jouneau, M. Rosina, G. Perillat, P. Ferret, and G. Feuillet
Nucleation of Metal Clusters on Carbon Nanotubes............................................... 141 X. Ke, A. Felten, D. Liang, S. Bals, J.J. Pireaux, J. Ghijsen, W. Drube, M. Hecq, C. Bittencourt, and G. Van Tendeloo
EDX and linescan modelling for core/shell GaN/AlGaN nanowire analysis ......... 143 L. Lari, R.T. Murray, T. Bullough, P.R. Chalker, C. Chèze, L. Geelhaar, and H. Riechert
Mo6S9-xIx nanowires: structure studies by aberration corrected high resolution TEM and STEM ....................................................................................... 145 V. Nicolosi, J.N. Coleman, D. Mihailovic, and P. Nellist
Discrete Atom Imaging in Carbon Nanotubes and Peapods Using Cs-Corrected TEM Operated at 100keV.................................................................. 147 Luca Ortolani, Florent Houdellier, and Marc Monthioux
Content
VII
Extended Defects in Semiconductor Nanowires ...................................................... 149 Peter Pongratz, Youn-Joo Hyun, Alois Lugstein, Aaron Andrews, and Emmerich Bertagnolli
Surface chemistry along a single silicon nanowire: Quantitative x-ray photoelectron emission microscopy (XPEEM) of the metal catalyst diffusion ..... 151 O. Renault, A. Bailly, P. Gentile, N. Pauc, T. Baron, L.–F. Zagonel, and N. Barrett
TEM characterization of metallic Ni nanoshells grown on gold nanorods and on carbon nanotubes........................................................................................... 153 J.B. Rodríguez-González, M. Grzelczak, M.A. Correa-Duarte, J. Pérez-Juste, and L.M. Liz-Marzán
Electron Irradiation Effects in Carbon Nanostructures: Surface Reconstruction, Extreme Compression, Nanotube Growth and Morphology Manipulation ................................................................................. 155 M. Terrones, L. Sun, J.A. Rodriguez-Manzo, H. Terrones, and F. Banhart
Crystallographic phase and orientation analysis of GaAs nanowires by ESEM, EDS, TEM, HRTEM and SAED............................................................. 157 A.M. Tonejc, S. Gradečak, A. Tonejc, M. Bijelić, H. Posilović,V. Bermanec, and M. Tambe
3-Dimensional Morphology of GaP-GaAs nanowires ............................................. 159 M.A. Verheijen, R. Algra,M.T. Borgström, G. Immink, E. Sourty, L.F. Feiner, W.J.P. van Enckevort, E. Vlieg, and E.P.A.M. Bakkers
Characteristics of Indium-Catalyzed Si Nanowires ................................................ 161 Z.W. Wang, Z.Y. Li, and F. Iacopi
M2
Nanomaterials and Catalysts
M2.1 Carbon-based HRTEM contribution to the study of extraterrestrial nanocarbons and some earth materials analogues ......................................................................... 163 J.N. Rouzaud and C. Le Guillou
Time resolved in-situ TEM observations of Carbon Nanotube growth................. 165 J. Robertson, S. Hofmann, R. Sharma, C. Ducati, and R. Dunin-Borkowski
Insulator-Metal transition: formation of Diamond Nanowires in n-type Conductive UNCD films ............................................................................ 167 R. Arenal, O. Stephan, P. Bruno, and D.M. Gruen
Field emission from iron-filled carbon nanotubes observed in-situ in the scanning electron microscope ......................................................................... 169 K.J. Briston, Y. Peng, N. Grobert, A.G. Cullis, and B.J. Inkson
Templated ordering of fullerenes on nanostructured oxide surfaces .................... 171 D.S. Deak, B.C. Russell, D.T. Newell, K. Porfyrakis, F. Silly, and M.R. Castell
VIII
Content
Carbon nanostructures produced by chlorination of Cr3C2 and Cr(acac)3 .......... 173 A. Gómez-Herrero, E. Urones-Garrote, D. Ávila-Brande, N.A. Katcho, E. Lomba, A.R. Landa-Cánovas, and L.C. Otero-Díaz
Structural peculiarities of carbon onions, formed by different methods .............. 175 B.A. Kulnitskiy, I.A. Perezhogin, and V.D. Blank
Electron Energy Loss Spectroscopy of La@C82 peapods........................................ 177 R.J. Nicholls, D.A. Eustace, D. McComb, G.A.D. Briggs, D.J.H. Cockayne, and D.G. Pettifor
HRTEM studies of Y-junction bamboo-like CN-nanotubes................................... 179 I.A. Perezhogin, B.A. Kulnitskiy, V.D. Blank, D.V. Batov, and E.V. Polyakov
EF-TEM observation of biological tissue for risk assessment of fullerene nanoparticles .......................................................................................... 181 K. Yamamoto, M. Makino, E. Kobayashi, and Y. Morimoto
M2.2 Nanoparticles and catalysts Looking at the surface of catalysts nanopowders .................................................... 183 J.C. Hernandez, A.B. Hungria, M. Lopez-Haro, J.A. Perez-Omil, S. Trasobares, S. Bernal, P. Midgley, O. Stephan, and J.J. Calvino
Gathering structural and analytical information on catalysts at sub-nanometer level with TEM............................................................................. 185 F.J. Cadete Santos Aires and M. Aouine
Size effect and influence of nanoparticles thickness on order/disorder phenomena in CoPt nanoparticles ............................................. 187 D. Alloyeau, C. Ricolleau, T. Oikawa, C. Langlois, Y. Le Bouar, and A. Loiseau
In situ L10 ordering of FePt nanoparticles ............................................................... 189 P. Bayle-Guillemaud, M. Delalande, V. Monnier, Y. Samson, and P. Reiss
Characterization of indium doped zinc oxide nanorods ......................................... 191 H. Burghardt, H. Schmid, and W. Mader
Adsorbate-induced restructuring on Pt nanoparticles studied by environmental transmission electron microscopy .............................................. 193 M. Cabié, S. Giorgio, and C.R. Henry
EELS in monochromated and Cs probe corrected TEM: ...................................... 195 M. Cheynet, S. Pokrant, and S. Ersen.
Atomic-resolution Electron Microscopy at Ambient Pressure............................... 197 J.F. Creemer, S. Helveg, A.M. Molenbroek, P.M. Sarro, and H.W. Zandbergen
Development of a system for TEM/STEM investigation of air-sensitive materials: Preliminary results on CeO2 reduction behaviour ................................ 199
J.J. Delgado, M. López-Haro, J.D. López-Castro, J.A. Pérez-Omil, S. Trasobares, and J.J. Calvino
Content
IX
Characterization of two new zeolites by combining Electron Microscopy and X-Ray Powder Diffraction analyses .................................................................. 201 E. Di Paola, E. Montanari, S. Zanardi, and A. Carati
Electron beam-induced effects on copper nanoparticles: coarsening and generation of twins........................................................................... 203 D. Díaz-Droguett, V. Fuenzalida, and G. Solórzano
Role of the catalyst and substrate in nucleation and growth of Single Wall Carbon Nanotubes in HFCVD ......................................................... 205 M.-F. Fiawoo, N. Brun, A.-M Bonnot, O. Stephan, J. Thibaultand, and A. Loiseau
PEMFC degradation phenomena studied by electron microscopy........................ 207 L. Guetaz, B. Vion-Dury, and S. Escribano
TEM investigation of magnetite nanoparticles for biomedical applications ......... 209 S. Gustafsson, A. Fornara, F. Ye, K. Petersson, C. Johansson, M. Muhammed, and E. Olsson
Catalytic soot oxidation studied by Environmental Transmission Electron Microscopy .................................................................................................. 211 S.B. Simonsen, S. Dahl, E. Johnson, and S. Helveg
Surface and interface structure of ceria supported ruthenium.............................. 213 J.C. Hernandez, S. Trasobares, J.M. Gatica, D.M. Vidal,M.A. Cauqui, J.J. Calvino, A.B. Hungria, and J.A. Perez-Omil
Characterisation of materials with applications in the photocatalytic activation of water .................................................................. 215 N.S. Hondow, R. Brydson, Y.H. Chou, and R.E. Douthwaite
Complementary EM study on highly active nanodendritic Raney-type Ni catalysts with hierarchical build-up.......................................................................... 217 U. Hörmann, U. Kaiser, N. Adkins, R. Wunderlich, A. Minkow, H. Fecht, H. Schils, T. Scherer, and H. Blumtritt
Structural properties of sol-gel synthesized Li+-doped titania nanowhisker arrays.................................................................................................... 219 U. Hörmann, J. Geserick, S. Selve, U. Kaiser, and N. Hüsing
Quantitative strain determination in nanoparticles using aberration-corrected HREM........................................................................... 221 C.L. Johnson, E. Snoeck, M. Ezcurdia, B. Rodríguez-González, I. Pastoriza-Santos, L.M. Liz-Marzán, and M.J. Hÿtch
Morphological characterization by HRTEM and STEM of Fe3O4 hollow nano-spheres.................................................................................... 223 A. Ibarra, G.F. Goya, J. Arbiol, E. Jr. Lima, H. Rechenberg, J. Vargas, R. Zysler, and M.R. Ibarra
Direct observation of surface oxidation of Rh nanoparticles on (001) MgO......... 225 N.Y. Jin-Phillipp, P. Nolte, A. Stierle, P.A. van Aken, and H. Dosch
X
Content
Characterization of catalyst poisoning in biodiesel and conventional diesel fuelled vehicles ................................................................... 227 T. Kanerva, K. Kallinen, Toni Kinnunen, M. Vippola, and T. Lepistö
TEM Characterisation of Highly Luminescent CdS Nanocrystals ........................ 229 H. Katz, A. Izgorodin, D.R. MacFarlane, and J. Etheridge
Structure and composition of dilute Co-doped BaTiO3 nanoparticles .................. 231 O.I. Lebedev, R. Erni, and G. Van Tendeloo
CoxFe3-xO4 catalytic materials for gaz sensors ......................................................... 233 L. Ajroudi, A. Essoumhi, S. Villain, V. Madigou, N. Mliki, and Ch. Leroux
(S)TEM investigation on the role of alumina dopants to prevent redox activity decay at high temperature in CePrOx /doped-Al2O3 catalysts .................. 235 M. López-Haro, K. Aboussaid, J.M. Pintado, J.J. Calvino, and S. Trasobares
Sulfated Zirconia Catalysts: Structure and Performance Relationship, a TEM Study............................................................................................................... 237 C. Meyer, D. Su, N. Hensel, F.C. Jentoft, and R. Schlögl
A novel procedure for an accurate estimation of the lattice parameter of supported metal nanoparticles from the analysis of plan view HREM images ....................................................................................... 239 C. Mira, J.A. Perez-Omil, J.J. Calvino, and S. Bernal
Microstructure of Pt particles and aggregates deposited on different carbon materials for fuel cells application ............................................................................ 241 D. Mirabile Gattia, E. Piscopiello, M. Vittori Antisari, S. Bellitto, S. Licoccia, E. Traversa, L. Giorgi, R. Marazzi, and A. Montone
Low-loss-energy EFTEM imaging of triangular silver nanoparticles ................... 243 J. Nelayah, L. Gu, W. Sigle, C.T. Koch, L. Pastoriza-Santos, L.M. Liz-Marzan, and P.A. van Aken
Microstructure of cobalt nanocluster arrays fabricated by solid-state dewetting.............................................................................................. 245 Y.-J. Oh, J. Kim, S. Hwang, C.A. Ross, and C.V. Thompson
Size Effect in Gold Nanoparticles Investigated by Electron Holography and STEM ................................................................................................................... 247 L. Ortolani, V. Morandi, and M. Ferroni
Post-Mortem investigation of Fischer Tropsch catalysts using cryo- transmission electron microscopy ......................................................... 249 D. Ozkaya, M. Lok, J. Casci, and P. Ash
TEM Investigations on Cu-impregnated Zeolite Y catalysts via chloride free preparation..................................................................................... 251 M.-M. Pohl, M. Richter, and M. Schneider
Content
XI
Coarsening of mass-selected Au clusters on amorphous carbon at room temperature .................................................................................................. 253 R. Popescu, R. Schneider, D. Gerthsen, A. Böttcher, D. Löffler, and P. Weiss
TEM investigations on Ni clusters electrodeposited on Carbon substrate............ 255 M. Re, M.F. De Riccardis, D. Carbone, D. Wall, and M. Vittori Antisari
Near-surface structure of FePt nanoparticles.......................................................... 257 B. Rellinghaus, D. Pohl, E. Mohn, and L. Schultz
Overgrowth of gold nanorods: From rods to octahedrons ..................................... 259 J.B. Rodríguez-González, E. Carbó-Argibay, I. Pastoriza-Santos, J. Pérez-Juste, and L.M. Liz-Marzán
Reactive Diffusion under Laplace Tension in Spherical Nanostructures.............. 261 C. Ene, C. Nowak, and G. Schmitz
Preparation and characterization of palladium nanoparticles with various size distributions................................................................................... 263 M. Slouf, H. Vlkova, and D. Kralova
Electron microscopy for the characterization of nanoparticles ............................. 265 D. Sommer and U. Golla-Schindler
Titanium dioxide nanoparticles prepared from TiOSO4 aqueous solutions ......... 267 J. Šubrt, J. Boháček, N. Murafa, and L. Szatmáry
Exploring nanoscale ferroelectricity in isolated and interacting colloidal ferroelectric nanocrystals using electron holography ............................................. 269 D. Szwarcman, Y. Lereah, G. Markovich, M. Linck, and H. Lichte
STEM investigation on the one-pot synthesis of nanostructured CexZr1-xO2-BaO·nAl2O3 catalytic materials ............................................................. 271 J.C. Hernandez, J.A. Perez-Omil, J.J. Calvino, S. Bernal, R. di Monte, S. Desinan, J. Kašpar, and S. Trasobares
Enhanced stability against oxidation due to 2D self-organisation of hcp cobalt nanocrystals ......................................................................................... 273 Isabelle Lisiecki, S. Turner, S. Bals, M.P. Pileni, and G. Van Tendeloo
Loaded porous Zn4O(bdc)3 (metal@MOF-5) frameworks characterised by TEM ....................................................................................................................... 275 S. Turner, O.I. Lebedev, F. Schröder, R.A. Fischer, and G. Van Tendeloo
Growth behaviour of sub-nm sized focused electron beam induced deposits ....... 277 W.F. van Dorp, C.W. Hagen, P.A. Crozier, P. Kruit, S. Zalkind, B. Yakshinskiy, and T.E. Madey
Ruthenium deposition on CO2-treated and untreated carbon black investigated by electron tomography........................................................................ 279 M. Wollgarten, R. Grothausmann, P. Bogdanoff, G. Zehl, I. Dorbandt, S. Fiechter, and J. Banhart
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Content
Size-dependent crystallinity and relative orientations of nano-Pt/γ-Al2O3 ............ 281 J.C. Yang, L. Li, S. Sanchez, J.H. Kong, Q. Wang, L.L. Wang, Z. Zhang, D.D. Johnson, A.I. Frenkel, and R.J. Nuzzo
Formation of nanometer-sized porous GaSb particles by vacancy clustering induced by electronic excitation ................................................................................ 283 H. Yasuda, A. Tanaka, N. Nitta, K. Matsumoto, and H. Mori
Structural investigations of membrane electrode assemblies in fuel cells via environmental scanning electron microscopy.................................................... 285 S. Zils, N. Benker, and C. Roth
M2.3 Nanostructured materials and Nanolab In situ TEM nanocompression testing ...................................................................... 287 A.M. Minor, J. Ye, and R.K. Mishra
Physical measurements on an individual nanostructure in a TEM nanolaboratory.......................................................................................... 289 M. Kobylko, S. Mazzucco, R. Bernard, M. Kociak, and C. Colliex
TEM study of nanostructured BZO templates in (001)-LAO and (001)-STO substrates for the growth of superconducting YBCO films.................................... 291 P. Abellan, M. Gibert, F. Sandiumenge, M.J. Casanove, T. Puig, and X. Obradors
Hydrothermal synthesis and characterisation of single crystal α-Fe2O3 nanorods ....................................................................................................... 293 T. Almeida, Y.Q. Zhu, and P.D. Brown
GaAs NWs and Related Quantum Heterostructures Grown by Ga-Assisted Molecular Beam Epitaxy: Structural and Analytical Characterization................ 295 J. Arbiol, S. Estradé, F. Peiró, J.R. Morante, C. Colombo, D. Spirkoska, G. Abstreiter, and A. Fontcuberta i Morral
A method for in-situ electrical measurements of thin film heterostructures using TEM and SEM.................................................................................................. 297 J. Börjesson, A. Kalabukhov, K. Svensson, and E. Olsson
Electron Beam Nanofabrication and Characterization of Iron Compounds ........ 299 K. Furuya, M. Shimojo, M. Takeguchi, M. Song, K. Mitsuishi, and M. Tanaka
TEM analysis of the chemical gradient in (Zn,Mn)Te/ZnTe nanowires .............. 301 H. Kirmse, W. Neumann, S. Kret, P. Dłużewski, E. Janik, W. Zaleszczyk, A. Presz, G. Karczewski, and T. Wojtowicz
Structural and morphological characterization of GaN/AlGaN quantum dots by transmission electron microscopy........................................................................ 303 M. Korytov, M. Benaissa, J. Brault, T. Huault, and P. Vennéguès
Structure and stability of core-shell AuAg nanopartciels....................................... 305 Z.Y. Li, R. Merrifield, Y. Feng, J.P. Wilcoxon, R.E. Palmer, A.L. Bleloch, M. Gass, and K. Sader
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XIII
In-situ studies on electrical and mass transport in multi-wall carbon and vanadium oxide nanotubes................................................................................. 307 M. Löffler, T. Gemming, R. Klingeler, and B. Büchner
Electron microscopy of nano-magnesium produced by Inert Gas Condensation for hydrogen storage ................................................... 309 E. Piscopiello, E. Bonetti, E. Callini, L. Pasquini, and M. Vittori Antisari
Two Different Structures of Crystalline Mesoporous Indium Oxide Obtained by Nanocasting Process.............................................................................................. 311 E. Rossinyol, E. Pellicer, M. Cabo, O. Castell, and M.D. Baro
Measuring electrical properties of carbon nanotubes using liquid metal immersion, an in situ scanning electron microscopy study..................................... 313 H. Strand, K. Svensson, and E. Olsson
TEM characterization of biogenic metal nanoparticles .......................................... 315 E.I. Suvorova, P.A. Buffat, H. Veeramani, J. Sharp, E. Schofield, J. Bargar, and R. Bernier-Latmani
On the structure of VxOy supported on multiwalled carbon nanotubes................ 317 D. Wang, J.-P. Tessonnier, M. Willinger, C. Hess, D.S. Su, and R. Schlögl
Hexahedral nano-cementites catalysing the growth of carbon nanohelices .......... 319 J.H. Xia, X. Jiang, C.L. Jia, and C. Dong
M2.4 Thin films and interfaces Investigation of organic/inorganic interfaces using nano-analytical transmission electron microscopy ............................................................................. 321 V. Jantou, M.A. Horton, and D.W. McComb
Cationic ordering and interface effects in superlattices and nanostructured materials ................................................................................... 323 P. Boullay, W.C. Sheets, W. Prellier, E.-L. Rautama, A.K. Kundu, V. Caignaert, B. Mercey, and B. Raveau
Strain in SrTiO3 layers embedded in a scandate/titanate multilayer system........ 325 D. Ávila, M. Boese, T. Heeg, J. Schubert, and M. Luysberg
Anisotropic growth of CGO islands on the (001)-LaAlO3 surface ........................ 327 A. Benedetti, M. Gibert, F. Sandiumenge, T. Puig, and X. Obradors
Diffraction contrast imaging and high resolution transmission electron microscopy of multiferroic thin films and heterostructures................................... 329 B.I. Birajdar, I. Vrejoiu, X.S. Gao, B.J. Rodriguez, M. Alexe, and D. Hesse
Imaging of compositional defects at silicide-silicon interfaces using aberration corrected HAADF ...................................................................................................... 331 M. Falke, U. Falke, P. Wang, and A. Bleloch
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Content
Characterization of nanometric oxide particles extracted from a steel surface onto a carbon replica ............................................................... 333 P. Haghi-Ashtiani, A. Ollivier, and M.-L. Giorgi
TEM characterization of textured silicon heterojunction solar cells..................... 335 A. Hessler-Wyser, C. Monachon, S. Olibet, and C. Ballif
Investigation of the change in the microstructure of thin p-type Bi-Sb-Te thermoelectric films after heat treatment ................................................................ 337 F. Heyroth, M. Schade, K. Rothe, H.S. Leipner, and M. Stordeur
EM study on forming Inorganic film with Periodically Organized Mesopores upon Polymer film...................................................................................................... 339 H. Wenqing, Z. Ying, Y. Fang, Zhaoxi, and Y. Wantai
Nanointerface analysis of hard coatings deposited by IBAD.................................. 341 D. Kakas, B. Skoric, A. Miletic, and L. Kovacevic
Transrotational crystals growing in amorphous Cu-Te film.................................. 343 V.Yu. Kolosov, A.V. Kozhin, L.M. Veretennikov, and C.L. Schwamm
TEM investigation of sputtered indium oxide layers on silicon substrate for gas sensors............................................................................................................. 345 Th. Kups, I. Hotovy, and L. Spieß
Microstructure of Sr4Ru2O9 thin films and Bi3.25La0.75Ti3O12/Sr4Ru2O9 bilayers................................................................... 347 R. Chmielowski, V. Madigou, M. Blicharski, and Ch. Leroux
Analysis of the LSM/YSZ interface on micro- and nano-scale by SEM, FIB/SEM and (S)TEM ............................................................................................... 349 Y. Liu, L. Theil Kuhn, and J.R. Bowen
ESI and HRTEM of chemical solution deposited (CSD) ........................................ 351 L. Molina, T. Thersleff, B. Rellinghaus, B. Holzapfel, and O. Eibl
CTEM diffraction contrast of biaxially-textured La2Zr2O7 buffer layers on nickel substrates .................................................................................................... 353 L. Molina, S. Engel, B. Holzapfel, and O. Eibl
TEM sample preparation of YBCO-coated conductors: conventional method and FIB........................................................................................................................ 355 L. Molina, T. Thersleff, C. Mickel, S. Menzel, B. Holzapfel, and O. Eibl
Nucleation and evolution of biepitaxial YBa2Cu3O7-δ thin film grown on SrTiO3 and MgO substrates ................................................................................. 357 H. Pettersson, K. Cedergren, D. Gustafsson, R. Ciancio, F. Lombardi, and E. Olsson
An investigation of Al-Pb interfaces using probe-corrected high-resolution STEM................................................................................................ 359 H. Rösner, S. Lopatin, B. Freitag, and G. Wilde
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Spectrometric Full-Color Cathodoluminescence Electron Microscopy Study of Grain Boundaries of ZnO Varistor ...................................................................... 361 H. Saijo, N. Daneu, A. Recnik, and M. Shiojiri
Study of structural properties of Mo/CuInS2/ZnS used in solar cells by TEM..... 363 J. Sandino, G. Gordillo, and H. Lichte
Texture analysis of silicide thin films: combining statistical and microscopical information.................................................................................. 365 H. Schletter, S. Schulze, M. Hietschold, K. De Keyser, C. Detavernier, G. Beddies, A. Bleloch, and M. Falke
Statistical Tomography of 3D Thin Film Structure using Transmission Electron Microscopy .................................................................................................. 367 E. Spiecker, V. Radmilovic, and U. Dahmen
Analytical TEM investigations of Pt/YSZ interfaces............................................... 369 V. Srot, M. Watanabe, C. Scheu, P.A. van Aken, E. Mutoro, J. Janek, and M. Rühle
Microstructure and self-organization of nano-engineered artificial pinning centers in YBa2Cu3O7-x coated conductors................................................. 371 T. Thersleff, E. Backen, S. Engel, C. Mickel, L. Molina-Luna, O. Eibl, B. Rellinghaus, L. Schultz, and B. Holzapfel
The determination of the interface structure between ionocovalent compounds: the general case study of the Al2O3-ZrO2 large misfit system........... 373 G. Trolliard, and D. Mercurio
Simple method to improve quantification accuracy of energy-dispersive X-ray spectroscopy in an analytical transmission electron microscope by specimen tilting...................................................................................................... 375 T. Walther
Comparison of transmission electron microscopy methods to measure layer thicknesses to sub-monolayer precision.................................................................... 377 T. Walther
Determination of interface structure of YBCO/LCMO by a spherical aberration- corrected HRTEM......................................................... 379 Z.L. Zhang, U. Kaiser, S. Soltan, and H.-U. Habermeier
HREM characterization of BST-MgO interface...................................................... 381 O.M. Zhigalina, A.N. Kuskova, A.L. Chuvilin, V.M. Mukhortov, Yu.I. Golovko, and U. Kaiser
M3
Structural and Functional Materials
M3.1 Alloys and Intermetallics TEM investigations on novel shape memory systems with Ni-depletions ............. 383 D. Schryvers, R. Delville, B. Bartova, and H. Tian
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Content
Crystalline-to-amorphous transformation in intermetallic compounds by severe plastic deformation .................................................................................... 385 K. Tsuchiya, T. Waitz, T. Hara, H.P. Karnthaler, Y. Todaka, and M. Umemoto
EELS quantification of complex nitrides in a 12 % Cr steel .................................. 387 M. Albu, F. Méndez Martin, and G. Kothleitner
Formation of ordered solid solution during phase separation in Cu-Ag alloy films.................................................................................................... 389 F. Misják, P.B. Barna, and G. Radnóczi
Precipitates and magnetic domains in an annealed Co38Ni33Al29 shape memory alloy studied by TEM................................................................................................. 391 B. Bartova, D. Schryvers, N. Wiese, and J.N. Chapman
On the gallium accumulation at the boundaries of Al layers in FIB prepared TEM specimens.......................................................................................... 393 P. Favia and H. Bender
Precipitation in an Al-Mg-Ge Alloy.......................................................................... 395 R. Bjorge, C.D. Marioara, S.J. Andersen, and R. Holmestad
Interaction between dislocations and oxide precipitates in an aluminium containing ferritic stainless steel ............................................................................... 397 L. Boulanger, S. Poissonnet, and F. Legendre
Voids Associated with Nano-Particles of Tin in Aluminium .................................. 399 L. Bourgeois, M. Weyland, and B.C. Muddle
Influence of thermal treatments in microstructure and recrystallization peak energy of P/M Al-Mg-X alloys................................................................................... 401 S.J. Buso, A. Almeida Filho, I.M. Espósito, J.R. Matos, and W.A. Monteiro
Analytical TEM of Nb3Sn Multifilament Superconductor Wires .......................... 403 M. Cantoni, V. Abächerli, D. Uglietti, B. Seeber, and R. Flükiger
3D Reconstruction of Ni4Ti3 Precipitates in Ni-Ti by FIB/SEM Slice-and-View...................................................................................... 405 S. Cao, W. Tirry, W. Van Den Broek, and D. Schryvers
Electron microscopy study of Mg78.5Pd21.5: aphase with nanothin 120° rotational twin domains..................................................................................... 407 W. Carrillo-Cabrera, J.P.A. Makongo, Yu. Prots, and G. Kreiner
Analysing small precipitates in a ferritic steel matrix............................................. 409 A.J. Craven and M. MacKenzie
Failure analysis of first stage land-based gas turbine blades.................................. 411 F. Delabrouille, F. Arnoldi, L. Legras, and C. Cossange
TEM investigation of microstructures in low-hysteresis Ti50Ni50-xPdx alloys with special lattice parameters .................................................................................. 413 R. Delville, D. Schryvers, Z. Zhang, S. Kasinathan, and R.D. James
Content
XVII
Evidence of silica layer at the interface between ferrite and the chromium oxide scale in oxidized Fe-Cr-Si alloys ..................................... 415 G. Bamba, P. Donnadieu, Y. Wouters, and A. Galerie
Applying a classical 2 beam diffraction contrast method for measuring nanoprecipitate misfit ....................................................................... 417 L. Lae and P. Donnadieu
Microstructure and interface composition of γ-phase in Co38Ni33Al29 shape memory alloy.......................................................................... 419 R. Espinoza, B. Bartova, D. Schryvers, S. Ignacova, and P. Sittner
Microstructural characterization of the aluminum alloy 6063 after work hardening treatments .............................................................................. 421 I.M. Espósito, S.J. Buso, and W.A. Monteiro
Microstructure- mechanical property relationships in dual phase automotive strip steels................................................................................................ 423 V. Tzormpatzdi and G. Fourlaris
Electron diffraction analysis of nanocrystalline Fe-Al............................................ 425 C. Gammer, C. Mangler, C. Rentenberger, and H.P. Karnthaler
Dual Beam and TEM characterisation of deformation structures in fatigued austenitic stainless steel........................................................................... 427 A. Garcia, L. Legras, M. Akamatsu, and Y. Bréchet
Microstructural characterisation of steel heat-treated by the novel quenching and partitioning process .................................................... 429 K. He, D.V. Edmonds, J.G. Speer, D.K. Matlock, and F.C. Rizzo
Martensite tempering behaviour relevant to the quenching and partitioning process ............................................................................................ 431 K. He, D.V. Edmonds, J.G. Speer, D.K. Matlock, and F.C. Rizzo
Chemical and structural analysis of NiAl-Al2O3 interface by FETEM and STEM ................................................................................................................... 433 W. Hu, T. Weirich, and G. Gottstein
TEM investigations of aluminum precipitate in eutectic Si of A356 based alloys ................................................................................................... 435 Z.H. Jia, L. Arnberg, P. Åsholt, B. Barlas, and T. Iveland
Microstructure of slow-cooled wedge-cast Cu58Co42 alloy with a metastable liquid miscibility gap ................................................................... 437 E. Johnson, S. Curiotto, N. Pryds, and L. Battezzati
TEM investigations of Elektron 21 magnesium alloy after long-term annealing .......................................................................................... 439 A. Kielbus
Microstructure of AJ62 magnesium alloy after long-term annealing.................... 441 A. Kielbus and J. Mizera
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Content
EELS characterisation of the interface between nanoscaled ODS particles and matrix in advanced fusion steels ........................................................................ 443 M. Klimenkov, R. Lindau, and A. Möslang
Microstructure-mechanical property relationships in a maraging 250 steel ........ 445 P. Kokkonidis, E. Papadopoulou, A. Rizos, T. Koutsoukis, and G. Fourlaris
Microstructure of Co-Ni based superalloys ............................................................. 447 T.J. Konno, T. Tadano, H. Matsumoto, and A. Chiba
Precipitation reactions in superferritic stainless steels ........................................... 449 T. Koutsoukis, K. Konstantinidis, P. Kokkonidis, E. Papadopoulou, and G. Fourlaris
Effect of ageing in cold rolled superaustenitic stainless steels ................................ 451 S. Zormalia, T. Koutsoukis, E. Papadopoulou, P. Kokkonidis, and G. Fourlaris
Structure and properties of P/M material of AlMg – SiO2 system processed by mechanical alloying............................................................................................... 453 A. Kula, L. Błaż, M. Sugamata, J. Kaneko, Ł. Górka, J. Sobota, and G. Włoch
Extrusion of rapidly solidified 6061 + 26 wt% Si alloy ........................................... 455 A. Kula, M. Sugamata, J. Kaneko, L. Błaż, G. Włoch, J. Sobota, and W. Bochniak
TEM and EELS study of carbide precipitation in low alloyed steels..................... 457 C. Leguen, M. Perez, T. Epicier, D. Acevedo, and T. Sourmail
Microstructural analysis of plastically deformed complex metallic alloy κ-AlMnNi ........................................................................................................... 459 M. Lipińska-Chwałek, M. Heggen, M. Feuerbacher, and A. Czyrska-Filemonowicz
Electron microscopy analysis of Mn partitioning in retained austenitemartensite- bainite islands......................................................................................... 461 A. Lis, J. Lis, and P. Wieczorek
TEM characterization of microstructures in a Ni2MnGa alloy.............................. 463 H. Maeda, E. Taguchi, K. Inoue, and A. Sugiyama
Nanocrystalline FeAl produced by high pressure torsion studied by TEM in 3D ............................................................................................................. 465 C. Mangler, C. Rentenberger, and H.P. Karnthaler
HRTEM study of precipitates in Al-Mg-Si-(Ag, Cu) alloys.................................... 467 K. Matsuda, J. Nakamura, T. Kawabata, T. Sato, and S. Ikeno
Martensite structure of non-stoichiometric Co2NiGa ferromagnetic shape memory alloy .............................................................................................................. 469 K. Prusik and M. Morawiec
Electron microscopy of Fe and FeB atomic clusters in the Fe-based amorphous alloys structure ............................................................ 471 E.V. Pustovalov, V.S. Plotnikov, B.N. Grudin, S.V. Dolzhikov, E.B. Modin, O.V. Voitenko, and E.S. Slabzhennikov
Content
XIX
Core/Shell Precipitates in Al-Li-Sc-Zr Alloys.......................................................... 473 V. Radmilovic, M.D. Rossell, A. Tolley, E.A. Marquis, R. Erni, and U. Dahmen
Analysis of basic mechanisms of hardening in ODS EUROFER97 steel using in-situ TEM ....................................................................................................... 475 A. Ramar and R. Schäublin
TEM investigation on the acicular ferrite precipitation in γ’-Fe4N nitride........... 477 X.-C. Xiong, A. Redjaïmia, and M. Gouné
Orientation Relationships between the δ-ferrite Matrix in a Duplex Stainless Steel and its Decomposition Products: the Austenite and the χ and R Frank-Kasper Phases ............................................. 479 A. Redjaïmia, T. Otarola, and A. Mateo
TEM study of localized deformation-induced disorder in intermetallic alloys of L12 structure........................................................................................................... 481 C. Rentenberger, C. Mangler, and H.P. Karnthaler
SEM and TEM study of dynamic recrystallisation of zirconium alloy.................. 483 L. Saintoyant, L. Legras, and Y. Brechet
Effects of solution treatment and test temperature on tensile properties of high strength high Mn austenitic steels ................................................................ 485 K. Phiu-on, W. Bleck, A. Schwedt, and J. Mayer
Microstructure evolution during Ni/Al multilayer reactions ................................. 487 S. Simões, F. Viana, A.S. Ramos, M.T. Vieira, and M.F. Vieira
TEM investigation of severely deformed NiTi and NiTiHf shape memory alloys .................................................................................................. 489 G. Steiner, M. Peterlechner, T. Waitz, and H.P. Karnthaler
TEM studies of nanostructured NiTiCo shape memory alloy for medical applications............................................................................................. 491 D. Stróż and Z. Lekston
TEM investigations of microalloyed steels with Nb, V and Ti after different treatments .......................................................................................... 493 G. Szalay, R. Grill, K. Spiradek-Hahn, and M. Brabetz
Initial Stages of the ω Phase Transformation .......................................................... 495 R. Tewari, K.V. Manikrishna, G.K. Dey, and S. Banerjee
TEM study of the Ni-Ti shape memory micro-wire ................................................ 497 H. Tian, D. Schryvers, and J. Van Humbeeck
Multi-scale observations of deformation twins in Ti6Al4V .................................... 499 W. Tirry, F. Coghe, L. Rabet, and D. Schryvers
Nd:YAG laser joining between stainless steel and nickel-titanium shape memory alloys .................................................................................................. 501 J. Vannod, A. Hessler-Wyser, and M. Rappaz
XX
Content
Focused Ion Beam application on the investigation of tungsten-based materials for fusion application.................................................. 503 L. Veleva, R. Schäublin, A. Ramar, Z. Oksiuta, and N. Baluc
HRTEM of NiTi shape memory alloys made amorphous by high pressure torsion............................................................................................. 505 T. Waitz, K. Tsuchiya, M. Peterlechner, and H.P. Karnthaler
Is the lattice structure of the martensite in nanocrystalline NiTi base centred orthorhombic? .............................................................................................. 507 T. Waitz
TEM study of the NiTi shape memory thin film...................................................... 509 B. Wang, A. Safi, T. Pardoen, A. Boe, J.P. Raskin, X. Wang, J.J. Vlassak, and D. Schryvers
Sub-nano analysis of fine complex carbide in high strength steel with probe Cs (S)TEM ............................................................................................... 511 K. Yamada, E. Hamada, K. Sato, and K. Inoke
Characterization of morphology and microstructure of different kinds of materials at NTNU Mater Sci EM Lab ................................................................ 513 Y.D. Yu, T. Nilsen, M.P. Raanes, J. Hjelen, and J.K. Solberg
Characterization of a Ti64Ni20Pd16 thin film by transmission electron microscopy........................................................................ 515 R. Zarnetta, E. Zelaya, G. Eggeler, and A. Ludwig
Analytical electron microscopy investigations of a microstructure of single and polycrystalline β-Mg2Al3 Samson phase............................................. 517 A. Zielińska-Lipiec, B. Dubie,l and A. Czyrska-Filemonowicz
M3.2 Ceramic materials Grain boundary interfaces in ceramics .................................................................... 519 D.J.H. Cockayne, S.-J. Shih, K. Dudeck, and N. Young
Structure and chemistry of nanometer-thick intergranular films at metal-ceramic interfaces........................................................................................ 521 W.D. Kaplan and M. Baram
Studying nanocrystallization behaviour of different inorganic glasses using Transmission Electron Microscopy ................................................................ 523 Somnath Bhattacharyya, Th. Höche, C. Bocker, C. Rüssel, A. Duran, N. Hémono, F. Muñoz, M.J. Pascual, K. Hahn, and P.A. van Aken
HRTEM and Diffraction Analysis of Surface Phases in Nanostructured LiMn1.5Ni0.5O4 Spinel.................................................................................................. 525 F. Cosandey, N. Marandian Hagh, and G.G. Amatucci
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The structural origin of the antiferroelectric properties and relaxor behavior of Na0.5Bi0.5TiO3 ..................................................................... 527 V. Dorcet, G. Trolliard, and P. Boullay
Electron beam probing of insulators ........................................................................ 529 H.-J. Fitting, N. Cornet, M. Touzin, D. Goeuriot, C. Guerret-Piécourt, D. Juvé, and D. Tréheux
Characterization of Ge-based clathrates oxidized in air by means of TEM and SEM...................................................................................................................... 531 C. Hébert, B. Bartova, M. Cantoni, U. Aydemir, and M. Baitinger
Microstructure analysis of thin Cr2AlC films deposited at low temperature by magnetron sputtering............................................................................................ 533 R. Iskandar, D.P. Sigumonrong, J.M. Schneider, and J. Mayer
Study of structural variation in YBaCo4O7+δ by electron diffraction .................... 535 Y. Jia, H. Jiang, M. Valkeapää, M. Karppinen, and E.I. Kauppinen
Exsolution phenomena in glass-ceramic systems..................................................... 537 I. Tsilika, Ph. Komninou, G.P. Dimitrakopulos, Th. Kehagias, and Th. Karakostas
Transmission Electron Microscopy Studies of Lead-Free Ferroelectrics in the System BNT-BT-KNN ..................................................................................... 539 H.-J. Kleebe, J. Kling, L. Schmitt, S. Lauterbach, and H. Fuess
ReO3-related aluminum tungsten oxides showing a novel type of crystallographic shear structure........................................................................... 541 F. Krumeich and G.R. Patzke
Structural Characterisation by TEM of a New Homologous Series Bi2n+4MonO6(n+1); n=3,4,5 and 6.................................................................................. 543 A.R. Landa-Canovas, J. Hernández-Velasco, E. Vila, J. Galy, and A. Castro
Structural characterisation of a new rich iron layered oxide TlεSr25-εFe30O76-ξ .... 545 C. Lepoittevin, S. Malo, S. Hebert, M. Hervieu, and G. Van Tendeloo
EBSD studies of stress concentrations in ferroelectrics .......................................... 547 I. MacLaren, M.U. Farooq, R. Villaurrutia, T.L. Burnett, T.P. Comyn, A.J. Bell, H. Kungl, and M.J. Hoffmann
High-resolution pictures of nucleation growth triangle of 180° ferroelectric domain wall in a thin film of LiTaO3 obtained by Lorentz DPC-STEM............... 549 T. Matsumoto, M. Koguchi, and Y. Takahashi
Size and structure of barium halide nano-crystals in optically active fluorozirconate-based glasses .................................................................................... 551 P.T. Miclea, B. Ahrens, C. Eisenschmidt, and S. Schweizer
Domain Structure And Microstructure Development of BaTiO3 Doped With Rare-Earth Dopants ......................................................................................... 553 V. Mitic, V.B. Pavlovic, V. Paunovic, M. Miljkovic, B. Jordovic, and Lj. Zivkovic
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SEM and EDS Analysis of BaTiO3 Doped With Yb2O3 and Ho2O3 ....................... 555 V. Mitic, V.B. Pavlovic, V. Paunovic, M. Miljkovic, B. Jordovic, and Lj. Zivkovic
Structure and superconductivity of Pr-Ba-Cu-O crystals prepared by ambient pressure synthesis using citrate pyrolysis method............................... 557 K. Nishio, T. Isshiki, T. Shima, and M. Hagiwara
Electron Diffuse Scattering in epitaxially grown SrTiO3 thin film ........................ 559 F. Pailloux and J. Pacaud
Investigation of the hydration of calciumsulfate hemihydrates with different microscopic methods.......................................................................... 561 C. Pritzel and R. Trettin
Investigation of holes in calciumsulfate-hemihydrate crystals by different microscopic methods ............................................................................. 563 C. Pritzel and R. Trettin
Analytical and high-resolution TEM investigation of Boron-doped CeO2 ............ 565 B. Rahmati, G. Gregori, W. Sigle, C.T. Koch, P.A. van Aken, and J. Maier
Accommodation of the compositional variations in the Ca1-xSrxMnO3-δ (0≤x≤1, 0≤δ≤0.5) system............................................................................................. 567 S. de Dios, J. Ramírez-Castellanos, A. Varela, M. Parras, and J.M. González Calbet
Evidence of SrO(SrTiO3)n Ruddlesden-Popper Phases by High Resolution Electron Microscopy and Holography .................................... 569 M. Reibold, E. Gutmann, A.A. Levin, A. Rother, D.C. Meyer, and H. Lichte
New Barium Antimony Aluminates evidenced by TEM techniques...................... 571 R. Retoux, A. Letrouit, M. Hervieu, and S. Boudin
(Multi-)ferroic domain walls– a combined ab-initio and microscopical investigation ................................................................................ 573 A. Rother, S. Gemming, D. Geiger, and N. Spaldin
Diagnostic of Li battery cathode by EELS, first principles calculation and spectrum-imaging with multi-variate analysis ................................................. 575 K. Tatsumi, Y. Sasano, S. Muto, T. Sasaki, Y. Takeuchi, K. Horibuchi, and Y. Ukyo
Local electronic structure analysis on Ce3+-containing materials by TEM-EELS and first principles calculations...................................................... 577 K. Tatsumi, I. Nishida, and S. Muto
Local chemical inhomogeneities in NaNb1-xTaxO3 as observed by HRTEM and HAADF-STEM.................................................................................................... 579 A. Torres-Pardo, E. García-González, J.M. González-Calbet, F. Krumeich, and R. Nesper
The influence of lanthanum doping on the structure of PbZr0.9Ti0.1O3 ceramics.......................................................................................... 581 R. Villaurrutia, I. MacLaren, and A. Peláiz-Barranco
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Anomalous absorption of electrons during electron diffraction on BaTiO3 single crystals near phase transition at 120°C ...................................... 583 A. Wall
Template-assisted synthesis and characterization of SrTiO3 nanostructures........................................................................................... 585 K. Žagar, S. Šturm, and M. Čeh
(S)TEM/EELS characterisation of a multilayer C/Cr PVD coating ...................... 587 Z. Zhou, W.M. Rainforth, M. Gass, A. Bleloch, and P.Eh. Hovsepian
M3.3 Magnetic materials High resolution imaging of magnetic structures in a TEM – what is possible? .... 589 J. Zweck
Phase segregation leading to spontaneous outcropping of (Sr,La)Ox dots in La1-xSrxMnO3 films ................................................................................................ 591 P. Abellan, F. Sandiumenge, C. Moreno, M.J. Casanove, T. Puig, and X. Obradors
Microstructure of epitaxially strained LaCoO3 thin films...................................... 593 L. Dieterle, D. Gerthsen, and D. Fuchs
Are the samples really flat? Influence of the supporting membrane on the magnetization of patterned micromagnets ................................................... 595 C. Dietrich and J. Zweck
HRTEM characterization of core-shell Fe@C and Fe@SiO2 magnetic nanoparticles prepared by the arc-discharge plasma method................................ 597 Rodrigo Fernández-Pacheco, Manuel Arruebo, Jordi Arbiol, Clara Marquina, Jesús Santamaría, and M. Ricardo Ibarra
Nanofabrication of ferromagnetic nanotips and nanobridges by 2D and 3D electron-beam cutting ................................................................................... 599 T. Gnanavel, Z. Saghi, Y. Peng, B.J. Inkson, M.R.J. Gibbs, and G. Möbus
An investigation into the crystallization of the MgO barrier layer of a magnetic tunnel junction .................................................................................... 601 V. Harnchana, A.P. Brown, R.M. Brydson, A.T. Hindmarch, and C.H. Marrows
FeCoAlN films with induced magnetic anisotropy .................................................. 603 A. Lančok, M. Klementová, M. Miglierini, F. Fendrych, K. Postava, J. Kohout, and O. Životský
The martensitic microstructure of 5M and NM martensites in off-stoichiometric Ni2MnGa ferromagnetic shape memory alloys..................... 605 Pallavi Sontakke, Amita Gupta, and Madangopal Krishnan
TEM characterization of nanometer-sized Fe/MgO granular multilayer thin films grown by pulsed laser deposition............................................................. 607 C. Magén, E. Snoeck, A. García-García, J.A. Pardo, P.A. Algarabel, P. Štrichovanec, A. Vovk, L. Morellón, J.M. De Teresa, and M.R. Ibarra
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Content
Structural modification and self-assembly of nanoscale magnetite synthesised in the presence of an anionic surfactant ................................................................... 609 S. Malik, I.J. Hewitt, and A.K. Powell
Electron microscopy phase retrieval of perpendicular magnetic anisotropy (PMA) FePd alloys ..................................................................................................... 611 A. Masseboeuf, C. Gatel, A. Marty, E. Snoek, and P. Bayle-Guillemaud
Magnetic domain wall propagation in nanostructures of alloys with perpendicular magnetic anisotropy.................................................................. 613 A. Masseboeuf, A. Mihaï, J.P. Attané, J.C. Pillet, P. Warin, A.L. Vila, G. Gaudin, M. Miron, B. Rodmacq, E. Gautier, A. Marty, and P. Bayle-Guillemaud
The effect of annealing in the microstructure and magnetic properties of NiCuZn ferrites ...................................................................................................... 615 D. Sakellari, V. Tsakaloudi, V. Zaspalis, and E.K. Polychroniadis
Microstructural and compositional analyses of nano-structured Co-Pt thin films ..................................................................................................................... 617 Z. Samardžija, K. Žužek Rožman, and S. Kobe
L10-type ordered structure of FePd nanoparticles studied by high-resolution transmission electron microscopy ............................................. 619 K. Sato, T.J. Konno, and Y. Hirotsu
Structural and chemical characterization of Co-doped ZnO layers grown on Si and sapphire ...................................................................................................... 621 R. Schneider, L.D. Yao, D. Gerthsen, G. Mayer, M. Fonin, and U. Rüdiger
TEM studies of cobalt-doped zinc oxide films ......................................................... 623 J. Simon, K. Nielsen, M. Opel, S.T.B. Goennenwein, R. Gross, and W. Mader
Nanocrystallization of amorphous Fe40Ni38B18Mo4 alloy ........................................ 625 D. Srivastava, A.P. Srivastava, and G.K. Dey
Structural and compositional properties of Sm-Fe-Ta magnetic nanospheres prepared by pulsed-laser deposition at 157 nm in a N2 atmosphere...................... 627 S. Šturm, K. Žužek Rožman, E. Sarantopoulou, and S. Kobe
Characterization of Ni-Mn-Ga magnetic shape memory alloys using electron holography and Lorentz microscopy ............................................... 629 K. Vogel, M. Linck, Ch. Matzeck, A. Rother, D. Wolf, and H. Lichte
Energy Loss Magnetic Chiral Dichroïsm (EMCD) for magnetic material............ 631 B. Warot-Fonrose, L. Calmels, C. Gatel, F. Houdellier, V. Serin, and E. Snoeck
M3.4 Dislocations, interfaces and other defects Determining the nanoscale chemistry and behavior of interfaces and phases in Al-Si(-Cu-Mg) nanoparticles using in-situ TEM................................................. 633 J.M. Howe, S.K. Eswaramoorthy, and G. Muralidharan
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Dislocations in AlPdMn quasicrystals: contrast in TEM and physical properties .............................................................................................. 635 D. Caillard and F. Mompiou
Characterization of a-plane GaN films grown on r-plane sapphire substrate by electron microscopy .............................................................................................. 637 Y. Arroyo Rojas Dasilva, T. Zhu, D. Martin, N. Grandjean, and P. Stadelmann
A method for atomistic/continuum analysis of defects in large HRTEM images ............................................................................................ 639 A. Belkadi, G.P. Dimitrakopulos, J. Kioseoglou, G. Jurczak, T.D. Young, P. Dluzewski, and Ph. Komninou
High resolution electron microscopy of interfaces in ultrafine microstructures of Zr and Ti based alloys ........................................... 641 G.K. Dey, S. Neogy, R.T. Savalia, R. Tewari, D. Srivastava, and S. Banerjee
Anisotropic strain relaxation in (110) La2/3Ca1/3MnO3 thin films .......................... 643 S. Estrade, I.C. Infante, F. Sanchez, J. Fontcuberta, J. Arbiol, and F. Peiró
Metadislocations in complex metallic alloys: core structures investigated by aberration-corrected scanning transmission electron microscopy ................... 645 M. Feuerbacher, L. Houben, and M. Heggen
TEM of high pressure torsion processed intermetallic Zr3Al................................. 647 D. Geist, C. Rentenberger, and H.P. Karnthaler
Multiscale characterisation of the plasticity of Fe-Mn-C TWIP steels .................. 649 H. Idrissi, L. Ryelandt, K. Renard, S. Ryelandt, F. Delannay, D. Schryvers, and P.J. Jacques
Misfit analysis of the InN/GaN interface through HRTEM image simulations ....................................................................................................... 651 J. Kioseoglou, G.P. Dimitrakopulos, Th. Kehagias, E. Kalessaki, Ph. Komninou, and Th. Karakostas
Application of TEM for Real Structure Determination of Rare Earth Metal Compounds.............................................................................. 653 L. Kienle, V. Duppel, Hj. Mattausch, M.C. Schaloske, and A. Simon
Quantitative Dislocation Analysis of 2H AlN:Si grown on (0001) Sapphire ......... 655 O. Klein, J. Biskupek, U. Kaiser, S.B. Thapa, and F. Scholz
Transrotational crystals in crystallizing amorphous films: new solid state order or novel extended imperfection............................................. 657 V.Yu. Kolosov
Determination of precise orientation relationships between surface precipitates and matrix in a duplex stainless steel................................................... 659 Y. Meng, G. Nolze, W.Z. Zhang, L. Gu, and P.A. van Aken
Interfaces in Cu(In,Ga)Se2 thin film solar cells ....................................................... 661 G. Östberg and E. Olsson
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Content
In-situ electron beam irradiation of nanopipes in GaN .......................................... 663 F. Pailloux and J.-F. Barbot
The atomic structure of an incommensurate (001)-(110) Si grain boundary resolved thanks to a probe Cs-corrector .................................................................. 665 J.L. Rouviere, F. Lançon, K. Rousseau, D. Caliste, and F. Fournel
Atomic structure and dopant segregation of [0001] tilt grain boundaries in ZnO bicrystals ........................................................................................................ 667 Y. Sato, T. Mizoguchi, J.P. Buban, N. Shibata, T. Yamamoto, T. Hirayama, and Y. Ikuhara
TEM study of strain and defect engineering with diluted nitride semiconductors ......................................................................... 669 J. Schöne, E. Spiecker, F. Dimroth, A.W. Bett, and W. Jäger
Investigation of the Co-Precipitation of Copper and Nickel in Silicon by Means of Transmission Electron Microscopy..................................................... 671 C. Rudolf, L. Stolze, and M. Seibt
Micro-structure analysis of a friction-stir welded 2024 aluminium alloy using electron microscopy.......................................................................................... 673 E. Sukedai, T. Maebara, and T. Yokayama
Deformation defects in a metastable β titanium alloy............................................. 675 H. Xing and J. Sun
Defect generation and characterization in 4H-SiC.................................................. 677 J.P. Ayoub, M. Texier, G. Regula, M. Lancin, and B. Pichaud
Investigation of defects in polymorph B enriched zeolite Beta............................... 679 D. Zhang, J. Sun, S. Hovmöller, and X. Zou
M3.5 Coatings and graded materials Optimizing electron diffraction and EDS for phase identification in complex structures: application to multilayered Ti-Ni-P coatings .................... 681 P.A. Buffat and A. Czyrska-Filemonowicz
Advanced analytical transmission electron microscopy to investigate the nano-graded materials properties ...................................................................... 683 M. Cheynet, S. Pokrant, L. Joly-Pottuz, and J.M. Martin
Characterisation of Nickel Nanocomposites by SEM, TEM and EBSD................ 685 D. Dietrich, Th. Lampke, B. Wielage, D. Thiemig, and A. Bund
Characterisation of Gold Nanocomposites by SEM, TEM and EBSD .................. 687 D. Dietrich, Th. Lampke, B. Wielage, P. Cojocaru, and P.L. Cavallotti
Alumina Coatings as Protection against Corrosive Atmosphere ........................... 689 I. Dörfel, R. Sojref, M. Dressler, D. Hünert, and M. Nofz
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Advanced Multilayer Systems for X-ray Optics: Quality Assessment by TEM ....................................................................................................................... 691 D. Häussler, W. Jäger, E. Spiecker, B. Ögüt, U. Ross, J. Wiesmann, and M. Störmer
Surface investigation of SU-8 by atomic force and scanning electron microscopy ............................................................................ 693 Th. Kups, Chr. Kremin, M. Hoffmann, and L. Spieß
SEM and TEM investigations of electrophoretical deposited Si3N4 and SiC particles in siloxane of steel substrate ........................................................ 695 Th. Kups, A. Knote, and L. Spieß
Contribution of electron microscopy techniques to the chemical and structural characterization of TiC/a-C nanocomposite coatings .................... 697 C. López-Cartes, D. Martínez-Martínez, J.C. Sánchez-López, and A. Fernández
TEM investigations of the Ti/TiN multilayered coatings deposited on the Ti-6Al-7Nb alloy.............................................................................................. 699 T. Moskalewicz, H.J. Penkalla, and A. Czyrska-Filemonowicz
Microstructural examination of Al and Cr alloyed zinc coatings on low carbon steels.................................................................................................... 701 D. Chaliampalias, G. Vourlias, E. Pavlidou, K. Chrissafis, G. Stergioudis, and S. Skolianos
Study of the structure and high temperature oxidation resistance of high alloyed tool steels ........................................................................................... 703 E. Pavlidou, D. Chaliampalias, G. Vourlias, and K. Chrissafis
A comparative study of NiCrBSi and Al coated steels with thermal spray process in different environments............................................................................. 705 D. Chaliampalias, G. Vourlias, E. Pavlidou, K. Chrissafis, G. Stergioudis, and S. Skolianos
Microscopical study of the influence of zinc addition on the structure of WO3.... 707 K. Nikolaidis, D. Chaliampalias, G. Vourlias, E. Pavlidou, and G. Stergioudis
Microstructural Studies by Electron Microscopy Techniques of TiAlSiN Nanostructured Coatings........................................................................................... 709 V. Godinho, T.C. Rojas, M.C. Jimenez, M.P. Delplancke-Ogletree, and A. Fernández
Structural and interface studies of a nano-scale TiAlYN/CrN/alumina coating ................................................................................... 711 I.M. Ross, W.M. Rainforth, C. Strondl, F. Papa, and R. Tietema
An investigation of SiC-fiber coatings ...................................................................... 713 T. Toplišek, Z. Samardžija, G. Dražić, S. Kobe, and S. Novak
HRTEM-EELS study of atomic layer deposited thin rare earth oxide films for advanced microelectronic devices ....................................................................... 715 S. Schamm, P.E. Coulon, and L. Calmels
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Content
New Fullerene like materials for tribological applications: TEM and EELS study................................................................................................ 717 Virginie Serin, Nathalie Brun, and Ch. Colliex
Phase determination of nanocrystalline Al-Cr-O coatings by analytical TEM...................................................................................................... 719 J. Thomas, J. Ramm, B. Arnold, B. Widrig, and T. Gemming
M3.6 Biomaterials Bioinspired synthesis of nanostructures based on S-layer lattices ......................... 721 D. Pum, N. Ilk, and U.B. Sleytr
Direct Imaging of Carbon Nanoparticles inside Human Cells ............................... 723 A.E. Porter, C. Cheng, M. Gass, K. Muller, J. Skepper, P. Midgley, and M. Welland
Micro- and Nano-Textured Surfaces on Ti-Implants Made by Various Methods ................................................................................................... 725 U. Beck, R. Lange, and H.-G. Neumann
Determination of the biocompatibility of biomaterials by scanning electron microscopy (SEM) .................................................................. 727 M. Bovi, N. Gassler, and B. Hermanns-Sachweh
Quantitative evaluation of the long-term marginal behaviour of filling restorations of human teeth using three-dimensional scanning electron microscopy.................................................................................................... 729 W. Dietz, S. Nietzsche, R. Montag, P. Gaengler, and I. Hoyer
The analysis of Si doped hydroxyapatite coatings using FIBSEM, TEM and RHEED ................................................................................................................ 731 H.K. Edwards, S. Coe, T. Tao, M.W. Fay, C.A. Scotchford, D.M. Grant, and P.D. Brown
Electron microscopic investigations of the polymer/mineral composite material nacre............................................................................................................. 733 K. Gries, R. Kröger, C. Kübel, M. Fritz, and A. Rosenauer
Studies on the microstructure of fresh-cut melon ................................................... 735 I. Hernando, L. Alandes, A. Quiles, and I. Pérez-Munuera
Ceramic-loaded mineralizing bioresorbable polymers for orthopaedic applications...................................................................................... 737 L.W. Hobbs, T. Lim, A. Porter, H. Wang, M. Walton, and N.J. Cotton
AFM and TEM study of Ag coated insulin-derived amyloid fibrils ...................... 739 M. Gysemans, J. Snauwaert, C. Van Haesendonck, F. Leroux, B. Goris, S. Bals, and G. Van Tendeloo
Transmission Electron Microscopy studies of bio-implant interfaces using Focused Ion Beam microscopy for sample preparation................................ 741 F. Lindberg, A. Palmquist, L. Emanuelsson, J. Heinrichs, R. Brånemark, F. Ericson, P. Thomsen, and H. Engqvist
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AFM and SEM of Wax Crystallisation on Artificial Surfaces Controlled by Temperature and Solvents ................................................................................... 743 A. Niemietz, W. Barthlott, K. Wandelt, and K. Koch
Characterization of layer-by-layer microcapsules made of hyaluronic acid by CLSM, SEM and TEM ......................................................................................... 745 I. Pignot-Paintrand, A. Szarpak, and R. Auzely-Velty
Atomic force microscopy analysis of crystalline silicon functionalization with oligonucleotides .................................................................................................. 747 A. Ponzoni, G. Faglia, M. Ferroni, G. Sberveglieri, A. Flamini, G. Andreano, and L. Cellai
Hidden hierarchy of microfibrils within fluorapatite gelatine nanocomposites induced by intrinsic electric dipole fields ..................................... 749 P. Simon and R. Kniep
M4
Soft Matter and Polymers
Self-assembled block copolymer structures studied by transmission electron microtomography ........................................................................................ 751 H. Jinnai, T. Kaneko, C. Abetz, and V. Abetz
Quantitative chemistry and orientation of polymers in 2-d and 3-d by scanning transmission X-ray microscopy............................................................ 753 A.P. Hitchcock, G.A. Johansson, D. Hernández Cruz, E. Najafi, J. Li and and H. Stöver
Characterization of cavitation processes in filled semi-crystalline polymers........ 755 F. Addiego, J. Di Martino, D. Ruch, A. Dahoun, and O. Godard
Quantitative analysis of protein gel structure by confocal laser scanning microscopy .................................................................................................. 757 K. Ako, L. Bécu, T. Nicolai, J.-C. Gimel, and D. Durand
Thermal stability of organic solar cells: the decay in photocurrent linked with changes in active layer morphology ................................................................. 759 S. Bertho, I. Haeldermans, A. Swinnen, J. D’Haen, L. Lutsen, T.J. Cleij, J. Manca, and D. Vanderzande
Determining absorptive potential variation in electron beam sensitive specimens using a single energy-filtered bright-field TEM image ......................... 761 S. Bhattacharyya and J.R. Jinschek
Elemental distribution of soft materials with newly designed 120kV TEM/STEM .................................................................................... 763 C. Hamamoto, N. Endo, H. Nishioka, T. Ishikawa, Y. Ohkura, and T. Oikawa
Preparation of titanate nanotubes and their polymer composites ......................... 765 D. Kralova, N. Neykova, and M. Slouf
XXX
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Nanometer size wear debris generated from ultrahigh molecular weight polyethylene in vivo .................................................................................................... 767 M. Lapcikova, M. Slouf, J. Dybal, E. Zolotarevova, G. Entlicher, D. Pokorny, J. Gallo, and A. Sosna
Analysis of nano-composites based on carbon nanoparticles imbedded in polymers.................................................................................................................. 769 Kangbo Lu, Joachim Loos, Sourty Erwan, and Dong Tang
New developments in SEM for in situ tensile tests on polymers............................. 771 P. Jornsanoh, G. Thollet, C. Gauthier, and K. Masenelli-Varlot
A study of the spatial distributions of the carbon blacks in polypropylene composites using TEM-Tomography and quantitative image analysis ................. 773 H. Matsumoto, H. Sugimori, T. Tanabe, Y. Fujita, H. Sano, and H. Jinnai
A study of the chain-folded lamellae structure and its array in the isotactic polypropylene spherulites by HAADF-STEM and HV-TEM Tomography techniques.................................................................... 775 H. Matsumoto, M. Song, H. Sano, M. Shimojo, and K. Furuya
Microstructural analysis of ultra-thin nanocomposite layers fabricated by Cu+ ion implantation in inert polymers............................................................... 777 G. Di Girolamo, E. Piscopiello, M. Massaro, E. Pesce, C. Esposito, L. Tapfer, and M. Vittori Antisari
In-situ experiments on soft materials in the environmental SEM – Reliable results or merely damage?......................................................................................... 779 P. Poelt, H. Reingruber, A. Zankel, and C. Elis
Structural studies on V-amylose inclusion complexes............................................. 781 J.L. Putaux, M. Cardoso, M. Morin, D. Dupeyre, and K. Mazeau
Multilamellar nanoparticles from PS-b-PVME copolymers .................................. 783 C. Lefebvre, J.-L. Putaux, M. Schappacher, A. Deffieux, and R. Borsali
TEM/SEM characterisation of hybrid titanoniobiates used as fillers for thermoplastic nanocomposites ............................................................................ 785 R. Retoux, S. Chausson, L. Le Pluart, J.M. Rueff, and P.A. Jaffres
Phase transitions and ordering in liquid crystals – a case study ............................ 787 A.K. Schaper
Study of degradation and regeneration of silicon polymers using cathodoluminescence........................................................................................ 789 P. Schauer, P. Horak, F. Schauer, I. Kuritka, and S. Nespurek
Orthogonal self-assembly of surfactants and hydrogelators: towards new nanostructures...................................................................................... 791 M.C.A. Stuart, A.M.A. Brizard, E.J. Boekema, and J.H. van Esch
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Structure of particles formed during Se redox process in aqueous polymer solutions .................................................................................... 793 E.I. Suvorova, V.V. Klechkovskaya, M. Cantoni, and P.A. Buffat
Exploring the 3D organisation of high-performance organic solar cells ............... 795 S. van Bavel, E. Sourty, B. de With, and J. Loos
Morphological study on three kinds of two-dimensional spherulites of PBT ....... 797 T. Yoshioka and M. Tsuji
Solution-Grown Crystals of Optically Active Propene–Carbon Monoxide Copolymer................................................................................................................... 799 T. Yoshioka, N. Kosaka, A. Nakayama, A.K. Schaper, W. Massa, T. Hiyama, K. Nozaki, and M. Tsuji
M5
Materials in Mineralogy, Geology and Archaeology
New insights into ultra-high pressure metamorphism from TEM studies ............ 801 F. Langenhorst and A. Escudero
Characterization of a (021) twin in coesite using LACBED and precession electron diffraction........................................................................... 803 P. Cordier, D. Jacob, and H.-P. Schertl
Rubens in the Prado National Museum: analytical characterization of ground layers.......................................................................................................... 805 M.I. Báez, L. Vidal, M.D. Gayo, J. Ramírez-Castellanos, J.L. Baldonedo, and A. Rodríguez
Development of the FIB-cryo-SEM approach for the in-situ investigations of the elusive nanostructures in wet geomaterials ................................................... 807 G. Desbois and J.L. Urai
TEM applied on the interface characterisation of the replacement reaction chlorapatite by hydroxyapatite ................................................................................. 809 U. Golla-Schindler, A. Engvik, H. Austrheim, and A. Putnis
Quantitative study of valence states of zirconolites ................................................. 811 U. Golla-Schindler and P. Pöml
Study of Organic Mineralogical Matter by Scanning Probe Microscopy ............. 813 Ye.A. Golubev and O.V. Kovaleva
Research of Nanoparticle Aggregates from Water Colloidal Solutions of Natural Carbon Substances and Fullerenes by Atomic Force Microscopy ...... 815 Ye.A. Golubev and N.N. Rozhkova
Diffusion in Synthetic Grain Boundaries ................................................................. 817 K. Hartmann, R. Wirth, R. Dohmen, G. Dresen, and W. Heinrich
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Content
An examination of Van Gogh’s painting grounds using analytical electron microscopy – SEM/FIB/TEM/EDX .......................................................................... 819 R. Haswell, U. Zeile, and K. Mensch
Amorphisation in fresnoite compounds – a combined ELNES and XANES study....................................................................................................... 821 Th. Höche, F. Heyroth, P.A. van Aken, F. Schrempel, G.S. Henderson, and R.I.R. Blyth
TEM study of Comet Wild 2 pyroxene particles collected during the stardust mission ....................................................................................... 823 D. Jacob, J. Stodolna, and H. Leroux
The mechanism of ilmenite leaching during experimental alteration in HCl-solution ........................................................................................................... 825 A. Janßen, U. Golla-Schindler, and A. Putnis
Microstructure and Texture from Experimentally Deformed Hematite Ore ....... 827 K. Kunze, H. Siemes, E. Rybacki, E. Jansen, and H.-G. Brokmeier
Identifying pigments in the temple of Seti I in Abydos (Egypt) ............................. 829 E. Pavlidou, H. Marey Mahmoud, E. Roumeli, F. Zorba, K.M. Paraskevopoulos, and M.F. Ali
Nanostructural study of ground layers of canvas of Rubens at “El Prado” National Museum ............................................................................... 831 J. Ramírez-Castellanos, J.L. Baldonedo, M.I. Báez, L. Vidal, M.D. Gayo, and M.J. García
Micro- and nano-diamond particles in carbon spherules found in soil samples ............................................................................................................. 833 Z. Yang, D. Schryvers, W. Rösler, N. Tarcea, and J. Popp
The use of FIB/TEM for the study of radiation damage in radioactive/non-radioactive mineral assemblages............................................... 835 A.-M. Seydoux-Guillaume, J.-M. Montel, and R. Wirth
Non-destructive 3D measurements of sandstone’s internal micro-architecture using high resolution micro-CT ................................................................................ 837 E. Van de Casteele, S. Bugani, M. Camaiti, L. Morselli, and K. Janssens
Author Index............................................................................................................... 839 Subject Index .............................................................................................................. 859
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Direct observation of atomic defects in carbon nanotubes and fullerenes K. Suenaga National Institute of Advanced Industrial Science and Technology (AIST) and the Japan Science and Technology Agency (JST), Tsukuba, 305-8565, Japan
[email protected] Keywords: defect, nanotube, fullerene
The diversified properties of carbon nano-structures (nanotubes, fullerenes and their derivatives) are related to their polymorphic arrangement of carbon atoms. Therefore the direct observation of carbon network, such as defects or chirality, is of great consequence in both scientific and technological viewpoints in order to predict the physical and chemical properties. In order to identify the local configuration of pentagons and hexagons in carbon nanostructures, an electron microscope with higher spatial resolution and higher sensitivity is definitively required. Since the spatial resolution of the conventional TEM is limited by the spherical aberration coefficient (Cs) of its objective lens and the wave length (λ) of incident electron beam, the Cs must be minimized to achieve the best performance because the reduction of the λ is detrimental to the high sensitivity to visualize individual carbon atoms. A highresolution transmission electron microscope (HRTEM, JEOL-2010F) equipped with a Cs corrector (CEOS) was operated at a moderate accelerating voltage (120kV). The Cs was set to 0.5 ~ 10 µm in this work. The HRTEM images were digitally recorded under a slightly under-focus condition (Δf = -2 to -7 nm) where the point resolution of 0.14 nm was achieved at 120kV. The spatial resolution of 0.14 nm (a typical C-C bond length) obtained at a moderate accelerating voltage provides us a great advantage because we can realize the visualization of carbon atomic arrangement without massive electron irradiation damage. Here we show some examples for atomic-level characterization of carbon nanostructures. The C60 and C80 fullerene molecule has been successfully identified its structure and orientation at a single-molecular basis (1, 2). Also the active topological defects have been eventually caught red-handed (3, 4). The technique can be widely applicable to visualize a biological activity, at an atomic level, for which any conformation change of the C-C bonds is responsible. The cis-/trans-isomerization of retinal molecules have been successfully visualized (5). 1. 2. 3. 4. 5.
Z. Liu, K. Suenaga and S. Iijima, J. Am. Chem. Soc., 129 (2007) 6666. Y. Sato, K. Suenaga, S. Okubo, T. Okazaki and S. Iijima, Nano Letters, 7 (2007) 3704. K. Suenaga, H. Wakabayashi, M. Koshino, Y. Sato, K. Urita and S. Iijima, Nature Nanotech. 2 (2007) 358. C.-H. Jin, K. Suenaga and S. Iijima, Nature Nanotech. 3 (2008) 17. Z. Liu, K. Yanagi, K. Suenaga, H. Kataura and S. Iijima, Nature Nanotech. 2 (2007) 422.
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6.
The supports of the JST-CREST, JST-ERATO, KAKENHI and JSPS are kindly acknowledged.
Figure 1. a, A 5–7 pair defect found in an SWNT after heat treatment at 2,273 K. b, An enlarged image of the area enclosed by the green line in a) in which the 5–7–7–5 defect can be more clearly seen. c, The black dots indicate the hexagons with six neighbors, the two red dots have seven neighbors, and the two black dots with circles have five neighbors.
Figure 2. a, The Stone-Wales (SW) transformation leading to the 5–7–7–5 defect, generated by rotating a C–C bond in a hexagonal network. b, HR-TEM image obtained for the atomic arrangement of the SW model. c, Simulated HR-TEM image for the model shown in b. (ref. 3)
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Atomic studies on ferroelectric oxides by aberration corrected transmission electron microscopy K. Urban and C.L. Jia Institute for Solid State Research, Research Centre Jülich, and Ernst Ruska Centre for Microscopy and Spectroscopy with Electrons, D-52428 Jülich, Germany
[email protected] Keywords: Oxides, Interfaces, Defects
The advent of aberration-corrected transmission electron microscopy in 1998 [1] has provided materials science with entirely new tools for quantitative investigations. Four key innovations have to be mentioned: (1) The possibility to operate the electron microscope as a variable-sphericalaberration instrument allows to derive a new phase contrast theory optimizing both resolution and point spread [2]. In classical Scherzer phase contrast theory the radius of the point spread disc amounts to three times the Scherzer resolution limit. Besides the fact that information is lost by placing an aperture in the diffraction plane to keep the contrast oscillations in the contrast transfer function from affecting the images this point spread is a second disadvantage of the classical Scherzer approach to phase contrast. Both limitations can be substantially reduced in a new theory in which by both the objective lens defocus Z as well as CS the spherical aberration parameter adopt specific values. As a result the resolution limit coincides with the information limit and the point spread gets reduced to about one half of the latter making it an uncritical parameter in practice. (2) The negative spherical aberration imaging (NCSI) technique leads to enhanced contrast of atoms with low nuclear charge number [3]. It relies on two advantages compared to the classical Zernike technique. The shift of the phase of the diffracted waves is, in contrast to the classical Scherzer technique, in clockwise direction leading to white atom contrast. Furthermore the contrast is enhanced by a dynamic non-linear effect. Oxygen, nitrogen and even boron can be imaged directly even when these atomic species occur in close distance to heavy cations. (3) Essentially point-spread-free atomic images allow to measure occupancies of atomic columns, i.e. local concentrations, with lateral atomic resolution evaluating atomically resolved intensity measurements [4]. This means that high-resolution is not only a structural technique. From now on also local composition maps can be derived which are forming an excellent starting point for ab-initio calculations of interface-, boundary- and defect structures. (4) Measurements of atomic distances can be carried out at an accuracy of a few picometers [5]. This promises to measure structure-dependent physical properties directly on the atomic scale [6]. The technique of choice in ultra-high resolution transmission electron microscopy is to take a focal series of images on the basis of the NCSI technique which is forming the
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basis of a reconstruction of the electron exit-plane wave function (EPW) [7]. However, in order to arrive at the EPW all major aberrations up to fifth order have to be known or must be compensated. This has to be done by proper diagnostics and adjustment software. The focal series reconstruction allows for additional small corrections of the aberrations but it is by principle not suited to replace aberration correction in general. It should be noted that ultra-high resolution requires modelling of the structure and composition on the computer. The reconstructed EPW is in general not(!) sufficient to carry out high precision measurements since neither the specimen illumination conditions nor the thickness is known a priori. Both seriously affect the EPW. While the latter may permit qualitative interpretation, provided that the projected potential approximation is valid, it is required to do the complete computer fit up to the eventual real structure in order to be able to carry out the picometer-accuracy measurements in the computer. A first example in which the enhanced accuracy of aberration correction has been successfully applied is the investigation of the core structure of Σ3{111} twin boundaries in BaTiO3 [4]. It could be shown that the occupancy of the oxygen sites in the boundary is only 68 %, i.e. 32 % of the sites are left vacant. The corresponding change in the Ti-Ti separation across the boundary of +35 pm and of the Ba-Ba-separation of 17 pm is well described by ab-initio calculations [8]. In a recent study of PbZr0 2Ti0 8O3 (PZT) a new inversion domain boundary was discovered [6]. This longitudinal boundary is charged and presumably formed by the dynamics of domain growth during cooling from above the critical temperature. The boundary could be characterised on the atomic level and the polarisation shifts were measured atom by atom at an accuracy of a few picometers. From these data the polarisation could be calculated as a function of distance from the core of the domain. This is a first example that ferroelectric properties can be measured by ultra-high resolution atomic transmission electron microscopy. 1. 2. 3. 4. 5. 6. 7. 8.
M. Haider, S. Uhlemann, E. Schwan, H. Rose, B. Kabius, and K. Urban, Nature 392 (1998) p. 768. M. Lentzen, Ultramicroscopy 99 (2004) p. 211. C.L. Jia, M. Lentzen and K. Urban, Science 299 (2003) p. 870. C.L. Jia and K. Urban, Science 303 (2004) p. 2001. L. Houben, A. Thust and K. Urban, Ultramicroscopy 106 (2006), p. 200. C.L. Jia, S.B. Mi, K. Urban, I. Vrejoiu, M. Alexe and D. Hesse, Nature mat. 7 (2008) p. 57 K. Tillmann, A. Thust, and K. Urban, Microsc. Microanal. 10 (2004) p. 185. W.T. Geng, Y.J.Zhang & A.J. Freemann, Phys. Rev. B 63 (2000) p. 060101 R
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Dark-field electron holography for the measurement of strain in nanostructures and devices M.J. Hÿtch, F. Houdellier, F. Hüe and E. Snoeck CEMES-CNRS, 29 rue Jeanne Marvig, 31055 Toulouse, France
[email protected] Keywords: strain, high-resolution, holography, semiconductors
We present a new method for measuring strain in nanostructures and electronic devices [1]. It is based on a combination of the moiré technique and off-axis electron holography. A hologram is created from the interference between the diffracted beam emanating from an unstrained region of crystal, which serves as the reference, and a beam from the region of interest containing strained crystal. A typical example for these two regions would be the substrate and an active region of a device. The aim is to measure geometric phase differences, from which the deformation can be calculated [2]. Naturally, any other phase contributions should be minimised, notably, dynamic phases due to thickness variations. For this reason, specimens should be prepared with suitably uniform thickness and regions exhibiting bend contours avoided. The technique has a number of advantages over geometric phase analysis (GPA) of high-resolution images for the study of transistors [3]. The specimens do not need to be so thin, being more like those of conventional TEM. Specimens are therefore easier to prepare and the effects of thin-film relaxation reduced. The major advantage, however, is the ability to analyse large regions of crystal at relatively low resolution. Results will be presented for different strained-silicon devices. TEM specimens are prepared by focussed ion beam (FIB) to thicknesses of about 200 nm. Observations are carried out on the SACTEM-Toulouse, a Tecnai (FEI) 200kV TEM equipped with a Cs corrector (CEOS), rotatable biprism and 2k CCD camera (Gatan). Strain fields are extracted using a modified version of GPA Phase 2.0 (HREM Research Inc.), a plug-in for DigitalMicrograph (Gatan). Typical fringe spacings are 1-2 nm and hologram widths from 300-400 nm allowing lengthwise fields of view of several microns. Figure 1 shows an example of a p-MOSFET with recessed Si80Ge20 source and drain [3]. Holograms were formed using the {111}, (004) and (220) diffracted beams (Figure 2a). The corresponding deformation map for the component parallel to the [220] direction, εxx, (Figure 2b) compares favourably with the result from finite element modelling (Figure 2c). In this case, the measurement precision is 0.2% for a spatial resolution of 4 nm. 1. 2. 3. 4.
M.J. Hÿtch, F. Houdellier, F. Hüe and E. Snoeck, Patent Application FR N° 07 06711. M. J. Hÿtch, E. Snoeck, and R. Kilaas, Ultramicroscopy 74 (1998), p. 131. F. Hüe, M.J. Hÿtch, H. Bender, F. Houdellier and A. Claverie, PRL (2008) accepted. F. Hüe is co-funded by the CEA-Leti. The authors thank the European Union for support through the projects PullNano (Pulling the limits of nanoCMOS electronics, IST: 026828)
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 5–6, DOI: 10.1007/978-3-540-85226-1_3, © Springer-Verlag Berlin Heidelberg 2008
6 and ESTEEM (Enabling Science and Technology for European Electron Microscopy, IP3: 0260019), and IMEC for the device material.
Figure 1. Bright-field image of an array of three dummy p-MOSFET strained-silicon channel transistors with Si80Ge20 sources and drains.
Figure 2. Experimental holographic dark-field: (a) hologram of (220) diffracted beam; (b) corresponding deformation map for εxx; (c) finite element modeling.
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Some device challenges towards the 22nm CMOS technology F. Andrieu, T. Ernst, O. Faynot, V. Delaye, D. Lafond, S. Deleonibus CEA-LETI Minatec, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France francois.andrieu@cea fr (invited abstract) Keywords: CMOS, integrated circuits
Since the 70’s, the transistor cost decreases exponentially thanks to the CMOS technology scaling down. However, this historical scaling is slowing down. Indeed, the IC’s manufacturers face different issues that the device engineers tend to solve thanks to a combination of both new materials and new architectures. The first (historical) challenge is to proceed the performance improvement while managing the dissipated power. Indeed, to achieve high performance devices, the supply voltage has not been reduced in the same proportion as the feature sizes. This has degraded the dynamic power consumption. The static power has increased even more [1]. At the same time, the performance enhancement has been limited by the difficult scaling of the gate oxide thickness (TOX). Strain has first been used, in order to maintain a good trade-off between performance and dissipated power. Stress Memorization Techniques, Nitride Contact Etch Stop Layers (CESL, see Figure 1), embedded SiGe source/drain [1] were integrated in the 65nm node. Moreover, wafer-level strain [2], (110) oriented substrates or Ge-based channels are currently evaluated in order to boost further the ON state currents of the sub-45nm technology. At the same time, new materials have been assessed to reduce the gate leakage. In particular, a combination of a Hf-based high-k dielectrics and a metal gate was introduced in the 45nm INTEL technology [3]. Finally, new architectures, like Fully Depleted thin films (planar, trigate or FINFETs) are evaluated as a solution to limit the source/drain leakage current (Figure 1-2) of sub-22nm devices. The second challenge is the scaling or, at least, the integration density growth, mainly because of the lithography limits. Moreover, even when the lithography techniques enable to draw aggressively scaled devices, it is found that strain is not necessarily as efficient as for longer ones. Finally, their OFF current is difficult to maintain because of the difficult scaling of all the other device dimensions (especially TOX and the junction depth). The historical scaling of the gate length thus tends to slow down. In the future, this trend will limit the integration density, unless new device architectures take over. Indeed, 3D layouts, like multi-channels or multi-fins structures already demonstrated very promising performance and density ([4], Figure 3-4). The third major concern is the variability issue. It is linked to the statistical technological variations (Line Edge Roughness of the gate, Random Dopant Fluctuation of the channel impurities…) that reduce the working window of the devices (e.g. the Static Noise Margin of the sub-45nm SRAM cells) [5]. To conclude, the slowing-down of the CMOS node shift (from 1.5 to 3 years per node) reflects 3 main technological issues: the more and more challenging trade-off S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 7–8, DOI: 10.1007/978-3-540-85226-1_4, © Springer-Verlag Berlin Heidelberg 2008
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between performance enhancement and dissipated power, the difficult increase of the integration density and the variability issues. Microscopy techniques could help device engineers characterizing new materials and new architectures (where thin films are mandatory, cf. Fig.5 and [6]). However, to be relevant, they have to extract local information (about strain, doping, structure, compound, potential…) because the nanometer-range properties will govern the overall 22nm device performance and variability. 1. 2. 3. 4. 5. 6.
S.E. Thompson et al., IEEE Transactions on Electron Devices, 18, 1, p. 26, 2005. F. Andrieu et al., Micoelec. Eng. 84, p. 2047-53, 2007. K. Mistry et al., IEDM Tech. Dig., pp. 24750, 2007. T. Ernst et al., Proc. of ICICDT, 2008. F. Boeuf et al., VLSI Symp., pp. 24-5, 2007. V. Barral et al., IEDM Tech Dig., pp. 61-4, 2007.
TiN/HfO2
TiN/HfO2
20nm
CESL
9nm thin strained Si channel
Burried Oxide
Figure 1. TEM cross section in the electron transport direction of 25nm short FDSOI MOS with CESL and TiN/HfO2 (TSi=9nm). Si 3.4nm
9nm thin strained Si channel
Figure 2. TEM cross section perpendicularly to the electron transport direction of 40nm narrow FDSOI MOS with TiN/HfO2.
SiO2
4.8nm
Figure 3. Cross section of a silicon nanowire obtained by a self-limited oxidation to obtain a sub-5nm channel.
5nm
TiN HfO2
2 5nm thin strained Si channel
Figure 4. The 3D configuration of staked nanowires compensates the pitch-limited current density observed in planar trigate structures.
Figure 5. Integration of a 2.5nm thin strained Si channel MOS with very good performance [6].
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Off-axis electron holography for the analysis of nm-scale semiconductor devices. D. Cooper1, R. Truche1, L. Clement2, S. Pokrant3, and A. Chabli1. 1. CEA LETI - Minatec, 17 rue des Martyrs, 38054 Grenoble, Cedex 9, France. 2. ST Microelectronics, 860 rue Jean Monnet, 38926 Crolles, France. 3. NXP Semiconductors, 860 rue Jean Monnet, 38926 Crolles, France. david.cooper@cea fr Keywords: off-axis electron holography, FIB, semiconductors
The reduction in the size of state-of-the-art semiconductors provides challenges for the characterisation of the doped regions during device development [1]. Off-axis electron holography is a promising TEM-based technique that can be used to provide 2D dopant maps with nm-scale resolution [2]. In this paper we will show how specimens containing nm-scale transistors are prepared using focused ion beam (FIB) milling for examination using off-axis electron holography. Parallel-sided specimens have been prepared using combinations of in situ lift out and back-side milling in order to avoid artefacts such as curtaining which can mask the phase measured in electrical junctions. Finally, low-energy FIB cleaning is used to reduce the thickness of the damaged surface regions on the specimens. Electron holograms have been acquired of state-of-the-art device specimens using a probe corrected FEI Titan electron microscope. The unprecedented electrical and mechanical stability of the Titan microscope allows electron holograms to be acquired for time periods of more than one minute allowing phase images of relatively thick, FIB prepared specimens to be reconstructed with a good signal-to-noise ratio [3]. Figure 1 shows reconstructed phase and amplitude images for a 45 nm gate nMOS device. In the amplitude image, no contrast is visible from the presence of the dopants, however, in the phase image the dopants can be clearly seen. The position of the gate is indicated by the white overlays. In this phase image, as well as the heavily doped regions (HDD), the lightly doped source and drain (LDD) regions can be clearly observed either side of the gate which can allow the electrical gate-width to be measured directly if all of the artefacts are understood. In this paper, we will discuss the suitability of using off-axis electron holography on FIB-prepared semiconductor specimens for dopant profiling. We will highlight many of the artefacts that are observed in phase images, including the effects of specimen thickness on the dopant concentration detection limit and the effects of strain in the doped regions. Finally we will show how electron holography has been applied to a range of samples in the semiconductor industry in order to support process development.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 9–10, DOI: 10.1007/978-3-540-85226-1_5, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3.
International Technology Roadmap for Semiconductors, 2005 ed. http://public.itrs net W.D. Rau, P. Schwander, F.H. Baumann, W. Hoppner and A. Ourmazd. Phys. Rev. Lett. 82, 2614 (1999). D. Cooper, R. Truche, P. Rivallin, J. Hartmann, F. Laugier, F. Bertin and A. Chabli. Appl. Phys. Lett. 91, 143501 (2007).
Figure 1. Shows a phase and amplitude image of a 450-nm-thick specimen prepared using FIB milling containing 45 nm gate nMOS devices.
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Influence of the oxide thickness on the SiO2/Si interface structure P. Donnadieu1, V. Chamard2, M. Maret1, J.P. Simon1 and P. Mur3 1. SIMAP, INPGrenoble-CNRS-UJF, BP 75, 38402 Saint Martin d'Hères – France 2. IM2NP, CNRS - Université Paul Cézanne, 13397 Marseille Cedex 20 France 3. CEA-DRT-LETI-Minatec CEA-GRE 17 rue des Martyrs, 38054 Grenoble Cedex 9, France
[email protected] Keywords: Interface, HRTEM, geometric phase analysis, x-ray grazing incidence diffraction
A key issue in the elaboration of nanodots deposited on substrates is to be able to control their size, density and organisation. In that perspective, major attention has been given to monitor the substrate strain which may be helpful in some case to induce organization. In that context, we studied a currently used substrate: a Si wafer covered by an oxide layer. The typical thickness for such oxide layer is usually in the nanometer range : namely 1 to 10 nm. To characterize the structure and local chemistry of the oxide layer as well as the Si/oxide interface, TEM provides a large number of possibilities. X-ray surface sensitive techniques like Grazing Incidence Diffraction (GID) and reflectivity can also provide information on strain and electron density profile in the near surface region. We report here on a combined TEM and x-ray study carried out on a series of Si wafers covered with oxide layer of different thicknesses. HRTEM associated to the Geometrical Phase Analysis (GPA) method [1] was used to study the substrate deformation as a function of the oxide layer thickness. The samples were prepared by oxidation of 8 inch (100) CZ P type silicon wafers [2]. Prior to oxidation, the wafers are cleaned with an ozone based process. The oxides were elaborated, using a N2/O2 atmosphere at 800°C in a rapid thermal processing machine. According to ellipsometry measurement, the oxide thickness varies from 1.2 nm to 7 nm within the series we have studied. For each Si wafer, cross section samples have been prepared and examined by HRTEM. The images were further analysed by the GPA method. In the numerical analysis, the g(200) reciprocal vector has been selected to measure the displacement of the (200) planes, i.e. planes parallel the interface. Hence the phase map reported here displays the displacement component normal to the substrate surface. Figure 1a and 1b show a HRTEM image and the related GPA map. The profile in the insert gives the displacement as a function of the position along the AB line indicated in Figure 1b. There is a significant displacement in the vicinity of the surface (about 1-1.5 nm) while at distances larger than 1.5 nm, the almost flat profile indicates a negligible displacement. This behaviour has been observed for all samples, regardless of the oxide thickness between 1.2 and 7 nm. Figure 1c gives the measured displacement amplitude as a function of the oxide layer thickness. Error bars have been estimated from the fluctuations of repeated measurements on the phase maps.
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Figure 1. HRTEM image (Fig. 1a) and phase image for (200) g vector (Fig. 1b) (here the 1.2 nm oxide layer). In insert, the phase profile from A to B. Fig. 1c. plot of the total displacement as a function of thickness (the reported oxide thicknesses are measured by HRTEM which slightly differs from the ellipsometry ones). It comes out that the oxide layer thickness strongly influences the strain in the vicinity of the substrate (approximately 1-1.5 nm). Besides, for increasing thickness, larger displacement are measured (Figure 1c). In terms of deformation, it gives about 1 % for the 1.2 nm oxide thickness up to ~ 3 % for the 7 nm thickness. The relation between the oxide layer thickness and the deformation state of the substrate has been confirmed by observations on a two other samples : one with an extremely thin oxide layer (0.6 nm according to ellipsometry) and one with a thick oxide layer (80 nm). For the 0.6 nm oxide, GPA measurements were within error bars because of a too low strain. For the thick oxide layer, numerous dislocations were observed in the substrate close to the oxide layer which is consistent with the relaxation of a high level of strains. The x-ray GID measurement exhibits a 4-fold modulation of the oxide diffraction peak, which follows the Si[011] and the 3 other equivalent surface directions. This modulation, which shows the preferred orientation of the SiO4 tetrahedra at the interface, decays with increasing oxide thickness. Besides the analysis of x-ray reflectivity measurements emphasizes the presence of a dense interfacial layer (density mismatch ~ 8% for the 0.8 nm layer), which disappears with increasing oxide thickness. This combined TEM and X-ray study points out the complex structure of the oxide layer and the strain in the substrate at the vicinity of the interface. Both the oxide layer structure and the subtrate strain state are sensitive to the oxide thickness. It suggests that further nanostructure deposition may be influenced by the oxide layer thickness. This work has been carried out, in the frame of CEA-LETI / CPMA collaboration, with PLATO Organization teams and tools 1. 2. 3. 4.
M. J. Hytch, E. Snoeck, R. Kilaas, Ultramicroscopy 74, (1998) p. 131 P. Mur, M.-N. Semeria, M. Olivier, A.M. Papon, Ch. Leroux, G. Reimbold, P. Gentile, N. Magnea, T. Baron, R. Clerc, G Ghibaudo, Applied Surface Science 175-176 (2001) p. 726 M. Castro-Colin, W. Donner, S. C. Moss, Z. Islam, S. K. Sinha, R. Nemanich, H. T. Metzger, P. Bösecke and T. Shülli, Phys. Rev. B 71, (2005) p. 045310. We kindly acknowledge the ESRF for allocating beamtime and the ESRF ID1 staff for their help during x-ray experiments.
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Challenges to TEM in high performance microprocessor manufacturing H.J. Engelmann, H. Geisler, R. Huebner, P. Potapov, D. Utess, E. Zschech AMD Saxony LLC & Co. KG, CCA, MS E23-MA, D-01109 Dresden, Germany
[email protected] Keywords: Electron Tomography, EELS, Strain Analysis, Dark-Field Diffraction
Smaller structures and new materials require the application of advanced TEM techniques for process control and failure analysis in 45 nm CMOS technology node and beyond. Both imaging TEM and analytical TEM techniques have to be modified or adapted to special questions. There are several reasons for that: A) Device structures that have to be characterized are often located completely within a TEM lamella. For example, typical gate lengths of 45 nm technology node transistors are in the range between 40 and 45 nm. The diameters of respective contacts are smaller than 80 nm. Assuming a standard lamella thickness of 60…80 nm, not all details of a 3D structure can be seen in the 2D projection image anymore. A possible approach to solve this problem is the application of Electron Tomography. While tomographic image acquisition and data treatment have already become a standard technique, sample preparation is still a challenge, especially in case of failure analysis. As an example, Figure 1 shows the 3D reconstruction of a defect in the contact area. Missing silicide caused an increased electrical resistance in that case. B) The application of new materials requires the characterization of their properties in dependence on deposition and treatment parameters. For example, low-k dielectric materials which are used to reduce the cross-talk between Cu interconnects show changes in chemical composition caused by plasma etch processes. The resulting kvalue increase has to be measured in the direct neighbourhood of etched structures like trenches and vias, with a spatial resolution better than 5 nm. While changes in chemical composition are analyzed by EELS, direct measurement of the k-value can be done by Valence EELS. A procedure was developed which allows determining the 1014 Hzfrequency dielectric permittivity [1]. Even though this is not the k-value corresponding to the GHz-frequency range used in microprocessors, relative changes in the dielectric constant can be detected very precisely (Figure 2). Ultra low-k (ULK) materials that are expected to be introduced for the 32 nm CMOS technology node will contain pores. Local pore size/pore distribution characterization will be another challenge for TEM. C) The introduction of ‘strained silicon’ into the channel region of transistors requires advanced characterization techniques. Mechanical stress results in a distortion of the silicon lattice which affects the electronic band structure, allowing improvements in carrier mobility. For process control and next technology node transistor development, local strain measurements in the Si MOSFET channel are needed. Nano Beam Diffraction (NBD) is an analysis technique that uses a small probe electron beam with reduced convergence angle to produce diffraction patterns with smaller spots than in CBED patterns [2]. The lattice parameter can be determined from the positions of the S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 13–14, DOI: 10.1007/978-3-540-85226-1_7, © Springer-Verlag Berlin Heidelberg 2008
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spots which allows strain quantification. The challenge in this technique is the NBD pattern analysis for very precise lattice spacing determination which is needed for strain quantification. Figure 3 shows the relative change of the Si lattice spacing in direction in the channel region of a tensile strained NMOS transistor. D) With shrinking of structure sizes new questions arise regarding product reliability. For example, the Cu microstructure becomes more and more important with decreasing dimensions of the interconnect lines. Grain size, grain orientation and twin formation can influence the stability against electromigration/stress migration to a high degree. So far, the EBSD technique has been used to characterize the Cu microstructure. Since agglomerates of small Cu grains are expected to be a reliability concern for the 32 nm CMOS technology, grains with sizes below 40 nm have to be analyzed which requires a TEM-based technique. Dark-Field Diffraction Circular Scanning with subsequent diffraction pattern reconstruction can be used to produce grain orientation maps. As an example, Figure 4 shows a [001] inverse pole figure map of a Cu interconnect stack. Further challenges exist in the field of sample preparation: Quality/precision, target preparation for failure analysis/defects and sample throughput. 1. 2. 3.
P. Potapov, H.J. Engelmann, E. Zschech, M. Stöger-Pollach, submitted to Micron. H.J. Engelmann, S. Heinemann, E. Zschech, Proc., IMC 16, Sapporo, Sept. 3-8, 2006. We kindly acknowledge the financial support by the German BMBF, FKZ 13N9431
Figure 1. 3D reconstruction of a defect in the contact area
Figure 2. Relative change in dielectric constant in surface-plasma treated low-k material (covered with a Cr layer).
Figure 3. Relative change of Si lattice spacing in NMOS channel region
Figure 4. Inverse pole figure map of a Cu interconnect stack
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Strain study in transistors with SiC and SiGe source and drain by STEM nano beam diffraction P. Favia1, D. Klenov2, G. Eneman1,3, P. Verheyen1, M. Bauer4, D. Weeks4, S.G. Thomas4 and H. Bender1 1. IMEC, Kapeldreef 75, 3001 Leuven, Belgium 2. FEI, Achtseweg Noord 5, 5651 GG Eindhoven, The Netherlands 3. K U Leuven, ESAT/INSYS, and Fund for Scientific Research-Flanders, Belgium 4. ASM America, 3440 E. University Dr., Phoenix, AZ 85034, USA paola
[email protected] Keywords: nano-beam diffraction, strain, SiGe, SiC
Strain is introduced in the fabrication of complementary metal-oxide-semiconductor devices to enhance their channel region carrier mobility [1]. Epitaxial Si1-xGex (1530at% Ge) or Si1-xCx (1-2at% C) are typical stressor materials. As Ge has a 4% larger lattice constant (0.566 nm) than Si (0.543 nm), Si1-xGex deposited in the source/drain (S/D) regions will induce compressive strain in the Si channel, while Si1-xCx in the S/D will induce tensile strain in the channel [2]. Nano-beam diffraction (NBD) is a TEM-based technique that allows to obtain a diffraction pattern from small regions and, as a result, to measure directly local lattice parameter and thus to quantify 2-D strain. NBD uses a small diameter electron probe which determines the lateral resolution (> Ce0 50Zr0 38Tb0 12O2 >>> Ce0 62Zr0 38O2). The electron microscopy studies (HREM,
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HAADF, XEDS and EELS) of the high-temperature treated catalysts showed that the Ru particles keep the same epitaxial relationship already observed after reduction at low temperature, defined by Ru(002) || Ce0 8Tb0 2O2(1-11) and Ru[-2-10] || Ce0 80Tb0 20O2 [2-11]. Likewise such studies also evidenced that the chemical composition in the pedestals, i.e. in the regions of the oxide close to the metallic nanoparticles, was comparable to that observed in other areas of the support, Figure 2. 1. 2. 3. 4.
D. L. Trimm, Z. I. Onsan, Catalysis Reviews-Science and Engineering 43 (2001), 31. R. Lanza, S.G. Järås and P. Canu, Applied Catalysis A: General 325 (2007), 57. S. Bernal, G. Blanco, J.J. Calvino, C. López-Cartes, J.A. Pérez-Omil, J.M. Gatica, O. Stephan and C. Colliex, Catalysis Letters 76 (3–4) (2001), 385. We acknowledge the financial support from Ministry of Education and Science of Spain (MAT2005-00333) and Junta de Andalucia (FQM334, FQM110). Electron microscopy imaging was carried out in the Central Service of Science and Technology from Universidad de Cadiz.
[010] (002) 61º (002)
61º
(101)
(100)
(101)
(101)
[110] (111)
2 nm Figure 1. HREM image of a Ru(1%)/Ce0 62Zr0 38O2 catalyst reduced under H2 at 1173K. The structural analysis shows an epitaxial relationship between the metallic phase and the support defined by Ru(002) || Ce0 62Zr0 38O2(-111) and Ru[010]|| Ce0 62Zr0 38O2[110]. 1000
800 Ru
Ce
O
600
Tb Tb 400
Ce
pedestal
Ce Tb
O
Cu
Tb
Ce Ce
Tb Tb
200
Cu
particle
Ce Tb
support
Cu
Tb
0 0
2
4
keV
6
8
10
Figure 2. (Left) HREM image of a Ru(1%)/Ce0 8Tb0 2O2 catalyst reduced under H2 at 1173K showing pedestal-like nanostructures. Note the epitaxial relationship between the metallic phase and the support. (Center) XEDS compositional analysis showing similar chemical composition in the pedestal areas and the support. (Right) HREM image showing the distance between the metal and the oxide in the interface.
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Characterisation of materials with applications in the photocatalytic activation of water N.S. Hondow1, R. Brydson1, Y.H. Chou2 and R.E. Douthwaite2 1. Institute for Materials Research, University of Leeds, Leeds, LS2 9JT, United Kingdom 2. Department of Chemistry, University of York, York, YO10 5DD, United Kingdom n
[email protected] Keywords: photocatalyst, TEM, STEM
The requirement for the development of alternative fuel sources is highlighted by the limited supplies of fossil fuels and the environmental impact caused by their extensive use. The conversion of solar energy is a particularly desirable option, with the photocatalytic conversion of water into hydrogen and oxygen representing an attractive source of fuel. An ideal material suitable for this type of reaction has yet to be reported, though several promising developments have been made. An ideal photocatalyst would exhibit certain characteristics, such as long term stability, optimum absorption of the solar spectrum, and the ability to oxidise and reduce water to O2 and H2 respectively. Stable oxide semiconductors have shown the best results, with the possibility of manipulating the valence and conduction bands of the materials through alteration of the composition and structure allowing the required redox reactions. However, the materials developed at present generally have an overall low efficiency as they only utilise the high energy UV periphery of the solar spectrum [1]. This therefore creates the need for either the development of new materials, or the further improvement of these known systems. In either case, the application of alternative synthetic methods may lead to the formation of new phases and novel materials. One such synthetic route currently being investigated is that of using microwave-induced plasma promoted dielectric heating [2]. Materials currently being made include titanates, niobates and tantalates. The morphology and crystallinity of the samples can directly affect the performance of the materials as photocatalysts. It is important that sites for the oxidation and reduction reactions are separated so as to prevent recombination. The morphology of the materials has been examined by field emission SEM and the crystalline structure of the materials has been confirmed using conventional high resolution phase contrast TEM and selected area diffraction. This has been examined at several points throughout the catalyst development, including before and after catalytic testing, enabling observations as to the stability of the materials. Attempts to increase the catalytic performance has led to the introduction of further elements into the systems being investigated, with particular interest in the formation of metal or nitrogen/oxynitride rich surface regions [3,4]. The key factors in how these materials will perform as photocatalysts for the splitting of water include the
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distribution and composition of these added particles. This has been analysed by high angle annular dark field STEM imaging in combination with EDX mapping. Initial studies have found that some of the conditions the materials are subjected to leads to phase separation of the metals present, rather than the desired surface doping (Figure 1). Further elemental composition, including lighter elements, has also been determined using EELS, and supporting chemical bonding information has been obtained by the surface sensitive technique XPS. 1. 2. 3. 4.
A. Kudo, H. Kato and I. Tsuji, Chemistry Letters 33 (2004), p. 1534. R.E. Douthwaite, Dalton Transactions (2007), p 1002. A. Kudo, R. Niishiro, A. Iwase and H. Kato, Chemical Physics 339 (2007), p 104. Y. Lee, H. Terashima, Y. Shimodaira, K. Teramura, M., Hara, H. Kobayashi, K. Domen and M. Yashima, Journal of Physical Chemistry C 111 (2007), p. 1042.
Figure 1. High angle annular dark field STEM image (left) and EDX maps (right) of NiTa2O6 after attempts at nitrogen doping by reduction in ammonia at 750 oC. EDX maps of Ni (top right) and Ta (bottom right) show that phase separation of the metals has occurred in some particles.
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Complementary EM study on highly active nanodendritic Raney-type Ni catalysts with hierarchical build-up U. Hörmann, U. Kaiser1, N. Adkins2, R. Wunderlich3, A. Minkow3, H. Fecht3, H. Schils3, T. Scherer4 and H. Blumtritt5 1. Ulm University, Electron Microscopy Group of Materials Science, Albert-Einstein-Allee 11, 89081 Ulm, FRG 2. Ceram, Queens Road, Penkhull, Stoke-on-Trent, ST4 7LQ, Great Britain 3. Ulm University, Institute of Micro- and Nanomaterials, Albert-Einstein-Allee 47, 89081 Ulm, FRG 4. Forschungszentrum Karlsruhe, Institute of Nanotechnology, PO Box 3640, 76021 Karlsruhe, FRG 5. Max-Planck-Institut für Mikrostrukturphysik, Weinberg 2, 06120 Halle/Saale, FRG
[email protected] Keywords: Raney-type Ni, dendrites, structure, Cs-corrected HRTEM, gas atomisation, slicing view
Nanostructured Raney-type Ni catalysts have been used in industry since the 1920s for the production of a wide range of chemicals. [1] In the EU supported project IMPRESS it has been shown that by using gas atomisation processing high surface area particles with significantly increased catalytic activity in hydrogenation reactions can be produced. [2,3] Structural investigations with complementary methods of electron microscopy in combination with X-ray powder diffractometry have enabled the link between processing, structure and catalytic activity to be explored. [4] Raney-type Ni catalysts were produced from alloy powder prepared by gas atomisation. After activation by leaching with NaOHaq and prior to the structural investigations the samples were passivated with oxygen. Size selected microparticles of ca. 100 µm size, grown from different melt compositions were chosen for this study. The microstructure of the samples was characterised in 2D by light microscopy and by SEM, see Fig. 1, and SEM EDX mappings. The nanostructure was investigated with HRTEM and Energy filtered TEM for elemental mappings (Ni, Al) using a Cs-corrected FEI 80-300 Titan microscope operated at 300kV. The use of a dual-beam FIB SEM for sample preparation allowed the investigation of one particular nanodendrite on different scales, first within the microparticle by SEM and hereafter as a single cut lamella in the TEM. In order to correlate the local structure with integral measurements, X-ray powder diffractometry was also carried out. The 3D interconnection of the nanodendrites, which build up the whole particle was imaged with slicing view by using a FIB SEM. The resulting porous particles were found to be built-up of nanodendrites. The thickness of the dendrites decreases with increasing Al content. The samples with the finest dendrites were obtained from Ni-75%Al alloy powder, i.e. from an alloy with a higher Al content than the one which is used to produce the standard commercial catalysts. The dendrites consist of two adjacent phases, from which one after leaching and passivation is transformed into NiO. This phase is located at the dendrite tips, and S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 217–218, DOI: 10.1007/978-3-540-85226-1_109, © Springer-Verlag Berlin Heidelberg 2008
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might offer the reactive sites for the catalytic reaction. The complex structure was characterised by Cs-corrected HRTEM. On the mesoscale it shows a polycrystalline framework structure with filled mesopores. The nanocrystals within the mesopores clearly reveal texture. The outer surface of the dendrite tips shows nanosteps, which increase significantly the surface area provided for the catalytic reaction, see Fig. 2. 1. 2. 3. 4.
M. Raney: US patent 1563587 (1925) A.M. Mullis, N.J. Adkins, Z. Huang, R.F. Cochrane, Proc. 3rd International Conference on Spray Deposition & Melt Atomization, 2006, Bremen, Germany, CD proceedings F. Devred et al., to be published. U. Hörmann, U. Kaiser, N.J.E. Adkins, R. Wunderlich, A. Minkow, H. Fecht, H. Schils, F. Devred, B. Nieuwenhuys, H. Blumtritt, submitted to 9th International Conference on Nanostructured Materials, Nano 2008, Rio de Janeiro, Brazil (2008)
Figure 1. Left: Light microscopy image of a microparticle from the 75 – 106 µm size fraction with a high inner porosity due to the nanodendritic structure. Right: SEM image of a single dendrite with capped tips (light grey), see arrows.
Figure 2. Left: Dark field micrograph of the interface between the caps and the dendrite backbone. Right: Cs-corrected energy-filtered HRTEM micrograph of the mesopores at the outer rim of the capped dendrite tips, showing the pores and the surface steps.
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Structural properties of sol-gel synthesized Li+-doped titania nanowhisker arrays U. Hörmann1, J. Geserick2, S. Selve1, U. Kaiser1, and N. Hüsing2 1. Ulm University, Electron Microscopy Group of Materials Science, Albert-Einstein-Allee 11, 89081 Ulm, FRG 2. Ulm University, Institute of Inorganic Chemistry I, Albert-Einstein-Allee 11, 89081 Ulm, FRG
[email protected] Keywords: titania, anatase, rutile, nanowhiskers, mesoporous material, HRTEM, sol-gel synthesis
Nanostructured titania is of particular interest for applications in photo-catalysis due to its high catalytic activity. Moreover, these structures are of particular interest for many applications due to their electronic properties, e.g. anti-reflection layers, sensors, vacuum microelectronics. The band gap of the nanoscaled semiconducting anatase is size dependent. The band gap increases in the size range of 15 nm to 3.9 eV [1], compared to the bulk value of 3.2 eV [2], suggesting already a quantum confinement effect. Nanowhiskers, grown in even smaller dimensions as in this study are prospective candidates for showing a transition to the quantum confinement effect. Sol-gel synthesis of mesoporous oxides relies on the self-assembly of the structure directing agents, the surfactants which rule the solidification or crystallisation of the inorganic oxide. Mesoporous ordered solids produced by sol-gel processing are e.g. monolithic SBA-15 type silica networks or the highly catalytically active mesoporous titania powders. In this study titania nanowhisker arrays with a high surface area were produced. Anatase nanowhiskers were grown in a sol-gel process [3] using ethylene glycol modified titanium(IV) (EGMT) as the titania precursor and Lithiumdodecyl sulphate (LDS) as a structure directing agent. The LDS simultaneously delivers the Li+ as a dopant. After synthesis the samples were dried and calcined in order to remove the surfactant. The samples were characterised by X-ray powder-diffractometry. The dried samples were found to consist of an approximately proportionate equal mixture of rutile and anatase. The high temperature phase of titania grows thus even under room temperature processing. After calcination at 400 °C for 4 h, the intensity of the anatase peak grew significantly, indicating a higher anatase ratio. After calcination nitrogen sorption measurements were performed in order to determine the average pore sizes as well as the specific surface areas. The resulting samples were investigated by HRTEM. The samples showed a strong growth anisotropy, i.e. the whiskers revealed a high aspect ratio. The diameter of the whiskers measures approximately 3 – 4 nm and the length up to 50 nm even after calcination. The titania crystals grew as whiskers with a preferential orientation. These needles are aggregated to radial bunches, thus forming nanowhisker arrays.
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Calcination caused the formation of roundish nanoparticles on top of the whisker tips. These crystalline particles are associated with an amorphous phase and were attributed to the anatase phase. 1. 2. 3.
T. Toyoda, Ikumi Tsuboya, Rev. Sci. Instrum. 74 (1) (2003), p. 782 W. Wunderlich, L. Miao, M. Tanemura, S. Tanemura, P. Jin, K. Kaneko, A. Terai, N. Nabatova-Gabin, R. Belkada, Int. J. Nanoscience 3 (4&5) (2004), p. 439. J. Geserick, N. Hüsing, R. Roßmanith, C.K. Weiß, K. Landfester, Y. Denkwitz, R.J. Behm, U. Hörmann, U. Kaiser MRS Spring Meeting 2007
Figure 1. Left: Survey of nanowhisker arrays in the sample after calcination for 4 h at 400 °C. Right: HRTEM micrograph showing the whiskers with some roundish particles.
d(101) = 0.35nm Figure 2. Left: Roundish particles observed after calcination for 4 h at 400 °C. Right: HRTEM micrograph. Detail of one particle of Fig. 2 left, identified as anatase with 0.35 nm d-spacing of the (101) plane.
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Quantitative strain determination in nanoparticles using aberration-corrected HREM C.L. Johnson1, E. Snoeck1, M. Ezcurdia1, B. Rodríguez-González2, I. Pastoriza-Santos2, L.M. Liz-Marzán2 and M.J. Hÿtch1 1. CEMES-CNRS, 29, rue Jeanne Marvig, 31055 Toulouse, France 2. Departamento de Quimica Fisica, CSIC, University of Vigo, 36310 Vigo, Spain
[email protected] Keywords: strain, nanoparticles, aberration correction, high-resolution electron microscopy
Metallic nanoparticles exhibit exceptional optoelectronic properties that are strongly size and shape dependant and locally variable. Recently, novel synthesis techniques have enabled precise control over the growth of metallic nanoparticles, occasionally resulting in morphologies that cannot be characterized using standard techniques [1]. One example is five-fold-twinned decahedral Au nanoparticles. Owing to the decahedral geometry, these nanoparticles must be strained or contain defects and models have been proposed to predict their strain states. We examined the internal structures of decahedral Au nanoparticles using a combination of aberration-corrected HREM, strain mapping, and finite-element analysis [2,3]. HREM images (Figure 1) were obtained using the SACTEM-Toulouse, a Tecnai F20 ST (FEI) equipped with an imaging aberration corrector (CEOS), rotatable electron biprism and a 2K CCD camera (GATAN). Strain analysis was done using DigitalMicrograph (GATAN) and the GPA Phase 2.0 (HREM Research) software. Microscope distortions were calibrated to obtain highly accurate (< 0.1% strain), highspatial-resolution (< 1 nm) maps of the lattice strain and rotation in the decahedral nanoparticle. Aberration correction provides high-contrast images necessary for accurate high-resolution strain determination. The strain mapping revealed that internal rigid-body rotations (Figure 2a) combined with shear strains (Figure 2b) accommodate the geometric constraints imposed by the decahedral geometry. Our measurements confirm, for the first time, the existence of a disclination. Furthermore, comparison of the results to finite-element analyses revealed that shear strains, which are not predicted by the commonly accepted strain models for decahedral particles, result from elastic anisotropy. The internal structure of these complex nanoparticles will determine their growth and stability as well as affect their surface structures, and, therefore, will be of great importance for engineering their electronic and optical properties. 1. 2. 3. 4.
A. Sánchez-Iglesias et al, Advanced Materials 18 (2006) p. 2529. M.J. Hÿtch et al, Ultramicroscopy 74 (1998), p. 131. C.L. Johnson et al, Nature Materials 7 (2008) p. 120. We thank the EU Integrated Infrastructure Initiative ESTEEM (Ref. 026019 ESTEEM) and the Spanish Ministerio de Educacion y Ciencia (Grants No. MAT2004-02991 and NAN2004-08843-C05-03) for support.
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Figure 1. Aberration-corrected HREM image of decahedral Au nanoparticle. (a) The image shows the 5-fold rotational symmetry marked by twin boundaries that intersect at the centre of the particle. (b, c) Enlarged views of the core and edge of the particle.
Figure 2. (a) Internal rigid-body rotation of the crystallographic lattice and (b) shearstrain distributions in the decahedral Au nanoparticle. The lattice rotation combined with the shear strains, which result from elastic anisotropy, accommodate the unique geometry of the decahedral particle.
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Morphological characterization by HRTEM and STEM of Fe3O4 hollow nano-spheres Alfonso Ibarra1, Gerardo F. Goya1, Jordi Arbiol3, Enio Lima Jr.4, Hercílio Rechenberg4, Jose Vargas5, Roberto Zysler5 and M. Ricardo Ibarra1,2 1. Aragon Nanoscience Institute (INA), 2. Materials Science Institute of Aragon (CSICZaragoza University), University of Zaragoza, 50009 Zaragoza, (Spain), 3. TEM-MAT, Serveis Cientificotecnics, UB, 08028 Barcelona, (Spain) 4. LMM, University of São Paulo (Brazil), 5. Centro Atomico Bariloche, 8400, S. C. Bariloche (Argentina)
[email protected] Keywords: Fe3O4 nanoparticles, HRTEM, STEM, HAADF, EELS, EFTEM
Morphology, surface and finite size effects in magnetic nanoparticles have been the subject of growing interest in recent years from both experimental and theoretical point of view [1]. The magnetic properties are strongly associated with the morphological and structural homogeneity of the nanoparticles [2]. Interparticle interactions also play an important role in the magnetic behaviour of an ensemble of nanoparticles, which differs from that of non-interacting systems [3]. The aim of this work is the characterization, by means of transmission electron microscopy (TEM), of the morphology and structure of Fe3O4 nanoparticles prepared by chemical route [4] in order to understand their magnetic behaviour. TEM specimens were prepared dispersing the nanoparticles in toluene and dropping this colloidal solution onto a carbon-coated copper grid. TEM analyses were performed in a JEOL 1010 (200 kV). Interesting enough is to point out that a deeper High Resolution TEM (HRTEM) and STEM analysis combined with Energy Filtered TEM (EFTEM) as well as high angular annular dark field (HAADF or Z-contrast) show that magnetite nanoparticles interact creating hollow nano-spheres, and thus affecting the magnetic behaviour of the sample. Figure 1 shows a general view of the sample where the projection reveals a toroidallike shape nanostructures with a weak interaction between them. Electron energy loss spectroscopy (EELS) analyses show that the nanostructures are constituted by Fe3O4, which is corroborated by a chemical analysis where the ~60 % wt. of the final powder corresponds to the Fe3O4 nanoparticles, and ~40 % wt. corresponds to the organic cap of oleic acid that covers the particles and avoid their agglomeration. A more detailed analysis by HRTEM, Figure 2, shows that depending on defocus; new crystallographic planes appear “inside” the projected toroids, indicating that the observed nanostructured may correspond to hollow spheres instead of toroids. In order to confirm this assumption, EFTEM Fe Maps (Figure 3), were obtained showing the homogeneity of the Fe around the whole surface of the sphere. The inferred morphology of hollow spheres is reported here for the first time in magnetic nanoparticles, and it is intimately related to novel magnetic properties displayed by these samples.
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1. 2. 3. 4. 5.
D. Fiorani., Surface effects in magnetic nanoparticles, Springer, New York (2005). C. B. Murray, C. R. Kagan and M. G. Bawendi, Annu. Rev. Mater. Sci. 30, 545 (2000). J. L. Dormann, E. D'Orazio, F. Lucari, E. Tronc, P. Prené, J. P. Jolivet, D. Fiorani, R. Cherkaoui and M. Nogués, Phys. Rev. B 53, 14291 (1996). J. M. Vargas and R. D. Zysler, Nanotech. 16, 1474 (2005); J. M. Vargas, W. C. Nunes, L. M. Socolovsky, M. Knobel and D. Zanchet, Phys. Rev. B 72, 184428 (2005). This work has been supported by the Spanish Projects Nanoscience Action NAN200409270C3-1/2 and Consolider Ingenio CSD2006-00012. GFG acknowledges support from the Spanish MEC through the Ramon y Cajal program
Figure 1. TEM micrograph of the sample where the projection of nanoparticles seems toroidal structures.
b)
c) 2 nm
a) 5 nm
Figure 2. a) HRTEM micrograph where crystalline planes are observed forming hollow spheres. b) FFT of the area. Welldefined rings show the polycrystalline character of the sample. c) Magnified image of a nanoparticle where the Fe3O4 microstructure is observed.
Figure 3. EFTEM Fe map. The presence of Fe forming the hollow sphere is observed.
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Direct observation of surface oxidation of Rh nanoparticles on (001) MgO N.Y. Jin-Phillipp, P. Nolte, A. Stierle, P.A. van Aken, and H. Dosch Max Planck Institute for Metal Research, Heisenbergstr.3, D-70569 Stuttgart, Germany
[email protected] Keywords: surface oxidation, nanoparticles, HRTEM, EELS, Rhodium
The late transition metals have been studied extensively for decades because of their catalytic applications. Understanding the oxidation behaviour and the structure of the oxides of these metals are essential in order to raise the efficiency of the catalysts [1]. In the present contribution we investigate surface structure of Rhodium (Rh) nanoparticles grown on (001) MgO and oxidized at oxygen (O2) pressure of 2x10-5 mbar by highresolution transmission electron microscopy (HRTEM) in both (110) and (100) crosssections, and by spatially-resolved electron energy-loss spectroscopy (EELS). Surface layer with a structure different from Rh fcc-structure may be clearly seen at the Rh (1 1 1) surface at the top-right side of the particle. As indicated in Figure 1 the measured spacing between the surface layer and the Rh (1 1 1) top layer is 0.28nm, markedly higher than d111,Rh of 0.220nm, measured from the core of the particle. The distance between the image points in the surface layer along Rh is 0.27nm. Similar measurements have been carried out for particles without any surface layer, and it is found that even for the very small Rh particles of a size of ~ 2nm the error is within ±0.005nm. This confirms, that our observation of larger spacing of the surface layer is not due to the deviation of the lattice spacing measurement found in the case of randomly oriented small particles [2]. Spatially-resolved EELS line-scans were performed across {111} surfaces of Rh particles free of epoxy. Figure 2(a) illustrates one of such line scans. Figure 2(b) shows the background subtracted energy-loss near-edge structure (ELNES) of the Rh-M edge of selected spectra. A small extra peak, marked with an arrow, is detected in the spectrum 3 taken at the surface. This small peak lies at the energy position of ~532eV, 11eV distant from Rh-M2 peak, and is therefore the O-K edge. This result suggests that the surface layer observed by HRTEM is surface oxide formed during oxidation. Image simulation using a theoretical model of the surface oxide obtained by density functional theory (DFT) [3] suggests a hexagonal trilayer of O-Rh-O at the Rh (111) surface of the nanoparticles. 1. 2. 3.
H. Over, Y.D. Kim, A.P. Seitsonen, S. Wendt, E. Lundgren, M. Schmid, P. Varga, A. Morgante, and G. Ertl, Science 287 1474 (2000). J.-O. Malm and M.A. O’Keefe, Ultramicroscopy, 68, 13 (1997). J. Gustafson, A. Mikkelsen, M. Borg, E. Lundgren, L. Köhler, G. Kresse, M. Schmidt, P. Varga, J. Yuhara, X. Torrelles, C. Quirós, and J.N. Andersen, Phys. Rev. Lett. 92, 126102 (2004).
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4.
Financial support from the European Union under Contract No. NMP3-CT-2003-505670 (NANO2) is acknowledged.
Figure 1. High-resolution micrograph of a Rh nanoparticle on the (110) cross-section, showing the surface oxide layer at Rh (1 1 1) surface.
Figure 2. (a) EELS line-scan across the {111} surface of a Rh nanoparticle with a spacing between the spectra of 0.3nm. The scan started at vacuum (spectrum 1) and ended inside the particle (spectrum 7), (b) Selected EELS spectra (3-5) are shown, and an O-K peak is found at the surface (spectrum3).
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Characterization of catalyst poisoning in biodiesel and conventional diesel fuelled vehicles T. Kanerva1, K. Kallinen2, Toni Kinnunen2, M. Vippola1 and T. Lepistö1 1. Tampere University of Technology, Department of Materials Science, P.O.Box 589, FIN-33101 Tampere, Finland 2. Ecocat Oy, Typpitie 1, FIN-90650 Oulu, Finland tomi kanerva@tut fi Keywords: TEM, catalyst, biodiesel, poisoning
Demand for lower and lower emissions in road transportation has promoted the development of more efficient exhaust emission catalysts. On the same time the fight against the impact of transportation on climate change has opened the way for the use of biofuels, e.g. biodiesel. Deactivation of catalytic surfaces is a serious problem in the design of more efficient automotive exhaust catalysts. Deactivation of catalysts can be classified in three types: chemical (e.g. poisoning), mechanical (e.g. fouling) and thermal (e.g. ageing). In the long run these deactivation processes can cause nearly total loss of catalytic activity in the catalyst material. In biodiesel fuelled vehicles these processes can lead to notably different effects in catalyst efficiency compared to those of conventional diesel vehicles [1]. In this study typical diesel catalyst with noble metals Pt and Pd was vehicle-aged using two different fuels: conventional diesel (EN590) and biodiesel (RME, rapeseed methyl ester). Samples were studied with analytical transmission electron microscope (TEM) and field emission scanning electron microscope (FEG-SEM), both equipped with energy dispersive x-ray spectrometer (EDS). Catalyst poisons from vehicle-ageing were analysed after conventional diesel and biodiesel use. Characterization included EDS-mapping, spot analyses and imaging. According to the results different types and contents of poisons were found in the samples depending on the fuel used. Poisons and their contents are presented in tables 1 and 2. In conventional diesel sample typical poisons were S and K with contents of around 0.5 wt%. In biodiesel samples the highest contents were for poisons S, K and Zn, with much higher proportions. Also overall number of poisons was higher in biodiesel sample. Locations of some EDS-spot analyses for RME and EN590 is presented in figures 1 and 2 respectively. In this vehicle-ageing microstructural effects were minor and no detectable effects were found. In this study the conventional diesel samples were less poisoned and there was not any effect on the performance of this catalyst. The vehicle-ageing using RME caused significant loss of efficiency in the catalyst. Higher contents of poisons and higher number of poisons in vehicle-aged samples using biodiesel rises questions of the quality of biodiesels. There is a lot of research going on to gain more knowledge on the properties and behaviour of biodiesels in automotive engines. Further studies in the
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effects of biodiesel use on the catalyst components is required to meet the future demands of international energy policies [2]. 1. 2.
J.B. Butt and E.E. Petersen, Activation, Deactivation and Poisoning of Catalysts (Academic Press Inc., 1988). M. Lapuerta, O. Armas and J. Rodriguez-Fernadez, Progress in Energy and Combustion Science 34 (2007), p. 198.
Table 1. Poisons detected in vehicle-aged RME-sample. Contents in wt%. RME
EDS 1 EDS 2 EDS 3 EDS 4 EDS 5 EDS 6 EDS 7 EDS 8
S
0.8
0.2
0.6
0.9
0.1
1.5
1.3
0.2
K
2.4
0.3
4.3
3.7
0.5
2.5
0.9
0.3
0.3
0.1
Ca
0.1
Cr
0.2
0.1
Fe
0.1
0.1
6.2
0.4
0.1
0.3
0.3
Zn
0.5
0.1
0.5
1.1
0.1
0.7
0.9
0.1
Table 2. Poisons detected in vehicle-aged EN590-sample. Contents in wt%. EN590 EDS 1 EDS 2 EDS 3 EDS 4 EDS 5 EDS 6 EDS 7 EDS 8 S
0.5
K Fe Zn
0.2
0.4
0.6
0.4
0.4
0.2
0.1
0.3
0.1
0.1 0.2
Figure 1. Area of RME EDS 1.
0.1
0.1 0.2
0.1
Figure 2. Locations of EN590 EDS 4,5 and 6.
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TEM Characterisation of Highly Luminescent CdS Nanocrystals Hadas Katz1, Alexey Izgorodin2, Douglas R. MacFarlane2, Joanne Etheridge1,3 1. Dept of Materials Engineering, Monash University, Clayton, Victoria, 3800, Australia 2. ARC Centre of Excellence for Electromaterials Science, Monash University, Clayton, Victoria, 3800, Australia 3. Monash Centre of Electron Microscopy, Monash University, Clayton, Victoria, 3800, Australia hadas katz@eng monash.edu.au Keywords: TEM, Cadmium Sulphide, nanocrystal, electroluminescence.
Luminescent II-VI semiconductor nanocrystals have been the focus of many studies in recent years due to their low energy consumption, wide variety of electroluminescence properties and a large number of combinations of core/shell materials that can be synthesized by the simple and cost efficient reverse micelle method [1]. The size of the nano-crystal determines its surface to volume ratio [2], which affects the band gap and hence the luminescence properties of the crystal. The crystal structure and defect structure, such as point defects and dislocations, also affect band gap energy [3] and hence wavelength of the emitted light as well as the stability of the luminescence materials over time. Characterizing composition, nanocrystals size, crystal structure, defect structure and atomic bonding in the atomic and even subatomic level by electron microscopy will enable us to understand how their size and atomic structure would affect the wavelength of the emitted light and stability of the luminescence materials over time. Those are important factors in the engineering of luminescence materials with desired properties. This work presents an electron microscopy and diffraction study of the crystal structure of highly luminescence CdS nanocrystals produced using the reverse micelles method. Energy-dispersive X-ray spectroscopy (EDX), selected area diffraction (SAD) and atomic resolution imaging using an analytical JEOL 2011 TEM fitted with a LaB6 filament were used to determine the CdS nanocrystal’s composition, crystallographic structure, defect structures and size with a view to understanding how these affect the stability, band gap energy and luminescence properties. Nanoparticles were observed both aggregated in clusters and distributed across the carbon film. The size of the nanoparticles is typically between 3-13nm and was determined by counting of atomic planes in atomic resolution images. Selected area diffraction (SAD) patterns taken from filtered solution indicate the presence of hexagonal CdS and cubic CdS only. However, in unfiltered solutions cubic
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CdO and cubic Na2S nanocrystals were also observed and are assumed to be byproducts. Careful measurement of a SAD pattern of numerous particles distributed across the carbon film shows a 4% difference in the nominal cubic 220 and -220 spacing, suggesting the cubic CdS structure is distorted. This could have an affect on band gap and electroluminescence properties. Using high resolution images and their Fourier transforms, it was conformed that both cubic and hexagonal CdS nanocrytals and a cubic CdO by–product are produced using the reverse micelle process (e.g. figure 1). In addition, a high resolution image of a small cluster containing 3 cubic CdS nanocrystals shows that these particles share a common atomic plane and might had inter-grown during growth process. 1. 2. 3.
D. R. Vij, Handbook of Electroluminescent Materials, Institute of Physics Publishing, (2004). S. J. Rosenthal, J. McBride, S. J. Pennycook, L.C. Feldman, Surface Science Reports, 62 (2007) 111-157. Ronghui Xu, Yongxian Wang, Guangqiang Jia, Wanbang Xu, Sheng Liang, Duanzhi Yin,. Journal of Crystal Growth, 299 (2007) p. 28-33.
Figure 1. (a) HRTEM Image of A small cluster of 4 CdS nanocrystals supported on a carbon film. (b) Fourier transform taken of the image. 3 nanoparticles have cubic structure and are oriented down the and zone axis and share the {31-1} atomic plane (fig. b1). The fourth nanoparticle has hexagonal CdS structure (fig. b2). (c) Inverse of the Fourier transform with diffuse background masked to enhance atomic contrast. The orientations of the 3 cubic CdS nanocryatls sharing the {31-1} atomic plane (fig c1) and of the hexagonal CdS nanocrystal atomic plane (fig. c2) are marked. The cluster diameter is ~7 nm.
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Structure and composition of dilute Co-doped BaTiO3 nanoparticles O.I. Lebedev, R. Erni and G. Van Tendeloo EMAT, University of Antwerp, Groenenborgerlaan 171, B2020 Antwerpen, Belgium
[email protected] Keywords: nanoparticles, edge dislocation, vacancies, EFTEM, HRTEM
Dilute ferromagnetism in semiconductors with Curie temperatures (Tc) above 300 K are materials of high technological interest due to their potential use in spin based electronic devices operable at room temperature. Additionally, the current technological trends towards device miniaturization are driving the development of materials research strongly in the direction of functional nanomaterials. Therefore, fabrication of well characterized nano dilute magnetic semiconductor systems is becoming increasingly important. Nanoparticles of 5% Co-doped BaTiO3 (Co-BTO) with nominal composition BaTi1¡xCoxO3 were synthesized following an established solvothermal drying route with additional cobalt(II) nitrate hexahydrate as Co precursor in stoichiometric quantities. Detailed TEM and ED studies confirm that the samples are indeed single phase Codoped BaTiO3, devoid of other impurity phases or Co metal clusters. HRTEM investigation indicates the presence of Ba vacancies in varying concentrations. Figure 1 shows a HRTEM image from a single Co-BTO nanoparticle along the [111]C zone axis. There is a clear reduction of contrast of the lattice fringes within 1-2 nm size areas. Such local variation in HRTEM contrast can be attributed to clusters of Ba vacancies. Moreover, exceeding a certain vacancy concentration leads to internal stress that can be reduced by the formation of edge dislocations (Fig.2a). The existence of dislocations is particularly surprising in case of nanoparticles where the energy stability of a dislocation is not a priori warranted. The closure failure of the Burgers circuit in the HRTEM image (Fig.2b) determines the Burgers vectors as b1 and b2 =a√2 [110]C (a being the lattice parameter of cubic BaTiO3). In order to clarify the origin of these contrast variations, EFTEM and Z- contrast imaging have been employed. The EFTEM Ti and Co map (Fig 2d,e) clearly confirms a quite narrow Ti-Co distribution inside the nanoparticles while the Z-contrast imaging (Fig.2c) indicates that the bright spots correspond to pore-like structures in the nanoparticles, and are not related to any chemical inhomogeneity or metal clustering effects. The elemental maps and the plasmon-loss image reveal that the regions with brighter contrast in the HRTEM image do not correlate with increased concentrations of Ti or Co. This leads to the conclusion that physical voids within the particles have to be present. Electron Microscopy results indicate that we have formed Co-doped cubic BTO nanocrystals in which the presence of vacancies and relative defect concentrations regulate the occurrence/absence of ferromagnetism.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 231–232, DOI: 10.1007/978-3-540-85226-1_116, © Springer-Verlag Berlin Heidelberg 2008
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Figure 1. [111] HRTEM image of Co-doped BaTiO3 nanoparticle and corresponding FT pattern. Filtered HRTEM image from selected region marked by white frame notice variation of the contrast ( marked by white arrow)
Figure 2. (a) - [111] HRTEM image of a highly defected Co-doped BaTiO3 nanoparticle and corresponding FT pattern; (b)-filtered HRTEM image showing presence of core dislocations. The associated Burgers circuits are indicated. (c) - Z contrast image of an agglomerate of nanoparticles exhibiting inhomogenities and EFTEM images of selected nanoparticles at Co-M edge(65 eV energy loss) (d) and TiM edge (45 eV energy loss) (e)
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CoxFe3-xO4 catalytic materials for gaz sensors L. Ajroudi1,2, A. Essoumhi1, S. Villain1, V. Madigou1, N. Mliki2, and Ch. Leroux1 1. IM2NP (UMR-CNRS 6242), South University Toulon-Var (USTV), Bat.R, B.P.20132, 83957 La Garde Cedex, France 2. LMOP, Physics Department, University Tunis El-Manar, 2092 Tunis, Tunisia
[email protected] Keywords: nanoparticles, ferrites, catalysis
Nanomaterials based on spinel ferrites have already numerous applications, mainly based on their magnetic properties. Recently, catalytic properties of nickel, copper and cobalt ferrites in the conversion of CO and CH4 were evidenced, opening a new field of applications for these materials [1]. Magnetic properties of nanoparticles, as well as catalytic properties, will depend on their size, but also on their shape and size distribution. These parameters are linked to the elaboration method. The properties of transition metal spinel oxides depend also on the nature of the transition metal and on his site occupation in the structure. The location of the cations in the spinel structure is related to their octahedral or tetrahedral sites preference, but also to the synthesis method. For magnetite Fe3O4, which adopts the inverse spinel structure, the tetrahedral sites are fully occupied by Fe3+ ions, whereas octahedral sites are occupied by Fe3+ and Fe2+. For Co3O4, the octahedral and tetrahedral sites are respectively occupied by Co3+and Co2+ (normal spinel structure). Intermediate situations occur for CoxFe3-xO4 ,. Thus, three different compositions were prepared CoFe2O4, Co0 6 Fe2 4O4, Co1 4 Fe1 6O4. In order to synthesize chemical homogeneous powders, we used a co-precipitation method, and a non aqueous elaboration technique developed by Pinna [2]. For the coprecipitation method, the starting iron and cobalt salts were FeCl3.6H2O, FeSO4.7H2O and CoSO4.7H2O. Two different co-precipitations were realised. The proportions and nature of the salts changed, but the rest of the procedure was the same. The salts were dissolved in distilled water, and the solution was then mixed to a solution of NaOH, heated at 70°C. Since Fe3O4 was obtained by co-precipitation of FeCl3.6H2O and FeSO4.7H2O, a co-precipitation was realised with FeCl3.6H2O, FeSO4.7H2O, and CoSO4.7H2O, (sample A FC 17). Another attempt with FeCl3.6H2O and CoSO4.7H2O, was also done (sample B FC16). The precipitates were annealed at different temperatures (250°C,300° and 500°C). For the second method, iron acetylacetonate and cobalt 2,4-pentanedionate were mixed in various proportion in benzyl alcohol. The mixture was stirred and put into a steel autoclave. After two days in a furnace, one obtains a dark suspension, which was ultrasonicated and centrifuged. The precipitates were thoroughly washed and subsequently dried in air. This procedure was applied to the elaboration of CoFe2O4 (sample C), Co0 6 Fe2 4O4,(sample D) and Co1 4 Fe1 6O4 (sample E). Whatever the elaboration method, CoxFe3-xO4 nanoparticles were obtained. The spinel nanoparticles, obtained by co-precipitation method, have irregular shapes (Figure 1a) and mean sizes ranging from 8 nm (sample B, 300° C) to 12 nm (sample A, 500 °C).
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 233–234, DOI: 10.1007/978-3-540-85226-1_117, © Springer-Verlag Berlin Heidelberg 2008
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Sample B is chemically homogeneous, but sample A contains Co3O4 in form of platelets and needles (Figure 1b). TEM studies of samples annealed at lower temperatures showed that the platelets result from the decomposition of cobalt hydroxide, as needles come from the initial cobalt sulphate. Spinel nanoparticles obtained by the non aqueous method are very regular in shape and well dispersed (Figure 2a). They are also chemically homogeneous and very homogeneous in size (Figure 2b). The chemical composition was tested by nanoprobe analysis. 1. 2.
D. Fino, S. Solaro, N.Russo, G. Saracco, V. Specchia, Topics in Catalysis 42, (2007), p.454 N. Pinna, S. Grancharov, P. Beato,| P. Bonville, M. Antonietti and M. Niederberger, Chem. Mater. 17 (2005), p. 3044.
Figure 1. a) HREM of one CoFe2O4 particle, with a [110] zone axis. b) Sample A, annealed at 500°C. The powder consists in a mixture of CoFe2O4 nanoparticles, and Co3O4, in form of platelet and needles.
Figure 2. a) CoFe2O4 nanoparticles (sample C). b) Histogram of the size distribution.
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(S)TEM investigation on the role of alumina dopants to prevent redox activity decay at high temperature in CePrOx /doped-Al2O3 catalysts M. López-Haro, K. Aboussaid, J.M. Pintado, J.J. Calvino, S. Trasobares Departamento de Ciencias de los Materiales, Ingeniería Metalúrgica y Química Inorgánica. Universidad de Cádiz, 11510-Puerto Real (Cádiz). Spain
[email protected] Keywords: CePrOx Catalysts, STEM, doped-alumina
CePrOx mixed oxides, which present extraordinary redox properties, have a wide range of potential applications in environmental catalysis [1]. A good dispersion of these oxides over a high surface area material, like alumina, increases their specific surface and, hence, their oxygen handling capabilities. Moreover if the alumina support is modified with a doping agent, their textural stability can be improved and their deactivation by solid-state reaction with the support inhibited [2]. In this work, Transmission Electron Microscopy Imaging Techniques (HREM and HAADF) have been combined with spectroscopic techniques, (EELS and X-EDS), to investigate the influence of the dopant nature (SiO2, La2O3) on the chemical and structural properties of CePrOx particles supported on modified aluminas. A mechanism to explain the differences observed [3] in the resistance against high temperature deactivation of a Ce0 8Pr0 2O2-x supported on Si-doped alumina (Ce0 8Pr0 2O2-x/Al2O3- 3.5% SiO2) and a La-doped alumina (Ce0 8Pr0 2O2-y/Al2O3 -4% La2O3) system is proposed. To reveal the effects of high temperature aging, samples of the two catalysts were studied after treatments under reducing conditions at low, 350ºC, and high temperature, 900ºC. HREM indicates the presence in both materials of a fluorite-like structure on the samples treated at 350ºC. Neatly different results are observed at the highest temperature. Thus, two different crystalline phases have been detected in the Si-doped material: a lanthanide hidroxycarbonate (Bastnesite-type) (figure 1.A) and a fluorite-like structure. In the case of La-doped sample, only a perovskite phase, LnAlO3 (Ln=Ce, Pr) (figure 1.B) is observed. The formation of the redox-inactive perovskite phase would explain the greater deactivation behaviour observed in the La-doped sample. To investigate further the role of the dopant in these solid state chemistry differences, the samples were characterised by STEM-XEDS and STEM EELS. These techniques provide a more accurate, high spatial resolution, description of the elements distribution. In the Si-doped samples (both at 350ºC and 900ºC) EELS studies indicated the simultaneous presence of small (a few nanometers) PrOx crystallites and, much larger (100-200 nm), Ce-rich, CePrOx crystals over the support. In contrast, in the La dopedsample, Pr is homogeneously distributed all over the catalyst, not only in the form of Ce-rich CePrOx mixed oxide particles (at low temperature) or as LnAlO3 (at high temperature) but also into/over the alumina particles (figure 2).
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 235–236, DOI: 10.1007/978-3-540-85226-1_118, © Springer-Verlag Berlin Heidelberg 2008
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These results suggest that the addition of SiO2 as dopant prevents the formation of the redox-inactive LnAlO3 perovskite phase by blocking the diffusion of the lanthanide ions into the alumina crystallites, which allows maintaining the lanthanide ions within crystalline phases which still present a high enough Ln4+/Ln3+ exchange capability.
Figure 1.- Experimental HREM images acquired on (a) the Si-doped material and (b) the La-doped material reduced at 900ºC. Ce M5,4 La M5,4
Pr M5,4
1050
0
950
50
Distance (nm)
100
850
Energy Loss (eV)
Figure 2.- EELS 3D representation of a collection of 50 spectra acquired on the Ladoped material after treatment at 900ºC. Similar results were observed at 350ºC. 1. 2. 3. 4.
M. Shelef, G.W. Graham, R.W. McCabe, Catalysis by Ceria and Related Materials A. Trovarelli,Imperial College Press, London 343-374 (2002). H. Schaper, E. B. M. Doesburg, L.L. van Reijen, Appl. Catal. 7, 211 (1983). K. Aboussaid, S. Bernal, G. Blanco, G.A. Cifredo, A. Galtayries, J.M. Pintado Mohamed Soussi el Begrani. Surface and Interface Analysis, 40, 3-4 ,250-253 (2008) We acknoledge the financial support from Ministry of Education and Science of Spain (Proyto MAT2005-00333) Junta de Andalucía (Grupos FQM-110 y FQM-334), and Programa Ramón y Cajal 2003. The electron microscopy work was carried out at the Electron Microscopy Division of Central Services of Science and Technology at the University of Cádiz.
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Sulfated Zirconia Catalysts: Structure and Performance Relationship, a TEM Study C. Meyer1, D. Su1, N. Hensel1, F.C. Jentoft1 and R. Schlögl1 1. Department of Inorganic Chemistry, Fritz-Haber Institute of the Max Planck Society, Faradayweg 4-6, 14195 Berlin, Germany
[email protected] Keywords: Sulfated zirconia, structure, heterogeneous catalysis, HR-TEM
Sulfated Zirconia (SZ) is an important and suitable catalyst for heterogeneous nalkane isomerization that experiences braod industrial application. Zirconia in tetragonal phase that is stabilized at low temperature in the form of nanosized particles [1] proved a high activity and selectivity for this reaction [eg. 2]. A key issue in elucidating parameters relevant for catalytic performance is the observation of processes occurring during the calcination process, in which the precursor material is transformed to the active catalyst [3]. During calcination exothermic crystallization of the amorphous zirconia precursor produces a specific overshoot in the temperature known as the glow phenomenon. Once this event has occurred, the material exhibits significant activity as opposed to the inactive precursor. This study is focused on the correlation of structural characteristics and activity or selectivity of SZ nano powder catalyst. A series of quenching experiments at different stages of the calcination process has been conducted. Sulfated hydrous zirconia (MEL Chemicals XZ0 682/01) was used as starting material. Calcination was performed in flowing air at 823 K in batches of 20 g. Temperature was held for 3 h, applying a temperature ramp for heating and cooling of 15 K / min. Quenching in liquid nitrogen was done before, during and after the glow phenomenon as well as at completion of the temperature program. Isomerisation of nbutane (1 % n-butane in He at atmospheric pressure) at 373 K was used as a test reaction. Activity and selectivity of the catalyst samples were monitored as a function of time by on-line gas chromatography. Further characterization included XRD, UV-VIS and BET measurements. Systematic TEM investigations reveal sample properties with a high degree of spatial resolution. Electron diffraction and Fourier transformation are used for phase characterization, EDX and EELS for chemical analysis and HR TEM for morphology and grain size measurements. Further, detailed structures of single grains are investigated. Figure 1 exhibits features of interest using a tetragonal grain of zirconia with a maximum diameter of 30 nm as an example. Line of sight is parallel to [100], (011) and (002) are projected. Surface steps (1), bending of projected lattice traces (2) and intra grain porosity (3) are observable. This sample is taken after the glow phenomenon and shows high activity. Figure 2 shows for comparison a sample with little catalytic activity. Lattice spacings
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 237–238, DOI: 10.1007/978-3-540-85226-1_119, © Springer-Verlag Berlin Heidelberg 2008
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prove the tetragonal nature of the zirconia. A difference lies for example in the smoothness of the surface termination denoting energetic differences.
Figure 1. HR TEM micrograph of SZ catalyst. Projection of (011) and (002), line of sight is [100]. The arrows indicate exemplarily structural features, for details see text.
Figure 2. HR TEM micrograph of SZ catalyst. (011) planes of tetragonal zirconia are projected. The arrows indicate porosity (1) and a smooth surface termination (2). 1. 2. 3.
R.C. Garvie, J. of Phys. Chem., 82 (1965), p. 218 – 224. M. Benaissa, J.G. Santiesban, G. Dias, C.D. Chang, M. Jose-Yacaman, J. of Catalysis, 161 (1996), p. 694 - 703. A.H.P. Hahn, R.E. Jentoft, T. Ressler, G. Weinberg, R. Schlögel, F.C. Jentoft, J. of Catalysis, 236 (2005), p. 324 - 334.
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A novel procedure for an accurate estimation of the lattice parameter of supported metal nanoparticles from the analysis of plan view HREM images C. Mira, J.A. Perez-Omil, J.J. Calvino and S. Bernal Dep. Ciencia de los Materiales e Ingeniería Metalúrgica y Química Inorgánica. Universidad de Cádiz. c/ Rep. Saharaui s/n. 11510 Puerto Real (Cádiz) - SPAIN cesar
[email protected] Keywords: HREM, plan view, nanoparticle lattice parameter
The lattice parameter of metal nanoparticles supported on a carrier material, as are those present in a large variety of catalysts, can suffer small modifications from the value expected for the bulk materials. Thus, dilatation or contraction of the lattice constant of supported metal has been related to effects like incorporation of support elements into the particles in the form of alloys; incorporation of small atoms of other elements, like H, O, or C, into the metal lattice; pseudomorphic growth; surface stress or encapsulation of the metal under a large compressive stress by the substrate. The observed modifications of the lattice parameter are frequently smaller than 5%. A precise quantification of this effect could help to understand their precise origin. The measure of lattice parameters by X-ray diffraction is quite extended. In spite of its accuracy this technique provides only an average value of all the particles under analysis. Likewise, the analysis becomes very unreliable with diffraction patterns with a low signal to noise ratio as are those obtained in systems with very small particles. In the case of HREM images, errors higher than 5% can be expected in direct measurement of lattice spacings [1]. In fact, the accuracy with this measurement technique depends on several experimental factors [2]. SAED patterns allow estimations with high relative errors (2-3%). An statistical approach, considering a large number of particles, improves the accuracy of HREM or SAED measurements [3], but in this case we are not characterising a single particle but an ensemble of them. We have developed a procedure to increase the accuracy in the determination of the lattice spacings of supported metal particles based on the detailed analysis of Moiré type fringes observed in plan view HREM images. These Moiré fringes correspond to linear combinations between the characteristic metal and support reflections. With particles of only a few nanometers a large number of Moiré spots can be detected in the corresponding image diffractograms. Using the lattice fringes of the bulk support as a reference, the measurement of each Moiré reflection in reciprocal space allows an estimation of metal lattice spacings. The error obtained in these individual determinations is lower than that expected for the direct measurement of the metal spacing, the exact value of the error depending on the specific Moiré reflection selected. Nevertheless, a much better estimation can be obtained if all the information present in the diffractogram is used simultaneously.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 239–240, DOI: 10.1007/978-3-540-85226-1_120, © Springer-Verlag Berlin Heidelberg 2008
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If we record an intensity profile of the image diffractogram we can try to fit it to a theoretical curve in which the metal structure is the parameter to be refined. We have developed a software to apply this approach. This program allows us fitting simultaneously all the visible peaks in the intensity profile, which includes not only the metal and support peaks but also a set of Moiré reflections. The positions of the different peaks in the theoretical curve are very sensitive to subtle modifications of the metal lattice spacings. Thus, the correlation coefficient between the theoretical and experimental curves shows a maximum for a precise metal spacing value, usually corresponding to situations of slight lattice contractions or expansions. In the figure below a planar view of a CeO2-supported Pd particle (a) and its diffractogram (b) are shown. Different Moiré reflections can be identified, which result from the combination of (111) CeO2 and (111) Pd reflections. The experimental intensity profile and its theoretical fitting are also shown. Two possible metal lattice modifications have been considered (c). A 0.6% contraction shows a better correlation (r) than that characteristic of a 0.4% expansion. By applying this procedure to the analysis of HREM images simulated for a set of models which consider exact metal lattice expansions/contractions we have estimated the precision of the method to be of the order of 0.2%. This new procedure has been applied to Pd/CeO2 images to detect, in single particles, variations in the metal lattice parameters lower than 1% with high reliability. 1. 2. 3. 4.
J.-O. Malm, M.A. O´Keefe; Ultramicroscopy 68 (1997) 13-23. W.J. DE Ruijter, R. Sharma, M.R. McCartney, D.J. Smith; Ultramicroscopy 57 (1995) 409422 S.-C.Y. Tsen, P.A. Crozier, J. Liu; Ultramicroscopy 98 (2003) 63-72 MEC/FEDER (MAT2005-00333) and JA (FQM-110, FQM-334) are acknowledged
a
b CeO2 Pd
Figure 1. (a) HREM image of a Pd/CeO2 catalyst ; (b) image diffractogram; (c) fitting of the experimental intensity profile to -0.6% and 0.4% variations in the lattice parameter. 300
300
-0,6%
r = 99,2%
250
r = 98,3%
250
+0,4%
Calc exptal.
200
intensity
intensity
200 150 100
100 50
50
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Microstructure of Pt particles and aggregates deposited on different carbon materials for fuel cells application D. Mirabile Gattia1, E. Piscopiello1*, M. Vittori Antisari1, S. Bellitto2, S. Licoccia2, E. Traversa2, L. Giorgi1, R. Marazzi1, A. Montone1 1. FIM Department, ENEA – C.R. Casaccia, Via Anguillarese 306, 00123 Rome, Italy *C.R.Brindisi, Via Appia Km 702, 72100 Brindisi, Italy 2. Department of Chemical Science and Technology, University of Roma “Tor Vergata”, Via della Ricerca Scientifica, 00133 Rome, Italy daniele
[email protected] Keywords: carbon nanostructures, fuel cell, catalyst
PEM (polymer electrolyte membrane fuel cells) and DMFC (direct methanol fuel cells) have demonstrated to be suitable devices in order to realize a widespread diffusion of electrical-H2 fed vehicles in the near future [1]. In recent years several research efforts were on the study of Pt clusters deposited on different kinds of nanostructured carbon, besides the classical carbon black, with the purpose of reducing the Pt loading by the optimization of the catalyst performances during electrochemical reactions at the cell electrodes. In this work Pt clusters have been deposited by an impregnation process on three carbon supports: Multi-Wall carbon Nanotubes (MWNT), Single-Wall carbon Nanohorns (SWNH) and Vulcan XC-72. MWNT and SWNH have been home synthesized by a DC [2] and an AC arc discharge process [3] respectively. The Pt particles, deposited on the three carbon supports, have been characterized by Scanning and Transmission Electron Microscopy, X-ray diffraction and cyclic voltammetry measurements. Electron microscopy investigations, revealed the presence of nanostructured aggregates with different diameters: 50-100 nm and 20-50 nm in Pt/SWNH and Pt/MWNT samples respectively, while in the case of Pt/Vulcan single nanoparticles were deposited. In this last sample the process resulted in a strongly inhomogeneous microstructure with several sample regions free from deposited particles. Electrochemical characterization showed that the Pt nanostructures deposited on MWNT were particularly efficient in the methanol oxidation reaction, even if the Pt active surface area on the Vulcan substrate is larger. This shows that particle aggregates can be more efficient respect to single particles, probably owing to the particular particle shape and to the presence of grain boundaries. The comparison between the two nanostructured substrates evidences furthermore a role for the small size of aggregates. The details of the microstructure, as evidenced by the high resolution TEM analyses, are reported in the insets (figure 1). The agglomerates deposited on the nanotubes appear to be constituted by single crystal region larger than the single leafs so that the structure appears quite complex and requires further analyses for a complete description.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 241–242, DOI: 10.1007/978-3-540-85226-1_121, © Springer-Verlag Berlin Heidelberg 2008
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a)
b)
c)
d)
Figure 1. Low magnification (a) and high magnification (b) TEM images of Pt particles deposited on Vulcan XC-72. In the inset a high resolution image is reported. In (c) and (d) Pt aggregates deposited on MWNT are shown with a high resolution detail in the inset. 1. 2. 3.
S. Gottersfield and T. Zawodzinski, Adv. Electrochem. Sci. Eng. 5 (1997), p. 195. D. Mirabile Gattia, M. Vittori Antisari, R. Marazzi, L. Pilloni, V. Contini and A. Montone, Materials Science Forum 518 (2006), p. 23. D. Mirabile Gattia, M. Vittori Antisari and R. Marazzi, Nanotechnology 18 (2007), p. 255604.
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Low-loss-energy EFTEM imaging of triangular silver nanoparticles J. Nelayah1, L. Gu1, W. Sigle1, C.T. Koch1, L. Pastoriza-Santos2, L.M. Liz-Marzan2, and P.A. van Aken1 1. Max Planck Institute for Metals Research, Stuttgart Center for Electron microscopy, Heisenbergstr. 3, D-70569 Stuttgart, Germany 2. Departemento de Quimica Fisica, Universidade de Vigo, 36310, Vigo, Spain nelayah@mf mpg.de Keywords: EFTEM, EELS, metallic nanoparticles, surface plasmon mapping
Understanding how light interacts with matter at the nanometer scale is a fundamental issue in optoelectronics and nanophotonics. It is known that the optical properties of nanoparticles are entirely dependent on collective excitations of their valence electrons, known as "surface plasmon resonances" (SPR´s), under electromagnetic illumination. Measuring these properties locally at the level of the individual nano-object constitutes a challenging issue for linking of the global response of the nanoparticles and the underlying structure and morphology. The high-energy electron beam in a transmission electron microscope (TEM) is an excellent tool for this application. In particular, it has been recently shown that low-loss electron energy-loss spectroscopy (EELS) in the context of a scanning transmission electron microscope (STEM) enables the optical properties of metallic nanoparticles in the ultraviolet–near-infrared (UV–NIR) domain to be probed with nanometer resolution via the mapping of SPR´s [1]. With the advent of recent instrumental improvements such as electron monochromators and in-column energy filters, it is expected that the understanding of these optical properties can be further improved through the gain in both energy and spatial resolution. In this contribution, we present energy-filtered transmission electron microscope (EFTEM) studies of the optical properties in the UVNIR regime of individual triangular silver nanoparticles.. Triangular silver nanoprisms have been synthesized as described in [2]. The prisms have edge lengths in the range between 100 and 300 nm and are typically between 5 and 10 nm thick. The instrument used for this investigation was the new 200 kV SESAM FEG-TEM (Zeiss) fitted with an electrostatic monochromator and a high-dispersion and high-transmissivity in-column MANDOLINE filter [3]. The energy resolution was set to 0.25 eV for these experiments. The EFTEM series were acquired on a 2k × 2k CCD camera (8 × 8 binning) using a 0.2 eV energy selecting slit. Energy-filtered images were recorded from 0.5 eV to 4 eV (17 images) with an acquisition time of 20 s/image. Figure 1 shows a series of EFTEM images of a triangular nanoparticle with 210 nm edge length recorded at energy losses of (a) 0 eV, (b) 1.0 eV, (c) 1.6 eV, and (d) 2.2 eV, respectively. The two-dimensional images in Figures 1(b)-(d) represent the intensity distributions of the first three main SPR´s of the silver nanoprism as already observed with STEM-EELS in [1]. But, compared to the latter data, our EFTEM images display
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increased spatial sampling and are obtained with shorter acquisition time and without post-acquisition data processing. These results clearly demonstrate that low-energy-loss EFTEM imaging in the SESAM microscope provides a fast and precise analytical tool for investigating the optical properties of single metallic nanoparticles in the visible range. This tool enables both unprecedented spatial and energy resolution which will provide new information about the optical properties of nanomaterials. 1. 2. 3. 4.
J. Nelayah, M. Kociak, O. Stephan, F.J.G. de Abajo, M. Tencé, L. Henrard, D. Taverna, I. Pastoriza-Santos, L.M. Liz-Marzan, and C. Colliex., Nature Physics 3 (2007) 348. V. Bastys, I. Pastoriza-Santos, B. Rodriguez-Gonzalez, R. Vaisnoras and L.M. Liz-Marzan, Adv. Funct. Mater. 16 (2006) 766. C.T. Koch, W. Sigle, R. Höschen, M. Rühle, E. Essers, G. Benner, and M. Matijevic, Micoscopy and. Microanalysis 12 (2006) 506. We acknowledge financial support from the European Union under the Framework 6 program under a contract for an Integrated Infrastructure Initiative. Reference 026019 ESTEEM
(a)
(b)
(c)
(d)
Figure 1. EFTEM images of a triangular silver nanoparticle with 210 nm edge length imaged at energy losses of (a) 0 eV, (b) 1.0 eV, (c) 1.6 eV and (d) 2.2 eV. Images were taken with a slit width of 0.2 eV centred on the specific energy loss. Scale bar and colour level are common for all images. The bright pixels indicate maximum intensity.
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Microstructure of cobalt nanocluster arrays fabricated by solid-state dewetting Yong-Jun Oh1, Junghwan Kim1, Sukhun Hwang1, Caroline A. Ross2, Carl V. Thompson2 1. Advanced Materials Science and Engineering Div., Hanbat University, Korea 2. Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, Massachusetts 02139
[email protected] kr Keywords: nanoclusters, dewetting, cobalt phases
Cobalt nanocluster arrays have recently attracted considerable attention due to their applications as patterned magnetic recording media as well as catalyst arrays for growing carbon nanotubes [1,2]. The processes to fabricate nanocluster arrays are mostly based on direct lithography of thin films using coatings of photo resist and their selective removal. Since the resultant clusters formed by this technique retain the microstructure of the original thin films, we need to control the microstructure of thin films before lithography to obtain clusters with specific crystal property. Recently, one of the authors developed a self-assembling technique to fabricate gold nanoparticle arrays by dewetting a thin metal film on topographic templates at elevated temperatures [3]. The technique was also characterized by crystal reorientation during agglomeration of the thin film by dewetting. In this study, we investigate the changes in phase, crystallography, and microstructure of cobalt before and after dewetting by thermal annealing. The topographic templates consisted of 200-nm-period square arrays of inverted pyramidal pits on (100) silicon wafers. Interference lithography using a laser beam was used to create patterns on the wafers. The templates were oxidized to prevent the reaction between the substrate silicon and the cobalt thin film. The cobalt films were deposited on the templates using ion-beam sputtering and were annealed in forming gas to induce dewetting. The ordered cobalt clusters at a high temperature were mostly in the fcc phase (Figure 1), while the films annealed at a low temperature showed a mixture of hcp and fcc cobalt phases (Figure 2). The dewetted clusters were almost single crystal when twin boundaries were disregarded. The orientation of more than a quarter of the observed particles was such that the (111) plane of the Co particles was parallel to a pair of the inverted pyramidal faces of the silicon template. The dewetting process is believed to be a promising technique to fabricate single-crystal nanocluster arrays for high-density magnetic recording media and carbon nanotube growth. 1.
J.I. Martin, J. Nogues, K. Liu, J.L. Vicent, I.K. Schuller, Journal of Magnetism and Magnetic Materials 256 (2003), p.449.
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3. 4.
C.A. Ross, M. Hwang, M. Shima, J.Y. Cheng, M. Farhoud, T.A. Savas, Henry I. Smith, W. Schwarzacher, F.M. Ross, M. Redjdal, F.B. Humphrey, Physical Review B, 65 (2002), p.144417. A.L. Giermann, C.V. Thompson, Applied Physics Letters 86 (2005), p.121903. We thank H.I. Smith and K. Berggren at MIT for interference lithography. (a)
(b)
Zfcc-Co=[110]
100 nm
100
Figure 1. (a) Scanning and (b) transmission electron microscope (SEM and TEM) images showing the dewetted clusters. The SAD pattern on the top right was taken from the particles in pits.
fcc[100]
fcc[110]
hcp[1213]
5 nm
Figure 2. High-resolution TEM image of the as-sputtered cobalt film. The arrows indicate the grains of fcc-Co and hcp-Co phases showing the atomic images along zone axis with low indices.
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Size Effect in Gold Nanoparticles Investigated by Electron Holography and STEM L. Ortolani1,2, V. Morandi2 and M. Ferroni3 1. University of Bologna, Dept. of Physics, v.le B. Pichat, 6/2, 40127 Bologna, Italy 2. CNR-IMM Bologna, v. Gobetti, 101, 40129 Bologna, Italy 3. INFM-CNR SENSOR and Dept. of Chemistry and Physics, Brescia University, v. Valotti, 9, 25133 Brescia, Italy
[email protected] Keywords: Gold Nanoparticles, Electron Holography, HAADF-STEM
Gold clusters are extremely interesting nanosystems with a high catalytic activity, exploited for sensing applications and to promote the growth of nanostructures [1,2]. All these potential applications motivated numerous structural studies on Au nano-clusters. The mean electrostatic potential (MIP) is a fundamental quantity and its value is crucial for accurate evaluation and simulation of experimental data from TEM imaging and electron diffraction. Recently, a dependence of the MIP on particle size has been reported, measured by electron holography (EH): the increase of the MIP over the bulk value for particle sizes lower than 5 nm has been shown, suggesting a correlation with the catalytic behaviour of gold [3,4]. Despite these very promising results, a similar effect was observed in amorphous carbon films as the result of a thickness independent phase shift [5]. EH is capable of providing a quantitative determination of this surface phase shift since it depends on the projected sample thickness. Unfortunately, gold particles are reported to change their contact angle with the substrate reducing their dimension, as a result of a size-dependent change in the particle-support interaction [3]. To overcome these limitations, and to determine unambiguously information on surface phase effects in gold clusters, a combination of High Angle Annular Dark Field STEM (HAADF-STEM) and electron holography has been used, exploiting the local sample thickness dependence of the HAADF-STEM signal. HAADF intensity depends also on the imaging conditions, on the atomic number and of the density of the observed material. By keeping fixed all these parameters, it is possible to directly correlate image intensity to sample thickness variations. From the holographic reconstructed phase map of gold clusters over an amorphous carbon film, like the one shown in Fig. 1a), it is possible to fit the induced phase shift and the projected radius of the particle, as reported in the linescan of Fig. 1b). From a HAADF-STEM image of the same clusters it is possible to fit the projected radius and the HAADF intensity, as shown in Fig. 2a). By keeping constant all the imaging parameters, intensity variation in the image only depends on changes in the particles shape. The fit of the data of Fig. 2b) shows that the smaller particles are thinner than the larger ones, as shown in the model of Fig. 3a). From these results it is possible to correct the electron holography phase shift measurements, obtaining the plot of Fig. 3b). Numerical fitting of the data reveals, for gold clusters dispersed over a carbon film, a thickness independent phase shift of 0.45
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rad. Additional studies are needed to find the origin of these surface electrostatic effects, which could be addressed to surface strain at cluster interfaces. Nonetheless, the synergy of EH and HAADF-STEM provides an insight into the interactions between the constituents of nanostructured systems. 1. 2. 3. 4. 5.
R Andres, T. Bein and M. Dorogi, Science, 272 (1996), p. 1323. M. Haruta, Catal. Today, 36 (1996) p.153. S. Ichikawa, T. Akita and M. Okamura, JEOL News, 38 (2003), p. 6. L. Ortolani et al., J. Eur. Ceram. Soc., 27 (2007), p. 4131 M. Wanner et al., Utramicroscopy, 106 (2006), p. 341.
Figure 1. a) Reconstructed phase map of three gold nanoparticles. b) Linescan profile and result of the numerical fit.
Figure 2. a) Linescan profile from HAADF image of a gold nanoparticles and result of the fit. b) Plot of the HAADF intensity for different particles and numerical fit.
Figure 3. a) Model for gold particles shape of decreasing dimension. b) Plot of the holographic phase shift and numerical fit showing a surface phase effect of 0.45 rad.
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Post-Mortem investigation of Fischer Tropsch catalysts using cryo- transmission electron microscopy Dogan Ozkaya1, Martin Lok2, John Casci2 and Peter Ash1 1. Johnson Matthey Technology Centre Blounts Court Sonning Common Reading RG4 9NH UK 2. Johnson Matthey Technology Centre, PO Box1 Belasis Avenue Billingham TS231LB UK
[email protected] Keywords: GTL, Fischer Tropsch, Cryo- microscopy
Co/Al2O3 catalysts are widely used in the Fischer-Tropsch, gas to liquids (GTL), catalytic reaction where syngas (CO2 and H2) is converted into higher hydrocarbon wax products. One important use of the wax product is through cracking to produce clean diesel fuel. In most production routes the catalyst is initially in the form of highly dispersed Co-oxide particles on a high surface area gamma alumina with up to 1 wt% addition of a precious metal promoter. The catalyst can then be reduced to its active state in-situ in the FT reactor or supplied in pre-reduced form. In the case of the prereduced catalyst the material is encapsulated in a wax product to prevent re-oxidation of the cobalt. Post reaction the catalyst is suspended in the wax product of the FT reaction. It is of paramount importance to analyse the initial state and the final state of a catalyst in order to understand how the reaction has progressed. Any pre or post-reactor analysis therefore needs to deal with the wax but leave the catalyst unchanged. However, the dewaxing procedures traditionally applied to the catalyst, (Soxhlet extraction and calcination at 350°C) before examination, not only partially oxidize the Co but also cause some changes in the microstructure. Consequently, the combination of Cryoelectron microscopy and cryo-microtomy offer a straightforward, but unique, route to analyse the catalyst within its original wax environment. The sample (spent catalyst had operated for 1000 hours in a slurry reactor at 210 C and 20 bar using a H2/CO ratio of 2.1) was prepared as follows: wax sample was stuck to a microtome stub using sucrose solution at liquid nitrogen temperatures and microtomed using a Leica FC-6 cryo-ultramicrotome. The slices were placed on a holey carbon coated Cu TEM grid and transferred to a Tecnai F20 field emission transmission electron microscope using a Gatan cryo-transfer system. Direct analysis of the catalyst is demonstrated with high-resolution images and analysis from the TEM. The catalyst analysed was produced using the high dispersion Catalyst (HDC) technique [1,2] rather than the more usual nitrate route. The material generated using HDC route was then reduced, generating Cobalt metal on the support, and encapsulated in wax An example of a microtomed wax on a TEM grid is shown in figure 1a. The marks from a diamond knife can be seen on the wax and this illustrates the difficulties of handling wax at low temperatures. Figure1b shows a part of the catalyst embedded
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within the wax. It shows that diamond knife marks do not necessarily prevent the analysis of the catalyst. Figure 2a shows a region of the catalyst close up. The particle size analysis from a region like this is shown in figure 2b with catalyst as prepared (fresh) and after it has been through the reaction (spent). The changes in the particle size reflect directly the conditions that the reactor has been through and possible to correlate it with reactivity. 1. 2.
1- Lok C. M. Studies in Surface Science and Catalysis 147 (2004) 283 2- Bonne R. and Lok C. M. US patent 5, 874 (1999) 381
Figure 1. Low magnification image of the wax slice on a TEM grid (fig 1a) and a catalyst piece within a wax surround.
Figure 2. The Co metal particles on alumina support (fig 2a) with particle size distributions for a fresh catalyst and spent catalyst.
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TEM Investigations on Cu-impregnated Zeolite Y catalysts via chloride free preparation M.-M. Pohl, M. Richter, M. Schneider Leibniz-Institute for Catalysis e.V.(branch Berlin), Richard-Willstätter-Str. 12, D-12489 Berlin, Germany
[email protected] Keywords: catalysis, impregnated zeolite, Cu dispersion
Copper-containing zeolites are potential catalysts for the oxidative carbonylation of methanol to DMC [1,2] but have a minor activity without chloride. The catalyst used for this study was a chloride-free system which combines high performance with the attractiveness of a process that abandons any halogens during catalyst preparation and process operation [3]. In all those zeolite catalyst systems the question of the distribution of the copper particulary at high loading is crucial. The investigation of the complex morphologies of zeolites by microscopy is a challenge due to their beam sensitivity. Ion exchanged zeolites decrease the beam stability further. The metals leave the zeolite matrix as particles and the zeolite structures are destroyed. As ultimate consequence nanowires can be formed, shown by Mayoral and Anderson [4]. For HRTEM on zeolites the use of Cs-corrected microscopes is a suitable way Tesche [5] showed this opportunity by imaging metal clusters inside ordered pores without any loss of the 3D-Structure. Since the access on Cs-corrected microscopes are limited most microscopists have to deal with conventional TEM. The here analyzed zeolite Y with 14wt% Cu load shows the typical growth of nanoparticles under electron beam at 200kV. Even in the first seconds with intact lattice planes Cu-seeds are visible which grow further under electron bombardment. Parallel the lattice plane disappear under shrinkage of the zeolite as shown in Figure 1. Additional Cu containing particles were detected both by XRD and TEM as CuO differ considerable from the beam grown Cu features (Figure 2.). Since damage could appear within the first seconds within the microscope, no secure interpretation on fine structures of these systems should be interpreted with care . 1. 2. 3. 4. 5. 6.
S. T. King, J. Catal. 161 (1996) 530. S. T. King, Catal. Today 33 (1997) 173 M. Richter, M. Faith, R. Eckelt, E. Schreier, M. Schneider, M.-M. Pohl, R. Fricke, Applied Catalysis B.: Environmental 73(2007) 269-281 A. Mayoral, PA Anderson, Nanotechnology 18 (2007)165708(6pp) B. Tesche, F.C. Jenthoft, R. Schlögl, S. R. Bare, L.T. Nemeth, S. Valencia, A. Corma, 41. Jahrestreffen deutscher Katalytiker, Weimar 2008, P43 We kindly acknowledge the financial support by the Federal Ministry for Education and Research of the FRG, the Senate of Berlin and the European Union (project 03X2002).The sample preparation is acknowledged to Mr. R. Eckelt.
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Within 20 sec
After 1 minute
After 3 minutes
Figure 1. High-resolution at 200kV of a Cu impregnated zeolite Y (1,4 nm lattice planes) with increasing beam damage and development of Cu nano particles
Relative Intensity (%)
Zeolite Y with 14wt%Cu
100.0
CuO
* 0.0
10.0
20.0
30.0
* 40.0
50.0
2Theta
Figure 2. CuO at high Cu load detected by XRD and TEM image with CuO and beginning electron damage within the zeolite Y matrix and first Cu seeds
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Coarsening of mass-selected Au clusters on amorphous carbon at room temperature R. Popescu1, R. Schneider1, D. Gerthsen1, A. Böttcher2, D. Löffler2 and P. Weiss2 1. Laboratorium für Elektronenmikroskopie and Center for Functional Nanostructures, Universität Karlsruhe, 76128 Karlsruhe, Germany 2. Institut für Physikalische Chemie, Universität Karlsruhe, 76128 Karlsruhe, Germany
[email protected] Keywords: Au clusters growths, transmission electron microscopy, surface Ostwald ripening.
Mass-selected Aun (n=4,6,13 and 20) clusters and clusters with an initial distribution of Aum (10≤m≤20) clusters were deposited on amorphous carbon (a-C) thin films by low-energy-beam cluster deposition. The samples were stored at room temperature under ambient conditions over more than two years to analyse the stability of the cluster sizes. The cluster-size distributions were investigated by transmission electron microscopy (TEM) in regular time intervals. Several hundred Au clusters were analysed for each sample and time interval. The cluster radii were assessed by measuring the projected cluster area, which is in a good approximation a circular one. Size histograms were derived and the average radii at a given time t R (t ) were determined. Figure 1 shows the measured R (t ) values which increase strongly over a period of more than two years although the samples were not exposed to elevated temperatures. The coarsening process is best described by a least-square fit of the experimental R (t ) based on R 4 (t ) = R 4 (0) + K d t (t: time, Kd: constant), which corresponds to surface diffusion-limited Ostwald ripening (OR) with the mass transport taking place through the cluster-substrate contact line [1,2]. Coalescence of clusters caused by Brownian motion can be excluded for the given experimental conditions. The values of the surface mass-transport diffusion coefficient Ds' can be calculated using the 45 K d ln( L) ϕ (θ ) k B T , where DS' is given by Kd values and the relation Ds' = 4 ω 2 γ n0 Ds' = Ds c Au with the surface-diffusion coefficient of single Au atoms Ds on a-C and the number of single Au ad-atoms on sites of the a-C substrate cAu. ω=1.7⋅10-29 m3 denotes the volume of gold atoms, γ=1.5 Jm-2 the Au surface energy [3], n0=1.1⋅1019 m-2 the density of sites on the cluster surface [4], kB the Boltzmann constant and T=298 K the temperature. L=2.5 is the screening distance, which is taken to be constant [1]. The parameter ϕ (θ ) = 0.45 depends on the contact angle θ between the Au cluster and the a-
C substrate [4]. The Ds' values are between (1.1±0.1) and (3.8±0.4)⋅10-25 m2s-1. Values for cAu between 0.8 and 2.4⋅1017 atoms⋅m-2 can be derived for our samples as outlined in detail elsewhere [5]. The surface-diffusion coefficient of single Au atoms on a-C is
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given by Ds (T ) = D0 exp(− Ed k BT ) with D0=1.6⋅10-8 m2s-1. Using the relation Ds' = Ds c Au with the values for cAu and Ds' , an activation energy for the surface diffusion of single Au atoms Ed=0.87±0.05 eV is determined in good agreement with a previous theoretical value of Ed=1.0 eV [6]. We show that the cAu values depend - besides on the temperature and the average distance between clusters - also on the initial Au-cluster size distribution on the substrate. They are not particularly sensitive (within our error limit) with respect to the initial size of the Aun (n=4,6,13 and 20) clusters in the case of mass-selected deposition. Moreover, the coarsening process for mass-selected Au clusters (even in case of an initial limited Au-cluster distribution) is quite different from that observed for the deposition of non mass-selected Au clusters on a-C at room temperature as reported in Ref. [7].
1. 2. 3. 4. 5. 6. 7. 8.
B. K. Chakraverty, J. Phys. Chem. Solids 28 (1967), 2401. M. Zinke-Allmang, L. C. Feldman and M. H. Grabow, Surf. Sci. 16(8) (1992), 377. W. R. Tyson and W. A. Miller, Surf. Sci. 62 (1977), 267. R. Popescu, E. Müller, M. Wanner, D. Gerthsen, M. Schowalter, A. Rosenauer, A. Böttcher, D. Löffler and P. Weis, Phys. Rev. B 76 (2007), 235411. R. Popescu, D. Gerthsen, M. Wanner, A. Böttcher, D. Löffler and P.Weiss, to be published. A. A. Schmidt, H. Eggers, K. Herwig and R. Anton, Surf. Sci. 349 (1996), 301. M. Wanner, R. Werner, G. Schneider and D. Gerthsen, Phys. Rev. B 72 (2005), 045426. This work has been performed within the project C4 of the DFG Research Center for Functional Nanostructures (CFN). It has been further supported by a grant from the Ministry of Science, Research and the Arts of Baden-Württemberg (Az: 7713.14-300).
Figure 1. a) Average radii of Au clusters as a function of storage time R (t ) . The
symbols represent the measured R for samples prepared by deposition of: Au4 (Y), Au6 ( ), Au13 (V), Au20 with 1.5⋅1012 clusters (U), Au20 with 0.75⋅1012 clusters ( ) and Aum (10≤m≤20) cluster distribution ( ). The solid lines with the corresponding colour represent fits of the data for diffusion-limited kinetics of surface OR under steady-state condition with the mass transport through the cluster-substrate contact line (see text); b) magnified section inside the dashed rectangle of a) up to 2.5 107 s .
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TEM investigations on Ni clusters electrodeposited on Carbon substrate M. Re1, M.F. De Riccardis1, D. Carbone1, D. Wall2 and M. Vittori Antisari1* 1. ENEA, FIM Department, C.R.Brindisi, Via Appia Km 702, 72100 Brindisi- Italy * C.R.Casaccia , Via Anguillarese 301, 00123 Roma – Italy. 2. FEI, Building AAE Achtseweeg Noord 5, Acht – Eindohven 5651GG marilena
[email protected] Keywords: catalysts, nanostructures, Ni clusters
The role of catalysts in the growth of carbon nanostructures by CVD is particularly critical, since the nano-carbon shape can depend on the catalyst composition and structure besides the deposition parameters [1-3]. In the synthesis of metal catalyzed carbon nanofilaments the performances of the catalyst particles can dramatically depend on both physical and chemical interaction with the substrate. In particular, a good adherence of the clusters to the substrate is necessary to avoid coalescence phenomena during the growth process generally carried out at relatively high temperature. The study of catalyst-substrate microstructure is particularly relevant for the optimization of the whole growth process. In this work Ni clusters were synthesized by electrodeposition, a versatile, rapid and inexpensive technique which, by a specific control of the process parameters, can allow the deposition of continuous metallic films or of nanoparticles, also on complex and convoluted substrates. The Ni clusters were electrodeposited on different C substrates, and have been successful used to assist the CVD growth of carbon nanofibres, having a particularly good adhesion with the substrate [4-5]. The experimental electrodeposition conditions are reported elsewhere [6]. Particularly interesting results were obtained in the case of Carbon Fibres (PAN fibres, produced by controlled pyrolysis of Polyacrylonitrile) [4-6] which were, by this method, decorated with carbon nanofibres grown at the electrodeposited Ni clusters. Transmission Electron Microscopy, with a TECNAI G2 F30 operated at 300 kV, was used to characterize the morphology and microstructure of the electrodeposited Ni and to study in detail the interface between the metal and the substrate in order to better understand the adhesion mechanism of Ni clusters to the C substrates. Considering the cylindrical shape of the Carbon Fibres, cross sectional samples for TEM observations were prepared by FIB with a FEI Strata 400 dual beam instrument. Conventional bright field and high resolution TEM images show that the Ni clusters have a globular shape, with a width in the range of 60–90 nm and a height between 50 and 80 nm (Figure 1 a). The clusters, on the contrary of a commonly observed situation, are polycrystalline and have a grain size of the order of 10 nm (Figure 1 c). The structure of the interface between the carbon substrate and the cluster is not particularly evident in the high resolution images (Figure 1 b), despite the favourable observation geometry. In order to characterize the carbon-cluster interface, EDS spectra were
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collected with a focused electron beam. The analyses, carried out on different clusters, showed systematically a higher O content at this interface. This finding, in agreement with SIMS experiments not reported, can bring another piece of information on the interface structure and indicating the way for the interpretation of the strong bonding of the Ni clusters with the substrate. 1. 2. 3. 4. 5. 6. 7.
I. Martin-Gullon, J. Vera, J. A. Conesa, J. L. . Gonzales, C. Merino, Carbon, 44 (2006), 1572-1580 F.H. Kaatz, M.P.Siegel, D.L. Overmyer; P.P. Provencio, D.R. Talland, Appl. Phys. Lett., 89 (2006), 241915 A. de Lucas, P. B. Garcia, A. Garrido, A. Romero, J.L. Valverde, Appl. Cat. A, 301 (2006), 123-132 Th. Dikonimos Makris, R. Giorgi, N. Lisi, L. Pilloni, E. Salernitano, M.F. De Riccardis and D. Carbone, Fullerenes, Nanotubes and Carbon Nanostructures, vol 13, supplement 1 , 2005, 383-392 M. F. De Riccardis, D. Carbone, Th. Dikonimos Makris, R. Giorgi, N. Lisi, E. Salernitano, Carbon, 44 (2005) 671 M. F. De Riccardis, D.Carbone, Appl. Surf. Sci. 252 (2006), 5403-5407, We kindly acknowledge the technical support of F. Tatti, Application Specialist of FEI Italy.
Figure 1. a: a BF TEM image of the longitudinal cross section of the sample; b: a HRTEM image of the interface between a Ni cluster and the substrate; c: a HRTEM image of a small area of the Ni cluster.
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Near-surface structure of FePt nanoparticles B. Rellinghaus1, D. Pohl1, E. Mohn1, and L. Schultz1 1. IFW Dresden, Institute for Metallic Materials, P.O. Box 270116, D-01171 Dresden, Germany b
[email protected] Keywords: FePt, Nanoparticles, Aberration corrected TEM.
Stoichiometric FePt nanoparticles in the chemically ordered tetragonal L10 phase have gained significant interest in the past decade, since their huge magneto-crystalline anisotropy makes them promising materials for future ultra-high density magnetic data storage media. However, simulations of the equilibrium structure of FePt nanoparticles imply that the formation of the L10 phase may be impeded by a segregation of Pt atoms to the particle surface. Recently, an increase of the lattice constant towards the particle surface – as expected for Pt-enriched surfaces – was reported for FePt nanoparticles [1]. We have therefore systematically investigated the structure of FePt nanoparticles by aberration corrected TEM utilizing a FEI Titan3 80-300 microscope equipped with an image CS corrector (CEOS). FePt nanoparticles were prepared from the gas phase [2,3] and deposited onto amorphous carbon support films. Owing to the clean preparation process, the particle surfaces are free of any stabilizing organics. Particle morphologies and structures range from icosahedral or deceahedral multiply twinned particles (MTPs) to truncated octahedra which are often single crystals. The mean particle diameter is roughly 7 nm. Two typical FePt icosahedra are depicted in Fig. 1 where the phase of the reconstructed exit wave as obtained from the evaluation of a focus series of TEM images is shown. A detailed analysis does not reveal any systematic increase of the lattice spacing upon approaching the particle surface, but a rather statistic variation of any inter-atomic distances (not only the radial spacings). Increased lattice parameters close to the particle surface were only observed, when the particles were terminated by incomplete layers of atoms, or in the vicinity of pronounced (near-surface) defects. The latter is illustrated in Fig. 2 which shows exemplarily a truncated FePt octahedron with a near-surface edge dislocation. As can be seen from the magnification of the defected area in Fig. 2b, the spacing between the last complete (though bent) surface layer to the atoms of an additional incomplete surface layer as manifested by the faint contrast marked by the white arrows is larger than in the depth of the particle. Such defects are more likely to occur in un-equilibrated particles (see, e.g., the twin boundary in the vicinity of the sintering neck to a neighboured particle). The effect of thermal equilibration of the particles by short-time in-flight annealing will be discussed. 1. 2.
R.M. Wang, O. Dmitrieva, M. Farle, G. Dumpich, H.Q. Ye, H. Poppa, R. Kilaas and C. Kisielowski, Phys. Rev. Lett. 100 (2007) 017205. S. Stappert, B. Rellinghaus, M. Acet and E.F. Wassermann, J. Cryst. Growth 252 (2003) 440.
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3.
B. Rellinghaus, E. Mohn, L. Schultz, T. Gemming, M. Acet, A. Kowalik, and B.F. Kock, IEEE Trans. Magn. 42 (2006) 3048.
Figure 1. (a) Reconstructed exit wave (phase image) of two icosahedral FePt nanoparticles lying with their two-fold symmetry axes parallel to the electron beam. (b) Magnification of the area marked by the dashed rectangle in fig. (a). (c) Line profiles as obtained from the areas indicated in fig. (b).
Figure 2. (a) Reconstructed exit wave of a [110]-oriented truncated FePt octahedron at a focus where a near-surface edge dislocation becomes visible. The dashed line indicates a twin boundary (TB) close the sintering neck to an adjacent particle (only partly shown, bottom left corner). (b) Blow-up of the area marked in fig. (a). Dashed lines highlight the dislocation. Arrows indicate incomplete layers of surface atoms.
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Overgrowth of gold nanorods: From rods to octahedrons J.B. Rodríguez-González, E. Carbó-Argibay, I. Pastoriza-Santos, J. Pérez-Juste and L.M. Liz-Marzán Departamento de Química Física, Unidad Asociada CSIC-Universidade de Vigo, 36310 Vigo, Spain.
[email protected] Keywords: TEM, gold, nanorods.
In this work, we use transmission electron microscopy (TEM), and selected area electron diffraction (SAED) to study the growth of previously synthesized monocrystalline gold nanorods (NRs) [1]. When HAuCl4 was reduced on the rods by DMF in the presence of PVP, preferential growth on the sides was obtained, together with sharpening of the tips. TEM images of one original NR, two intermediate particles, and the final stage particle, as well the corresponding SAED patterns, are shown in Figure 1. The starting single-crystal Au NRs, are enclosed within eight {110} and {100} alternating lateral facets, whereas the rod tips are terminated by {100}, {110}, and {111} facets [2]. The SAED analysis (Figure 1) reveals that the tips of the intermediate particles are composed of four {111} facets and indicates that the lateral facets are {110}. The final particles display an octahedral shape whit all facets type {111}, while maintaining a single-crystalline structure [3]. We propose a growth mechanism as sketched in Figure 2. Transformation of the initial rods (Figure 2a) into particles with four {110} lateral facets and sharp tips enclosed by four {111} facets should involve disappearance of the {100} side facets through preferential addition of Au atoms on them (Figure 2b). The morphological transition produced by further growth of the sharp NRs will accordingly be determined by the ratio between the growth rates along the [110] and the [111] directions, this is schematically shown in Figure 2c by the red spheres closing the {110} facet and forcing the {111} facets to join, one with each other, in the final octahedral structure. The described mechanism, based on preferential growth of certain crystalline facets should correlate with a sequence of surface energies in the order {100}>{110}>{111}, which is not in full agreement with the general sequence of surface energies for the different crystallographic Au fcc planes γ(111)< γ (100)20 nm and so in the FFT spots a clearly visible, for the higher oxygen case the particle size decreased into the range of ~10 nm and the polycrystalline structure was increased. For both samples the majority of reflexes correspond to rhombohedral In2O3 whereas some single reflexes could associate with cubic phase!
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d b
Figure 1. (HR)TEM images of In2O3 layers formed by using low (a and b) and high (c and d) oxygen flow while reactive magnetron sputtering. The differences in crystallinity presented in the FFT inset in b and d show smaller particle sizes for the high oxygen flow case. The majority of the FFT reflexes correspond to the rhombohedral In2O3 phase. 1.
C. Falcony, J. R. Kirtley, D. J. Dimaria, T. P. Ma, and T. C. Chen, J. Appl. Phys. 58, 3556 (1985) 2. C. G. Granqvist, Appl. Phys. A 57, 19 (1993) 3. S. Kasiviswanathan and G. Rangarajan, J. Appl. Phys. 75, 2572 (1994) 4. M. Bender, N. Katsarakis, E. Gagaoudakis, E. Hourdakis, E. Douloufakis, V. Cimalla and G. Kiriakidis, J. Appl. Phys. 90, 5382 (2001) 5. Ch. Li, D. H. Zhang, B. Lei, S. Han, X. L. Liu, and Ch. W. Zhou, J. Phys. Chenm. B 107, 12451 (2003) 6. ICDD PDF-2 Data base, JCPDS-Int. Center for Diffraction Data, Pennsylvania, USA (2006). 7. C. A. Pan, and T. P. Ma, Appl. Phys. Lett. 37, 163 (1980). 8. S. Kasiviswanathan and G. Rangarajan, J. Appl. Phys. 75, 2572 (1993). 9. H. Imai, A. Tominaga, H. Hirashima, M. Toki, and N. Asakuma, J. Appl. Phys. 85, 203 (1998). 10. Ch. Y. Wang, V. Cimalla, H. Romanus, Th. Kups, G. Ecke, Th. Stauden, M. Ali, V. Lebedev, J. Pezoldt, and O. Ambacher, Thin Solid Films 515 (2007) 6611–6614 11. This project is supported by the German Academic Exchange Services (DAAD) within project D/06/07398
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Microstructure of Sr4Ru2O9 thin films and Bi3.25La0.75Ti3O12/Sr4Ru2O9 bilayers R. Chmielowski1,2, V. Madigou1, M. Blicharski2, Ch. Leroux1 1. IM2NP (UMR-CNRS 6242), South University Toulon-Var (USTV), Bat.R , B.P.20132, 83957 La Garde Cedex, France 2. AGH - University of Science and Technology, Al. Mickiewicza 30, 30-059 Kraków, Poland
[email protected] Keywords: Sr4Ru2O9, pulsed laser deposition, oxide electrode
It has been reported that using oxide bottom electrodes reduces the fatigue phenomenon in ferroelectric random access memories (FeRAM). For that aim, strontium ruthenate SrRuO3 was already used as bottom oxide electrode for ferroelectric thin films [1], but compounds with less ruthenium are more interesting from the economical and environmental point of view. We elaborate Sr4Ru2O9 thin films and bilayers Bi3 25La0 75Ti3O12 /Sr4Ru2O9 by pulsed laser deposition. Sr4Ru2O9 compound differs from the other strontium ruthenates. It has a hexagonal perovskite structure, contains pentavalent ruthenium atoms, with octahedral sharing faces, thus building dimmers Ru2O9 [2]. All the other strontium ruthenates belong to the Ruddlesden Poppers homologous series, Srn+1RunO3n+1 (Sr2RuO4 corresponds to n=1 and SrRuO3 to n=∞). The common features of these compounds are an orthorhombic structure, tetravalent ruthenium atoms, and an octahedral surrounding by oxygen of ruthenium atoms. The films were grown by pulsed laser deposition on Si [100] substrate, using a Sr2RuO4 target with an excimer laser (KrF λ=248 mm, COMPex 301, Lambda Physik). The laser beam had a fixed size of 2mm x 5 mm and the fluence of the laser on the target was 1.5 Jcm-2. The substrate was heated up to 700°C and the films were deposited in two different oxygen pressures, 50 10-3 Torr and 300 10-3 Torr [3]. The thin films and the targets were characterized by X-ray diffraction; in order to identify the strontium ruthenate phases. The surface morphology of the target and the thin films was observed with a JEOL JSM-6320F high resolution scanning microscope. A Tecnai G2 transmission electron microscope, operating at 200 kV with a LaB6 filament, equipped with a 1k x 1k Slow Scan CCD camera, was used for microstructural characterizations. All electron diffraction patterns obtained with various zone axes were indexed in the Sr4Ru2O9 hexagonal structure and each time, diffuse streaks parallel to a* or b*, were present (Figure 1). The occurrence of forbidden spots could be explained by an ordering of Ru atoms. In the case of Sr4Ru2O9 films, transmission electron microscopy revealed, on cross section samples, the existence of a supplementary strontium ruthenate phase, in form of nanograins, layered on top of the Sr4Ru2O9 columnar grains, for all the thin films deposited with 50 mTorr oxygen pressure in the deposition chamber. This is at the
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origin of the intermediaire layer observed the ferroelectric film and the oxide electrode. For thin films Bi3 25La0 75Ti3O12 /Sr4Ru2O9, this layer is single phase and has a SrTiO3 structure. 1. 2. 3.
H.N.Lee, D.Hesse, N.Zakharov,U.Gösele, Science 296 (2002), 2006 C.Dussarat, J;Fompeyrine, J.Darriet, Eur.J. Solid State Inorg. Chem. 32, (1995), 3 R. Chmielowski, V. Madigou, Ph. Ferrandis, R. Zalecki, M. Blicharski, Ch. Leroux, Thin Solid Films, 515 (2007), 6314
Figure 1. Cross sectional view of Bi3 25La0 75Ti3O12 /Sr4Ru2O9 Diffraction patterns taken on the Sr4Ru2O9 films, corresponding to [010] (a) and to [201] zone axes (b).
BLT Intermediate layer Sr4Ru2O9 100 nm
Figure 2. Cross sectional view of a Bi3 25La0 75Ti3O12 /Sr4Ru2O9 bilayer, showing the occurrence of a well crystallized intermediate layer.
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Analysis of the LSM/YSZ interface on microand nano-scale by SEM, FIB/SEM and (S)TEM Yi-lin Liu, Luise Theil Kuhn and Jacob R. Bowen Fuel Cells & Solid State Chemistry Department Risø National Laboratory for Sustainable Energy, Technical University of Denmark DK-4000 Roskilde, Denmark
[email protected] Keywords: interface, nano-structure, FIB/SEM, (S)TEM
8 mol% Y2O3-stabilized ZrO2 (8YSZ) and lanthanum strontium manganite (LSM) are well developed materials for electrolyte and cathode, respectively, in solid oxide fuel cells (SOFCs). Thermodynamically the 8YSZ/LSM interface is not stable and solid state reactions will result in formation of zirconates (La2Zr2O7, SrZrO3) which are electrically insulating phases. For cells operated at 1000˚C, zirconates were found covering large areas of the interface: this was readily characterized by XRD, SEM, TEM and identified as a degradation mechanism in SOFC [1,2]. However, the newly developed SOFCs are operated at much lower temperatures (700-800oC). Due to the slow reaction kinetics, the zirconate formation is localized and its growth is limited. In this case, identification of microstructural degradation has become a great challenge. This paper presents how such an interface has been studied by FEGSEM, FIB and (S)TEM combined with EDS from the micro- to nanometre scales. The interface sample used in this work was taken from a cell that has been tested at 750˚C for 1500h. The cathode was removed by HCl etching (LSM dissolves quickly whereas 8YSZ and La2Zr2O7 are not affected in HCl) and then the electrolyte surface was analyzed using a Zeiss Supra FEGSEM. A low accelerating voltage of ~2 kV was used which allowed direct imaging of the 8YSZ surface without carbon coating. The contact points left over by (dissolved) LSM grains (a few hundred nanometre) appear as craters, and many nanoparticles are found in association with these craters (Figure 1). To identify these nanoparticles lining the craters, a site-specific TEM lamella was made using a Zeiss 1540XB CrossBeam microscope. Prior to the FIB milling several protective layers of W were deposited on the area where the lamella later was extracted. Firstly, to avoid excessive Ga ion damage to the nanoparticles a 60 nm thick W layer was locally electron beam deposited using the in-situ chemical vapour deposition gas injection system (GIS). Subsequently, the Ga ion beam was used with the GIS to deposit successively thicker W protection layers with increasing Ga ion beam probe currents. After building up a 1.5 µm thick protective layer the lamella was extracted using the standard in-situ TEM lamella preparation technique. The lamella dimension was 1 x 30 x 8 µm3 (thickness, length and depth, respectively) prior to mounting on a Cu Omniprobe™ TEM sample holder. Final thinning was performed in a 24 x 2 µm window until the region with the crater surfaces was approximately 50 nm thick suitable for high resolution TEM, STEM and EDS investigation (Figure 2).
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The (S)TEM analysis was performed using a 300kV Jeol 3000F TEM equipped with a Gatan Imaging Filter (GIF), a STEM unit and an Oxford Inca EDS system. The craters were studied in normal TEM and in STEM mode, and their chemical composition and environment were analysed by site-specific nanobeam EDS with an average electron beam diameter of 10 nm. Nanoparticles with sizes in the range 3-20 nm (see Figure 3) were only observed in the crater regions confirming the SEM studies. EDS analysis was performed in several crater regions and also at the crater environment to detect any changes in chemical composition. The EDS analysis of nanoparticles showed that there are clear observations of La- and Sr-phases further confirming the hypothesis that the nanoparticles are a result of the cathode degradation and are closely connected with the contact points of the LSM grains. The TEM and EDS studies were supplemented by STEM imaging where the enhanced contrast favoured imaging of the crater rims and the nanoparticles. 1. 2.
E. Ivers-Tiffee, A. Weber, K. Schmid and V. Krebs, Solid State Ionics, 174 (2004) p. 223. A. Mitterdorfer and L.J. Gauckler, Solid State Ionics, 111(1998) p. 185.
Figure 1. SEM image showing LSM craters and nano-particles on the 8YSZ electrolyte surface.
Figure 2. SEM image of the TEM lamella after extracted and thinned.
Figure 3. TEM image showing the cross section of a crater region with nanoparticles. The light and dark region to the left of the diagonal is the protective layer of W and the dark contrast is caused by implanted Ga. Only a few % of Ga was detected in the nanoparticles, so we assume that they were not affected by the Ga-ion beam. The circles indicate some of the larger nanoparticles at the electrolyte (dark contrast to the right) interface.
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ESI and HRTEM of chemical solution deposited (CSD) YBa2Cu3O7-δ coated conductors L. Molina1, T. Thersleff2, B. Rellinghaus2, B. Holzapfel2, O. Eibl1 1. Institut für Angewandte Physik, Eberhard-Karls Universität Tübingen Auf der Morgenstelle 10, D-72076, Tübingen, Germany 2. IFW Dresden, P.O. Box 270116, D-01171, Dresden, Germany leopoldo
[email protected] Keywords: thin films, interfaces, coated conductors
Coated conductors are currently of great interest for technological applications which involve transportation of large currents. TEM Sample preparation techniques and CTEM analysis are presented in other papers at this conference. Control of the stoichiometry of CSD prepared films is a key issue for coated conductor technology. Performance limiting secondary phases have to be eliminate or minimized during film deposition. Thus, a reliable nano-scaled chemical analysis is crucial for coated conductor fabrication. The chemical composition and distribution within YBCO-coated conductor samples of Ni5at%W(substrate)/La2Zr2O7(CSD)/CeO2(CSD)/YBCO(CSD) layer architecture was studied by energy-filtered TEM (ESI) in a Zeiss 912 OMEGA operating at 120 kV. By electron spectroscopic imaging (ESI), elemental maps with high lateral resolution of YBCO-coated conductor samples in cross-section were obtained. Fig.1 shows an RGB overlay of Ce, Ba, Y and La elemental maps. The used ionization edges are outlined in table 1. The Ni substrate, the La2Zr2O7 and the YBCO thin film appear in strong contrast. An intermediate layer is found between the Ni and the LZO. Two YBCO layers were deposted for this sample by CSD and these layers and their interface are clearly imaged, e.g. 100 nm Yttria secondary phases (blue in fig. 1) which were identified by energy dispersive X-ray microanalysis in the TEM. Understanding the Ni/oxide interface is of crucial importance for the deposition of biaxially textured LZO buffer layers. The small lattice parameters of the Ni substrate requires a point resolution of significantly better than 0.2 nm. Therefore, the crosssection sample of fig.1 was analysed in a FEI-TITAN 300 kV transmission electron microscope equipped with a Cs corrector yielding a point resolution of better than 0.1 nm. Fig.2 shows a cross-section of a Ni/oxide interface. 1.
Molina L, Knoth K, Engel S, Holzapfel B, Eibl O 2006 Supercond. Sci. Technol. 19 12001208
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Figure 1. ESI chemical maps of a full YBCO-coated conductor sample. R(Ce)G(Ba)B(Y) overlay image and ESI R(Ce)G(Ba)B(La) overlay image Table 1. a.) Ionization edges used. Ce (N 4,5)
La (N 4,5)
La (M 4,5)
Ba (N 4,5)
(eV)
(eV)
(eV)
(eV)
ΔE1
76
76
127
60
ΔE2
102
95
149
80
ΔE3
137
118
184
111
Edges
Pre-edge 1 Pre-edge 2 Post-edge
Figure 2. HRTEM of the Ni /oxide interface. Lattice spacing in Ni is 0.2 nm.
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CTEM diffraction contrast of biaxially-textured La2Zr2O7 buffer layers on nickel substrates L. Molina1, S. Engel2, B. Holzapfel2, O. Eibl1 1. Institut für Angewandte Physik, Eberhard-Karls Universität Tübingen Auf der Morgenstelle 10, D-72076, Tübingen, Germany 2. IFW Dresden, P.O. Box 270116, D-01171, Dresden, Germany leopoldo
[email protected] Keywords: thin films, interfaces, coated conductors
Chemically deposited La2Zr2O7 (LZO) buffer layers are of crucial importance for YBCO-coated conductor fabrication. They transfer the texture from a highly biaxially textured nickel tungsten substrate to the final YBCO superconducting layer and act as nickel diffusion barriers. The misfit of LZO with respect to the nickel tungsten substrate is 7.6% (compressive). LZO films were grown by Chemical Solution Deposition (CSD) and were annealed at temperatures between 900°C and 1000°C [1]. The films are highly biaxially textured as suggested by the XRD pole figures shown in figure 1 [2]. Due to the XRD pole figures an epitaxial growth of LZO is claimed in the literature, which was the motivation for this investigation. Conventional TEM two-beam diffraction contrast imaging is sensitive to small misorientation of grains and allows to image small-angle grain boundary networks present in the LZO thin films. Carefully tilting the sample under two-beam conditions reveals grains of 100-200 nm in size that are slightly titled with respect to each other and to the underlying Ni grain in the substrate, (figure 3 plan-view LZO sample). Dark areas are grains that are in a strong diffraction condition. The contrast of these images are particularly sensitive to the out-of-plane tilt of the crystallites. Figure 2 shows a similar sample prepared as cross-section and LZO grains of 100-200 nm can be identified. The average Ni substrate grain size is 40 μm and the LZO grain size is 100200 nm, therefore epitaxial growth of LZO on Ni does not occur . Figure 4 shows a dark-field image at the Ni-LZO interface under two-beam diffraction conditions using the g(400) reflection of LZO. Bright areas show LZO grains that are in a strong diffraction contrast condition denoted in the figure as grain 1 to 5. Also observed in the images is the strong contrast from the nanovoids present in the films [3]. At the Ni /LZO interface a roughness with a wavelength of 60 nm and an amplitude of 10 nm is found. LZO grain size was beyond 100 nm and agrees with plan view results. The LZO grains are shown to extend over the full film thickness. The tilting of the LZO grains and the roughness at the interface contribute to strain relief in the LZO film. 1. 2.
Molina L, Engel S, Knoth K, Hühne R, Holzapfel B, Eibl O 2008 Journal of Physics: Conference Series 97 (2008) 012108 Molina L, Knoth K, Engel S, Holzapfel B, Eibl O 2006 Supercond. Sci. Technol. 19 12001208
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Figure 1. (a-b) XRD pole figures. Figure 2. TEM bright-field images under two beam conditions of an LZO buffer layer sample in crosssection (a-c). Nano-voids 10-20 nm in size can be seen. (d-f) corresponding dark-field images. Note that all images are from the same sample area. The white arrow indicates a LZO grain.
Figure 3. TEM bright-field image tilt series of an LZO buffer layer sample in planview. Scale bar is 100 nm
Figure 4. Dark-field image under two beam conditions of an LZO buffer layer sample in cross-section. The horizontal white arrow indicates the interface. Vertical arrows show roughness at the interface.
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TEM sample preparation of YBCO-coated conductors: conventional method and FIB L. Molina1, T. Thersleff2, C. Mickel2, S. Menzel2, B. Holzapfel2, O. Eibl1 1. Institut für Angewandte Physik, Eberhard-Karls Universität Tübingen Auf der Morgenstelle 10, D-72076, Tübingen, Germany 2. IFW Dresden, P.O. Box 270116, D-01171, Dresden, Germany leopoldo
[email protected] Keywords: thin films, interfaces, coated conductors
YBa2Cu3O7-δ (YBCO)-coated conductors prepared by the chemical solution deposition process (CSD) are currently of great interest for industrial applications. A flexible nickel tungsten substrate is dip-coated in a buffer layer and a YBCO precursor solution. The resulting thin films systems have a high misfit. TEM sample preparation is challenging and is a fundamental issue for the characterization of YBCO-coated conductors. Conventional sample preparation involves gluing the sample of interest between oxide and silicon dummies followed by mechanical polishing and grinding [1]. Further thinned is done using Ar+ ions in a Baltec Res 100 ion milling machine operating at 4.5 kV and 3.5 mA with etching angles ranging from ±12° to ±6° for up to 30 hrs. The insitu lift-out focused ion beam (FIB) allows to prepare as TEM lamellas previously choosen sample areas. Figure 1 describes the preparation procedure. A Pt protection layer is deposited on top of the area of interest in the YBCO film. Trapezoid shaped cuts area made with a Ga-ion beam. The in-situ lift-out is done with a nanomanipulator. The sample is attached to a sample holder and further ion milled until reaching ~100 nm thickness. Samples prepared by this method are highly homogenous which makes them especially suitable for electron spectroscopic imaging (ESI) and HRTEM. 1.
Eyidi D, Eibl O 2002 Micron 33 499
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Figure 1.- a.) Pt protection layer deposited on the area of interest on top of the YBCO film. Scale bar is 1 μm b.) Trapezoid shaped cuts made with the Ga-ion beam. Scale bar is 1 μm c.) In-situ lift out of the TEM lamella done with a nano-manipulator. Scale bar is 10 μm d.) Attachment of the TEM lamella to the half-grid sample holder Scale bar is 100 μm e.) Encircled area shows were the sample was attached. Scale bar is 100 μm f.) Overview of the polished final TEM lamella. Scale bar is 1 μm.
Figure 2.- a.) Zero-loss dark-field image of FIB prepared TEM lamella. b.) Zero-loss dark-field image of a conventionally prepared sample.
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Nucleation and evolution of biepitaxial YBa2Cu3O7-δ thin film grown on SrTiO3 and MgO substrates H. Pettersson1, K. Cedergren2, D. Gustafsson2, R. Ciancio1, F. Lombardi2 and E. Olsson1 1. Dept. of Applied Physics, Chalmers University of Technology, Göteborg, Sweden 2. Dept. of Microtechnology and Nanoscience, Chalmers University of Technology, Göteborg, Sweden
[email protected] Keywords: Nucleation, thin films, YBa2Cu3O7-δ, SrTiO3, MgO
The importance of the microstructure for the transport properties of grain boundaries in high-temperature superconducting YBa2Cu3O7-δ (YBCO) (e.g. SQUID-based devices) is well established. Basal-plane grain boundaries can be used as high-quality Josephson junctions [1,2]. In this work artificial grain boundaries (AGB) of (103)/(001) YBCO films produced by pulsed laser deposition (PLD) have been investigated, in order to obtain knowledge of the YBCO film nucleation and the subsequently grain boundary evolution. For this purpose ultra thin films (50 and 100 pulses) and full films have grown on two different substrate/template-layer combinations (110) MgO / (110) SrTiO3 (STO) and (110) STO / (110) CeO2 (CEO) respectively. The (103)/(001) YBCO biepitaxial artificial grain boundaries were obtained on substrates with a vicinal cut, in order to achieve pure (103) YBCO growth on top the STO regardless if STO is used as substrate or template-layer. The YBCO grow with (001) orientation and 45º in-plane rotation on the MgO and CEO. The microstructural investigation has been performed with scanning electron microscopy (SEM) using a Leo Ultra 55 FEG SEM. The transmission electron microscopy (TEM) investigation was carried out using a Philips CM 200 FEG and JEOL JEM 2100-F both operated at 200kV. The results show that the nucleation site (see Fig. 1) of the YBCO depends on the morphology of the substrate close to the step (See Fig. 2). Both (103) and (-103) YBCO are found at the YBCO/STO interface due to surface roughness, exposing both (100) and (010) facets. The nucleation of YBCO film and subsequently growth at the step edge, produced by the etching procedure, affects the grain boundary evolution. By growth of ultra thin YBCO films the nucleation sites of YBCO could be determined. 1. 2.
Granozio, F.M., et al., Structure and Properties of a Class of CeO2-based Biepitaxial YBa2Cu3O7-™ Josephson Junctions. Physical Review B, 2003. 67(18). Tafuri, F., et al., Microstructure and Josephson phenomenology in 45 degrees tilt and twist YBa2Cu3O7-™ artificial grain boundaries. Physical Review B, 1999. 59(17): p. 11523-11531.
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a)
(103) YBCO c)
b)
(001) YBCO (001) YBCO
(103) YBCO
(103) YBCO
(001) YBCO
d)
(001) YBCO
(103) YBCO
Figure 1. SEM micrographs of (103)/(001) YBCO films grown on STO substrates with a) 50 pulses, b) 100 pulses and c) full thickness, (-103) oriented YBCO growth is marked (red), and d) full thickness on MgO substrate
Figure 2. TEM micrograph of (103)/(001) YBCO films grown on MgO substrate.
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An investigation of Al-Pb interfaces using probe-corrected high-resolution STEM Harald Rösner1, Sergei Lopatin2, Bert Freitag2 and Gerhard Wilde1 1. Westfälische Wilhelms-Universität Münster, Institut für Materialphysik, Wilhelm-Klemm-Str. 10, D-48149 Münster, Germany 2. FEI Company, Achtseweg Noord 5, 5600 KA Eindhoven, The Netherlands
[email protected] Keywords: high-resolution STEM, probe Cs-corrector, misfit dislocation
Nanometre-sized Pb inclusions embedded in an Al host matrix serve as a model system for size-dependent melting studies [1-3]. For this purpose, an Al-Pb composite containing 1 at.% Pb was prepared by high-energy ball milling. Former TEM investigations revealed that the metal/metal interfaces were found to display a cube-on-cube orientation relationship with a large lattice mismatch of 22% with respect to the Al lattice [14]. The accommodation of the lattice misfit has remained a puzzle until quite recently due to Moiré contrasts that veiled a direct observation of the interface structure. Recently, this mismatch has been visualized by high-resolution TEM of uncovered Pb inclusions [4,5] displaying a hetero-interface with the Al matrix remaining on one side only. Such Pb particles were located in the amorphous edge area of the TEM specimen. In this study, partially embedded Pb inclusions located in the thicker part of the TEM foil, having hetero-interfaces on all lateral sides with the Al matrix, have been analyzed. For this purpose, high-resolution HAADF-STEM investigations have been carried out in a FEI Titan 80-300 (field-emission gun, super-twin lens, Cs= 1.2 mm, monochromator) operated at 300 kV and equipped with a CEOS Cs probe-corrector. The spherical aberration of the probe has been corrected down to the level of a few microns. Fourier-filtering was performed on the high-resolution micrograph using the Digital Micrograph Software Package (Version 3.7.4, Gatan). Circular masks with a diameter of 0.5 nm have been used for the Fourier-filtering in Figure 2. In Figure 1 two Pb inclusions embedded in an Al matrix are shown. The larger inclusion appears faceted but without a Moiré pattern. The smaller one has a spherical shape and a strong Moiré contrast. The insert in Figure 1 is the Fourier transform of the whole image. The Al and Pb reflections of the [011] zone axis confirm the cube-oncube orientation relationship. The observation of a partially encased faceted Pb inclusion in the thicker part of the TEM foil allows a direct comparison with the results of former studies [4,5] where the Pb inclusions investigated were located in the thinnest parts of the TEM foil. Here by use of high-resolution STEM imaging instead of HRTEM, one may discern that even in the thicker parts of the TEM foil the lattice misfit at the Al/Pb interface is localized in the form of interfacial dislocations. In the absence of a Moiré pattern, the faceted Pb inclusion shown in Figure 2 can be seen to display a semicoherent interface with the Al matrix. A misfit on about every fifth Al plane is observed; white lines indicating their positions have been added to guide the eye. In addition, the current observation provides the opportunity to learn more about the particle S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 359–360, DOI: 10.1007/978-3-540-85226-1_180, © Springer-Verlag Berlin Heidelberg 2008
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shapes with respect to matrix strains [6]. Since most of the Pb inclusions in ball-milled Al-Pb composites exhibit curved non-faceted shapes, the question arises why in this case one of the Pb inclusions shows a faceted morphology. The answer may be found in the fact that this Pb inclusion was only partially embedded in the Al matrix without an overlapping Al layer above or underneath. As a result the strain relief is thought to be accomplished via these two ‘free’ surfaces by the high mobility of the Pb atoms [1]. In contrast, the fully embedded Pb inclusion nearby, displaying a curved morphology, is totally governed by strains of the heavily distorted Al matrix. A strain relief in the distorted Al matrix of ball-milled material achieved by annealing also led to a change in particle shapes [6]. Thus, faceted interfaces account for a lower total excess free energy of the nanoparticle and hence, a melting point increase. 1. 2. 3. 4. 5. 6. 7.
E. Johnson, H.H. Andersen, and U. Dahmen, Microsc. Res. Techniq. 64 (2004), p. 356. K. Chattopadhyay and R. Goswami, Prog. Mater. Sci. 42 (1997), p. 287. Q.S. Mei and K. Lu, Prog. Mater. Sci. 52 (2007), p. 1175. H. Rösner, J. Weissmüller, and G. Wilde, Phil. Mag. Lett. 86 (2006), p. 623. H. Rösner, B. Freitag, and G. Wilde, Phil. Mag. Lett. 87 (2007), p. 341. H. Rösner, P. Scheer, J. Weissmüller, and G. Wilde, Phil. Mag. Lett. 83 (2003), p. 511. The experiments have been performed during a FEI Demo at the Ernst-Ruska Centre in Jülich.
Figure 1. High-resolution HAADF-STEM micrograph showing two Pb inclusions embedded in an Al matrix projected along the [011] zone axis. The larger particle shows clear facets and no Moiré pattern. The smaller Pb inclusion appears spherically shaped and with a Moiré pattern due to the overlapping Al matrix. The inset shows the Fourier transform revealing the cube-on-cube orientation relationship between the two lattices.
Figure 2. Detail of Figure 1 showing the faceted Pb inclusion after Fourier-filtering. Masks were set around all reflections in order to remove the noise. The lattice mismatch is indicated by the white lines. No Moiré pattern is visible.
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Spectrometric Full-Color Cathodoluminescence Electron Microscopy Study of Grain Boundaries of ZnO Varistor H. Saijo1, N. Daneu2, A. Recnik2, and M. Shiojiri3 1. School of Biology-Oriented Science and Technology, Kinki University, 930 Nishimitani, Kinokawa, Wakayama 649-6493, Japan 2. Jozef Stefan Institute, Jamova 39, SI-1001 Ljubljana, Slovenia 3. Prof. Emeritus of Kyoto Institute of Technology, 1-297 Emmyoji-Wakiyama, Ohyamazaki, Kyoto 618-0091, Japan (home) hirsaijo@waka kindai.ac.jp Keywords: cathodoluminescence, ZnO, varistor
Scanning electron microscopy (SEM) with cathodoluminescence (CL) detection can draw a map of electronic states and chemical bonding in the specimen. The energy of CL photons gives information of Band Gap of semiconductors or Highest Occupied Molecular Orbital (HOMO)-Lowest Unoccupied MO (LUMO) of organic materials. Our new spectrometric full-color CL microscope [1] obtains CL spectra of each observing point and draws full-color CL micrograph and SEM image of 512 x 512 pixels in 8 s. The block diagram of the system is shown in Fig. 1. Spatial resolution of CL micrographs is worse than that of SEM due to bulb-shape spread of incident electrons in solids. Delayed CL emission is another cause of poor resolution. Both result practical CL resolution around 100 nm. However, lower incident energy and slow image scan improve the resolution significantly, and in our system 25 nm resolution was verified with 6 kV beam and 40 s/frame. With this high resolution, we observed grain boundaries of ZnO varistor ceramics with 1 mol % of Bi2O3. The specimen also contains SnO2 from 0.1 to 10 mol %. The grain size of the ZnO decreased with increasing amount of SnO2. [2] However, electronic properties did not show significant difference with SnO2 concentration. We examined CL spectra of boundaries, surface and cutout plane of the ZnO grains as shown in Figs. 2 and 3. The cutout plane can be regarded as the inside bulk of the grain. The grain surfaces emit strong bluegreen CL, and the bulk emits weak or no luminescence. CL observation across the boundary revealed that several peaks appear and disappear depending on the position from the boundary. 500 nm peak appears first (Fig. 2-7), and then 510 nm peak rises if mouse-pointer moved slightly inward the grain. Then, 490 nm peak appears forming twin with 510 nm peak. Further inside shows various CL peaks as shown in Fig. 3-4 and 3-5. CL peaks between 490 and 510 nm observed across the boundary correspond to the layers with different composition, and all of them appear in every specimen of different SnO2 concentrations in the same order, which explains a non-linearity parameter α the same. The layer thickness of each peak becomes thin with SnO2 concentration to result the threshold voltage decrease and leak current increase with increasing SnO2 [2]. 1. 2.
H. Saijo and M. Shiojiri, in “Proc. 16th Intern’l Micros. Congr. Sapporo”, Vol. 2 (2006), p. 883. 2. N. Daneu, A. Recnik, S. Bernik and D. Kolar, J. Am. Ceram. Soc., 83 3165 (2000).
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Figure 1. Block Diagram of CL Microscope
Figure 2. CL observation of SnO2 0.1% and Bi2O3 1% doped ZnO varistor.
Figure 3. CL observation of SnO2 10% and Bi2O3 1% doped ZnO varistor. 1; SEM image, 2; CL micrograph of 1, 3-8; CL spectra of points marked in 1.
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Study of structural properties of Mo/CuInS2/ZnS used in solar cells by TEM J. Sandino1, G. Gordillo2, H. Lichte1 1. Triebenberg Laboratory, Institute of Structure Physics, Technische Universitaet Dresden, D-01062 Dresden, Germany 2. Departamento de física, Universidad Nacional de Colombia, Bogotá, Colombia
[email protected] Keywords: TEM, Thin films solar cells, structural properties, TEM
One of the steps to improve the efficiency of thin film solar cells is the exhaustive comprehension of the physical aspects present in the bulk of materials and their interphases. TEM is a tool that provides significant information about the specimen, by means of e.g. analysis of crystalline structure, defects visualization, grain boundaries and interfaces, and composition mapping. Additionally, Electron Holography allows measuring the inner potentials and electric fields [1], which will be decisive for the performance of thin films solar cells. Here, solar cells consisting of Mo/CuInS2/ZnS are investigated (Figure 1). CuInS2 has potential advantages in comparison with the most widely investigated selenium containing chalcopyrites. These include the significantly reduced toxicity of sulfur compared to selenium, which minimizes environmental impacts and facilitates mass production handling; furthermore, fabrication can be performed with short processing times and low thermal budget, quite appropriate for commercial application. The energy gap of the CuInS2 is close to the ideal one (1.5 eV) for a single junction solar cell [2]. For the performance of the solar cells, grain size distributions, composition of each layer, and structural mismatch between the different layers have a fundamental importance. For example, in the XRD spectra of a CuInS2-layer [4], obtained under variation of the incidence angle between 1.5 and 4º (see Fig. 2), only reflections corresponding to the In2S3 and CuInS2 phases appear. This indicates that a Cu2S phase is formed predominantly at the bottom of the sample, whereas In2S3 and CuInS2 phases grow in the whole volume [2]. For a more detailed analysis, a TEM investigation is needed. For this work, Mo/CuInS2/ZnS thin films were prepared with different deposition methods of the ZnS-layer, in order to analyze their influence on structural and electrical properties under TEM and Electron holography. The multilayer Mo/CuInS2/ZnS system in both samples was sequentially prepared on a substrate consisting of soda lime glass. First, Mo was deposited by DC magnetron sputtering with a S-gun configuration electrode. Next, for CuInS2, two sequential steps using elemental sulphur evaporated from an effusion cell, and metallic precursors of Cu and In evaporated from tungsten boats were used. Finally, ZnS was deposited by CBD (chemical bath deposition) and by co-evaporation, respectively. The CBD-deposited ZnS was grown from the following chemical bath composition [zinc acetate] = 15×10-3 M; [sodium citrate] = 7.5×10-3 to
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60×10-3 M; [ammonia] = 300×10-3 M; [thiourea] = 150×10-3 M. Alternatively, a coevaporated ZnS-layer was grown from elemental evaporation sources; a detailed description of the process of deposition of each layer is given in [3]-[6]. Finally, the samples were prepared by traditional cross-sectioning and by FIB, and the TEM and HTEM measurements were accomplished using a CM200 ST/Lorentz Philips TEM. First results are shown in Figures 3-4. [7]. 1. 2. 3. 4. 5. 6. 7.
P.Simon, H. Lichte, D. Mönter, W. Reschetilowski, A. Valera, and W. Carrillo-Cabrera, 631, ZAAC (2005), p.983. N. Meyer, I. Luck, U. Ruehle, R. Klenk, M. C. Lux-Steiner, R. Scheer, Proc. 19th European Photovoltaic Solar Energy Conference, Munich, 2004, p.1698 G. Gordillo; F. Mesa; C. Calderón, Brazilian Journal of Physics 36 (2006), p. 982. J. Clavijo, E. Romero, J. S. Oyola and G. Gordillo, Proc. 22nd European Photovoltaic Solar Energy Conference, 2007, p. 2279. M. Ladar, E.J. Popovici, I. Boldea, R. Grecu, E. Indrea, Journal of Alloys and Compounds, 434 - 435 (2007) 697-700 G.Gordillo, E. Romero, Thin Solid Films 484 (2004) p.352 We kindly acknowledge supply of samples by the group of Materiales Semiconductors y Energía Solar de la Universidad Nacional de Colombia, financial support of the agreement ALECOL and DAAD. Discussions within the Triebenberg lab were indispensable.
InS CuS
ZnS CuInS2 Molybdenum Subtract (soda-lime glass)
Figure 1. Scheme of the Mo/CuInS2/ZnS system. Drawing is not to scale.
Figure 2. XRD spectra measured under small angle incidence for a CuInS2 film ZnS
CuInS2
Mo
Figure 3. TEM micrograph of the CuInS2 thin films analyzed.
0 1µm 1 0µm
1 1µm
Figure 4. TEM micrograph of the Mo/CuInS2/ZnS system prepared by CBD.
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Texture analysis of silicide thin films: combining statistical and microscopical information H. Schletter1, S. Schulze1, M. Hietschold1, K. De Keyser2, C. Detavernier2, G. Beddies1, A. Bleloch3, and M. Falke1 1. Institute of Physics, Chemnitz University of Technology, 09107 Chemnitz, Germany 2. Department of Solid State Physics, Ghent University, 9000, Belgium 3. SuperSTEM Laboratory, Daresbury, WA4 4AD, Cheshire, UK
[email protected] Keywords: EBSD, texture, SEM
The statistical distribution of grain orientations in a poly-crystalline thin film, called texture, is an important characteristic. The presence of preferred orientations is important for the macroscopical behaviour of the film due to an anisotropy of certain properties, like resistivity, carrier mobility, etc. Therefore it is important to know, which types of texture evolve for a certain combination of substrate and film material and in which way they depend on the sample preparation. In our measurements, electron backscatter diffraction (EBSD) was used to reveal the texture of the films. On one hand side, this technique provides good statistical data, because a large number of grains (in the order of 10³) is examined. On the other hand, because the measurement is carried out in an SEM, the information about grain orientation is available for every single point on the sample surface. Therefore, direct correlations between texture components (statistical data) and film morphology (microscopical data) can be drawn. As an additional source of information, high resolution TEM investigations were carried out. These measurements verify the EBSD results and reveal the interface structure at an atomic scale. The samples we investigated consisted of thin films (approximate thickness d = 50 nm) of CrSi2 and MnSi1 7 on Si(001) single crystal substrates. The films were grown by reactive codeposition on the heated substrate under UHV conditions. For some samples, a very thin metal layer was deposited as a template prior to the film growth. We found different kinds of texture present in all samples. Besides the various known epitaxial relations [1], we could identify new crystal orientations which belong to epitaxy, fiber texture and axiotaxy. The statistical data are visualized in pole figures. An example for CrSi2 is shown in Figure 1. The lines of high intensity, which are not concentric in the pole figures, are of special interest for us, since they refer to axiotaxial alignments. This orientation was identified as a new type of texture only a few years ago [2]. For CrSi2, the axiotaxial relation is CrSi2{100}||Si{110}. For MnSi1 7, two different axiotaxial orientations were found, namely MnSi1 7{118}||Si(110) and MnSi1 7{110}||Si(110). By analysing grain orientation maps (shown in Figure 2) and the corresponding SEM images, a correlation between texture components and grain size was found. The
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pictures show, that the largest grains have an epitaxial orientation while randomly oriented grains have the smallest size on average. Therefore we conclude that the different texture components correspond to different growth conditions for the crystallites. 1. 2.
O. Filonenko et al., Journal of Crystal Growth 262 (2003), 281 C. Detavernier et al., Nature 426 (2003), 641
Figure 1. Statistical pole figure of CrSi2 (left) + theoretical lines for axiotaxy and spots for epitaxy (right).
Figure 2. Corresponding grain orientation map (different colours / grey levels represent different CrSi2 crystal planes being parallel to the silicon substrate surface).
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Statistical Tomography of 3D Thin Film Structure using Transmission Electron Microscopy E. Spiecker1,2, V. Radmilovic2 and U. Dahmen2 1. Faculty of Engineering, Christian-Albrechts-University Kiel, Kaiserstrasse 2, 24143 Kiel, Germany 2. National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, USA
[email protected] Keywords: thin films, transmission electron microscopy, tomography
Thin films play a crucial role in many modern technologies, like microelectronics, solar cells, sensors, and coatings. A key to understanding and controlling thin film growth processes and properties is the knowledge of the variation in structure with distance from the substrate. Therefore experimental techniques for quantitative characterization of the three-dimensional (3D) structure of thin films are highly demanded. Transmission electron microscopy (TEM) is well established as a powerful tool for investigation of thin film structures down to the nanometer scale. However, cross-section and plan-view geometries commonly used in TEM studies of thin films are not suited for quantitative evaluation of 3D data since these geometries represent only two particular 2D sections through the film. On the other hand most film structures are much too complicated for application of standard 3D tomography techniques. We have recently developed a new sample preparation technique that overcomes these limitations by a double wedge geometry that probes different depths in the film and separates depth information in a continuous series of thin slices that are spread over a lateral distance [1,2]. The key steps of the preparation are (i) the formation of an ultrashallow dimple (slope ~ 0.1º) in the film surface, (ii) the determination of the dimple profile, and (iii) the formation of a small-angle wedge whose edge intersects the center of the dimple (Fig. 1). Along the edge of the double-wedge sample quasi-horizontal 2D sections at different film depths can be analyzed by plan-view TEM over large areas, thus enabling a statistical description of the 3D film structure. Fig. 2 depicts results of an application of this technique to a 3C-SiC film with columnar grain structure [2]. By evaluating more than 1800 grains in TEM images of a double-wedge sample the grain size distributions at six different height levels in the film have been determined with good statistics (Fig. 2b). From these data the evolution of the mean grain size could be determined (Fig. 2c) and, for the first time, the self-similarity of the grain size distribution, predicted earlier by computer simulations of faceted film growth [3,4], could be confirmed experimentally (Fig. 2d, for details see [2]). An alternative way of representing the results of the statistical analysis is shown by Fig. 2e which depicts an “average grain”. This kind of statistical tomography reconstructs average film morphologies from statistical analysis of large-area 2D sections in different film depths.
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1. 2. 3. 4. 5.
E. Spiecker, Phil. Mag. 86, 4941 (2006). E. Spiecker, V. Radmilovic, U. Dahmen, Acta Mater. 55, 3521 (2007) M. Grujicic and S.G. Lai, J. Mater. Synt. Proc. 8, 73 (2000). P. Smereka, X. Li, G. Russo, and D.J. Srolovitz, Acta Mater. 53, 1191 (2005). This work is supported by the AvH-Foundation and by the DOE (DE-AC02-05CH1123).
Figure 1. Schematic showing the geometry of a double-wedge sample for statistical 3D analysis of thin film structure. Along the sample edge the thin film structure can be investigated at each depth over large areas by plan-view TEM (see Ref. [2]).
Figure 2. Statistical analysis of the 3D grain structure of a polycrystalline SiC-film with columnar grain structure by large-area evaluation of TEM images of a double-wedge sample. Figure part e) depicts an “average grain” which has been reconstructed from the statistical grain size data (see Ref. [2]).
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Analytical TEM investigations of Pt/YSZ interfaces V. Srot1, M. Watanabe2, C. Scheu3, P.A. van Aken1, E. Mutoro4, J. Janek4 and M. Rühle1 1. Max Planck Institute for Metals Research, Heisenbergstr. 3, Stuttgart, Germany 2. National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, 1 Cyclotron Rd., Berkeley, CA94720, USA 3. Department Physical Metallurgy and Materials Testing, University of Leoben, FranzJosef-Straße 18, Leoben, Austria 4. Institute of Physical Chemistry, Justus-Liebig-University Gießen, Heinrich-Buff-Ring 58, Gießen, Germany
[email protected] Keywords: metal-ceramic interfaces, EDXS, EELS
Metal-ceramic interfaces are of fundamental importance for a variety of industrial applications and are also of scientific interest. The physical and chemical properties of several technologically relevant nano-structural materials and devices are strongly affected and controlled by the presence of interfaces between the components. Already small amounts of impurities at the interface can dramatically change the properties of the system. Zirconia (ZrO2) is an industrially important ceramic material with a wide range of applications. Due to large volume changes during phase transitions, pure zirconia cannot be used as a high-temperature structural ceramic without stabilisation. Yttria (Y2O3)-stabilised zirconia (YSZ) is an extremely tolerant material with many interesting properties and applications [1]. In solid state electrochemistry platinum (Pt) is a highly important material for electrodes and as potential catalyst films [2]. The aim of our work was the investigation of Pt/YSZ interfaces using analytical transmission electron microscopy (TEM) techniques in order to detect any possible changes (diffusion, segregation…) across the interfaces. For our study Pt films were deposited on the (111) and (100) surfaces of YSZ single crystals using pulsed laser deposition (PLD). The dense and well oriented Pt layers were formed during the thermal treatment at 1023 K [3]. Scanning transmission electron microscopy (STEM) images revealed that the Pt films deposited on the (111) surface of YSZ have a thickness of about 100 to 150 nm and the Pt films on (100) YSZ surface are ca. 1 μm thick (Figure 1). The analytical TEM investigations of Pt/YSZ interfaces included energy-dispersive X-ray spectroscopy (EDXS) and electron energy-loss spectroscopy (EELS). In order to detect any possible variations across the Pt/YSZ interfaces we have performed several EDXS line-scan measurements in the samples with both substrate orientations. At every beam position intensities of the Zr-K, Y-K and Pt-L X-ray emission lines were measured. For the quantitative analysis of EDXS measurements, the new ζ-factor method introduced by Watanabe and Williams [4] was applied. According to our experimental results neither diffusion nor segregation could be detected across the S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 369–370, DOI: 10.1007/978-3-540-85226-1_185, © Springer-Verlag Berlin Heidelberg 2008
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Pt/YSZ interfaces for both investigated YSZ orientations (Figure 2). Additionally the line-scan measurements across the domain boundaries in Pt films with the (111)Pt / (100)YSZ orientation showed that no detectable amounts of elements except Pt were present. The Pt/YSZ interfaces were additionally examined using EELS. Spectra were recorded in the energy ranges of O-K, Y-L2,3, Zr-L2,3 and Pt-M4,5 edges. In order to predict possible atomic models of the investigated interfaces ab initio full multiple scattering FEFF [5] calculations were performed and will be discussed. 1. 2. 3. 4. 5.
A. Navrotsky, Journal of Materials Chemistry 15 (2005), 1883-1890 B. Luerßen et al., Angew. Chem. Int. Ed. 45 (2006), 1473-1476 G. Beck et al., Solid State Ionics 178 (2007), 327-337 M. Watanabe & D.B. Williams, J. Microscopy 221 (2006), 89-109 A.L. Ankudinov et al., Physical Review B 58 (1998) 7565-7576
Figure 1. Bright field STEM images of Pt/YSZ interfaces of a) (111)Pt / (111)YSZ and b) (111)Pt / (100)YSZ investigated orientations.
Figure 2. EDXS line-scan measurement across the (111)Pt / (100)YSZ interface is shown. The values of Zr, Y and Pt atomic fractions are plotted versus the distance from the interface (IF).
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Microstructure and self-organization of nano-engineered artificial pinning centers in YBa2Cu3O7-x coated conductors T. Thersleff1, E. Backen1, S. Engel1, C. Mickel1, L. Molina-Luna2, O. Eibl2, B. Rellinghaus1, L. Schultz1, B. Holzapfel1 1. IFW – Dresden, Institute for Metallic Materials, Postbox 270116, 01171 Dresden 2. University of Tübingen, Auf der Morgenstelle 10, D-72076 Tübingen, Germany
[email protected] Keywords: TEM, FIB, artificial pinning centers, YBCO, 2411, PLD, Moiré
The introduction of nano-sized Artificial Pinning Centers (APCs) to YBa2Cu3O7-x (YBCO) coated conductors immobilizes flux lines at higher fields, thus increasing their usefulness and commercial applicability. Moreover, careful nano-engineering of these APCs facilitates the fine-tuning of the superconducting properties of coated conductors such as enhanced pinning along specific crystallographic orientations or an overall reduction in anisotropy [1]. Understanding the self-organizational behavior of these APCs and their effect on the superconducting properties of YBCO thin films is the focus of this work. TEM lamellae from samples prepared on single crystal SrTiO3 and LaAlO3 substrates using both Pulsed Laser Deposition (PLD) and Chemical Solution Deposition (CSD) methods incorporating APCs were produced using a Carl Zeiss 1540XB Focused Ion Beam (FIB) employing the in-situ lift-out method. TEM investigations on a FEI Tecnai T20 were carried out to elucidate the effect of the processing parameters on the organizational behavior of APCs and subsequently correlate this to the macroscopic properties of these films. The first pinning centers to be introduced are the so-called 2411 phase (Y2Ba4CuMOx where M = metal) into the YBCO matrix using the quasi-multilayer PLD approach described in [2]. Initial TEM investigations looked at the microstructure of YBCO / Y2Ba4CuZrOx (2411-Zr) quasi-multilayers deposited using an off-axis geometry described in [3]. The high resolution image shown in figure 1 indicates that a strong Moiré contrast is present with a spacing of about 2 nm, which may be indicative of a columnar defect structure caused by the presence of either 2411-Zr or BaZrO3 (BZO) nano-particles. Dark-field images were also taken in various orientations. Some results from the (100) reflection are shown in figure 2, which reveals a mixed rotational and lattice constant Moiré. A detailed chemical analysis of the sample was not completed at the time of abstract submission. 1. 2. 3.
S. R. Foltyn et al., Nature Materials 6 (2007) p. 631-642 J. Haenisch et al., Appl. Phys. Lett. 86 (2005), p. 122508/1-3 B. Holzapfel et al., Appl. Phys. Lett. 61 (1992) p. 3178
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YBCO (001)
Moiré
YBCO (100)
2 nm-1
c-axis
Figure 1. High-resolution micrograph of YBCO / Zr-2411 quasi-multilayers in the [100] pole with a defocus close to zero Scherzer focus. The sample surface is located in the upper direction. The FFT subset reveals both the YBCO structure as well as the presence of Moiré reflexes, corresponding to a spacing of approximately 2.0 nm in real space.
Figure 2. Dark field image in the YBCO (100) reflection revealing the presence both mixed and lattice constant Moiré contrast.
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The determination of the interface structure between ionocovalent compounds: the general case study of the Al2O3-ZrO2 large misfit system G. Trolliard, D. Mercurio SPCTS - UMR6638, 123 Av. Albert THOMAS, 87060 Limoges Cedex – France gilles.trolliard@unilim fr Keywords: thin films, interface structure, epitaxy
Electron microscopy is known to be a powerful tool to investigate the structure of crystalline interfaces. Many studies have been carried out with great success to analyze interfaces and defects at the atomic scale in metals or covalent semi-conducting materials. The atomic structure of oxide-oxide interfaces is however far much unknown. HREM observations of interfaces in aligned eutectics demonstrate that complex oxides, such as Al2O3 and ZrO2, can fit along particular crystallographic planes perpendicular to the interface [1]. More recently, we presented the results of a study devoted to the establishment of the interface crystallographic models of such oxides [2]. In this communication we propose an approach to determine the interface models in oxides, based on an HREM approach. The method is proposed in the case of large misfit systems which represents the general case study and should therefore be applicable to all kinds of complex oxide-oxide interfaces. With this aim, the ZrO2-Al2O3 system has been chosen as no intermediate phase could be formed within the interface. In addition, the solid solution occurring between these two compounds is usually very limited in composites produced by conventional methods. Thus the zirconia-alumina system can be considered as non-reactive interface. The studied samples are synthesized by the sol-gel route, which allows growing thin monoclinic zirconia islands of pure zirconia showing a heteroepitaxial growth on (112 0)sa rhombohedral single crystal sapphire substrates. Their heteroepitaxial relationship is first established by electron diffraction on plane view samples (Fig. 1). Different orientation variants were determined. Specific transverse sections where then elaborated by careful selection of the substrate orientation on the base of XRD experiments. They were then prepared normal to the in-plane common directions of the two crystals. In a second step, the metric of the coincidence super-cell between the two lattices was obtained by these HREM images. Indeed, whatever the orientation variant, the HREM images of the interfaces reveal the occurrence of super-periodicities corresponding to common planes across the interface, that account for a full structural continuity between the two phases (Fig. 2). Finally, the geometry of the interfacial coordination polyhedra is deduced and presented as a 3D interface model (Fig. 3). This communication will present the results obtained both on the main orientation variant, corresponding to the most commonly observed variant, and on a secondary orientation variant which is more seldom seen.
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Figure 1. Monoclinic zirconia island showing heteroepitaxial relationship with the alumina substrate. 3d (-101) ZrO2
ZrO2
Al2O3
8d(11 08)Al2O3
Figure 2. Bragg filtered image obtained on the secondary variant and showing the supercell periodicity.
Figure 3. Interface model of the main orientation variant. Specific polyhedra in which the zirconium cation is 7-fold coordinated are involved to build up the interface between alumina and zirconia.
It is shown that the orientation relationship is governed by the occurrence of the common planes crossing the interface, observed in the HREM images (Figs. 2, 3). These common planes are clearly attributed to common cationic planes of both ZrO2 and Al2O3. They correspond to ‘structural walls’ within which the cationic lattices of the two phases are similar and thus perfectly coherent, explaining the continuity across the interface. Alumina-zirconia represents a case of large misfit system. However, except the expected geometrical dislocations and ledges (Fig. 2), the interfaces present very few physical defects. The overall interfacial polyhedra of coordination however suffer high distortions. The lattice mismatch could thus be locally accommodated by an important elastic strain linked to fluctuations of the cation-anion interatomic distances. The predominance of this relaxation mechanism based on the adaptability of the ionocovalent bonds seems a characteristic of oxide-oxide interfaces. 1. 2.
L. Mazerolles, D. Michel, M. Hÿtch, J. Eur. Ceram. Soc.; 25, (2005), 1389-1395 G. Trolliard, R. Benmechta, D. Mercurio and O.I. Lebedev; J. Mater. Chem., 16, 36, (2006), 3640 – 3650
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Simple method to improve quantification accuracy of energy-dispersive X-ray spectroscopy in an analytical transmission electron microscope by specimen tilting T. Walther Department of Electronic & Electrical Engineering, University of Sheffield, Sir Frederick Mappin Building, Mappin Street, Sheffield S1 3JD, UK
[email protected] Keywords: microanalysis, X-ray, energy-dispersive X-ray spectroscopy (EDXS), specimen tilt
The accuracy of measurements of elemental concentrations by energy-dispersive Xray spectroscopy (EDXS) depends on the knowledge of the ionisation cross-sections of the corresponding elements (Z-effect) [1, 2] and proper corrections for absorption (A) [3] and for fluorescence (F) [4]. These effects are now routinely combined for given specimen thickness and density in the so-called ZAF-correction. More recently, fully self-consistent spectrum modelling has become available which models X-ray generation as a function of mass-thickness and then integrates the result over the specimen thickness, e.g. with the PROZA96 software package [5]. If a structure to be analysed is much smaller than the sample thickness, however, two further factors become relevant, namely electron beam broadening [6, 7] and the top-bottom effect, i.e. the depth position of the feature within the foil, as X-rays going through a thicker part of the sample before reaching the detector will be absorbed more strongly. The further a feature is located within the foil towards the bottom surface, the fainter the X-ray signal recorded from it. If this depth position could be measured experimentally it could be taken account of in the absorption correction. Here I show this can be determined indirectly, taking two measurements for different specimen tilts. For a thin layer embedded at depth d within the foil, Figure 1 shows that X-rays from the thin layer have to travel a distance D=d/sin θ through the sample to reach the detector. Hence, for two take-off angles θ1 and θ2 the X-ray path length difference is (1). D2–D1 = d (1/sinθ2–1/sinθ1) The X-ray intensity I for any given element will decay exponentially with the effective path length D, with some attenuation wavelength λ, i.e. for measurements at two different take-off angles θ1 and θ2 and thus different path lengths D1,2: (2) I1,2 = A exp(–D1,2/λ) The intensity ratio then is given by I1/I2 = exp[(D2–D1)/λ] (3) which can be solved for D2–D1 = λ ln(I1/I2) (4). A comparison of equations (1) and (4) with some trigonometry finally yields [8]: d = 4λ ln(I1/I2) sin[(θ1–θ2)/2] cos[(θ1+θ2)/2] /{[cos[(θ1–θ2)] – cos[(θ1+θ2)]} (5) This means one can calculate the depth of the feature in a foil from the intensity ratio of
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two X-ray measurements using the attenuation wavelength λ, the angular difference (θ1–θ2) that can directly be read from the goniometer stage and the average value (θ1+θ2)/2 that can be obtained from the manufacturer of the X-ray detector. With the knowledge of d, the absorption/fluorescence correction for the top-bottom effect can then be refined if either experimental or calculated values of λ are available. As an example, Monte Carlo simulations of electron scattering and X-ray generation as a function of foil thickness have been carried out using the CASINO software (version 2.4.2) [9, 10] to model the absorption effect to a first (non-relativistic) approximation and determine some tentative values of typical attenuation lengths. Figure 2 depicts plots of the calculated X-ray intensities of lines for the case of 2nm thin layers at different depths d within 200nm of GaAs. Elements that can occupy substitutional lattice sites in the III/V semiconductor GaAs and are technologically relevant (i.e. N, Al, P, In and Sb) have been included. Linear regression analysis to logarithmic plots of I(d) yielded the following attenuation lengths: λ(NK)=122nm; λ(AlK)=319nm; λ(PK)=618nm; λ(InL)=2323nm; λ(SbL)=3009nm; λ(InK)=λ(SbK)=∞ for hard X-rays.
intensity [counts]
12 11
2nm thin film in 200 nm GaAs (x10)
10 9 8
Al_K P_K Sb_K
7 6 5
Sb_L In_K In_L N_K
4 3 2 1 0 0
Figure 1. Basic sketch of geometry, angles and distances used in the equations and their analytical solution
50
100 depth d [nm]
150
200
Figure 2. Plot of signal attenuation for 2nm thin layers of of N, Al, P, In or Sb if there is GaAs of thickness d between them and the Xray detector. U=200kV, take-off angle θ=90°
1. H.A. Bethe, Ann. Phys. 397 (1930) 325 2. P. Duncumb and S.J.B. Reed, in “Quantitative Electron Probe Microanalysis”, ed. K.F.J. Heinrich (NBS Spec. Publ., Washington) 298 (1968) 133 3. J. Philibert, Proc. Int. Symp. X-ray Optics and X-ray Microanalysis, ed. H.H. Pattee, V.E. Cosslett and E. Engström (Academic Press, New York) (1963) 379 4. S.J.B. Reed, Brit. J. Appl. Phys. 16 (1965) 13 5. G.F. Bastin, J.M. Dijkstra and H.J.M. Heijligers, X-ray Spectrom. 27 (1998) 3 6. J.I. Goldstein, J.L. Costley, G.W. Lorimer and S.J.B. Reed, Proc. Workshop Anal. Electron Microsc., ed. O. Johari (IIT Research Institute, Chicago) 1 (1977) 315 7. P.A. Crozier, M. Catalano and R. Cingolani, Ultramicroscopy 94 (2003) 1 8. T. Walther, Proc. EMAG conference, Glasgow (2007) in print 9. P. Hovington, D. Drouin and R. Gauvin, Scanning 19 (1997) 1 10. P. Hovington, D. Drouin, R. Gauvin, D.C. Joy and N. Evans, Scanning 19 (1997) 29
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Comparison of transmission electron microscopy methods to measure layer thicknesses to sub-monolayer precision T. Walther Department of Electronic & Electrical Engineering, University of Sheffield, Sir Frederick Mappin Building, Mappin Street, Sheffield S1 3JD, UK
[email protected] Keywords: Energy-dispersive X-ray spectroscopy (EDXS), wetting layer, grain boundary width
Energy-dispersive X-ray spectroscopy (EDX) is an analytical tool to measure the chemistry of thin layers in cross-sectional transmission electron microscopy (TEM). Two methods are compared here using Monte Carlo simulations of electron scattering and X-ray generation as a function of foil thickness by the CASINO software [1]: 1. scanning TEM (STEM) line profiles recorded by scanning a focused electron beam across the layer oriented edge-on: the effective density of atoms within the layer is given by the integral of the measured profile which can be directly calculated up to a certain extent (method 1a) or modelled using Gaussian distributions (method 1b). 2. series of nano-probe TEM illuminations with different beam diameters. The effective layer thickness can be determined from a least-squares fit of atomic ratios plotted as a function of beam diameter; in the following this is called method 2 [2]. The model consists of ½ unit cell of InAs (i.e. 1 monolayer (ML) of In atoms or d=0.14nm) sandwiched between GaAs, imaged edge-on using 200kV electrons and an X-ray detector at 20° take-off angle. For fair comparison the number of electrons in the simulations have been the same for both methods (dose: 1.2×107 e–), distributed either over 100 STEM probes (diameters 2r=0.14, 0.5 or 2nm) or a dozen TEM illumination spots of radii r = 7, 10, 13,…, 40nm. Specimen thicknesses considered were t = 48, 100 and 197nm. From thin bulk InGaAs k-factors for the relative intensities of the X-ray lines have been determined as kInK/GaK=0.351 and kInL/GaL=1.284. These were used, after elimination of some program bugs [3], to convert intensity ratios into atomic % [4]. The apparent widths (FWHM) of the scanned profiles were 0.4nm, 1.5nm and 5nm. Table 1 demonstrates that methods 1a and 1b yield inconsistent results and that their output values scatter by a factor of ~2, without any possibility to guess how reliable a particular measurement would be. Table 2 shows that in method 2, large thicknesses introduce non-linearities due to the effect of beam broadening, which degrades the fit quality. For small and intermediate thicknesses the linear fit is excellent, however. Here, if R2>0.9 then the output of both K- and L-lines is highly accurate. In this case the average d=136±19pm, where the error bar contains statistical and systematic errors [3], corresponds to 0.97±0.14 ML and is a reliable measure of the input value of 1 ML. 1.
P. Hovington, D. Drouin and R. Gauvin, Scanning 19 (1997) 1
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T. Walther, J. Microsc. 215 (2004) 191 and 223 (2006) 165 T. Walther, Proc. EMAG2007, Glasgow (2007) in print S.J.B. Reed, Ultramicroscopy 7 (1982) 405 10000
InL @ t=48nm
intensity [counts]
InL @ t=100nm
1000
InL @ t=197nm
Figure 1. Profiles of In and Ga L-line intensities vs. STEM probe position for 0.14nm wide electron beam and different thicknesses. Please note the logarithmic intensity scale.
GaL @ t=48nm 100
GaL @ t=100nm GaL @ t=197nm
10
-2.2
-1.2
1 -0.2
0.8
1.8
probe position [nm]
Table 1. Effective chemical layer width d [pm] determined from STEM profiling for three probe sizes (diameter 2r) and sample thicknesses (t) using k-factors for K- and Llines as described. Input chemical width: dtrue=140 pm. Labels (a) and (b) refer to methods 1a and 1b, respectively. Mean values for each row are given in the last column. t [nm] 48 100 197
2r [nm] X-ray lines K L K L K L
0.14 (a) (b) 199.8, 181±30 187.2, 175±29 170.4, 149±25 159.1, 142±24 139.2, 119±20 131.0, 113±19
0.5 (a) (b) 193.9, 119±40 175.2, 108±36 173.0, 167±33 154.7, 152±30 140.7, 160±27 128.3, 123±25
GaK / InK ratio * 0.351
t = 48 nm 100 nm 197 nm
400
200
0 0
5
10
15
∅ 173±30 173±34 166±16 148±24 138±29 124±26
Table 2. Effective chemical layer widths d [pm] from method 2 for various sample thicknesses t and ratios of Ga/In K- or L-line intensities after k-factor correction. R2 is linear correlation coefficient.
800
600
2 (a) (b) 182.2; 160±40 204.6; 189±38 189.4, 145±29 176.3, 104±26 176.1, 95±24 163.2, 84±21
20
25
30
35
40
beam radius r [nm]
Figure 2. Plot of the k-factor corrected Ga/In ratio of the K-lines as a function of beam radius. The results for L-lines are very similar.
t R2 0.88 0.91 0.96 [nm] X-ray lines 48 K 143±14 L 127±12 100 K 147±21 L 129±19 197 K 102±18 L 93±16
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Determination of interface structure of YBCO/LCMO by a spherical aberration- corrected HRTEM Z.L. Zhang1, U. Kaiser1, S. Soltan2 and H.-U. Habermeier2 1. Materials Science Electron Microscopy, University of Ulm, 89069 Ulm, Germany 2. Max-Planck-Institut für Metallforschung, D-70569, Stuttgart, Germany
[email protected] Keywords: Cs-corrector, HRTEM, interface
Heterostructure interface plays an important role in modern microelectronics, which could be effectively used to control various electronic and magnetic properties by tuning the atomic structure and chemical composition of the interfaces. The interface also creates new and unexpected phenomena produced by the strong interaction among the electrons. Heterostructure interfaces consisting of superconducting and ferromagnetic manganese oxide attract much attention recently [1-3]. Generally, a chemically pure and atomically sharp interface is needed for achieving particular properties. However, to visualize the atomic structure of interface consisting of light and strong scatters with high precision in detail is only possible by applying spherical-aberration correction techniques [4-5]. Figure 1 is an representative HRTEM image of one interface of a bi-layer heterostructure YBCO(YBa2Cu3O7-δ)-LCMO(La2/3Ca1/3MnO3)-YBCO, which was acquired by an objective-lens Cs-corrected Titan FEI 80-300 microscope under a CS of -1.0 µm and small over-focus. The image indicates a perfect epitaxial relationship along the caxis due to extreme low mismatch. The orientation relationships in between are: YBCO [001] //LCMO [001] and YBCO [100]//LCMO [110]. Under this condition atoms in the adjacent two layers and the interface are clearly visible and imaged as white atoms, which enable to readily determine the atom types combined with atomic model. Further analysis reveals that the interface is composed of CuO plane and –LaCaO plane. Oxygen atom column contrast is clearer and stronger in the LCMO layer than in YBCO layer. Oxygen contrast is not pronounced in CuO2 chain (but, detectable by line-profile) whereas it is remarkable in BaO chains. Using multislice method [6], simulated HRTEM images based on the atomic model are inserted, which demonstrates reasonable fitting with experimental contrast under similar experimental condition, therefore, corroborating the experimental contrast observed and atom configurations at the interface. Figure 2 shows the variations of plane spacing over a rectangular area about 3.0 nm by 1.5nm from position A to B. Closely examination on the spacing reveals that the interface spacing (denoted by arrow 1 and 2 ) is larger, 0.227nm, than the averaged value (denoted by arrow 5 and 6), 0.192 nm, in LCMO layer. Atom planes at the interface slightly expand compared to the adjacent atom planes. Further HRTEM study, combined with exit wave reconstruction, the whole three layer recorded in one image, reveals that such heterostructure usually exhibits two dis-
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tinct interfaces, which are either CuO-LaCaO (as shown in Fig.1.) or BaO-MnO termination layers, one on each side of the LCMO layer. 1 2 3 4 5 6
J. Chakhalian, et al, Science 318(2007)1114. M. Varela, A.R. Lupini, S.J. Pennycook, et al, Solid-State Electronic. 47 (2003) 2245. T. Holden, H.-U. Habermeier, et al , Phy.Rev.B 69(2004)064505. M. A. O’Keefe, Ultramicroscopy 108 (2008) 196. L. Houben, A. Thust, K. Urban, Ultramicroscopy 106 (2006) 200. A. Chuvilin and U.Kaiser, Ultramicroscopy 104 (2005) 73.
Figure 1 (a) High-resolution image of the interface of YBCO/LCMO recorded under a negative CS (-1.0 µm) and defocus, where the LCMO terminates in a LaCa-O plane while YBCO terminates in CuO plane. The simulated image based on the atomic model (left) is inserted. Note that oxygen atom column contrast in CuO plane is weak.
Figure 2. Intensity traces across the interface over a rectangular areas (in Fig.1) illustrating the variation of plane spacing. It is clearly seen that a large spacing exist in the interface, i.e. 0.237 nm. The first neighbor atom planes both in YBCO and LCMO are somehow compressed, which reflect a smaller spacing relative to the averaged values.
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HREM characterization of BST-MgO interface O.M. Zhigalina1, A.N. Kuskova1, A.L. Chuvilin2, V.M. Mukhortov3, Yu.I. Golovko3 and U. Kaiser2 1. Institute of Crystallography, Russian Academy of Sciences, 119333, Moscow, Russia 2. University of Ulm, Albert-Einstein-Alee 11, 89069, Ulm, Germany 3. South Scientific Centre, Russian Academy of Sciences, 344006, Rostov-on-Don, Russia
[email protected] ras ru Keywords: films, ferroelectric, interface
It is known the properties of thin ferroelectric films can be different from bulk materials, that is concerned with an influence of mechanical stress (up to several GPa) at the film – substrate interface [1]. The stress can be partially (or fully) relaxed by means of misfit dislocation formation. So it is necessary to investigate both thin films and heterostucture interfaces at the atomic level to understand relationships between the microstructure and the electrical properties. High resolution electron microscopy (HREM) is a powerful method for the study of the film – substrate interface at the nanometer scale. Combined with geometric phase analysis, useful information can be obtained concerning local strains, variations in lattice parameters in the region of the film – substrate interface. Image analysis was carried out using especially written scripts for Digital Micrograph 3.5 (Gatan) [2]. For TEM and HREM investigations cross sections were prepared by both ion milling in Gatan PIPs 691 and focussed ion beam (FIB) technique in Quanta 200 3D (SMA Company). All samples were characterized in a Tecnai G2 30ST and a FEI Titan 80-300 at accelerating voltage of 300kV, using imaging, electron diffraction and highangle-annular dark-field (HAADF) STEM detector. (Ba0 8Sr0 2)TiO3 (BST) thin epitaxial films were deposited on [001]-oriented MgO substrate by rf sputtering. Recently it has been shown [3] the degree of stress in the epitaxialy grown thin films is a function of the film thickness. The method of geometric phase analysis was used to visualize local strains and extrinsic dislocations in the BSTMgO interface for films with different thickness (5 – 1000 nm). BST thin films revealed a monocrystalline structure with low-angle blocks boundaries (θ < 20) (Figure 1). The main reason of the block structure was a surface geometry (holes and hills) of MgO substrate. Digital analysis of cross sections HREM images (Figure 2) allowed us to visualize and compare misfit dislocations and displacement fields around their cores at the heterostructure interface for films with different thickness. It has been shown by Z-contrast STEM images and image simulation there are two possible ways of film-substrate atomic bonds at the BST-MgO interface: titan-oxygen or barium-oxygen. Analysis of dark field high resolution STEM images has indicated both variants of bonds can be observed.
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1. 2. 3. 4.
Y.S. Kim, D.H. Kim, J.D. Kim, et. al., Appl. Phys. Lett., 86 (2005), p.102907 A.K. Gutakovskii, A.L. Chuvilin, Se Ahn Song, Izvestiya RAS, ser. phys., 71 (2007) p.1464 P.-E. Janolin, Bo-Kuai Lai, Y.I. Yuzuk et. al., Book of abstracts EMF-2007 Bled, Slovenia p.74 The work was supported by RFBR grant № 07-02-12259-ofi and the State Program for Support of Leading Scientific schools, project № NSh-1955.2008.2. We kindly acknowledge SMA Company for the help in sample preparation.
Figure 1. High-resolution dark field image of the block boundary (white arrows) in BST thin film. Cross-section.
b) a)
c)
Figure 2. a) HREM image of BST-MgO interface; b) corresponding FFT; c) maps of variations of lattice parameters combined with (020) – filtered image displaying the (020) planes ending the interfacial dislocations.
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TEM investigations on novel shape memory systems with Ni-depletions D. Schryvers, R. Delville, B. Bartova* and H. Tian EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium * now at Swiss Federal Institute of Technology (EPFL), CH-1015 Lausanne, Switzerland
[email protected] Keywords: martensite, microstructure, precipitation
Many of today’s shape memory systems contain a substantial amount of Ni. The most typical system is of course Nitinol® based on the binary Ni-Ti alloy used near to its equiatomic composition. Another well known system is Ni-Al, which has been very thoroughly studied more for its fundamental physical characteristics rather than its practical applications. Recently, however, new ternary systems have attracted a lot of attention attempting to, e.g., introduce magnetic transitions and driving forces, such as in Ni-Mn-Ga or Co-Ni-Al or lower the hysteresis and increase the transformation temperatures in Ni-Ti-(Au,Pd,Pt), replacing mainly Ni with the ternary compound. At the same time, new studies have revealed unwanted Ni release in commercial wires for medical use. Focusing on Ni thus remains important for understanding the behaviour of martensitic transformations and shape memory applications. In the present lecture we will present results on the ternary system Ni50-xTi50Pdx in which different amounts of Pd substitution on Ni positions lead to special ratios between the austenite and martensite lattice parameters. As a result, the hysteresis of the transformation becomes very narrow and the amount of microtwinning, necessary to yield an invariant plane strain, decreases drastically. Also, when changing the amount of Pd, the type of stable microtwinning changes from Type I & II into compound. A first result is shown in Figure 1 demonstrating the change of microstructure as the content of Pd is decreased from 23 at.% (1.a) to 20 at.% (1.b) and 11 at.% (1.c) where the compatibility condition between austenite and martensite is satisfied. Detwinning and rearrangement of the martensite plates lead to a radically different microstructure which can account for the decrease of hysteresis [1]. Another system is the magnetic Co38Ni33Al29 with a Curie temperature above 90°C due to the replacement of mainly Ni by Co. Depending on the thermal treatment, nanoscale precipitates of fcc or hcp Co grow in the B2 matrix [2], as shown in Figure 2. Due to the large amount of these precipitates, they can seriously change the composition of the matrix thus affecting martensitic and magnetic transformation temperatures. In medical devices an oxide film is formed on the surface inhibiting leakage of toxic Ni into the human body. Even without special thermal treatments, a natural Ti-oxide films forms, which normally should not contain Ni. Recent work, however, has shown that in some cases pure metallic Ni particles can exist inside this Ti-oxide surface film, as seen in Figure 3. Also a Ni3Ti intermediate layer was observed. The composition and structure of the film is investigated by HRTEM, EELS and EFTEM techniques, with samples being prepared by FIB.
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1. 2. 3.
J. Cui, Y.S. Chu, O.O. Famodu et al., Nature Materials 5 (2006) 286-290 B. Bartova, D. Schryvers, Z. Yang, S. Ignacova and P. Sittner, Scripta Materialia 57 (2007) 37-40 We kindly acknowledge support of the projects FWO G.0465.05, FWO G.0180.08N, MULTIMAT MRTN-CT-2004-505226
Figure 1. Example of the changing microstructures observed in Ni50-xTi50Pdx samples.
Figure 2. (a) fcc Co precipitates observed in as-cast Co38Ni33Al29 material (1530°C, 10ms-1) and (b) hcp Co precipitates observed after a 4 hour annealing at 1275°C.
Figure 3. EFTEM revealing (a) metallic Ni particles (black) in the Ti-oxide film (whitegrey) of a cold-rolled Ni-Ti microwire and (b) an intermediate Ni3Ti layer between the metallic core and oxide surface.
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Crystalline-to-amorphous transformation in intermetallic compounds by severe plastic deformation K. Tsuchiya1, T. Waitz2, T. Hara1, H.P. Karnthaler2, Y. Todaka3, M. Umemoto3 1. National Institute for Materials Science, Sengen 1-2-1, Tsukuba, Ibaraki 305-0047, Japan 2. Physics of Nanostructured Materials, University of Vienna, Boltzmanngasse 5, 1090 Vienna, Austria 3. Dept. of Production Systems Engineering, Toyohashi University of Technology, Tempaku-cho Hibarigaoka 1-1, Toyohashi, Aichi 440-8580, Japan tsuchiya
[email protected] Keywords: shear band, twin, martensite
Process of nanostructure formation and crystalline-to-amorphous transformation (CTAT) by high pressure torsion have been studied for various intermetallic compounds, such as, TiNi, ZrCu and Ni3Al. These compounds have been chosen to elucidate the effect of crystal structure on CTAT. Crystal structures of the samples were B2 or B19’ martensite for TiNi, Cm martensite for ZrCu and L12 for Ni3Al. Discs (10 mmφ x 0.85 mm) were cut from a hot-rolled sheet (TiNi) or a casted ingot (TiNi, ZrCu and Ni3Al). After an adequate heat treatment for homogenization, the discs were deformed by high pressure torsion (HPT) apparatus up to 50 turns under an applied pressure of 5 GPa. For TEM observations, 3mmφ discs were cut from the deformed samples and electropolished by Tenupol. HRTEM observations were carried out using a CM 30ST operating at 300 kV. Nanobeam diffraction (NBD) study was done on a JEM-2010F operated at 200 kV. Complimentary X-ray diffractometry (XRD) and optical microscopy (OM) were also carried out. TiNi has been known to undergo CTAT by various SPD processes, e.g., cold rolling[1, 2], shot peening [3] and HPT[4]. We have pointed out that a drastic increase in lattice defects in shear bands is one of the dominant mechanisms of CTAT by cold rolling [2]. Fig.1(a) is a TEM bright field image of B2/amorphous lamellar often seen in HPT sample as well as in cold rolled samples. HRTEM observations and FFT analysis (Fig. 1(b)) indicate that the B2 crystals at the both sides of the amorphous band are in the same orientation. Such structures suggest that the martensitic twin boundaries are preferred site for defect accumulation and the CTAT starts at these boundaries. Long range order was retained even just before the CTAT. Fig.2(a) is an OM image of ZrCu after HPT deformation of 50 turns. Numerous shear bands can be seen running nearly parallel to the shear direction (denoted as SD) with many branching. Fig.2(b) shows a corresponding TEM image. The shear bands, which have very similar morphology to those observed in OM, are seen running nearly parallel to each other, but some of them also exhibit complex branching. Thicker bands exhibit some diffraction contrast but thinner ones seem to be rather featureless. It should be also noted that, compared to the case to TiNi the grains remain to be coarse and the dislocation density is not so high in the surrounding area of the deformation S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 385–386, DOI: 10.1007/978-3-540-85226-1_193, © Springer-Verlag Berlin Heidelberg 2008
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bands. Fig.2(c) exhibits HREM of the deformation bands of a few nm width (denoted by the arrows). Detailed HREM observations and NBD study revealed that these nanoscale deformation bands are amorphous. On the contrary CTAT was not observed for Ni3Al even after HPT deformation of 30 turns; the sample remained to be nanocrystalline of about 50 nm. XRD results indicated significant decrease in the intensity of superlattice reflection was observed in the first few turns of HPT deformation as reported previously [5]. The process and mechanism of CTAT will be compared and discussed for these intermetallic compounds. 1. 2. 3. 4. 5.
J. Koike, D. M. Parkin, Journal of Meterials Research 5 (1990) 1414-1418. H. Nakayama, K. Tsuchiya, Z.-G. Liu, M. Umemoto, K. Morii, K. Shimizu, Materials Transactions 42 (2001) 1987-1993. D. M. Grant, S. M. Green, J. V. Wood, Acta Metallurgica et Materialia 43 (1995) 10451051. Y. V. Tat'yanin, V. G. Kurduymov, V. B. Fedorov, Physics of Metals and Metallography 62 (1986) 133-137. C. Rentenberger, H. P. Karnthaler, Acta Materialia 53 (2005) 3031-3040.
Figure 1. lamellar structures of Ti-50.2Ni HPT deformed for 2 turns. (a) TEM bright field micrograph and selected area electron diffraction (inset). (b) HREM image and corresponding FFT (inset).
Figure 2. microstructures of ZrCu HPT deformed for 50 turns. (a) optical micrograph. (b) TEM bright field image. (c) HREM image.
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EELS quantification of complex nitrides in a 12 % Cr steel M. Albu1, F. Méndez Martin2 and G. Kothleitner1 1. Institute for Electron Microscopy, Graz University of Technology, 8010 Graz, Austria 2. Institute for Materials Science, Welding and Forming, Graz University of Technology, 8010 Graz, Austria
[email protected] Keywords: nitride precipitates, EELS, quantification
In 9–12 % Cr steels, nitrides (MX, M2X and modified Z-phase ((Cr,V,Nb,Fe)N)) are of special interest because of their different contribution to the creep strength of the material. The changes in the chemical composition and their crystallography were investigated using transmission electron microscopy (TEM). The different phases have been identified by using the energy filtered TEM (EFTEM) technique [1] and the corresponding bivariate scatter diagrams of chromium and vanadium jump ratios (Fig. 1a-b). Their elemental composition has been established both with electron energy loss spectroscopy (EELS) and energy dispersive X-ray spectroscopy (EDX). Nevertheless the light elements are not easily quantifiable with EDX, for which reason EELS was involved. Furthermore, the energy-loss near edge structure (ELNES) of the nitrogen ionisation K edge has been used to differentiate between different metastabile nitride phases [2]. Analysis was carried out on a total of 60 thin particles from three differently treated samples: as-received (after tempering at 780 °C), thermally aged and creep loaded at 600 °C for 24 639 h at 115 MPa. The specimens were prepared by extracting the precipitates from the matrix into an amorphous carbon film. In the EEL spectra, edges coming from Nb(M45,M23), N(K), V(L23), Cr(L23) and Fe(L23) were recognised in addition to amorphous carbon from the extraction film. Since the edges from Nb, N, V and Cr are energetically too closely spaced and a conventional edge intensity extraction is not possible, the multiple linear least squares (MLS) fit deconvolution has been employed, fitting suitable references to all overlapping edges. Such an MLS fit together with the references used is shown in Fig. 2 for a mod. Z-phase particle. Computing the relative fit weights and integrating the references over an energy range of 100 eV, allows calculation of the atomic percentages of each element quite accurately, provided experimentally determined cross-sections [3,4] are available, as in this case. The mean elemental concentrations of precipitates under study i.e. the M2X, MX and modified Z-phase shows a good agreement comparing with the outcomes of the thermodynamic model implemented in the software package MatCalc (Table I) [5]. 1. 2.
I. Letofsky-Papst, P. Warbichler, F. Hofer, E. Letofsky, H. Cerjak: Z. Metallk. 95 (2004) 18. F. Hofer, P. Warbichler, A. Scott, R. Brydson, I. Galesic, B. Kolbesen: Journal of Microsc. 204 (2001) 166.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 387–388, DOI: 10.1007/978-3-540-85226-1_194, © Springer-Verlag Berlin Heidelberg 2008
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M. Albu, F. Méndez Martin, B. Sonderegger, G. Kothleitner: Int. Journal Mat. Res., in print F. Hofer: Microsc. Microanal. Microstruct. 2 (1991) 215. J. Svoboda, F.D. Fischer, P. Fratzl, E. Kozeschnik: Mater. Sci. Eng. A 385 (2004) 157.
200 nm
Figure 1. Phase diagram of a representative position from a creep loaded sample and the respective diffraction pattern from the mod. Z-phase –green particle. M23C6 –blue and MX(VN) –yellow were identified as well.
arb. units
V L23
Cr L23
NK
Nb M45
Nb M23
eV
Figure 2. MLS fit using Nb, N, V and Cr -references for a mod. Z-phase particle. With black colour is represented the experimental spectrum and with grey the MLS fit. Table I. Mean concentrations and their standard deviations from EELS measurements of M2X, MX and mod. Z-phase particles compared with the MatCalc simulations. EELS Cr (at.%) M2X 49.8 ± 3.4 MX 11.6 ± 4 mod. Z-phase 31.4 ± 4.2 Simulations M2X MX mod. Z-phase
Cr (at.%) 50.69 0.05 31.31
V (at.%) 17.2 ± 2.1 40.0 ± 4.5 30.1 ± 2.4
Nb (at.%) 6.9 ± 2.7 3.1 ± 1.6
Fe (at.%) 2.6 ± 1.7 1.7 ± 1.1 3.5 ± 1.2
N (at.%) 30.4 ± 4.2 41.2 ± 4.8 32.3 ± 3.4
V (at.%) 14.75 47.39 31.20
Nb (at.%) 1.04 5.40 3.45
Fe (at.%) 0.09 0.02 3.30
N (at.%) 32.87 46.61 30.70
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Formation of ordered solid solution during phase separation in Cu-Ag alloy films F. Misják, P.B. Barna, G. Radnóczi Research Institute for Technical Physics and Materials Science (MFA) of the Hungarian Academy of Sciences, H-1525 Budapest, P.O.Box 49, Hungary
[email protected] hu Keywords: coating, Ag-Cu alloy, phase separation, ordered solid solution
The Cu-Ag binary system belongs to the two component systems of practically non mixing components at thermodynamic equilibrium. At non equilibrium conditions however, substantial mixing could be observed. Cu–Ag thin films prepared by vapour [1] or DC sputter [2-4] deposition show solubility of both components (approx. 10 at.%) in each other as observed by XRD measurements. Besides of enhanced solubility formation of hexagonal (hcp) domains was also observed [4]. These results make probable the formation of metastable (ordered solid solution) phases in this system. Cu-Ag films were prepared by thermal evaporation and co-deposition of Cu and Ag at 10-5-10-6 mbar onto amorphous thin carbon foils at room temperature. The composition of the films covered a wide range of 10-60 at % of Ag. Deposition rates were around 1 nm/s. A CM20 TEM at 200 kV with a Noran EDS was used to determine the morphology, texture and chemical composition. High resolution TEM analysis was performed with a JEOL 3010 operated at 300kV. Films of eutectic composition (60 at% of Ag) are constituted of nanosized grains of slide solutions of Cu and Ag. The grain size is around a few nm as shown in the bright field and dark field images in Fig. 1. Electron diffraction patterns (Figs. 1b and 1c) show that the film has a pronounced one axis texture [5]. The HREM analysis discovered the existence of domains with ordered solid solutions in the grains. Lattice fringe image of an Ag domain is shown in Fig. 2a. The FFT corresponding to the image is presented in Fig. 2b. Indexing the diffraction pattern (Fig. 2b) shows, that this crystal is oriented with its [110] direction parallel to the electron beam. The fitting of the 111 reflections (printed in italic in fig. 2b) shows the epitaxial relation between Cu and Ag crystals as evidenced by the moiré fringe in Fig. 2a. The occurrence of 100 and 110 type forbidden reflections proves that besides the enhanced solubility limit [1-4] and hexagonal ordering [4] the Cu-Ag solid solutions could arrange into ordered solid solution during the structure evolution of the film. The structure of this ordered solid solution is demonstrated in Fig. 2c obtained by Fourier filtering of Fig. 2a. These results suggest that the phase separation occurring during film growth could start by ordering processes. Ordering in solid solution can be one of the routs through the metastable states occurring between supersaturated random solid solutions and phase separated Ag and Cu grains as a final equilibrium state.
S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 389–390, DOI: 10.1007/978-3-540-85226-1_195, © Springer-Verlag Berlin Heidelberg 2008
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1. 2. 3. 4. 5.
A.G. Dirks, J.J. van den Broek, P.E. Wierenga, J. Appl. Phys. 55 (1984) 4248 H.W. Sheng, G. Wilde, E. Ma, Acta Mater. 50 (2002) 475 Hao Chen, Jian-Min Zuo, Acat Materialia 55 (2007) 1617-1628 Smita Gohil, Rajarshi Banerjee, Sangita Bose and Pushan Ayyub, Scripta Materialia (2008) , doi:10.1016/j.scriptamat.2007.12.043 F. Misják, P.B. Barna, A.L. Tóth, T. Ujvári, I. Bertóti, Gy. Radnóczi, Thin Solid Films 516 (2008) 3931–3934
Figure 1. Bright field (a) and dark field (b) as well as electron diffraction patterns (c,d) of eutectic Cu-Ag film co-deposited at room temperature. The diffraction patterns reveal, the textured nature of the film in on axis ((c) and tilted (d) position.
Figure 2. High resolution image of a Ag domain (a) showing superlattice reflections of (001) and (1-10) type in the FFT pattern indicating the presence of an ordered solid solution in the Cu-Ag alloy (b). The index printed in italic in (b) is common for Cu and Ag reflections. Filtered version of the image indicating the position of the unit cell and the ordered structure in the insert (c). x stands for only Ag positions, ● stands for common positions of Ag and Cu.
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Precipitates and magnetic domains in an annealed Co38Ni33Al29 shape memory alloy studied by TEM B. Bartova1,*, D. Schryvers1, N. Wiese2 and J.N. Chapman2 1. EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium 2. Department of Physics and Astronomy, University of Glasgow, Glasgow G12 8QQ, United Kingdom * now at LSME, Ecole polytechnique Fédérale, Station 12, CH-1015 Lausanne, Switzerland
[email protected] Keywords: CoNiAl shape memory alloy, precipitates, magnetic domains, Lorentz microscopy
The Co-Ni-Al system undergoes a martensitic transformation from β-phase (cubic) austenite to L10 (tetragonal) martensite in a temperature range between 93 and 393 K depending on composition. Furthermore, the martensitic start temperature TMs and Curie temperature Tc can be independently controlled by the composition. TMs decreases with increasing content of Co and Al whereas Tc increases with increasing Co content and decreasing amounts of Al [1]. This work presents a detailed study of the microstructure of austenite and an investigation of the relation between magnetic and crystallographic structure. Since TMs for the material studied here is below room temperature, in-situ cooling experiments were performed involving conventional transmission electron microscopy and Lorentz microscopy. The Fresnel mode of Lorentz microscopy was used to study the magnetic domain structure of the sample. The morphology of the sample, following annealing and subsequent quenching, consists of the B2 matrix and a γ-phase. A dark field TEM image reveals small precipitates present in the B2 matrix see Fig.1. The rod-like precipitates have dimensions ranging from 10 to 60 nm for the longest axis. In the high resolution image of Fig. 2a, a single precipitate in the B2 matrix can be seen. From the fast Fourier transform (FFT) pattern it can be concluded that the B2 matrix is viewed along a [110] zone axis. The FFT of the precipitate clearly reveals hexagonal [001] or cubic symmetry. The measured value d100=0.219 nm fits with the values for hcp ε-Co. Figure 2b shows the FFT pattern from the entire area of Fig. 2a. It corresponds to the Burgers orientation relationship which is in good agreement with simulation for the same case shown in Fig. 2c. In some regions close to the hole, the phase transformation is incomplete. At these positions, the habit plane between the austenite and martensite regions can be observed. In Fig. 3, a Fresnel image of such a region is shown. The magnetic domain structure has to accommodate across the interface, an inevitable consequence of the difference between the preferred domain spacing in the two phases. Moreover, in this very thin region, it is clear that there is heavy twinning within the martensite and locally the domain walls adopt irregular zig-zag structures with walls running for variable distances along either of two preferred directions. S. Richter, A. Schwedt (Eds.): EMC 2008, Vol. 2: Materials Science, pp. 391–392, DOI: 10.1007/978-3-540-85226-1_196, © Springer-Verlag Berlin Heidelberg 2008
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K. Oikawa, L. Wulff, T. Iijima et al. APL 79 (20), (2001) 3290-3292 We gratefully acknowledge financial support from the MULTIMAT Marie Curie Research Training network (MRTN-CT-2004-505226).
Figure 1. (a) SEM image of the microstructure of Co38Ni33Al29 annealed alloy consisting of B2 matrix with dispersed γ-phase. (b) Dark-field image of the Co precipitates present in the B2 matrix taken from the reflection marked with a circle. (c) Diffraction pattern taken from the [110] zone axis of B2 matrix (a=0.287 nm)
Figure 2. (a) HRTEM image shows hcp precipitate in B2 matrix. (b) FFT plot taken from whole area of Fig. 2a. (c) Simulated diffraction pattern combining ε-Co precipitate along the [001] zone axis with B2 matrix in the [110] zone.
Figure 3. (a) Fresnel image of a partially transformed region of the specimen. (b) Detail close to the interface between austenite and martensite as indicated.
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On the gallium accumulation at the boundaries of Al layers in FIB prepared TEM specimens P. Favia and H. Bender IMEC, Kapeldreef 75, 3001 Leuven, Belgium
[email protected] Keywords: FIB, EDS, specimen preparation, Ga in Al
Focused ion beam milling in a dual beam FIB/SEM is nowadays the most common TEM specimen preparation method for semiconductor device structures. Numerous studies investigated the damage induced by the ion beam near the outer faces of the TEM lamellae. Methods to reduce this damage are for example low energy Ga ion milling as final milling step or low energy Ar ion milling (ex-situ or in-situ). In many materials the ion beam damage manifests itself as amorphisation of the outermost layer (e.g. ~ 23 nm in Si for 30 kV Ga). In metals the damage can result in an increase of the dislocation density [1]. An estimate of the amount of implanted Ga for different materials is discussed in [2] predicting for example the presence of 4 and 9 at % Ga in the outer 1, martensite plates contain fine internal twins (microtwins) which are the result of stress accommodation at the austenite-martensite habit plane. Electron diffraction shows that the fine twinning occurs along a (1-11) type I mode. Martensite plates are also found to be related to each other along a type II twinning mode. Calculations derived from the Geometrically Non-Linear theory of Martensite (GNLTM) [1] predict the two observed twins along with a (011) compound twin. It can also be demonstrated theoretically that as λ2>1, compound twins are forbidden. Conversely, as λ2
